Electrochemical surface science: basics and applications

Granozzi, Gaetano (Ed.)
Alonso-Vante, Nicolas (Ed.)

October 2019

398

MDPI - Multidisciplinary Digital Publishing Institute

Electrochemical surface science (EC-SS) is the natural advancement of traditional surface science (where gas–vacuum/solid interfaces are studied) to liquid (solution)/electrified solid interfaces. Such a merging between two different disciplines—i.e., surface science (SS) and electrochemistry—officially advanced ca. three decades ago. The main characteristic of EC-SS versus electrochemistry is the reductionist approach undertaken, inherited from SS and aiming to understand the microscopic processes occurring at electrodes on the atomic level. A few of the exemplary keystone tools of EC-SS include EC-scanning probe microscopies, operando and in situ spectroscopies and electron microscopies, and differential EC mass spectrometry (DEMS). EC-SS indirectly (and often unconsciously) receives a great boost from the requirement for rational design of energy conversion and storage devices for the next generation of energetic landscapes. As a matter of fact, the number of material science groups deeply involved in such a challenging field has tremendously expanded and, within such a panorama, EC and SS investigations are intimately combined in a huge number of papers. The aim of this Special Issue is to offer an open access forum where researchers in the field of electrochemistry, surface science, and materials science could outline the great advances that can be reached by exploiting EC-SS approaches. Papers addressing both the basic science and more applied issues in the field of EC-SS and energy conversion and storage materials have been published in this Special Issue.

Science (General)
Chemistry (General)
Inorganic Chemistry

English

9783039216420; 9783039216437

10.3390/books978-3-03921-643-7

surfaces

Electrochemical
Surface Science
Basics and Applications

Edited by

Gaetano Granozzi and Nicolas Alonso-Vante
Printed Edition of the Special Issue Published in Surfaces

www.mdpi.com/journal/surfaces

Electrochemical

Surface

Science

Electrochemical
Surface
Basics and Applications

Special Issue Editors
Gaetano Granozzi
Nicolas Alonso-Vante

MDPI • Basel • Beijing • Wuhan • Barcelona • Belgrade

Science:

Special Issue Editors
Gaetano Granozzi

Nicolas Alonso-Vante

Universita degli Studi di Padova

University of Poitiers

Italy

France

Editorial Office
MDPI
St. Alban-Anlage 66
4052 Basel, Switzerland

This is a reprint of articles from the Special Issue published online in the open access journal Surfaces
(ISSN 2571-9637) from 2018 to 2019 (available at: https://www.mdpi.com/journal/surfaces/special
issues/Electrochemical Surface Science)

For citation purposes, cite each article independently as indicated on the article page online and as
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Contents
About the Special Issue Editors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

ix

Preface to ”Electrochemical Surface Science: Basics and Applications” . . . . . . . . . . . . . .

xi

Nicolas Alonso-Vante and Gaetano Granozzi
Electrochemical Surface Science: Basics and Applications
Reprinted from: Surfaces 2019, 2, 33, doi:10.3390/surfaces2030033 . . . . . . . . . . . . . . . . . .

1

Elizaveta A. Konstantinova, Alexander S. Vorontsov and Pavel A. Forsh
Investigation of Photoelectron Properties of Polymer Films with Silicon Nanoparticles
Reprinted from: Surfaces 2019, 2, 28, doi:10.3390/surfaces2020028 . . . . . . . . . . . . . . . . . .

4

Yongan Tang and Shouzhong Zou
Formic Acid Oxidation on Pd Thin Film Coated Au Nanocrystals
Reprinted from: Surfaces 2019, 2, 27, doi:10.3390/surfaces2020027 . . . . . . . . . . . . . . . . . . 12
Gustav W. Sievers, Anders W. Jensen, Volker Brüser, Matthias Arenz and Marı́a
Escudero-Escribano
Sputtered Platinum Thin-films for Oxygen Reduction in Gas Diffusion Electrodes: A Model
System for Studies under Realistic Reaction Conditions
Reprinted from: Surfaces 2019, 2, 25, doi:10.3390/surfaces2020025 . . . . . . . . . . . . . . . . . . 27
Kiyotaka Abe, Hiroyuki Uchida and Junji Inukai
Electro-Oxidation of CO Saturated in 0.1 M HClO4 on Basal and Stepped Pt Single-Crystal
Electrodes at Room Temperature Accompanied by Surface Reconstruction
Reprinted from: Surfaces 2019, 2, 23, doi:10.3390/surfaces2020023 . . . . . . . . . . . . . . . . . . 40
Benedetto Bozzini, Danjela Kuscer, Matteo Amati, Luca Gregoratti, Patrick Zeller, Tsvetina
Dobrovolska and Ivan Krastev
Spatially Resolved XPS Characterization of Electrochemical Surfaces
Reprinted from: Surfaces 2019, 2, 22, doi:10.3390/surfaces2020022 . . . . . . . . . . . . . . . . . . 51
Nolwenn Mayet, Karine Servat, K. Boniface Kokoh and Teko W. Napporn
Probing the Surface of Noble Metals Electrochemically by Underpotential Deposition of
Transition Metals
Reprinted from: Surfaces 2019, 2, 20, doi:10.3390/surfaces2020020 . . . . . . . . . . . . . . . . . . 71
Jing Wang, Shuwei Zhang, Haihong Zhong, Nicolas Alonso-Vante, Dianqing Li, Pinggui
Tang and Yongjun Feng
Nitrogen-Doped Ordered Mesoporous Carbons Supported Co3 O4 Composite as a Bifunctional
Oxygen Electrode Catalyst
Reprinted from: Surfaces 2019, 2, 18, doi:10.3390/surfaces2020018 . . . . . . . . . . . . . . . . . . 91
Maximilian Eckl, Steve Zaubitzer, Carsten Köntje, Attila Farkas, Ludwig A. Kibler and Timo
Jacob
An Electrochemical Route for Hot Alkaline Blackening of Steel: A Nitrite Free Approach
Reprinted from: Surfaces 2019, 2, 17, doi:10.3390/surfaces2020017 . . . . . . . . . . . . . . . . . . 103
v

Anna Testolin, Stefano Cattaneo, Wu Wang, Di Wang, Valentina Pifferi, Laura Prati, Luigi
Falciola and Alberto Villa
Cyclic Voltammetry Characterization of Au, Pd, and AuPd Nanoparticles Supported on
Different Carbon Nanofibers
Reprinted from: Surfaces 2019, 2, 16, doi:10.3390/surfaces2010016 . . . . . . . . . . . . . . . . . . 116
Vladyslav Mishyn, Patrik Aspermair, Yann Leroux, Henri Happy, Wolfgang Knoll, Rabah
Boukherroub and Sabine Szunerits
“Click” Chemistry on Gold Electrodes Modified with Reduced Graphene Oxide by
Electrophoretic Deposition
Reprinted from: Surfaces 2019, 2, 15, doi:10.3390/surfaces2010015 . . . . . . . . . . . . . . . . . . 127
Mohammad Ali Kamyabi, Ricardo Martı́nez-Hincapié, Juan M. Feliu and Enrique Herrero
Effects of the Interfacial Structure on the Methanol Oxidation on Platinum Single
Crystal Electrodes
Reprinted from: Surfaces 2019, 2, 14, doi:10.3390/surfaces2010014 . . . . . . . . . . . . . . . . . . 139
Svitlana Yarova, Deborah Jones, Frédéric Jaouen and Sara Cavaliere
Strategies to Hierarchical Porosity in Carbon Nanofiber Webs for Electrochemical Applications
Reprinted from: Surfaces 2019, 2, 13, doi:10.3390/surfaces2010013 . . . . . . . . . . . . . . . . . . 155
Yvonne Grunder, Jack Beane, Adam Kolodziej, Christopher A. Lucas and Paramaconi
Rodriguez
Potential Dependent Structure and Stability of Cu(111) in Neutral Phosphate Electrolyte
Reprinted from: Surfaces 2019, 2, 12, doi:10.3390/surfaces2010012 . . . . . . . . . . . . . . . . . . 173
Marco Favaro, Fatwa Firdaus Abdi, Ethan Jon Crumlin, Zhi Liu, Roel van de Krol and David
Edward Starr
Interface Science Using Ambient Pressure Hard X-ray Photoelectron Spectroscopy
Reprinted from: Surfaces 2019, 2, 8, doi:10.3390/surfaces2010008 . . . . . . . . . . . . . . . . . . . 187
Alessandro Zana, Gustav K. H. Wiberg and Matthias Arenz
Oxygen Reduction Reaction on Polycrystalline Platinum: On the Activity Enhancing Effect of
Polyvinylidene Difluoride
Reprinted from: Surfaces 2019, 2, 7, doi:10.3390/surfaces2010007 . . . . . . . . . . . . . . . . . . . 209
Alicja Szczepanska, Gary Wan, Mattia Cattelan, Neil A. Fox and Natasa Vasiljevic
Surface Investigation on Electrochemically Deposited Lead on Gold
Reprinted from: Surfaces 2019, 2, 6, doi:10.3390/surfaces2010006 . . . . . . . . . . . . . . . . . . . 217
Leonardo Girardi, Luca Bardini, Chiara Maurizio, Gian Andrea Rizzi and Giovanni Mattei
Co3 O4 Nanopetals on Si as Photoanodes for the Oxidation of Organics
Reprinted from: Surfaces 2019, 2, 4, doi:10.3390/surfaces2010004 . . . . . . . . . . . . . . . . . . . 230
Tomasz Kosmala, Matı́as Blanco, Gaetano Granozzi and Klaus Wandelt
Potential Driven Non-Reactive Phase Transitions of Ordered Porphyrin Molecules on
Iodine-Modified Au(100): An Electrochemical Scanning Tunneling Microscopy (EC-STM) Study
Reprinted from: Surfaces 2018, 1, 3, doi:10.3390/surfaces2010003 . . . . . . . . . . . . . . . . . . . 243
Gladys Arteaga, Luis M. Rivera-Gavidia, Sthephanie J. Martı́nez, Rubén Rizo, Elena Pastor
and Gonzalo Garcı́a
Methanol Oxidation on Graphenic-Supported Platinum Catalysts
Reprinted from: Surfaces 2019, 2, 2, doi:10.3390/surfaces2010002 . . . . . . . . . . . . . . . . . . . 260
vi

Juan Carlos Calderón Gómez, Verónica Celorrio, Laura Calvillo, David Sebastián, Rafael
Moliner and Marı́a Jesús Lázaro Elorri
Electrochemical Behavior of Pt–Ru Catalysts Supported on Graphitized Ordered Mesoporous
Carbons toward CO and Methanol Oxidation
Reprinted from: Surfaces 2019, 2, 1, doi:10.3390/surfaces2010001 . . . . . . . . . . . . . . . . . . . 276
Ebru Özer, Ilya Sinev, Andrea M. Mingers, Jorge Araujo, Thomas Kropp, Manos Mavrikakis,
Karl J. J. Mayrhofer, Beatriz Roldan Cuenya and Peter Strasser
Ir-Ni Bimetallic OER Catalysts Prepared by Controlled Ni Electrodeposition on Irpoly and Ir(111)
Reprinted from: Surfaces 2018, 1, 13, doi:10.3390/surfaces1010013 . . . . . . . . . . . . . . . . . . 291
Ebru Özer, Zarina Pawolek, Stefanie Kühl, Hong Nhan Nong, Benjamin Paul, Sören Selve,
Camillo Spöri, Cornelius Bernitzky and Peter Strasser
Metallic Iridium Thin-Films as Model Catalysts for the Electrochemical Oxygen Evolution
Reaction (OER)—Morphology and Activity
Reprinted from: Surfaces 2018, 1, 12, doi:10.3390/surfaces1010012 . . . . . . . . . . . . . . . . . . 313
Momath Lo, Rémi Pires, Karim Diaw, Diariatou Gningue-Sall, Mehmet A. Oturan,
Jean-Jacques Aaron and Mohamed M. Chehimi
Diazonium Salts:
Versatile Molecular Glues for Sticking Conductive Polymers to
Flexible Electrodes
Reprinted from: Surfaces 2018, 1, 5, doi:10.3390/surfaces1010005 . . . . . . . . . . . . . . . . . . . 326
Emanuele Salvietti, Walter Giurlani, Maria Luisa Foresti, Maurizio Passaponti, Lorenzo
Fabbri, Patrick Marcantelli, Stefano Caporali, Stefano Martinuzzi, Nicola Calisi, Maddalena
Pedio and Massimo Innocenti
On the Contrasting Effect Exerted by a Thin Layer of CdS against the Passivation of Silver
Electrodes Coated with Thiols
Reprinted from: Surfaces 2018, 1, 4, doi:10.3390/surfaces1010004 . . . . . . . . . . . . . . . . . . . 342
Francesco Giannici, Giuliano Gregori, Alessandro Longo, Alessandro Chiara, Joachim Maier
and Antonino Martorana
X-ray Absorption under Operating Conditions for Solid-Oxide Fuel Cells Electrocatalysts: The
Case of LSCF/YSZ
Reprinted from: Surfaces 2019, 2, 3, doi:10.3390/surfaces1010003 . . . . . . . . . . . . . . . . . . . 356
Long Zhang, Ruth Belinda Domı́nguez Espı́ndola, Benjamin B. Noble, Vinicius R. Gonçales,
Gordon G. Wallace, Nadim Darwish, Michelle L. Coote and Simone Ciampi
Switchable Interfaces: Redox Monolayers on Si(100) by Electrochemical Trapping of
Alcohol Nucleophiles
Reprinted from: Surfaces 2018, 1, 2, doi:10.3390/surfaces1010002 . . . . . . . . . . . . . . . . . . . 365
Martina Fracchia, Paolo Ghigna, Alberto Vertova, Sandra Rondinini and Alessandro
Minguzzi
Time-Resolved X-ray Absorption Spectroscopy in (Photo)Electrochemistry
Reprinted from: Surfaces 2018, 1, 11, doi:10.3390/surfaces1010011 . . . . . . . . . . . . . . . . . . 374

vii

About the Special Issue Editors
Gaetano Granozzi is Full Professor of Inorganic Chemistry at University of Padova since 1990,
where he teaches the courses Inorganic Chemistry of Materials, and Catalysis and Surface
Chemistry. He has been Director of the PhD Course in Science and Engineering of Materials and
Nanostructures of the University of Padova from 2005 to 2016. He is Leader of the Surface Science
and Catalysis Group (SSCG) of the Department of Chemical Science of the University of Padova
(http://www.chimica.unipd.it/surfacescience/). Granozzi’s research interests are in the areas of
materials chemistry, with particular emphasis of heterogeneous catalysis and surface chemistry.
Granozzi has authored over 320 scientific publications and his current h-index is 40, representing
approximately 7541 citations (Google Scholar).
Nicolas Alonso-Vante has served as Professor at the University of Poitiers since September
1997, where he has been teaching the following graduate courses: Specific Analysis of Solids;
Electrocatalysis, Photocatalysis; Conversion and Storage of Chemical Energy. From 1985 to 1997,
he was a Senior Scientist at the Hahn-Meitner-Institut-Berlin (now Helmholtz-Zentrum Berlin).
His current main research interests at IC2MP UMR CNRS 7285 include (photo)electrochemistry
and (photo)electrocatalysis of novel materials using various ex situ and in situ techniques, fuel
generation, interfacial characterization, and surface analytical techniques. He is the author of over
250 publications and book chapters, editor of a two-volume e-book on electrochemistry in Spanish,
and the author of 2 books and 6 patents. His current h-index 46 corresponding to ca. 6750 citations
(RG source).

ix

Editorial

Electrochemical Surface Science:
Basics and Applications
Nicolas Alonso-Vante 1, * and Gaetano Granozzi 2, *
1
2

*

IC2MP-UMR CNRS 7285, University of Poitiers, 86022 Poitiers, France
Department of Chemical Sciences, University of Padova, 35131 Padova, Italy
Correspondence: nicolas.alonso.vante@univ-poitiers.fr (N.A.-V.); gaetano.granozzi@unipd.it (G.G.)

Received: 12 June 2019; Accepted: 1 July 2019; Published: 4 July 2019

The great success of the Surfaces Special Issue entitled “Electrochemical Surface Science (EC-SS):
Basics and Applications” reflects the great vitality and relevance of the addressed topic. EC-SS stems from
the merging of two different disciplines, i.e., surface science (SS) and electrochemistry (EC), which dates
back to ca. four decades ago. The two separate disciplines mainly contributed either a methodological
approach (SS, aiming at understanding the microscopic processes occurring at surfaces/interfaces at
an atomic level) or a selection of the objects to study (EC, aiming at describing the phenomena at
electrified interfaces). The combination of the two is nowadays making possible the challenge of
unraveling the complex mechanisms of the electron transfer processes occurring at electrodes, i.e., truly
interface-driven phenomena which encompass both solid/liquid, solid/solid, and solid/gas interfaces.
This hybrid approach has had an enormous impact on fields such as electrocatalysis [1–3], solar energy
harvesting [4–6], corrosion, electrochemical energy storage and conversion devices [7], and sensors [8].
Among them, electrocatalysis is omnipresent and plays a key role. Actually, processes at electrodes
are often kinetically limited to efficiently run multi-charge transfer reactions. An electrocatalyst is
usually needed, i.e., a substance that can reduce the overall activation barrier height of the redox
chemical reaction via complex surface-chemistry steps (adsorption/desorption of reactants and products,
low kinetic barriers for charge transport) and determine the product selectivity distribution. The figures
of merit of an electrocatalyst are nowadays determined following the standard parameters of catalysis,
i.e., turnover frequency and number. Determining such parameters and correlating them with the
electrocatalyst structure is a task highly facilitated by the adoption of the hybridized EC-SS method.
A historical perspective is useful to better understand the evolution of electrocatalysis. At the
beginning of the 20th century (1905), Julius Tafel [9], in Switzerland, reported on the hydrogen
evolution reaction (HER) on various electrode materials, thus establishing a quantitative method for
HER electrocatalysts benchmarked through the “Tafel equation” [10]. The HER two-electron process,
which started to be academically studied in the 1950s, is still under development in many laboratories
in the world [11]: the main goal is to provide a sustainable route for the preparation of molecular
hydrogen through the electrochemical splitting of water (water splitting; WS). Actually, WS is expected
to promote the envisioned hydrogen economy [12], based on molecular hydrogen as an energy vector
for the development of a sustainable energy infrastructure established on the efficient interconversion
of chemical energy into electricity and vice versa.
One of the current key concepts in electrocatalysis is the replacement of noble-metal-based
electrocatalysts with those based on elements that are abundant on Earth [13,14]. The role played by
the synergetic EC-SS approach in such a paradigmatic revolution is similar to that it already played
in the 1980s, when platinum-based electrocatalysts were optimized [15,16]. In addition, nowadays,
two more relevant key concepts have appeared, i.e., in situ and operando techniques. In situ or
operando characterization tools aim to monitor the electrochemical reaction while the electron transfer
process is occurring [17,18]. The difference between the two terms is subtle. The in situ term relates
to experiments where the experimental conditions (e.g., pressure, atmosphere, potential, current,
Surfaces 2019, 2, 455–457; doi:10.3390/surfaces2030033

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Surfaces 2019, 2

electrolyte, etc.) are controlled during acquisition, but no temporal discrimination is explicitly taken
into account. On the other hand, operando tools are related to the study of the system in real life
applications. When applied to electrocatalysis, this means obtaining more detailed electrochemical
information while monitoring the working electrodes with other techniques, such as X-ray diffraction
(XRD), transmission electron microscopy (TEM), atomic force microscopy (AFM), Raman spectroscopy,
ultraviolet visible (UV-vis) absorption spectroscopy, X-ray absorption near-edge structure (XANES),
nuclear magnetic resonance (NMR), X-ray photoelectron spectroscopy (XPS), etc.
In this Special Issue, a total of 27 scientific papers (one of which is a review) report some of the
latest advances in the field of EC-SS. It was a particular target of the Guest Editors to demonstrate the
importance and the large scope of EC-SS through examples from a variety of systems and applications.
Four papers are related to innovative methods for characterization. As expected, electrocatalysis papers
make up the lion's share, but other interesting topics such as sensors, switchable interfaces, potential
driven self-assembly on surfaces, and flexible electrodes are also addressed. The Guest Editors hope
that the readers will appreciate the different contributions closest to their own field of research.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

3

Article

Investigation of Photoelectron Properties of Polymer
Films with Silicon Nanoparticles
Elizaveta A. Konstantinova 1,2,3, *, Alexander S. Vorontsov 1 and Pavel A. Forsh 1,2,3
1
2
3

*

Faculty of Physics, M.V. Lomonosov Moscow State University, Leninskie Gory 1-2, Moscow 119991, Russia;
as.vorontsov@physics.msu.ru (A.S.V.); phorsh@mail.ru (P.A.F.)
Department of Nano-, Bio-, Information Technology and Cognitive Science, Moscow Institute of Physics and
Technology, Institutskij 9, Dolgoprudny, Moscow 141701, Russia
National Research Center Kurchatov Institute, Akademika Kurchatova 1, Moscow 123182, Russia
Correspondence: liza35@mail.ru

Received: 21 February 2019; Accepted: 11 May 2019; Published: 13 May 2019

Abstract: Hybrid samples consisting of polymer poly-3(hexylthiophene) (P3HT) and silicon
nanoparticles were prepared. It was found that the obtained samples were polymer matrixes
with conglomerates of silicon nanoparticles of different sizes (10–104 nm). It was found that, under
illumination, the process of nonequilibrium charge carrier separation between the silicon nanoparticles
and P3HT with subsequent localization of the hole in the polymer can be successfully detected
using electron paramagnetic resonance (EPR) spectroscopy. It was established that the main type of
paramagnetic centers in P3HT/silicon nanoparticles are positive polarons in P3HT. For comparison,
samples consisting only of polymer and silicon nanoparticles were also investigated by the EPR
technique. The polarons in the P3HT and Pb centers in the silicon nanoparticles were observed.
The possibility of the conversion of solar energy into electric energy is shown using structures
consisting of P3HT polymer and silicon nanoparticles prepared by different methods, including the
electrochemical etching of a silicon single crystal in hydrofluoric acid solution and the laser ablation
of single-crystal silicon in organic solvents. The results can be useful for solar cell development.
Keywords: polymer; silicon nanoparticles; EPR spectroscopy; photoconversion

1. Introduction
Solar cells based on crystalline or polycrystalline silicon are now the dominant technology
globally [1–3]. However, the cost of such batteries is quite high, which is leading to the rapid
development of technologies using amorphous silicon [4,5]. In addition to the low cost of devices,
the use of amorphous silicon can decrease the thickness of solar cells, as well as their weight and
material consumption, due to its higher absorption ability. However, the efficiency of solar cells based
on amorphous silicon remains quite low, at 14%, compared to crystalline solar cells (approximately
25%) [6]. Increasing the efficiency of amorphous silicon-based solar cells is potentially possible
using semiconductor nanocrystals. However, this is a very complex problem because it requires
the development of methods for improving the injection and transport of charge carriers in such
structures. Therefore, the development of cheap photo-electric converters has become a subject
of great interest in the last years [1–4]. One of the perspective directions of the depreciation of
phototransformation is the development of solar elements on the basis of polymers [7,8]. However, at
the moment, the effectiveness of the phototransformation of such solar elements and their stability
are low. The available literary data [7,9] demonstrate the prospects of use of hybrid structures on
the basis of silicon nanoparticles (nc-Si) and organic compounds for transformation of solar energy.
Such devices have the advantage of solution preparation but at the same time are characterized by
a much broader spectral range of absorption because of inorganic semiconductors. It should be
Surfaces 2019, 2, 0; doi:10.3390/surfaces2020028

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borne in mind that the absorption coefficient of nc-Si with a diameter of about 10 nm and more is
comparable with that of bulk Si. For small nc-Si with quantum confinement effects, the band gap
strongly increases and the absorption coefficient grows. There have only been a few investigations
of silicon/organic semiconductor heterojunctions for solar cells [7,9–11]. However, many questions
concerning the correlation between the conditions of preparation of nc-Si and the electrophysical and
optical properties of such hybrid structures still have no definite answer. Since point defects in solids
are the centers of capture of charge carriers and limit the transport of charge carriers, their study is an
important problem [12]. A powerful tool for studying the nature and properties of defects is electron
paramagnetic resonance (EPR) spectroscopy [13]. The structure and properties of defects in nc-Si and
other semiconductor nanoparticles have been successfully studied with EPR spectroscopy, taking
advantage of the high specific area of these materials [13,14]. In spite of the different types of nc-Si
structures, different preparation and storage conditions, the dominant type of defect (paramagnetic
centers) in nc-Si is an Si dangling bond at the Si/SiO2 interface or Pb center [14,15]. There are two
types of Pb centers, Pb0 (at the (111), (100) Si/SiO2 interface) and Pb1 (at the (100) Si/SiO2 interface),
which are characterized by different parameters of EPR spectra [14]. The Pb center concentration is
very sensitive to vacuum heating and oxidation [14]. During vacuum heating (at temperatures higher
than 300–400 ◦ C), an Si dangling bond similar to that in amorphous silicon is detected in nc-Si [14].
A free electron EPR signal is observed in some samples of nc-Si [14]. EX center defects are detected
in high-temperature oxidized nc-Si [14]. In poly-3(hexylthiophene) (P3HT), the main type of defect
center is a positive polaron [9]. Therefore, in this work, we investigated the properties of defect states
in hybrid samples consisting of polymer P3HT and silicon nanoparticles and in samples consisting
only of polymer and silicon nanoparticles for comparison. We also measured the current–voltage
characteristics of structures containing silicon nanocrystals in a polymeric matrix.
2. Experimental
Silicon nanoparticles were prepared by two methods [15,16]. The first method was the
electrochemical etching of monocrystal plates of silicon of n-type of conduction in a solution of
hydrofluoric acid (HF) and ethanol (C2 H5 OH). After etching, the samples were dried, and then silicon
particles were removed from the plate mechanically (nc-Si(1)). The second method of manufacturing
silicon nanoparticles (nc-Si(2)) was the pulse laser ablation of a substrate of monocrystal silicon in organic
solvents (C6 H5 Cl, benzene chloride, and CHCl3 , chloroform). For the formation of the structures,
the received nanoparticles were immediately added to the P3HT solution. Then, the resulting mixture
was applied on the glass substrate containing an ITO layer as one of electrical contacts. The second
was the aluminum contact which was sprayed from above onto the P3HT film. The current–voltage
characteristics of the samples were measured using Keithley 6487.
The investigation of structure of the prepared samples was performed using scanning electronic
microscopy (using a Carl Zeiss Supra 40-30-87 microscope, Oberkochen, Germany). The features
of the internal structure of the polymer and the characteristics of the obtained silicon nanoparticles
were determined by means of the Raman technique. For the Raman spectra registration, the Horiba
Jobin Yvon HR800 spectrometer was used. A helium–neon laser (λ = 632.8 nm) was the source of
light excitation. Raman signal recording was made in a configuration of reflection. A digital camera
with a charge-coupled device (CCD) matrix was used as the detector. For the direct detection of
separation processes of photoinduced charges in structures of nc-Si/P3HT, EPR spectroscopy was
used. Measurements were performed using the Bruker EPR-spectrometer ELEXSYS-500, Rheinstetten,
Germany (with a frequency of 9.5 GHz and a sensitivity of 5 × 1010 spin/mT). The EPR spectra simulation
was carried out using the EasySpin MATLAB toolbox [17]. Illumination was carried out immediately
in the spectrometer cavity by means of a mercury lamp with high pressure (Δλ = 270–900 nanometers,
power is 50 W).

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Surfaces 2019, 2, 0

3. Results and Discussions
Figure 1 shows an example of the structure of the nc-Si(1)/P3HT. Similar images were observed
also for the other samples, nc-Si(2)/P3HT. According to Figure 1, the samples under investigation
represent conglomerates of nanoparticles of various sizes (10–104 nm) dissolved in a polymeric matrix.

Figure 1. SEM image of nc-Si(1)/poly-3(hexylthiophene) (P3HT) samples.

Raman spectra were similar for both types of samples. The Raman spectrum of the nc-Si(2)/P3HT
samples is shown in Figure 2. According to the literature data, the peak near 430 and 520 cm−1
in the Raman spectrum corresponds to the C–C deformation mode of the polymer and silicon
nanocrystals [11,18], respectively; the peak near 728 cm−1 corresponds to an antisymmetric C–S–C
ring skeleton deformation in the thiophene ring of the polymer [19]; and the peak near 1390 and
1450 cm−1 can be assigned to the C–C skeletal stretch mode and the C=C symmetric stretch mode of
the polymer, respectively [20,21]. The other less intense peaks observed in the wavenumber range
of 570, 850, 1000, 1100, 1150, 1250 and 1525, 1560 cm−1 can be assigned to the C–C rotational mode,
C=C deformation mode, C–C stretching mode [21], the C–H bending mode [22], the C–C symmetric
stretching mode [23], a combination of the C–C stretching and the C–H stretching mode [21,24,25] and
the C = C antisymmetric stretch mode [20,23], respectively. Thus, according to Figure 2, the Raman
spectrum consists of lines caused by a polymeric matrix and also a poorly resolved peak, which is
characteristic of silicon nanocrystals. The results of a Raman investigation confirmed the data of
electronic microscopy, showing that all samples under investigation represent the non-uniform structure
of a polymeric matrix with the silicon nanoparticles dissolved in it.

Figure 2. Raman spectrum of the nc-Si(2)/P3HT samples.
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Surfaces 2019, 2, 0

For the direct detection of the separation processes of photoinduced charges in the nc-Si/P3HT
structures, the EPR-spectroscopy technique was used. In Figure 3a, the EPR spectra of nc-Si(1)/P3HT
samples in the dark and under illumination are presented.

Figure 3. (a) EPR spectra of nc-Si(1)/P3HT samples in the dark (1) and under illumination (2).
(b) Experimental and simulated EPR signals of nc-Si(1)/P3HT samples.

The EPR spectra are characterized by two overlapped signals: one is poorly resolved, with a
g-factor equal to 2.0050 originating from the silicon dangling bonds [26], while the second one has an
anisotropic form and is described by the following parameters: g1 = 2.0030, g2 = 2.0022, g3 = 2.0013,
extracted from computer simulation. The result of the computer simulation is shown in Figure 3b.
According to the literature data, this EPR signal can be attributed to the positive polarons in P3HT [9,21].
Under illumination, an amplitude of a polaron EPR signal increases, pointing to the separation process
of photoexcited charge carriers. To be sure that the positive polarons are localized in P3HT, we have
investigated these structures separately. In Figure 4a, the EPR spectra for P3HT samples in similar
conditions are shown, both in the dark and under illumination.

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Surfaces 2019, 2, 0

Figure 4. (a) EPR spectra of P3HT samples in the dark (1) and under illumination (2). (b) Experimental
and simulated EPR signals of P3HT samples.

According to the computer simulation, the EPR spectrum consists of two EPR signals with
parameters that differ slightly: EPR(I) g1 = 2.0028, g2 = 2.0019, g3 = 2.0009 and EPR(II) g1 = 2.0031,
g2 = 2.0021, g3 = 2.0012. The result of the computer simulation is shown in Figure 4b. The values of
the g-factors show the polaron nature of the EPR-centers [9,21]. Probably, we observe the polarons in
two different orientations. Thus, data from the EPR spectroscopy are direct proof of the formation of
polarons in nc-Si/P3HT structures as a result of the separation of charge carriers with the subsequent
localization of a hole in P3HT. Note that this process takes place under the conditions of indoor light
and are amplified under additional illumination.
The silicon nanoparticles was measured for a comparison of the defects’ natures and properties.
Figure 5 shows the EPR spectra measured under dark conditions and in the presence of illumination.

8

Surfaces 2019, 2, 0




,(35DX



















%P7
Figure 5. EPR spectra of nc-Si(1) in the dark (1) and under illumination (2).

As follows from the presented data, the EPR spectrum has an anisotropic structure. The following
values of the g-tensor were obtained from computer simulation: g1 = 2.0021, g2 = 2.0088. The EPR
signal with these anisotropic parameters can be attributed to the paramagnetic centers, which are
Pb centers (a silicon dangling bond at the Si/SiO2 interface) according to the literature data [27,28].
Therefore, an oxidation of silicon nanoparticles takes place in air. Under illumination, the intensity of
the EPR spectrum increases. The effect of illumination was reversible, showing recharge processes in
the samples.
For the determination of the photovoltaic parameters of the samples and qualitative evaluation
of the effectiveness of solar energy conversion, the current density–voltage (abbreviated as
current–voltage [8,9]) characteristics were measured (Figure 6). White light illumination was used
with an intensity of 50 mW/cm2 . As can be seen from Figure 6, nc-Si(1)/P3HT samples have a higher
open-circuit voltage (0,2 V) than nc-Si(2)/P3HT samples (0,16 V). Although these open-circuit voltage
values are not a record, it is possible in principle to use nc-Si/P3HT samples for photoconversion.







-$VP

























89
Figure 6. Current–voltage characteristics of the studied samples. Dependence 1—nc-Si(1)/P3HT
samples; dependence 2—nc-Si(2)/P3HT samples.

Although the efficiency of the solar energy conversion of the samples in our study was low (not
exceeding 5%), we have demonstrated that EPR spectroscopy is an effective tool for the detection of
polarons in nc-Si/P3HT prepared by different techniques.
9

Surfaces 2019, 2, 0

4. Conclusions
The structures consisting of P3HT polymer and silicon nanoparticles, in which silicon nanoparticles
are prepared by electrochemical etching and the laser ablation of single-crystal silicon, have the
possibility of converting solar energy into electric energy. In nc-Si samples, Pb centers are present
according to the EPR data. A lack of Pb centers in the nc-Si/P3HT structures evidences an absence of
oxygen in the samples. Positive polarons are detected in P3HT samples. Using EPR spectroscopy, it has
been found that it is possible to detect the separation of photoexcited charge carriers in nc-Si/P3HT
samples prepared both by the electrochemical etching of a silicon single crystal in hydrofluoric acid
solution and by using the laser ablation of single-crystal silicon in organic solvents. The obtained data
is useful for solar energy application since they demonstrate the possibility of the preparation of solar
cell samples using simple and cheap methods of sample synthesis (for example, the electrochemical
etching of silicon monocrystals).
This work was supported by the Russian Foundation for Basic Research (Grant No 18-29-23005).
Author Contributions: Conceptualization, E.A.K. and P.A.F.; Methodology, E.A.K.; Software, A.S.V.; Validation,
E.A.K. and A.S.V.; Formal Analysis, P.A.F.; Investigation, E.A.K., A.S.V. and P.A.F.; Resources, A.S.V.; Data
Curation, E.A.K.; Writing-Original Draft Preparation, E.A.K.; Writing-Review & Editing, E.A.K., A.S.V. and P.A.F.;
Visualization, P.A.F.; Supervision, E.A.K.; Project Administration, A.S.V.; Funding Acquisition, P.A.F.
Funding: This research was funded by the Russian Foundation for Basic Research Grant number 18-29-23005.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

11

Article

Formic Acid Oxidation on Pd Thin Film Coated
Au Nanocrystals
Yongan Tang 1,† and Shouzhong Zou 2, *,†
1
2

*
†

DOD Technologies Inc., Cary, IL 60013, USA; tangy3@miamioh.edu
Department of Chemistry, American University, Washington, DC 20016, USA
Correspondence: szou@american.edu
Department of Chemistry and Biochemistry, Miami University, Oxford, OH 45056, USA.

Received: 5 April 2019; Accepted: 6 May 2019; Published: 10 May 2019

Abstract: Cubic, octahedral, and rhombic dodecahedral gold nanocrystals enclosed by {100}, {111}, and
{110} facets, respectively, were prepared by a seed-mediated growth method at the room temperature.
Palladium thin films were coated on these Au nanocrystals by a redox replacement approach to
explore their catalytic activities. It is revealed that formic acid and carbon monoxide oxidation in
0.1 M HClO4 on Au nanocrystals coated with one monolayer (ML) of Pd are facet-dependent and
resemble those obtained from corresponding Pd single crystals and Pd films deposited on bulk
Au single crystals, suggesting epitaxial growth of Pd overlayers on the Au nanocrystal surfaces.
As the Pd film thickness increased, formic acid oxidation current density decreased and the CO
oxidation potential moved to more negative. The catalytic activity remained largely unchanged after
3–5 MLs of Pd deposition. The specific adsorption of (bi)sulfate was shown to hinder the formic
acid oxidation and the effect decreased with the increasing Pd film thickness. These observations
were explained in the framework of the d-band theory. This study highlights the feasibility of
engineering high-performance catalysts through deposition of catalytically active metal thin films on
facet-controlled inert nanocrystals.
Keywords: formic acid oxidation; Au nanocrystals; Pd thin films; electrocatalysis; d-band theory

1. Introduction
Direct formic acid fuel cells (DFAFCs) have attracted much attention in the past decade because
they have the advantages of high power density, fast oxidation kinetics, high theoretical cell potential,
and mild fuel crossover as compared to direct methanol fuel cells [1–5]. As the anode reaction of
DFAFCs, electrooxidation of formic acid has been extensively studied [2,5]. Although the nature of
the reactive intermediate is under active debate, it is generally accepted that there are two reaction
pathways in formic acid oxidation (FAO) [2,6]. The direct pathway (dehydrogenation) involves the
adsorption of reactive intermediates and their direct oxidation to carbon dioxide, while the indirect
pathway goes through the dehydration of formic acid and forms strongly adsorbed carbon monoxide
(COads ) that poisons catalysts [2,6]. It has been long recognized that the direct pathway is dominant
on Pd while the indirect pathway is more prevalent on Pt, which suggests that Pd is a better catalyst
than Pt for DFAFCs [6,7]. Indeed, using unsupported and carbon-supported Pd as anode catalysts in
DFAFCs, Masel and coworkers obtained better fuel cell performance than Pt-based catalysts at high
power densities [8,9]. Despite the higher activity observed on Pd, a significant performance decay was
observed after only a few hours of operation [9]. Different adsorbed species, including COads , have
been proposed to explain the deactivation of the Pd catalysts [2]. In addition, it has been shown that
Pd dissolution occurred in an acidic media, resulting in a decreased active surface area [10,11].
Significant efforts have been devoted to develop Pd-based catalysts with high activity as well as
long-term durability and stability [2,5,12]. A commonly employed approach is to introduce a second
Surfaces 2019, 2, 372–386; doi:10.3390/surfaces2020027

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Surfaces 2019, 2

metal either to form an alloy with Pd or as a support for Pd overlayers. Of particular interest in
the present work is the Pd–Au combination. Zhou and Lee reported that Au core Pd shell (Au@Pd)
nanoparticles showed significantly improved FAO activity and durability, which was attributed to the
Au and Pd electronic interactions [13]. More detailed studies by others revealed that the improved
FAO activity depends on the Au/Pd ratio [14] and the degree of alloying [15]. Lee et al. found that FAO
activity and durability of Pd3 Au nanoparticles can be significantly improved by increasing the surface
Pd content through CO-induced surface aggregation [16]. Hong et al. studied FAO on structurally
well-controlled PdAu nano-octahedra, and revealed that PdAu alloy octahedra are about 30% more
active than similar size Au@Pd octahedra, but both are more active than pure Pd octahedra [17].
By using a seed-mediated growth method, Fermín and coworkers prepared Au@Pd nanoparticles with
varying Pd shell thickness from 1 to 10 nm [18]. FAO activity on these core-shell particles was shown
to increase with the Pd shell thickness, which was attributed to the lattice strain effect. More recently,
by combining differential electrochemical mass spectrometry and in situ Fourier transform infrared
spectroscopy with the density functional theory (DFT) calculations, the same group concluded that a
faster HCOOads oxidation step on the thicker Pd shell, induced by the lattice strain effect, might be
responsible for the more facile FAO [19].
In addition to the core shell or alloyed nanoparticles, PdAu catalyst can also be prepared by
depositing Pd thin films on Au surfaces. Hsu et al. deposited Pd thin films on hollow Au nanospheres
through a spontaneous galvanic replacement deposition [20]. Depending on the Pd film’s thickness,
the Pd-coated Au nanospheres showed enhanced FAO peak current density compared to Pd black,
which was ascribed to the electronic coupling between Au and Pd. Obradović and Gojković showed
film-thickness-dependent FAO activity on electrochemically deposited Pd films on polycrystalline Au
electrodes with the film thickness ranging from 1 to 17 monolayers (MLs) [21]. The maximum peak
current density was observed on four ML Pd, but the peak potential is most negative at one ML Pd. All
of the Pd films had a higher FAO activity than Pd black, which was attributed to the strain effect from
the Au substrate. In an earlier study of FAO on Pd films deposited on various metal surfaces with (111)
orientation, Kibler et al. demonstrated nicely the correlation between the d-band center position tuned
by the strain effect and the FAO catalytic activity [22]. In their work, Pd monolayer deposited on Au
(111) exhibited a higher peak current density than Pd (111), albeit at a higher potential.
One limitation in most of the previous studies of nanoparticles is that the particle surface structure
is not controlled and the observed Pd film thickness dependent activity can at least partly arise
from the surface structure difference. It has been well-demonstrated that formic acid oxidation is
structure-sensitive [7,23–26]. Therefore, controlling the exposed facets of nanoparticles is critical when
comparing catalytic activities, and can be used as a tool to improve catalytic performance. In this paper,
Pd thin films were deposited on cubic, octahedral, and rhombic dodecahedral (RD) Au nanocrystals
that were enclosed by six {100}, eight {111}, and twelve {110} facets, respectively, by a redox replacement
approach. The surface-limiting nature of this deposition method ensures the precise control of the
film thickness to a monolayer level [27], and pseudomorphic growth of Pd films on Au (111) surface
up to several tens of MLs has been demonstrated [28]. The use of nanocrystals with well-defined
surface structure minimizes the complication from the surface structural difference of the substrates
in film-thickness-dependent studies. Our focus of this study is to explore whether the Pd thin films
on Au nanocrystals mimic the electrochemical behaviors of the bulk Pd single crystals and how the
FAO activity varies with Pd film thickness. The results reveal that cyclic voltammetric features of
hydrogen adsorption/desorption, surface oxidation/reduction, and CO stripping on Au nanocrystals
covered with a monolayer of Pd resemble those from the corresponding bulk Pd single crystals in
0.1 M HClO4 . FAO on the monolayer Pd-coated Au nanocrystals showed peak shapes in the cyclic
voltammograms similar to those of bulk Pd single crystals and Pd film-covered bulk Au single crystals.
FAO and CO oxidation were also conducted on Au nanocrystals covered with Pd films with increasing
thickness up to 10 MLs. The results show that the FAO peak current density and CO oxidation onset
potential decreased with increasing Pd film thickness. These observations were explained with the

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Surfaces 2019, 2

strain and ligand effects from the d-band theory. This study demonstrates that depositing catalytically
active materials as a thin film on facet-controlled nanocrystals is an effective means of obtaining
high-performance catalysts because it combines the advantages of optimizing the d-band position by
the substrate and the facet-dependent catalytic activity of the overlayer.
2. Experimental Section
2.1. Materials
Hydrogen tetrachloroaurate(III) trihydrate (HAuCl4 ·3H2 O) (99.9%), silver nitrate (AgNO3 )
(≥99.0%), sodium borohydride (NaBH4 ) (≥99%), potassium bromide (KBr) (≥99.0%), cetyl
trimethylammonium bromide (CTAB) (≥98%), and cetylpyridinium chloride (CPC) (99.0%) were
purchased from Sigma-Aldrich. L-Ascorbic acid (AA) (99%) was obtained from Acros Organics. Formic
acid (88%) was obtained from Pharmco. Sodium hydroxide (98.5%), and hydrochloric acid (36.5–38%),
were received from Fisher Scientific. Double-distilled perchloric acid (70%) and sulfuric acid (95–98%,
ultra trace metal grade) were from GFS. Carbon monoxide (99.997%) was from Airgas. All of the
chemicals were used without further purification. Ultrapure water from a Milli-Q system (18.2 MΩ·cm)
was used in this work.
2.2. Synthesis of Au Seeds
The synthesis of 38 nm Au seed followed previously reported multi-step methods by others with
some modifications [29–32]. First, 125 μL of 10 mM HAuCl4 was added to 5 mL of 100 mM CTAB
solution at room temperature (22 ± 1 ◦ C), followed by the addition of 300 μL of 10 mM ice-cold NaBH4 .
The mixture was stirred for 5 min and stored at room temperature for further uses. Second, 24 μL
of the above solution was added to a mixture of 20 mL of 100 mM CTAB solution, 1 mL of 10 mM
HAuCl4 , 120 μL of 10 mM AgNO3 , and 160 μL of 100 mM ascorbic acid. The solution was kept at
room temperature for 2 h without stirring. The Au nanoparticles were centrifuged at 6500 rpm and
redispersed in 20 mL of 100 mM CTAB. Third, 1 mL of 10 mM HAuCl4 and 200 μL of 100 mM ascorbic
acid were added in sequence into the solution at 40 ◦ C. The solution was left undisturbed for 1 h
then centrifuged at 6500 rpm and redispersed in 20 mL of 10 mM CTAB solution. Lastly, 400 μL of
10 mM HAuCl4 was added into the solution and kept undisturbed at 40 ◦ C overnight. The obtained
Au nanoparticles were washed three times with a 100 mM CPC solution by centrifugation at 6500 rpm
and then redispersed in 30 mL of 100 mM CPC. The solution was marked as the Au seed solution for
further uses.
2.3. Synthesis of Au Nanocrystals
The shape-controlled synthesis of Au nanocrystals mostly followed a reported method [29]. In the
synthesis of cubic Au nanocrystals, 5 mL of 100 mM KBr solution, 1 mL of 10 mM HAuCl4 solution,
150 μL of 100 mM ascorbic acid solution, and 2 mL of the Au seed solution were added into 50 mL of
100 mM CPC solution in sequence. The solution was kept at room temperature for 2 h without stirring.
The obtained cubic Au nanocrystals were washed five times by centrifugation at 3500 rpm with warm
water and then redispersed in 1 mL ethanol. The synthesis of octahedral Au nanocrystals is the same
as that of cubes, except that no KBr was used. In the synthesis of RD Au nanocrystals, besides the fact
that no KBr was used, the volume of AA with the same concentration was increased to 2 mL and the
concentration of CPC was reduced to 10 mM with the same volume.

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2.4. Electrochemical Measurement
For a typical measurement, a whole batch of Au nanocrystals were further washed three times with
warm water and redispersed in 500 μL ethanol before electrochemical measurement. Ten microliters
(10 μL) of Au nanocrystals ethanol suspension was drop-coated on a glassy carbon (GC) disk electrode
and used as the working electrode. Cyclic voltammograms (CVs) and chronoamperograms (CAs) were
recorded in a conventional two-compartment three-electrode glass cell using a CHI 700C electrochemical
analyzer. A Pt wire served as the counter electrode and a KCl-saturated Ag/AgCl electrode was
used as the reference electrode. The counter electrode and the reference electrode were in the same
compartment that is separated from the working electrode compartment by a fine porous glass frit.
The cell resistance was compensated for with the iR compensation function in the analyzer. Before any
electrochemical measurements, the Au nanocrystals were subject to an electrochemical cleaning process,
which entails electrochemical potential cycling between −0.7 to +0.7 V versus Ag/AgCl for 50 cycles at
a scan rate of 0.5 V s−1 in 0.1 M NaOH. To prepare Pd-monolayer-coated Au nanocrystals, an atomic
layer of Cu was first deposited by underpotential deposition (UPD) in 0.1 M H2 SO4 + 1 mM CuSO4
solution at 0.05 V for 10 min. Then, the Cu-layer-coated Au nanocrystals were immersed in a deaerated
0.1 M HClO4 + 5 mM PdCl2 for 20 min to replace the atomic layer of Cu. Because of the surface-limiting
nature of the Cu UPD, nominally one ML of Pd is deposited after one redox replacement cycle [27,28].
By repeating this procedure, multi-monolayers of Pd were obtained. The electrochemical surface
area of Pd films was measured from CO stripping charges by assuming 320, 315, and 315 μC cm−2
for complete removal of a saturated CO adlayer on the Pd overlayer deposited on RD, cubic, and
octahedral Au nanocrystals [33], respectively. The CO adlayer was formed by purging the gas in the
solution for 5 min with the electrode potential being held at -0.10 V versus Ag/AgCl followed by 10 min
N2 purging to remove solution CO. The stripping voltammograms were recorded in 0.1 M HClO4 , and
the potential was first scanned negatively to -0.17 V and reversed. Typically, two additional cycles
between −0.17 and +0.98 V were recorded after CO oxidation, but for clarity, only the CO stripping
segment was shown.
2.5. Instrumentation
Scanning electron microscopy (SEM) images were obtained on a Zeiss Supra35 scanning electron
microscope operating at 5 kV. High-resolution transmission electron microscopy (HRTEM) and
selected-area electron diffraction (SAED) images were taken using a JEOL2100 transmission electron
microscope (JEOL USA Inc., Peabody, Massachusetts, USA) operating at 200 kV. X-ray diffraction (XRD)
measurements were performed on a Scintag X1 powder diffractometer (Scintag Inc., San Francisco,
CA, USA) with Cu Kα radiation (λ = 0.154 nm) operated at 40 kV and 25 mA by step scanning with a
step size of 0.02◦ /step. UV-Vis spectra were recorded by an Agilent 8453 spectrophotometer (Agilent,
Santa Clara, CA, USA).
3. Results and Discussion
3.1. Synthesis and Characterization of Au Nanocrystals
Cubic, octahedral, and RD gold nanocrystals enclosed by {100}, {111}, and {110} facets were
synthesized following a seed-mediated growth method developed by Niu et al. with some modifications
(see experimental section) [29]. The Au seeds used in this approach have a relatively large diameter
(38 ± 4 nm), as shown by the SEM image in Figure S1 in the Supplementary Material. Figure 1a–c
displays the SEM images of the cubic, octahedral, and RD Au nanocrystals, respectively. The shape
difference of Au nanocrystals is clearly observed. The edge length of the Au nanocubes is 61.5 ± 3.4 nm
(average length and standard deviation based on measurements from 100 particles). The TEM image
and square SAED pattern shown in Figure 1d confirm that these Au nanocubes are enclosed by
well-defined {100} facets [34,35]. The edge lengths of Au octahedra and RD are 63.2 ± 4.2 nm and
58.7 ± 2.8 nm, respectively. TEM images and SAED patterns of octahedral and RD Au nanocrystals are
15

Surfaces 2019, 2

presented in Figure 1e,f. The hexagonal dot array in the SAED pattern indicates that the octahedral
nanocrystal is enclosed by {111} facets [35]. The RD Au nanocrystals show the elongated hexagonal
dot array, which is also consistent with the reported diffraction pattern of Au single crystals along the
[011] zone axis [29,36,37]. Compared to those reported in [29], the surfaces of our RDs are smoother,
suggesting higher quality of the crystals.

Figure 1. SEM images of (a) cubic, (b) octahedral, and (c) rhombic dodecahedral (RD) Au nanocrystals.
Scale bar: 100 nm. SAED patterns of (d) cubic, (e) octahedral, and (f) RD Au nanocrystals. Inset:
corresponding TEM images, scale bar: 20 nm.

The structure of the nanocrystals was further examined with powder X-ray diffraction (XRD).
The cubic, octahedral, and RD Au nanocrystals show dominant (200), (111), and (220) peaks respectively,
confirming the single crystal structure of these Au nanocrystals (Figure 2). The XRD pattern for
RD Au nanocrystals showed a much more intense (220) peak compared to the previously reported
results by others [29], suggesting that the RD Au nanocrystals synthesized in the present work have
a much higher quality, which ensures more accurate activity comparison between different facets.
The structural improvement may arise from the lower synthesis temperature (22 ◦ C here versus 30 ◦ C
in [29]). Synthesizing the particles at room temperature simplifies the process by eliminating the use of
temperature control devices. In addition, we increased the amount of reactants by 10 times, which
produces a larger quantity of Au nanocrystals that are needed in the catalytic activity investigations.
The UV-Vis spectra in Figure S2 show shape-dependent surface plasmon resonance peaks.
The peak wavelengths of cubic, octahedral, and RD Au nanocrystals are 539 nm, 559 nm, and 598
nm, respectively, which are similar to those reported in ref. 29. The differences in UV-Vis spectra
mainly result from the differences in the shape of these Au nanocrystals. It is interesting to note
that the peak width of the RD is significantly broader than those of cubes and octahedra. The full
width at half maximum (FWHM) for RD is 110 nm, but only about 55 and 60 nm for the cubes and
octahedra, respectively.
One of the advantages of the seed-mediated synthesis is that the crystal size can be conveniently
tuned through varying the amount of the seed. To demonstrate this point, we decreased the volume of
Au seed solution to 100 μL in the growth step, and larger Au nanocrystals were obtained, as shown
in Figure S3. The facet edges of these larger nanocrystals are easier to see compared to the smaller

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ones. The size distributions of these larger cubic, octahedral, and RD Au nanocrystals are 112 ± 13 nm,
117 ± 12 nm, and 135 ± 8 nm, respectively.

Figure 2. XRD pattern of (a) cubic, (b) octahedral, and (c) RD gold nanocrystals.

Before electrochemical studies, the residual organic adsorbates must be removed from the Au
nanocrystal surfaces. This cleaning step is crucial for catalytic applications in order to obtain true
shape-dependent activity [38–40]. Several methods have been developed to clean nanocrystal surfaces.
However, the ozone removal method has been confirmed to cause severe surface structure distortion
due to its high oxidation power [41]. Although carbon monoxide (CO) has been used for effective
cleaning of surfactants on Pd nanocrystals [24,42], it does not adsorb strongly on Au surfaces [43,44].
An electrochemical fast potential cycling in a strong basic solution was used here for cleaning these Au
nanocrystals, which is similar to the surfactant removal methods reported by others [39,45]. In this
cleaning procedure, the as-prepared Au nanocrystals supported on a glassy carbon electrode were
scanned in the potential range of −0.70 to +0.70 V versus Ag/AgCl at 0.5 V s−1 in 0.1 M NaOH for about
50 cycles until stable CVs, such as those in Figure 3a, were obtained. Before the electrochemical potential
cycling treatment, the surface oxidation/reduction charge density was small and the surface oxidation
current rise was sluggish as exemplified by the CV from Au nanocubes in Figure 3a (dotted trace),
indicating that, even after several times of washing with water, some surfactants remained on the
crystal surfaces. After the potential cycling, the surface oxidation current increased sharply between
0.2 and 0.35 V, and the surface oxidation/reduction current peaks showed facet-dependent features.
The RD Au nanocrystals showed a surface oxidation peak at around 0.33 V, while on cubes it appeared
at 0.4 V and two peaks located at 0.38 and 0.44 V were observed on the octahedra. On the cathodic
scan, the surface oxide reduction peak of octahedra appeared at ca. 0.15 V, which is sharper and more
positive than those from the cubes and RD. The adsorption peak observed around 0.0 V on the bulk
Au (100) single crystal surface, which is a signature of surface reconstruction [46], was not seen on
the cube surfaces, suggesting that surface reconstruction does not occur on the nanocube surfaces.

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Surfaces 2019, 2

Facet-dependent voltammogram features were also observed in 0.1 M H2 SO4 (Figure 3b), albeit the
differences are more subtle than in the alkaline solutions. These voltammetric features are similar
to those observed on bulk Au (111), Au (100), and Au (110) single crystal surfaces [46–48], further
confirming that the nanocrystals are enclosed by the three low index facets.

Figure 3. Cyclic voltammograms (CVs) of cubic, octahedral, and RD Au nanocrystals in (a) 0.1 M
NaOH (b) 0.1 M H2 SO4 . The dotted line in (a) is the CV of cubic Au nanocrystals before the potential
cycling process. Scan rate: 0.100 V s−1 .

3.2. Electrochemical Properties of Pd-Monolayer-Coated Au Nanocrystals
As stated in the outset, a central task in this study is to examine whether Pd thin films coated on Au
nanocrystals have facet-dependent electrochemical properties, especially formic acid oxidation activity,
that mimic those of bulk Pd single crystals or Pd thin film coated bulk Au single crystals [24,33,49].
After electrochemical cleaning in 0.1 M NaOH, the Au nanocrystals were coated with a layer of Pd film
by the surface limited redox replacement approach, which nominally deposits one atomic layer of Pd
in one deposition-replacement cycle. Figure 4a shows the CVs of different Au nanocrystals coated with
a monolayer of Pd recorded in 0.1 M HClO4 . Different nanocrystals exhibited different electrochemical
features both in the hydrogen adsorption/desorption and Pd oxidation/reduction regions. The Pd
monolayer on cubic Au nanocrystals showed hydrogen adsorption/desorption peaks at around 0.0 V,
while these peaks appeared at around −0.10 V on RD nanocrystals. A clear shoulder at ca. 0.0 V on
the cathodic scan was also observed on the latter nanocrystals. These observations agree with those
reported on Pd cubic and RD nanocrystals, respectively [24]. On Pd-monolayer-coated octahedral
Au nanocrystals, the hydrogen adsorption/desorption peaks were more reversible than the other two
crystals, in accordance with that on bulk Pd (111) single crystal electrodes [33]. At potentials above
0.40 V, a pair of redox peaks appeared at around 0.5 V on Pd-monolayer-coated Au cubes and RDs,
which can be assigned to the surface oxidation/reduction of Pd on the basis of similar observations
on bulk Pd (100) and (110) surfaces [33]. Interestingly, on Pd-monolayer-coated Au octahedra, the
surface oxidation/reduction peaks are much smaller and less reversible, which also agree with the
results from bulk Pd (111) single crystals [33,50–52], and can be explained by the relatively inert nature
of the Pd {111} facet that has a much higher oxidation potential compared to the Pd {100} and Pd
{110} facets [33,49]. The CO stripping voltammograms from monolayer-Pd-coated Au nanocrystals
displayed in Figure 4b also show facet-dependent features. The CO oxidation peak potential for the
cubic shape nanocrystals is about 50 mV more negative than the other two nanocrystals, which has
been observed on Pd (100) bulk single crystal electrodes [33]. In addition to the main CO oxidation
peak at 0.75 V, the octahedral nanocrystals show an additional peak at 0.9 V, which has also been
observed on a bulk Pd (111) single crystal electrode [51]. The slanted voltammogram observed on
the octahedra is from the background current of the supporting GC electrode and only discernable
when the nanocrystal coverage is low. The similarities of the observed features in CVs and CO

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stripping voltammograms between Pd-monolayer-coated Au nanocrystals and the bulk single crystal
Pd electrodes/Pd nanocrystals suggest the epitaxial growth of a Pd monolayer on Au nanocrystals.

Figure 4. (a) CVs; (b) CO stripping voltammograms of cubic, octahedral, and RD Au nanocrystals
coated with one monolayer of Pd recorded in 0.1 M HClO4 . Scan rate: 0.100 V s−1 .

Figure 5a shows CVs of FAO on Pd-monolayer-coated Au nanocrystals obtained in 0.5 M HCOOH
and 0.1 M HClO4 . The FAO on these Pd monolayers exhibits shape-dependent features similar to the
bulk Pd single crystal electrodes [23] and corresponding Pd nanocrystals [24,26,53]. The oxidation
of formic acid on different Pd-coated Au nanocrystals started at around −0.10 V, but the monolayer
Pd deposited on cubic Au nanocrystals showed the highest oxidation peak current density at about
0.4 V, which is about double that of the other two nanocrystals. On the bulk electrodes, the formic
acid oxidation current density on Pd (100) is about 4 times of the other two low-index single crystal
electrodes in H2 SO4 [7], but about double that of Pd (111) in HClO4 [23]. In addition, the CVs of
Pd-monolayer-coated octahedral Au nanocrystals show similar behavior at the beginning of formic
acid oxidation with cubic nanocrystals, but it reaches its maximum current at around 0.3 V, which is
much more negative than that of cubic nanocrystals. The RD nanocrystals show a broader maximum
current density peak at the potential range of 0.2~0.6 V. These results are consistent with the reported
formic acid oxidation on different bulk Pd single crystal electrodes, in which Pd (111) showed a more
negative oxidation peak potential than the other two surfaces, Pd (100) had the highest peak current
density, and Pd (110) presented a broader peak [7,23]. It is interesting to note that CV features of FAO
resembling those from the bulk Pd single crystals were only observed after 2 MLs of Pd deposited on the
corresponding bulk Au single crystals [7]. The CVs of formic acid oxidation from these Pd-ML-coated
Au nanocrystals also resemble those obtained on the corresponding Pd nanocrystals. Thus, the current
density is the highest on cubic Pd nanocrystals and the peak potential is the most negative on Pd
octahedra [24,26,53]. The anodic and cathodic potential scans for these Pd-ML-coated Au nanocrystals
show nearly identical current densities, which signify the absence of CO poisoning on the time scale of
the CV measurements.
To further compare the catalytic activity of Pd monolayers on different Au nanocrystals, formic
acid oxidation current densities at 0.38 V (around the peak potential of the cubic particles) and 0.08 V
read from the CVs in Figure 5a were plotted in Figure 5c,d. The plots reveal potential dependent
catalytic activities of the Pd monolayers. At near peak potential (0.38 V), the Pd monolayer on cubic Au
nanocrystals has the highest catalytic activity, which is more than 2 times of the other two nanocrystals.
The Pd monolayer on RD nanocrystals is only slightly higher than that of octahedra. At the lower
potential (0.08 V), however, the catalytic activity of the Pd monolayer deposited on RD nanocrystals is
1.5 times of the octahedral and cubic nanocrystals (Figure 5d). Moreover, the Pd monolayer on RD
nanocrystals shows a better durability with 50% loss of the initial current density at 1000 s, in contrast
to 70% in cubic and 90% in octahedral nanocrystals, as shown by the chronoamperometric curves
(CAs) in Figure 5b. The higher durability on Pd-coated RD suggests that the surface poisoning is less
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Surfaces 2019, 2

prominent on this surface at this potential. The surface poisoning species cannot be identified based
solely on these results. It has been reported that, at this potential, adsorbed CO can be formed on
Pd black and Pd-coated Au electrodes through electroreduction of CO2 [21,54,55], which is a major
product of FAO. It is therefore conceivable that accumulation of adsorbed CO may be responsible for
the decay of the activity over time and the CO2 reduction is structure-sensitive.

Figure 5. (a). CVs of cubic, octahedral, and RD Au nanocrystals coated with one monolayer
of Pd recorded in 0.5 M HCOOH + 0.1 M HClO4 , scan rate: 0.100 V s−1 . (b). Corresponding
chronoamperometric curves recorded in the same solution at 0.08 V with an initial potential set at
−0.10 V. (c,d). Comparison of catalytic activities of Pd monolayer at (c) 0.38 V and (d) 0.08 V.

3.3. Thickness-Dependent Formic Acid Oxidation on Pd Thin Layer Coated Au Nanocrystals
The layer-by-layer growth nature of the redox replacement method for preparing the Pd films
allows for the examination of film-thickness-dependent catalytic activity. We examined FAO on Au
nanocrystals coated successively with up to 10 MLs of Pd. CO stripping voltammograms were recorded
after the formic acid oxidation experiments to assess the electrochemically active surface area and
examine Pd film thickness dependence of the CO oxidation. Figure 6 displays thickness-dependent
formic acid oxidation activity in terms of current density at 0.38 V in CVs. To facilitate the comparison,
the current density was normalized to that of the first layer. The error bars represent the standard
deviations of results from three separately prepared Pd films. For each nanocrystal, the highest activity
was observed at the first monolayer of Pd. The thickness-dependent activity of all of the nanocrystals
examined followed the same trend. As the number of Pd layers increased, the formic acid oxidation
activity decreased, and became largely unchanged after five monolayers of Pd. Kibler and Kolb did
not observe a clear trend of thickness-dependent FAO on Pd films coated on bulk Au single crystal
surfaces, but showed that the FAO current density on the first layer of Pd was in general much lower
than the thicker films [7]. Similarly, Obradović and Gojković showed that the first monolayer of Pd
film coated on a polycrystalline Au electrode had the lowest FAO activity, but after four monolayers of
Pd, the activity remained largely the same [21]. The distinction of the first layer of Pd in the activity
series between these previous studies and the present work may come from the different deposition
method employed and the use of H2 SO4 in their studies (vide infra).

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Surfaces 2019, 2

Figure 6. Normalized formic acid oxidation current density at 0.38 V as a function of the Pd film
thickness measured in 0.5 M HCOOH + 0.1 M HClO4 . The error bars represent the standard deviations
of results from three separately prepared Pd films.

Two effects are likely contribute to the thickness-dependent catalytic activity of these Pd overlayers
on Au nanocrystals. First, the lattice constant of Au is 5% larger than that of Pd. Pseudomorphic growth
of Pd on Au results in an expansion of Pd lattice, which increases the d-band center and therefore
the adsorption energy of adsorbates [56]. This is the so-called strain effect. Second, the electronic
interaction between two different metals results in electronic effect (the so-called ligand effect), which
will cause the energy renormalization of the d-band [57]. A stronger interaction between the overlayer
and the substrate atoms will result in an upshift of the d-band center of the overlayer [58]. Decoupling
the two effects experimentally is challenging. In density functional theory (DFT) calculations, the two
effects could be separated by artificially enlarging the lattice constant of the Pd top layer to account
for the tensile strain while changing the underneath substrate from Au to Pd to eliminate the ligand
effect [19,59]. All of the calculations show that the tensile strain elevates the d-band center and plays
a dominant role [19,56,59,60], but there is no consensus on the ligand effect. Roudgar and Groß
predicted that the ligand effect further increases the d-band center position [59], while others showed
the opposite [19]. Nevertheless, the net result is an upshifted d-band center for Pd deposited on Au.
As the number of Pd layers on Au nanocrystals increases, the lattice constant gradually relaxes back
to the value of pure Pd, which has been demonstrated by both scanning tunneling microscopy and
surface X-ray scattering [61–64]. The lattice relaxation releases the strain effect. Meanwhile, as the film
thickness increases, the electronic interaction between the Pd top layer and the Au substrate becomes
weaker and therefore the ligand effect diminishes. Our formic acid oxidation results are in general
agreement with these arguments.
The thickness-dependent catalytic activity is also manifested in CO oxidation. Figure 7 displays
CO stripping onset potential as a function of Pd overlayer thickness obtained in 0.1 M HClO4 . The onset
potential here was taken as the potential where CO oxidation current is at 10% of the peak current. In the
case where there are multiple peaks, the most negative peak was used. All of the three nanocrystals
show a negative potential shift of CO oxidation onset potential when more than one ML of Pd was
deposited. After three layers of Pd, the onset potential of CO oxidation remains largely the same.
The RD nanocrystals coated with two Pd MLs showed a smaller negative shift compared with the
other two Au nanocrystals. The trend of CO oxidation onset potential shift is consistent with catalytic
activities shown in Figure 4 and can again be explained by the diminishing strain and ligand effects
with increasing film thickness. The thickness-dependent film growth mode may be revealed from the
CO stripping voltammograms (Figure S4). As the Pd film thickness increases, one obvious feature

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of CO oxidation is the appearance of a second oxidation peak at around 0.9 V. This second oxidation
peak was observed on octahedral nanocrystals when only one ML of Pd was deposited, but after the
deposition of five MLs of Pd it can be clearly seen in all three nanocrystals, as shown in Figure S4.
The second CO oxidation peak does not exist on the bulk Pd (100) and Pd (110) electrodes [50,52].
This peak might come from the defect sites formed through three-dimensional (3D) growth on the
thicker Pd films, which may play a minor role in the thickness-dependent FAO activity given its smaller
magnitude on the cubes and RDs.

Figure 7. CO oxidation onset potential as a function of the Pd film thickness on different Au nanocrystals
in 0.1 M HClO4 .

It has been demonstrated that formic acid oxidation rates are higher in HClO4 than in
H2 SO4 [7,23,26,65], because the strong adsorption of (bi)sulfate anion on Pd surfaces above 0.2 V
(versus RHE) [66,67] hinders the formic acid oxidation in H2 SO4 . It is therefore interesting to examine
whether the thickness-dependent adsorption will manifest itself in (bi)sulfate adsorption and hence
influences formic acid oxidation. Because of their higher catalytic activity of formic acid oxidation,
Pd-film-coated cubic Au nanocrystals were chosen for this study. Figure 8 displays the formic acid
oxidation current density at 0.38 V as a function of film thickness in 0.1 M HClO4 or 0.1 M H2 SO4 with
0.5 M formic acid. As expected, formic acid oxidation is more facile in 0.1 M HClO4 than in 0.1 M
H2 SO4 due to the adsorption of (bi)sulfate. Interestingly, the activity disparities in the two electrolytes
gradually decreased as the film thickness increased. On the first monolayer of Pd, the current density
of FAO in 0.1 M HClO4 is 1.6 times of that in 0.1 M H2 SO4 , but it decreased to only 1.2 times on
10 monolayers of Pd. This observation can be explained by thickness-dependent (bi)sulfate adsorption
on Pd film surfaces. As discussed above, the d-band center upshift is the highest on the first monolayer
of Pd deposited on Au nanocubes. As a result, the adsorption energy of (bi)sulfate is the strongest.
The strong adsorption of (bi)sulfate anions competes with formate adsorption, and results in a lower
catalytic activity compared to the nonspecifically adsorbed supporting electrolytes, such as HClO4 .
This effect decreases as the thickness of the Pd film increases because the d-band center gradually
returns to the value of bulk Pd. While the d-band upshift also increases the adsorption energy of
formate, the effect is likely stronger on the (bi)sulfate adsorption. Given that (bi)sulfate adsorption
even competes with CO adsorption on Pd [33], this assertion is not unreasonable.

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Figure 8. Formic acid oxidation current density at 0.38 V in 0.5 M HCOOH + 0.1 M HClO4 or H2 SO4
from cubic Au nanocrystals coated with Pd films of different thicknesses.

4. Conclusions
In summary, Pd-monolayer-covered high-quality Au nanocrystals showed facet-dependent formic
acid and carbon monoxide oxidation activities that resemble those observed on bulk Pd single crystal
surfaces and Pd thin film covered bulk Au single crystal surfaces, suggesting pseudomorphic growth
of Pd MLs on these Au nanocrystals. As the Pd film thickness increased, the formic acid oxidation
peak current density decreased and the decrement leveled off after about five MLs of Pd. Similarly,
the CO oxidation onset potential also gradually decreased with the increase of Pd film thickness and
remained largely unchanged after three MLs of Pd. The effects of bi(sulfate) specific adsorption on
FAO also decreased with the increasing Pd film thickness. These observations are explained in terms
of the d-band center shift with the Pd film thickness as a result of the tensile strain effect and the
ligand effect. This study demonstrates that, by combining the facet dependence of catalytic activity
with the fine-tuning of d-band position, deposition of catalytically active materials as a thin film on
facet-controlled nanocrystals can be an effective approach for engineering high-performance catalysts
for a variety of reactions.
Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/2/2/27/s1,
Figure S1: SEM image of Au seed; Figure S2: UV-Vis spectra of Au nanocrystals; Figure S3: SEM images of larger
Au nanocrystals, and Figure S4: CO stripping voltammograms of Au nanocrystals coated with different layers of
Pd films.
Author Contributions: Conceptualization, S.Z.; Data curation, Y.T.; Formal analysis, Y.T. and S.Z.; Funding acquisition,
S.Z.; Methodology, Y.T.; Project administration, S.Z.; Supervision, S.Z.; Validation, Y.T.; Writing—original draft, Y.T.;
Writing—review & editing, S.Z.
Acknowledgments: This work was partially supported by the US National Science Foundation through
CHE-1156425 and CHE-1559670.
Conflicts of Interest: The authors declare no conflicts of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Sputtered Platinum Thin-films for Oxygen Reduction
in Gas Diffusion Electrodes: A Model System for
Studies under Realistic Reaction Conditions
Gustav W. Sievers 1,2,† , Anders W. Jensen 1,† , Volker Brüser 2 , Matthias Arenz 3, * and
María Escudero-Escribano 1, *
1
2
3

*
†

Nano-Science Centre, Department of Chemistry, University of Copenhagen, Universitetsparken 5, 2100
Copenhagen Ø, Denmark; sievers@inp-greifswald.de (G.W.S.); awj@chem.ku.dk (A.W.J.)
Leibniz Institute for Plasma Science and Technology, Felix-Hausdorff-Strasse 2, 17489 Greifswald, Germany;
brueser@inp-greifswald.de
Department of Chemistry and Biochemistry, University of Bern, 3006 Bern, Switzerland
Correspondence: matthias.arenz@dcb.unibe.ch (M.A.); maria.escudero@chem.ku.dk (M.E.-E.);
Tel.: +41-31631-5384 (M.A.); +45-35-32-83-90 (M.E.-E.)
These authors contributed equally to this work.

Received: 20 March 2019; Accepted: 24 April 2019; Published: 28 April 2019

Abstract: The development of catalysts for the oxygen reduction reaction in low-temperature fuel cells
depends on efficient and accurate electrochemical characterization methods. Currently, two primary
techniques exist: rotating disk electrode (RDE) measurements in half-cells with liquid electrolyte
and single cell tests with membrane electrode assemblies (MEAs). While the RDE technique allows
for rapid catalyst benchmarking, it is limited to electrode potentials far from operating fuel cells.
On the other hand, MEAs can provide direct performance data at realistic conditions but require
specialized equipment and large quantities of catalyst, making them less ideal for early-stage
development. Using sputtered platinum thin-film electrodes, we show that gas diffusion electrode
(GDE) half-cells can be used as an intermediate platform for rapid benchmarking at fuel-cell relevant
current densities (~1 A cm−2 ). Furthermore, we demonstrate how different parameters (loading,
electrolyte concentration, humidification, and Nafion membrane) influence the performance of
unsupported platinum catalysts. The specific activity could be measured independent of the applied
loading at potentials down to 0.80 VRHE reaching a value of 0.72 mA cm−2 at 0.9 VRHE in the GDE.
By comparison with RDE measurements and Pt/C measurements, we establish the importance of
catalyst characterization under realistic reaction conditions.
Keywords: electrocatalysis; oxygen reduction; ORR; gas diffusion electrode; platinum; fuel cells;
thin-films; benchmarking; mass transport

1. Introduction
Fuel cell technologies, which convert chemical energy directly into clean electricity, are expected to
play a key role in environmentally friendly energy conversion schemes [1]. Polymer exchange membrane
fuel cells (PEMFCs) are very promising both for transportation and stationary applications [2,3].
Platinum-based catalysts typically catalyze both the oxidation of hydrogen at the anode and the
reduction of oxygen at the cathode of PEMFCs; as a result, the overall cost of the PEMFC technology is
closely linked with the platinum loading [4]. The slow kinetics of the oxygen reduction reaction (ORR)
causes significant losses in cell potential. Consequently, the loading at the cathode accounts for the
majority of platinum usage in PEMFC. In order to reduce the Pt loading, most scientific and industrial
research focuses on developing new ORR catalysts that present both high activity and stability for the
ORR [5–7].
Surfaces 2019, 2, 336–348; doi:10.3390/surfaces2020025

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Surfaces 2019, 2

The ORR activity can be enhanced by modifying the geometric [8] or electronic structure [9] of
the Pt-based electrocatalytic surfaces. To improve the catalyst mass activity, we can enhance: i) the
electrochemically active surface area (ECSA) [10] and/or ii) the specific catalytic activity by modifying
the electronic properties of platinum by alloying [11–13]. As a result, numerous catalyst concepts have
been developed with activities far exceeding commercial Pt catalysts using thin-film rotating disk
electrode (RDE) measurements in liquid half cells [10]. However, the impressive activity enhancements
obtained with RDE have not been translated to real devices [14]. To the best of our knowledge, only a
single concept based on dealloyed bimetallic Pt catalysts have exceeded the Department of Energy
(DOE) target of activity normalized by mass of precious group metal (PGM) of 0.44 A/mgPGM at 0.9
VRHE in a membrane electrode assembly (MEA) [15].
Benchmarking Pt-based catalysts using the RDE method allows a fast investigation of the trends in
electrocatalytic activity of nanoparticles supported on high surface area carbon materials and sputtered
thin films [16–18]. However, there are many differences between RDE and MEA measurements. These
include different catalyst loadings; in RDE, typical 5–20 μgPt cm−2 are used; on the other hand, MEAs
typical employs loadings in the range of 0.1-0.5 mgPt cm−2 [19]. Secondly, different testing parameters
are utilized. In RDE measurements, the potential is typically swept with a rate between 5 and 100 mV
s−1 at room temperature [17]. In contrast, MEA testing is generally carried-out at elevated temperatures
with constant currents or potentials [20]. A third difference is the reaction environment at the point
of the active site. In MEAs, the oxygen reduction occurs at the triple phase boundary where O2 gas,
proton transport mechanisms and electrical conductive particles coexist [21]. Conversely, the catalytic
layer in the RDE configuration is immersed in an oxygen-saturated aqueous electrolyte, enabling
superior access of oxygen, protons and electrons [22].
Another drawback of RDE benchmarking is the low diffusion and solubility of molecular oxygen
in aqueous electrolytes. Consequently, the rate of the reaction is under full mass transport limitation
at current densities orders of magnitude lower than actual fuel cells. Using the RDE technique, it is
therefore, only possible to measure activities at low overpotentials (typically 0.9 V versus the reversible
hydrogen electrode, VRHE ), far from the working potential in real fuel cells (0.6–0.8 VRHE ). While trends
in ORR rate can be investigated at wider potential ranges with use of the Koutecky-Levich analysis [23],
it is associated with great uncertainties in weakly adsorbing electrolytes. In MEAs, the catalyst layers
are applied on porous gas diffusion electrodes in order to enable diffusion of O2 from the gas phase.
However, MEA benchmarking requires large quantities of catalyst and specialized equipment; as a
result, it is less accessible for early stage catalysts development.
Since Zatilis et al. showed that superior, near mass transport free current could be achieved
with a floating electrode configuration using an artificial gas diffusion interface [24], several gas
diffusion electrode (GDE) half-cells designs have been reported for ORR benchmarking [25–28].
These works include measurements in concentrated phosphoric acid at elevated temperature [25],
methanol and ethanol oxidation [26] and Pt/C in HClO4 showing comparable activity to MEA, both
at room temperature [27] and at elevated temperatures using a Nafion separating membrane [28].
Furthermore, different GDE setups have recently been used for both CO [29] and CO2 reduction [30–32].
In this study, we utilize a GDE setup that enables testing commercial gas diffusion layers
with an interchangeable Nafion membrane resembling actual PEMFCs configuration [28]. Using
sputtered platinum as a model catalyst, we demonstrate the influence of different testing parameters
(electrolyte, scan rate, membrane and humidification) on the performance of the unsupported catalysts.
By investigating the effect of different catalyst loadings in both GDE and RDE, we identify key design
principles and establish the importance of benchmarking under realistic reaction conditions.

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Surfaces 2019, 2

2. Materials and Methods
2.1. Preparation of Pt Thin-Film Electrodes
For the Pt thin film deposition, the magnetron electrode was equipped with a planar target of
99.95 % Pt (Junker Edelmetalle, Waldbüttelbrunn, Germany). It was located at the superior part of the
recipient. The RF generator (Advanced Energy, Fort Collins, CO, USA) had a driving frequency of
13.56 MHz. The reactor chamber was evacuated to a base pressure of 5 × 10−3 Pa. An argon plasma
was ignited in the chamber at a working pressure of 5 Pa. The platinum loadings of the sputtered
Pt TFs were estimated by inductively coupled plasma mass spectrometry (ICP-MS) measurements
(Aurora Elite, Bruker, Billerica, MA, USA). For this purpose, the catalysts were digested in aqua regia
freshly mixed with 30% HCl (Suprapur, Merck, Darmstadt, Germany) and 65% HNO3 (Suprapur,
Merck, Darmstadt, Germany) in a volumetric ratio of 3:1, respectively. A Dektak 3ST profilometer
(Veeco, Plainview, NY, USA) was used to measure the film thickness. Corresponding to the Pt loadings
of 11.7 μgPt cm−2 , 16.2 μgPt cm−2 , 29.4 μgPt cm−2 and 66.2 μgPt cm−2 , the films had a thickness of
15 nm, 20 nm, 37 nm and 83 nm, respectively.
2.2. Physical Characterization
The morphology of the sputtered Pt thin-film was investigated with scanning electron microscopy
(JSM 7500F, JEOL, Tokyo, Japan) with a field-emission gun, a semi-in-lens conical objective lens and a
secondary electron in-lens detector for high-resolution and high-quality image observation of structural
features of the deposited films at a maximum specified resolution of 1.0 nm at 15 keV. The technique
enables imaging the surface without any preparative coatings. The GDE was directly placed onto the
SEM unit holder. The phase identity and crystallite size of the Pt catalyst was investigated by means of
x-ray diffraction (XRD) using a Bruker D8 Advance Diffractometer (Bruker, Billerica, MA, USA), with
measurements performed over a 2Theta range from 20◦ to 80◦ , step width 0.05◦ and 5 s per step. Cu
was used as the anode target.
2.3. Chemicals, Materials and Gases
Deionized ultrapure water (resistivity >18.2 MΩ·cm, total organic carbon (TOC) < 4 ppb) from a
Milli-Q system (Millipore, Burlington, MA, USA) was used for acid dilutions, catalyst ink formulation,
and the GDE cell cleaning. Perchloric acid (70% HClO4, Suprapur, Merck) were used for electrolyte
preparation. The Nafion membrane (Nafion 117, 183 μm thick, Chemours, Wilmington, DE, USA)
and the gas diffusion layer (GDL) with a microporous layer (MPL) (H23 C2, Freudenberg, Weinheim,
Germany) were employed in the GDE cell measurements. The following gases (Air Liquide, Taastrup,
Denmark) were used in electrochemical measurements: Ar (99.999%), O2 (99.999%), and CO (99.97%).
2.4. Gas Diffusion Electrode (GDE)
Electrochemical measurements were carried out using a GDE setup [28]. In contrast to RDE,
benchmarking is limited to low overpotentials, the GDE setup allow testing of ORR activity at high
current densities under realistic mass transport conditions. The GDE half-cell consists of two cell
components (see Figure 1): (i) a lower component made of body of stainless steel with a flow field
and gas supply (ii) an upper cell body of polytetrafluoroethylene (PTFE) with electrolyte and counter
electrode (platinum mesh) and reference electrode (reversible hydrogen electrode, RHE). For cleaning,
the Teflon upper part was soaked in mixed acid (H2 SO4 :HNO3 = 1:1, v:v) overnight. Subsequently, it
was rinsed thoroughly by ultrapure water, and boiled three times.

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Surfaces 2019, 2

Figure 1. Schematic illustration of the gas diffusion electrode (GDE) half-cell setup, showing the catalyst
layer and membrane sandwiched between the upper PTFE with electrolyte and the lower stainless
steel body with the gas flow field.

The geometric surface area of the working electrode was 3 mm. The GDE was used with and
without membrane between electrolyte and working electrode. When the membrane was used, it was
pressed during assembling of the cell to the catalyst layer. For measurements using humidification, the
gas was bubbled through a gas humidifier.
2.5. Electrochemical Measurements
All electrochemical measurements were performed using a computer-controlled potentiostat (ECi
200, NordicElectrochemistry, Copenhagen, Denmark). The measurements were performed using 4 M
HClO4 and 1 M HClO4 aqueous solutions at room temperature. The high electrolyte concentrations
reduce the solution resistance between working electrode, counter electrode and reference electrode.
Prior to the measurements, the working electrode was purged from the backside (through the gas
diffusion layer) with O2 gas and the catalyst was conditioned by potential cycles between 0.1 and 1.0
VRHE at a scan rate of 100 mV s−1 until a stable cyclic voltammogram could be observed (ca. 30 cycles).
The ECSA of the catalyst was determined by conducting CO stripping voltammetry. The working
electrode was held at 0.05 V during a CO purge through the GDL for 2 min followed by an Ar purge
for 10 min. The ECSA was determined from the CO (QCO ) oxidation charge recorded at a scan rate of
50 mV s−1 and the respective Pt loading (LPt ) using a fixed conversion coefficient of 390 μC cm−2 Pt [16];
ECSA = QCO /(LPt × 390 μC cm−2 Pt ). To determine the ORR activity, linear sweep voltammetry (LSV,
anodic scan) was conducted by purging the electrode with O2 from below and scanning the potential
at a scan rate of 50 mV s−1 or 100 mV s−1 . The polarization curves were corrected for the non-faradaic
background by subtracting the cyclic voltammograms (CVs) recorded in Ar-purged electrolyte at the
identical scan rate. Furthermore, the resistance between the working and reference electrode (~10 Ω)
was determined using an AC signal (5 kHz, 5 mV) and thereafter compensated for using analogue
positive feedback scheme of the potentiostat. The resulting effective resistance was 1 Ω or less for
each experiment.
3. Results and Discussion
3.1. Structural Characterization of Pt Thin-Films
Representative SEM measurements of Pt thin films (TF) are displayed in Figure 2a–c. It can be seen
that small Pt domains with a size of ca. 5 nm agglomerate to form cauliflower-like structures with a size
between 50 nm and 150 nm. The cauliflower morphology is related to a fractal structure. The resulting
surface activity and surface area is determined by this structure. The cross-section in Figure 2b shows
30

Surfaces 2019, 2

the structure of the top-surface layer to be porous and the attaching part to the microporous layer of
the gas diffusion electrode to be denser. XRD of Pt TF shows broad diffraction peaks of the indices
(111), (200) and (220) crystal structure (see Figure 2d). The average crystallite size estimated by the
Scherrer equation is 7.2 nm, which fits to the Pt domain size identified with the electron microscope.

Figure 2. (a–c) SEM micrographs of ~500 nm Pt thin films (TF) on a gas diffusion electrode (GDE) with
microporous layer (MPL) H23 C2 from Freudenberg prepared by RF unbalanced magnetron sputtering
at 15 W and a working pressure of 5 Pa: top view (a), cross-section view (b) high-magnification top-view
(c). (d) XRD measurements of Pt TF deposited onto a gas diffusion electrode (Freudenberg, H23 C2)
with Pt (111), (200) and (220) diffraction pattern (PDF 00-001-1194).

3.2. Electrochemical Characterization
Figure 3 shows the base cyclic voltammograms in argon-saturated electrolyte as well as
CO-stripping measurements in order to determine the electrochemically active surface area. Figure 3a
displays in detail the cyclic voltammograms of Pt TF with increasing Pt characteristic features for
higher loadings, as observed in the hydrogen under potential deposition (0.1–0.4 VRHE ) region as well
as the region where O-containing species are adsorbed (0.7–1.0VRHE ). The linear sweep voltammogram
of the CO oxidation charge transfer is shown in Figure 3b for the different loadings. The sweep shows
two CO oxidation peaks, one peak between 0.75 and 0.77 VRHE and a second around 0.82–0.85 VRHE , in
contrast to the CO oxidation sweep measured in the RDE with one peak at 0.73 VRHE . The integrated
CO oxidation charge is plotted against the Pt loading with the ideal linear slope crossing 0 and the
second lowest loading; see Figure 3c. The oxidation charge increases with loading. However, it can be
seen that the CO oxidation charge as function of the loading deviates from an ideal straight line for the
highest loading. This could indicate that there is a mass transport problem due to blocking of CO as a
reactant or trapped product; however, no CO oxidation peaks were visible in sub sequential sweeps.
Thus it can be assumed that the ECSA measurement with CO adsorption is valid.

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Surfaces 2019, 2

Figure 3. (a) Loading dependent electrochemical behaviour of Pt TF in GDE setup (298 K, 4 M HClO4 ,
0% RH), cyclic voltammograms recorded with 50 mV s-1 and Ar saturated flow field. (b) CO stripping
curves recorded with 50 mV s-1 . (c) Plot of CO oxidation charge as function of the Pt loading on the
GDE and RDE. The linear slope indicates the ideal correlation between oxidation charge and Pt loading.

Our results from CO-stripping experiments therefore, show that the decrease in ECSA for higher Pt
loadings is a result of decreasing Pt utilization. This is supported by the cross-section SEM micrograph
displayed in Figure 2b with a higher loading. The graph shows a higher porosity at the near surface
region, compared to the region near the Microporous Layer (MPL). Increasing the loading/film thickness
leads to lower Pt utilization due to lack of long-range porosity for higher loadings of sputtered Pt
thin-films. Consequently, the ECSA measurements have to be carefully done for each loading, especially
with loadings above 30 μgPt cm−2 . The ECSA of the Pt TF GDE is determined to be 20 m2 g−1 Pt for the
low-loaded 16.2 μgPt cm−2 gas diffusion electrodes.
3.3. Performance of Pt Thin-Films in the GDE Setup
Figure 4 shows a comparison between the linear sweep voltammograms during the oxygen
reduction reaction for both RDE and GDE. As stated in our previous work [28], the GDE enables the
measurement of realistic mass transport conditions and high current densities. In contrast, the RDE is
mass transport limited due to low solubility of oxygen in the condense phase/liquid electrolyte, but it’s
well defined hydrodynamics allow extraction of intrinsic kinetic currents [16,17]. Therefore, it is only
possible to measure the kinetic current densities at very low overpotentials.
The ORR activity from RDE measurements in acidic electrolyte is typically reported at 0.9 VRHE
for less active platinum catalysts and for highly active catalysts it is often only possible to measure
at even higher potentials up to 0.95 VRHE [33]. However, fuel cells typically operate at a working
potential of in the range of 0.6–0.8 VRHE , as indicated in Figure 4. At these potentials, the current in the

32

Surfaces 2019, 2

RDE is under full mass transport limitation. While in the GDE the current density is not free of mass
transport limitations, GDE measurements reveal more realistic current densities [28].
Our GDE setup allows measuring a maximum oxygen reduction current density of ~1 A cm−2
using only 16.2 μgPt cm−2 ; this is over two orders of magnitude higher than what is observed in the
RDE measurements. The maximum current density is reached at ~0.3 VRHE . At potentials below
0.3 VRHE in the GDE it can be seen that the current decreases. This correlates to the region where
hydrogen under potential adsorption occurs. Therefore, it is presumed that the decrease in the current
density observed between 0.1 VRHE and 0.3 VRHE is due to blocking by protons on the Pt surface which
would limit the full oxygen reduction to water. In this region, the catalyst reduces oxygen to hydrogen
peroxide, which corresponds to a two-electron transfer compared to a four-electron transfer for the
reduction of oxygen to water [34]. It is notable that the maximum current density is reached in the
same potential region as the maximum in the extended Koutecky-Levich analysis [23] as well as the
potential of zero total charge (PZTC) [35]. The maximum oxygen reduction rate, therefore, correlates
with the minimum ion coverage; most likely a competition between protons and oxygen species at the
surface and, to a lesser extent, perchlorate anions.

Figure 4. Comparison of the current density-potential curves as measured in a GDE setup and a
conventional RDE setup. The activity measurement in the GDE setup with oxygen saturated flow field
was conducted at 298 K, 4 M HClO4 , Nafion 117 membrane, 0% RH and 50 mV s−1 . The measurement
of the activity in the RDE setup was conducted at 1600 RPM, 298 K, 0.1 M HClO4 and 50 mV s−1 .
The catalyst was a Pt thin film prepared by magnetron sputtering with a Pt loading of 16.2 μgPt cm−2
for both RDE and GDE. The right axis shows the current density enhancement as compared to the
diffusion limited current observed in RDE.

3.4. Benchmarking Different Pt Loadings in the GDE Setup
The Pt TF catalyst was tested with different Pt loadings on the GDE, as shown in Figure 5.
The obtained linear sweep voltammograms are plotted as kinetic current densities normalized by the
catalyst geometric surface area, the specific activities (normalized by the roughness factor) and mass
activities. The activities normalized by geometric surface area show the expected increase of current
density over the full potential range. At 0.3 VRHE , the maximum activity is achieved and reaches 0.75 A
cm−2 for 11 μgPt cm−2 and nearly 1 A cm−2 for the highest loading 66 μgPt cm−2 . The onset potential for
the ORR is shifting to lower overpotentials. The specific activities were calculated using the roughness
factor determined by CO oxidation; it can be seen that the onset potential stays constant with a value
of 0.71 mA cm−2 Pt at 0.90 VRHE . The specific activity starts to differ in the region from 0.75 VRHE –0.8
VRHE for the different loadings. Furthermore, the lowest Pt loading leads to the highest specific activity
measured. The Pt TF catalyst with 11 μgPt cm−2 achieves a specific activity of 94.0 mA cm−2 Pt at 0.65
33

Surfaces 2019, 2

VRHE . While the GDE with the highest loading of 66 μgPt cm−2 reaches only a value of 38.6 mA cmPt −2
at 0.65 VRHE , corresponding to a “loss” of 59% at 0.65 VRHE . The reason for the potential-dependent
specific activity is most likely related to mass transport, which can be distinguished due to oxygen
transport, proton transport, H2 O transport or electronic conductivity. In order to distinguish between
the different influences, several operation parameters were changed and discussed later in this paper.

Figure 5. Influence of different Pt loadings on the catalyst layer activity (11.7 μgPt cm−2 , 16.2 μgPt cm−2 ,
29.4 μgPt cm−2 , 66.4 μgPt cm−2 ). (a) Activity normalized by the geometric surface area, (b) surface-area
specific activity and (c) mass specific activity. The measurement of the was conducted in the GDE setup
with oxygen saturated flow field (298 K, 4 M HClO4 , 0% RH, 100 mV s−1 ).

Comparing the mass activity for the different loadings, the highest activity can again be achieved
with the GDE with the lowest loading of 11 μgPt cm−2 . The activity of the GDE with 11 μgPt cm−2
reaches 17 A mg−1 Pt at 0.65 VRHE and 55 A mg−1 Pt at 0.3 VRHE . The highest loaded GDE achieves only
a mass activity of 5 A mg−1 Pt at 0.65 VRHE and 19 A mg−1 Pt at 0.3 VRHE . At 0.65 VRHE this is a loss of
75% and at 0.3 VRHE the “loss” is 70% of the mass activity compared to the lowest loading.
3.5. Systematic Change of Operation Parameters
The Pt TF GDE was evaluated with different operational parameters to understand the activity
determining factors of the unsupported catalyst layers (see Figure 6). First, Figure 6a shows the ORR
polarization curves of the GDE with a separating membrane, as compared with a non-separated GDE
setup with 4 M perchloric acid. The separating membrane increases the geometric activity over the full
potential window and at 0.65 VRHE by 104%. To clarify the cause of this drastic increase in activity, we
performed a set of measurements with 1 M HClO4 instead of 4 M HClO4 (see Figure 6a). Reducing the
molarity of the electrolyte greatly enhanced the activity. Interestingly, the polarization curve for Pt
in 1 M HClO4 resembles the one obtained in 4 M HClO4 using a Nafion membrane. This suggests
that specific anion adsorption is inhibiting the rate of ORR at high electrolyte concentrations. While
perchloric acid is generally considered a non-adsorbing electrolyte, it is known to inhibited ORR

34

Surfaces 2019, 2

on Pt at high concentrations [36]. Thus, in the absence of a separating membrane, an electrolyte
concentration of 1 M HClO4 would be preferable. While no differences were observed in the pure
kinetic region (until ~0.8 VRHE ), the membrane slightly enhanced the activity as compared to 1 M
HClO4 in the mass-transport-dominated region (0.3–0.8 VRHE ), possibly due to enhanced proton
transport or reduced flooding of the hydrophilic catalyst layer. At high current densities, oxygen
transport is critical. Recently, Kongkanand el. al showed that local O2 resistance in the ionomer film
is essential for the performance supported catalysts [4]. While unsupported catalysts are generally
free of ionomer, the degree of wetting would influence the oxygen transport in the catalyst layer. It is,
therefore, possible that the membrane reduces flooding as a result of less penetrating electrolyte, as
compared to measurements without the Nafion membrane. However, from the presented data, it
cannot be concluded if adsorbed water on the surface hinders the reaction.

Figure 6. Influence of: (a) the separating Nafion 117 membrane, (b) the humidification on the activity
of Pt TF 0% vs 100% relative humidity (RH), (c) the scan rate change between 50 mV s−1 and 100 mV
s−1 . The measurement of the activity was conducted in the GDE setup with oxygen saturated flow field
(16.2 μgPt cm−2 , 298 K, 4 M HClO4 , 0% RH, 50 mV s−1 ). (d) Influence of catalyst system on the catalyst
layer performance measured in the GDE with Nafion 117 as membrane in 4 M HClO4 at RT at 50 mV
s−1 with Pt TF (16.2 μgPt cm−2 , in red) and Pt/C (10 μgPt cm−2 , in black) [28].

Secondly, the scan rate was varied to 50 mV s−1 compared to the 100 mV s−1 , the results showed
negligible decrease in activity (see Figure 6b). This indicates that the current densities are defined
by faradaic currents instead of capacitive currents. Capacitive currents increase as the current is
proportional to the applied scan rate. Furthermore, it is a sign for a reaction in equilibrium, where
mass transport should not be the significant driver limiting the reaction. If mass transport would limit
the reaction rate, as the reactant supply drops, the geometric current density should decrease with
lower scan rate.
The third operation parameter studied was the humidification of oxygen. It was previously
shown that a change of oxygen humidification can have a significant impact on the gained activity
of platinum nanoparticles supported on carbon (Pt/C) [28]. However, there is nearly no change in
the activity of Pt TF after humidifying the oxygen with 100% relative humidity (RH); see Figure 6c.
35

Surfaces 2019, 2

This indicates that proton transport is not a limiting step with unsupported Pt catalysts. This is
in contrast to Pt/C in GDE, where we observed a doubling in activity with humidification at room
temperature, indicating a limiting behaviour; see Figure 6d. In our previous study, Pt/C catalyst was
mixed with Nafion [28]. The Nafion ionomer has to be humidified to work as a sufficient proton
conductor. In the unsupported Pt TF catalyst, protons seem to be transported on the metallic surface
by either Pt-OH or Pt-H. The transport of protons on platinum was first investigated by McBreen
et al. [37] and recently, Zenyuk et al. suggested that the transport mechanism might be potential
dependent above 0.7 V, where Pt-O and Pt-OH are present on the surface. At potentials below the point
of zero charge, the proton is believed to be transported by surface migration [38]. This speculation
was already introduced earlier for oxygen reduction catalysts by Debe et al. for Pt on non-conductive
organic whiskers [39], the so-called nano-structured thin films (NSTF). This is further supported by the
fact that the catalyst layer is also more hydrophobic compared to Pt/C. Thus, at room temperature,
sufficient wetting of catalyst layer can be obtained only from water generated by the ORR reaction.
Thus, our results support the claim that proton transport on bare Pt is not a limiting step.
The specific activities and ECSAs for the RDE and GDE benchmarking measurements on Pt TFs are
summarised in Table 1, as well as compared to commercial Pt/C in the RDE, GDE and MEA. The specific
activity of Pt/C in the GDE cell at 60◦ C and 0.9 VRHE was found to be comparable to the specific activity
of the MEA with commercial Pt/C and decreased at room temperature. In general, the specific activity
at 0.9 VRHE of GDE and MEA of Pt/C was found to be smaller compared to the RDE measurements.
This is also the case for Pt TF where the specific activity in the RDE was benchmarked with 1.82 mA
cm−2 Pt at 0.9 VRHE and in the GDE with 0.72 mA cm−2 Pt with membrane or 0.71 mA cm−2 Pt without
membrane. The decrease in specific activity for commercial Pt/C is nearly 50% compared to a loss
approximately 60% for Pt TF. In the higher current density region, at 0.65 VRHE , the specific activity
with Pt TF is significantly higher than with Pt/C. In contrast, the current density at 0.65 VRHE by
mass is only slightly increased for Pt TF - with 24.1 A/mgPt for Pt/C and 28.6 A/mgPt for Pt TF. This
indicates that unsupported Pt catalysts have an advantage in activity compared to commercial Pt/C
catalysts, especially without humidification at low temperatures. In contrast, at higher loadings, where
most devices operate, the low Pt utilization inhibits the performance of the unsupported Pt catalysts.
The investigations show that with the help of simple benchmarking of different catalyst systems in the
GDE, new ways of developing possible applications can be explored. It turns out that different catalyst
systems also require different environmental conditions and can play to their strengths under different
operating conditions.
Table 1. Comparison of the characterization under different conditions for sputtered Pt thin-film
catalyst in GDE and RDE setup and as well as a commercial 46.5 wt % Pt/C catalyst (TEC10E50E,
Tanaka) in RDE, GDE and MEA measurements. SA, ECSA and RH stands for ORR specific activity,
electrochemical active surface area and relative humidity, respectively.
Catalyst
Layer
(Reference)
Pt TF RDE
(this work)
Pt TF GDE
(this work)
Pt TF GDE
(this work)
Pt/C RDE
[40]
Pt/C GDE
[28]
Pt/C GDE
[28]
Pt/C MEA
[41]

Electrolyte

Loading
[μgPt cm−2 ]

Temperature/Humidity
[◦ C]

SA
@0.9VRHE
[mA cm−2 Pt ]

SA
@0.65VRHE
[mA cm−2 Pt ]

ECSA [m2
g−1 Pt ]

0.1 M HClO4

16

rt

1.82

-

21

Nafion 117/4
M HClO4

16

rt/0% RH

0.72

144

20

4 M HClO4

16

rt/0% RH

0.71

67.5

20

0.1 M HClO4

14

rt

0.49

-

76

5

rt/100% RH

0.14

25.9

93

5

60/100% RH

0.18

47.9

81

90

80/100% RH

0.21

-

80

Nafion 117/4
M HClO4
Nafion 117/4
M HClO4
Nafion 117

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Surfaces 2019, 2

4. Conclusions
In this paper, we present the electrochemical characterization of platinum thin-films in a GDE
setup. Unlike the commonly used RDE method, this half-cell method facilitates investigations under
triple phase boundary conditions. Moreover, the GDE setup allows achieving current density values
comparable to operating current densities in actual PEMFCs.
Herein, we investigate sputtered Pt with very low loadings, achieving high current densities
(close to 1 A cm−2 ) in GDE measurements. Furthermore, we show that different parameters such as
loading, electrolyte concentration, humidification and Nafion membrane influence the performance of
the Pt catalyst layer. By carefully measuring the activity with different loadings in the GDE, we observe
that specific activities at potentials down to 0.75–0.80 VRHE can be determined independently of the
applied loading. The values obtained in the GDE were two-and-half times lower than those measured
with the RDE technique using the same catalyst. Our results confirm that activity values obtained in
liquid half-cells are not directly transferable to real devices. This is mainly due to transport resistance
at the triple phase interface and not a result of higher loadings or different testing parameters.
Interestingly, our results show that at fuel-cell relevant overpotentials, the activity is highly
dependent on the loading and use of the membrane. Furthermore, we observed that humidification
had no influence on the activity of catalyst layer for the unsupported catalyst. This is in contrast to
Pt/C, where the use of humidification nearly doubled the catalytic activity. These findings highlight
the potential impact of GDE half-cell benchmarking. While RDE remains a valuable tool for observing
trends in intrinsic catalytic performance, the properties of the catalyst layer have to be investigated
under more realistic conditions. Using sputtered platinum as a model catalyst, we show that the
catalyst layer can be efficiently evaluated in the fuel cell operation region using a GDE half-cell setup.
It is our hope that this enables accelerated development of fuel cell catalysts by shortening expensive
and time-consuming MEA optimization.
Author Contributions: Conceptualization, A.W.J. and G.S.; formal analysis, A.W.J. and G.S.; investigation, A.W.J.
and G.S.; data curation, A.W.J. and G.S.; project administration, M.A. and M.E.-E.; resources, V.B.; supervision,
M.E.-E. and M.A.; writing—original draft preparation, A.W.J. and G.S.; writing—review and editing, M.A. and
M.E.-E.
Funding: This research was funded by the Danish Innovation Fund (4 M center) and the Danish 350 DFF through
(grant number 4184-00332). M.E.-E. and A.W.J. gratefully acknowledge the Villum Foundation V-SUSTAIN grant
9455 to the Villum Center for the Science of Sustainable Fuels and Chemicals.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

39

Article

Electro-Oxidation of CO Saturated in 0.1 M HClO4 on
Basal and Stepped Pt Single-Crystal Electrodes at
Room Temperature Accompanied by
Surface Reconstruction
Kiyotaka Abe 1,2 , Hiroyuki Uchida 1, * and Junji Inukai 1, *
1
2

*

Clean Energy Research Center, University of Yamanashi, 4 Takeda, Kofu, Yamanashi 400-8510, Japan
Present address: GS Yuasa Corporation, 555 Sakaijuku, Kosai-shi, Shizuoka 431-0452, Japan;
kiyotaka.abe@jp.gs-yuasa.com
Correspondence: h-uchida@yamanashi.ac.jp (H.U.); jinukai@yamanashi.ac.jp (J.I.); Tel.: +81-55-220-8185 (J.I.)

Received: 28 December 2018; Accepted: 11 April 2019; Published: 16 April 2019

Abstract: The electro-oxidation of CO on Pt surface is not only fundamentally important in
electrochemistry, but also practically important in residential fuel cells for avoiding the poisoning of
Pt catalysts by CO. We carried out cyclic voltammetry on Pt(111), (110), (100), (10 10 9), (10 9 8), (10 2 1),
(432), and (431) single-crystal surfaces using a three compartment cell to understand the activity and
durability towards the electro-oxidation of CO saturated in 0.1 M HClO4 . During the potential cycles
between 0.07 and 0.95 V vs. the reversible hydrogen electrode, the current for the electro-oxidation of
CO at potentials lower than 0.5 V disappeared, accompanied by surface reconstruction. Among the
electrodes, the Pt(100) electrode showed the lowest onset potential of 0.29 V, but the activity abruptly
disappeared after one potential cycle; the active sites were extremely unstable. In order to investigate
the processes of the deactivation, potential-step measurements were also conducted on Pt(111) in
a CO-saturated solution. Repeated cycles of the formations of Pt oxides at a high potential and Pt
carbonyl species at a low potential on the surface were proposed as the deactivation process.
Keywords: CO electro-oxidation; Pt single-crystal electrodes; potential cycling; potential stepping;
surface reconstruction

1. Introduction
The electrochemical oxidation (electro-oxidation) of CO, particularly on the Pt surface, is one
of the most fundamental and most studied electrochemical reactions [1–3]. Many papers have been
devoted to the oxidation of a pre-adsorbed monolayer of CO, often referred to as CO stripping, which
usually occurs at a relatively high potential, in the 0.7 to 1.0 V range, vs. the reversible hydrogen
electrode (RHE), within a narrow potential range, far from the equilibrium potential for the reaction,
−0.104 V vs. RHE [4,5]. On the other hand, a current can often be observed in the lower potential
region, from 0.3 to 0.5 V vs. RHE, depending both on the characteristics of the Pt surface and also to
some extent on the potential at which CO was initially adsorbed. This current has been referred to as
“pre-peak,” “pre-wave,” “pre-oxidation”, or “pre-ignition” current. The origins of this current are still
controversial [6–15]. With the use of an electrolyte solution saturated with CO (bulk CO), the current
can flow steadily, since CO is being supplied continuously from the solution, in contrast to the case of
pre-adsorbed CO, which is only a monolayer of CO adsorbed on the surface.
Strmcnik et al., using ex situ scanning tunneling microscopy (STM), presented results that have
provided insight into the electro-oxidation of bulk CO in the lower potential region and found that
the presence of Pt islands and other low coordination structures on the Pt(111) surface is associated
with high activity for CO oxidation at low potentials [16]. They found that, during the potential
Surfaces 2019, 2, 315–325; doi:10.3390/surfaces2020023

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Surfaces 2019, 2

cycling process, these islands were removed. Thus, the authors proposed that the steps associated with
islands, particularly at the nanometer level, are important catalytic sites, due to their low coordination
number (CN). We reported the electro-oxidation of bulk CO saturated in 0.1 M HClO4 on stepped
Pt(111) surfaces: Pt(10 10 9) with (111) steps and Pt(20 19 19) with (100) steps (ball models in Figure 1),
combined with imaging by operando STM during the electro-oxidation of bulk CO [15]. Our results,
in agreement with those of Strmcnik et al. [16], suggest that the activity increases in the order: Terraces
< steps < kinks < step-adatoms (single adatoms at steps), i.e., increasing activity with decreasing
CN [17]. One of the most striking findings of ours was that potential cycling in the presence of CO
tends to strongly favor the transformation of steps to maximize the occurrence of (111) steps, so that,
for example, initially disordered (111) steps become perfectly straight, and initially straight (100)
steps become zigzagged with (111) microsteps [15]. Both steps become highly inert, again due to the
superb ordering of the CO adlayer across the step. The atomic-level insight, with a direct link between
surface structure and electrocatalytic activity, provided a step forward in guiding the design of new,
more active catalysts and electrocatalysts for CO oxidation and CO-tolerant hydrogen oxidation.

Figure 1. Unit triangle of stereographic projection, showing the location of various Pt surfaces.
Ball models of (20 19 19), (10 10 9), (10 9 8), (10 2 1), (432), and (431) planes are depicted in blue.

In spite of the fundamental and practical importance of the electro-oxidation of bulk CO in
electrolyte solutions in the lower potential region, the reaction on single-crystal electrodes was studied
only on some Pt(111)-related surfaces [12,15–18], Pt(110) [16], and Pt(100) [16,19]. However, in order to
increase the reaction rate of the bulk CO electro-oxidation in the lower potential region, the reaction on

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Surfaces 2019, 2

surfaces with “atomically-controlled” defects also must be elucidated. The information thus obtained
could be applied to the development of catalysts for residential fuel cells by avoiding the poisoning
of Pt catalysts with hydrogen-rich fuel gas (reformate) produced from hydrocarbon fuels. In this
work, we carried out potential cycling in a CO-saturated HClO4 solution not only on basal Pt(111),
(110), and (100), but also on surfaces with atomically zigzag steps of Pt(10 10 9), (10 9 8), (10 2 1),
(432), and (431). On all the surfaces, the electro-catalytic activity in the pre-peak region decreased and
disappeared after the potential cycling between 0.07 and 0.95 V in CO-saturated HClO4 . By examining
cyclic voltammograms (CVs) in N2 -saturated HClO4 before and after the potential cycling in the
CO-saturated solution, we propose the surface reconstruction on all the Pt electrodes. Potential-step
experiments were also carried out on Pt(111) to elucidate the possible mechanism of the Pt deactivation
during the potential cycles.
2. Materials and Methods
Single-crystal beads of Pt, approximately 3 mm in diameter, were made by crystallization of a
molten ball formed at the ends of Pt wires in a hydrogen/oxygen flame as reported previously [15,20–22].
The reflection of a laser beam from the crystal facets were used to determine the orientation of the
single-crystal bead for exposing either Pt(111) or the (100) plane (Figure S1, Supporting Information).
Each plane was then mechanically polished with successively finer-grade diamond pastes down
to 0.25 μm with an accuracy of <0.2◦ . For Pt(110), (10 10 9), (10 9 8), (10 2 1), (432), and (431)
planes, single-crystal beads were installed in a 4-axis X-ray diffractometer (Rigaku, Tokyo, Japan) for
determining those planes prior to the mechanical polishing [15,23]. All of the samples were treated
in a hydrogen/oxygen flame for two hours to eliminate surface damage caused by the mechanical
polishing. Each sample was placed in an infrared image furnace filled with hydrogen [15,22], heated
up to 1450 K, and gradually cooled down to 473 K, at which the gas was exchanged to Ar. The sample
was taken out and placed in ultrapure water (18.2 MΩ, Milli-Q, Merck Millipore, Billerica, MA,
USA). The Pt electrode covered with ultrapure water was carefully and immediately transferred into
an electrochemical cell filled with a 0.1 M HClO4 (ultrapure grade, Kanto Chemical, Tokyo, Japan)
solution saturated either with N2 or CO. The continuous potential cycling in a CO-saturated solution
was carried out between 0.07 and 0.95 V at 50 mV s−1 (35.2 s for each cycle), whereas the CVs in
a N2 -saturated solution were recorded between 0.05 and 0.95 V at 50 mV s−1 before and after the
30 potential cycles in a CO-saturated solution. On Pt(111), the potential-step measurements were
carried out. The sample was first exposed to a CO-saturated solution at 0.05 V for 2 min. Then, the
potential was stepped and kept at 0.60 V for 6 min, then at 0.95 V for 2 min, and back at 0.60 V for 6 min.
This 16-min procedure was repeated 26 times, and the current density was simultaneously measured.
All electrochemical measurements were carried out at 293 K using a potentiostat (PGSTAT128N with an
analog scan module, Metrohm Autolab, Herisau, Switzerland) with the hanging meniscus method in a
three-compartment electrochemical cell (Figure S2, Supporting Information). All electrode potentials
in this work are referred to the RHE.
3. Results and Discussion
3.1. CVs on Pt Single-Crystal Surfaces in CO-Saturated HClO4
The continuous potential cycling was carried out on Pt samples in a 0.1 M HClO4 solution
saturated with CO. After 30 cycles, the steady-state CV curves showed no change on all samples.
Figure 2(a-1)–(c-1) show CVs obtained in N2 -saturated HClO4 before (black line) and after (red line)
the 30 potential cycles in CO-saturated HClO4 on Pt(111), (10 9 8), and (432), respectively. Those three
planes are related to Pt(111) (Figure 1). On the Pt(111) surface used in this study, a small amount
of (111) steps existed as reported in our previous study [15]. On (10 9 8) and (432), ideally, steps
are composed of alternating (111) and (100) steps with one-atomic width (Figure 1). The average
step density was estimated to be 1/9 on Pt(10 9 8) and 1/4 on Pt(432). The hydrogen and hydroxide

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Surfaces 2019, 2

adsorption/desorption regions are seen at potentials lower and higher than 0.5 V, respectively. A subtle
growth of the spikes at around 0.8 V was seen on Pt(111), which is indicative of a small increase in the
flatness of the terraces [15,23]. On (10 9 8) and (432), very small changes are observed at the peaks
observed in both hydrogen and hydroxide adsorption/desorption regions. In H2 SO4 solutions, the
peaks around 0.13 and 0.28 V have been reported to be associated with (110) and (100) planes [23].
If this is also the case in HClO4 , both surfaces had characteristics of (110) and (100) planes, which
is evident from the ball models shown in Figure 1, although detailed structural analysis should be
carried out based on techniques, such as in situ STM [15]. Figure 2(a-2)–(c-2) show CVs on Pt(111),
(10 9 8), and (432), respectively, obtained in a 0.1 M HClO4 solution saturated with CO. The potential
was swept repetitively from 0.07 to 0.95 V with 30 cycles at a scan rate of 50 mV s−1 . Only the CVs
for positive-going scans are shown in the figures. Figure 2(a-3)–(c-3) are enlargements of the lower
potential range in 2(a-2), 2(b-2), and 2(c-2), respectively. The current due to CO oxidation initially
commenced at 0.37 V on Pt(111), at 0.43 V on (10 9 8), and at 0.42 V on (432). A broad anodic oxidation
feature, often termed a “pre-peak”, but referred to more properly as “low-potential CO oxidation
current”, at about 0.5 to 0.7 V was seen, followed by a sharp spike at around 0.8 V, known as the main
peak. After 30 cycles, practically no current was seen at the lower potential region than the main peak
on all surfaces. Because the CVs in the N2 -saturated HClO4 solution (Figure 2(a-1)–(c-1)) were similar
before and after the potential cycles, the surface structures might not be drastically changed; only
atomic structures of steps might have been changed as reported on Pt(111) and stepped Pt(111) with
either (111) or (100) steps [15]. In the case of (111) steps on Pt(111), the increase of the step density
increase the “low-potential CO oxidation current” [15]. However, for Pt(10 9 8) and (432), the increase
in the step density by a factor of 2.3 rather decreased the current (Figure 2). This could be explained
by the difference in the actual configurations at the steps at Pt(10 9 8) and (432), although in the ideal
models shown in Figure 1, the steps are identical on both surfaces. The ratio of the peak currents at 0.13
and 0.28 V are different on Pt(10 9 8) and (432) in an N2 -saturated solution, which also indicates that
the step structures on Pt(10 9 8) and (432) are different. Interestingly, the decrease of the “low-potential
CO oxidation current” was similar on Pt(111) and (10 9 8). Therefore, the step structure on Pt(10 9 8)
might be similar to that on Pt(111).
Figure 3 shows CVs on Pt(110) and (431). Pt(431) is oriented with (110), ideally with only a
2-atomic width of the (111) terraces (see Figure 1). On Pt(110), the CV in N2 -saturated HClO4 before the
potential cycles (black line in Figure 3(a-1)) was identical to those reported previously [24–26]. By in
situ STM, the surface prepared in the same manner as in this study was reported to consist of parallel
atomic rows aligned along the [110] direction, with an atomically-resolved (1 × 1) structure [22]. After
the potential cycles in CO-saturated HClO4 , many characteristic peaks in the CV were lost (red line in
Figure 3(a-1)). Therefore, the surface structure must be largely reconstructed. The lost peaks might
indicate a reconstruction into “dim” surface structures with no distinct steps and terraces after the
potential cycles. The CVs on Pt(431) in N2 -purged 0.1 M HClO4 obtained before and after the potential
cycles in a CO-saturated solution showed a large peak at 0.1 V in the hydrogen adsorption/desorption
region. Little change was observed before and after the potential cycles (Figure 3(b-1)) in contrast to
the results on Pt(110) (Figure 3(a-1)). This is indicative of little change in the surface structure after the
potential cycles on Pt(431), because of the unique surface morphology of Pt(431); the defect sites were
densely and uniformly arranged on the surface, which cause the surface energy to be rather small.
Figure 3(a-2),(b-2) show CVs on Pt(110) and (431) obtained in a 0.1 M HClO4 solution saturated with
CO, respectively. Figure 3(a-3),(b-3) are enlargements of the lower potential ranges. On Pt(110), a small
“low-potential CO oxidation current” with an onset of 0.35 V was seen to decrease as the potential
was cycled. On Pt(431), the “low-potential CO oxidation current” commencing at 0.40 V during the
first scan was comparable to those on Pt(111)-related surfaces (Figure 2), which abruptly decreased
even at the second cycle. Because we expected a large “low-potential CO oxidation current” on this
highly-defected surface, the results were surprising. In our previous paper, we concluded that bulk
CO electro-oxidation at potentials lower than 0.5 V proceed at surface sites not fully covered by CO

43

Surfaces 2019, 2

molecules [15]. If this reaction model is correct, a small amount of sites not covered by CO existed
before the potential cycles on Pt(431), but the sites quickly disappeared and the surface was efficiently
covered by CO during the first cycle. To conclude, bulk CO electro-oxidation at low potentials is not
necessarily high at a surface with many defects; the absorption structure of CO molecules must be
simultaneously taken into account, too.

Figure 2. (a-1), (b-1), and (c-1) show CVs obtained in N2 -saturated HClO4 before (black line) and after
(red line) the 30 potential cycles in CO-saturated HClO4 on Pt(111), (10 9 8), and (432), respectively.
Scan rate = 50 mV s−1 . (a-2), (b-2), and (c-2) show the results of the CVs on Pt(111), (10 9 8), and (432),
respectively, obtained in a 0.1 M HClO4 solution saturated with CO. Scan rate = 50 mV s−1 . (a-3), (b-3),
and (c-3) are enlargements of the lower potential range in (a-2), (b-2), and (c-2), respectively.

Figure 4 shows CVs on Pt(100) and (10 2 1). On (10 2 1), the steps are ideally all (111), but kinks
are located on the four-fold sites of Pt(100) (CN = 6), which are unique to this surface. On Pt(100), the
CV in N2 -saturated HClO4 before the potential cycles (black line in Figure 3(a-1)) was identical to
those reported previously [26]. In situ STM images with Pt atoms forming (100)–(1 × 1) structure were
already reported on the surfaces after the similar sample preparation procedures [27,28]. On Pt(100), the
electro-oxidation of bulk CO commenced at 0.29 V, the lowest potential among the surfaces investigated
in this study. On Pt(10 2 1), the onset potential for CO oxidation was 0.35 V, higher by 0.06 V than on
Pt(100). After the potential cycles in CO-saturated solution, the shape of the CVs changed on both
surfaces [15], but not in an extinctive manner as on Pt(110). Although the onset potential was very low
on Pt(100), the “low-potential CO oxidation current” became non-observable even at the second cycle.
Therefore, the sites for the electro-oxidation of bulk CO were extremely active, but unstable on Pt(100).
Since the onset potential of Pt(100) was lower than on Pt(10 2 1) with kinks, the active sites on Pt(100)
might be those with lower CN than kinks, or isolated Pt adatoms (CN = 4) or clusters composed of a
few Pt atoms (CN = 4–6) as proposed previously [15,16,27,29]. However, as already mentioned, the
reaction rate of the electro-oxidation of bulk CO is not determined only by the numbers of sites with
smaller CNs, but also by the structure of the CO adlayer [15,28].
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Surfaces 2019, 2

Figure 3. (a-1) and (b-1) show CVs obtained in N2 -saturated HClO4 before (black line) and after
(red line) the 30 potential cycles in CO-saturated HClO4 on Pt(110) and (431), respectively. Scan
rate = 50 mV s−1 . (a-2) and (b-2) show the results of the CVs on Pt(110) and (431), respectively, obtained
in a 0.1 M HClO4 solution saturated with CO. Scan rate = 50 mV s−1 . (a-3) and (b-3) are enlargements
of the lower potential range in (a-2) and (b-2), respectively.

Figure 4. (a-1) and (b-1) show CVs obtained in N2 -saturated HClO4 before (black line) and after
(red line) the 30 potential cycles in CO-saturated HClO4 on Pt(100) and (10 2 1), respectively. Scan
rate = 50 mV s−1 . (a-2) and (b-2) show the results of the CVs on Pt(100) and (10 2 1), respectively,
obtained in a 0.1 M HClO4 solution saturated with CO. Scan rate = 50 mV s−1 . (a-3) and (b-3) are
enlargements of the lower potential range in (a-2) and (b-2), respectively.

In Table 1, the onset potentials during the first potential scans are listed on all the Pt electrodes.
It is clear that the numbers of defects shown in the ideal structures (Figure 1) are not related to the order
of the onset potentials. This might be indicating that the structures of actual active sites are mostly not
as those shown in Figure 1. In our previous paper, it was reported that the active sites on Pt(111) existed
preferentially at steps and were very small in number [15]. Even by using single-crystal surfaces with
zigzag steps in this study, the active sites for bulk CO electro-oxidation were not increased. Interestingly,
45

Surfaces 2019, 2

the onset potentials were always lower on the basal planes rather than on stepped surfaces on Pt(111)-,
Pt(110)-, and Pt(100)-oriented surfaces. This might imply that those highly-active sites were adatoms
or clusters, having low CNs, on the basal planes or at steps, as discussed on Pt(100). The instability of
those sites is discussed in the following section.
Table 1. Onset potentials for bulk CO electro-oxidation in 0.1 M HClO4 saturated with CO.
Pt(111)
Onset potential/V vs. RHE

Pt(111)-Oriented
Pt(10 9 8)
Pt(432)

0.37

0.43

0.42

Pt(110)-Oriented
Pt(110) Pt(431)
0.35

0.40

Pt(100)-Oriented
Pt(100)
Pt(10 2 1)
0.29

0.35

3.2. Potential Stepping on Pt(111) in CO-Saturated HClO4
On all Pt surfaces with different types of defects, the current for the CO electro-chemical oxidation
in the lower potential region disappeared after potential cycles. By ex situ [16] and operando [15]
STM, it has been reported that the loss in the activity for the CO electro-oxidation was related to the
transformation of Pt surface structures by a reduction of the surface defects [16]. In our previous report,
the deactivation of the electrodes was connected with the formation of the single atomic Pt species [15]:
Anodic scan (up to 0.95 V):
Pt(OH)2 → PtO2 + 2H+ + 2e−

(1)

PtO + H2 O → PtO2 + 2H+ + 2e

(2)

Cathodic scan (lower than 0.80 V):
PtO2 + 2H2 O + 4H+ + 2e− → Pt(H2 O)4 2+

(3)

Pt(H2 O)4 2+ + 4CO → Pt(CO)4 2+ + 4H2 O

(4)

Pt(CO)4

2+

−

+ 2e → Pt(CO)2 + 2CO

(5)

At 0.95 V, bulk CO electro-oxidation and the formation of Pt oxides must competitively proceed
on surface, where Pt atoms with low CNs are expected to be easily oxidized. When the potential was
scanned in the cathodic direction, those (single or several atomic) oxides were reduced to Pt(CO)4 2+
or Pt(CO)2 on the surface. The single atomic Pt species, either as Pt(CO)4 2+ or Pt(CO)2 , would then
have the opportunity to redeposit at the nearby site in an energetically favorable configuration. In the
previous section, the active sites on Pt(100) were proposed as adatoms or atomically small clusters.
Pt(CO)4 2+ or Pt(CO)2 could be preferentially formed from those isolated Pt adatoms with small CNs,
therefore, those active sites should be very unstable as shown in Figure 4.
In addition to CV measurements (Figures 2–4), we carried out potential stepping [18,30]. The Pt(111)
surface was used because the atomically-resolved structural information was obtained only on
Pt(111)-related single-crystal electrodes during bulk CO electro-oxidation [15]. The sample was
exposed to a CO-saturated solution at 0.05 V for 2 min, stepped to 0.60 V and left for 6 min, to 0.95 V for
2 min, then back to 0.60 V for 6 min. The current density obtained at 0.60 V was chosen as a measure
for the CO electro-oxidation in the lower potential region. This 16 min procedure was cycled 26 times.
The current density was simultaneously recorded. Figure 5 shows the current density during the
potential-step cycles. The current density was initially 0 mA cm−2 at 0.05 V, then abruptly increased
to 0.17 mA cm−2 as the potential was stepped to 0.60 V at 2 min (black line in Figure 5). The current
density slowly decreased to 0.06 mA cm−2 during the 6-min interval. Subsequently, the current
density increased sharply to 0.5 mA cm−2 upon the potential step to 0.95 V at 8 min and decreased to
0.25 mA cm−2 at 10 min. When the potential was switched down to 0.60 V at 10 min, after a cathodic
spike, the current density very slowly decreased down to 0.02 mA cm−2 in 6 min. Interestingly, after the
sample was again kept at 0.05 V for 2 min, the current density jumped to 0.14 mA cm−2 accompanied
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Surfaces 2019, 2

by the potential step to 0.60 V, slightly lower than the first jump (0.17 mA cm−2 ). The current decreased
to 0.04 mA cm−2 at 0.60 V after 6 min (or at 8 min) in the second cycle. Indeed, the current density at
0.60 V between 2 and 8 min in the (N + 1)th (N = 1 to 25) cycle was always larger than that at 0.60 V
between 10 and 16 min in the Nth cycle, showing the partial recovery of the activity in the lower
potential region. Between 2 and 8 min and between 10 and 16 min at 0.6 V, the current density at the
corresponding period in the (N + 1)th cycle was always smaller than that in the Nth cycle (Figure 5),
showing a continuous deactivation. Between 8 and 10 min at 0.95 V, the current density was not in an
order. This discrepancy at 0.95 V might be due to the diffusion of the solution not controlled in the
experimental configuration shown in Figure S2. Rotation disk electrode measurements are needed for
detailed analyses.

Figure 5. Current density obtained on Pt(111) during the potential steps in CO-saturated HClO4 .

Operando STM measurements were carried out on Pt(111) in CO-saturated HClO4 at 0.05, 0.60,
and 0.95 V. At 0.95 V, the atomic arrangements were not able to be obtained because of the continuous
electro-oxidation of CO, although the terraces were observed to be flat. At 0.05 and 0.60 V, the (2 × 2)-3
CO structure was steadily observed as shown in Figure 6. Therefore, the CO adlayer was rigid and
stable between 0.05 and 0.60 V. It is also understood that the bulk CO electro-oxidation at 0.6 V
proceeded at steps. The reconstruction of steps without potential cycling either at 0.05 or 0.60 V was
not observed by STM, indicating that the mobile sites were very limited in number.

Figure 6. (a): Operando STM image of a CO adlayer on a Pt(111) terrace forming (2 × 2)-3CO obtained
in 0.1 M HClO4 saturated with CO at 0.60 V vs. RHE during the CO electro-oxidation. A unit cell is
depicted as a rhombus. The length of the sides was 0.54 nm, and the angle was 120◦ . (b): Structural
model of a (2 × 2)-3CO adlayer on Pt(111)–(1 × 1). The red and purple circles represent CO molecules
and Pt atoms, respectively.

Based on the results obtained above, the deactivation processes during the potential steps
were modeled.
1.

At 0.95 V, the surface was continuously deactivated by the simultaneous formation of surface
oxides (Figure 5, Equations (1) and (2)). Pt adatoms or Pt clusters with low CNs were preferentially
oxidized. Those oxidized species were immobile on the surface.
47

Surfaces 2019, 2

2.

3.

4.

The current density at 0.60 V, after being treated at 0.95 V, became much smaller (Figure 5) along
with the formation of a stable CO adlayer on the terraces (Figure 6). Pt(CO)4 2+ (Equation (4)) or
Pt(CO)2 (Equation (5)) species were also formed after the potential step by the reduction of the
(single or several atomic) Pt oxides. We believe that, at least partially, Pt(CO)4 2+ remained on
the surface at 0.60 V after the potential step from 0.95 V, because Pt oxides were not completely
reduced to metal Pt at 0.60 V after going through 0.95 V, as seen in the CVs obtained in pure
HClO4 (Figures 2–4). Because the reactivity did not change very much at 0.60 V, the surface
morphology did not change either. Therefore, the planar and electronically charged Pt(CO)4 2+
species, inactive towards the bulk CO electro-oxidation, might be immobile on the Pt surface.
At 0.05 V, all Pt(CO)4 2+ species became Pt(CO)2 . The current density at 0.60 V after treatment at
0.05 V became larger than that after treatment at 0.95 V (Figure 5), because of the formation of
Pt(0)(CO)2 from Pt(II)(CO)4 2+ ; metallic Pt species can be active towards bulk CO electro-oxidation.
The reaction rate, on the other hand, was continuously lowered at 0.60 V. This is because of
the deposition of mobile Pt(CO)2 at the energetically favorable adjacent site on the surface,
thus lowering its CN and eventually its reaction activity.
Repeating the treatments at 0.05 and 0.95 V gradually decreased the reaction rate of the bulk CO
electro-oxidation in the lower potential region (Figure 5). The potential cycling between 0.07 and
0.95 V much enhanced the deactivation process on Pt electrodes (Figures 2–4) compared with the
potential-step treatment (Figure 5), possibly because of the continuous morphological change
during the potential cycling.

4. Conclusions
The electro-oxidation of CO was investigated on Pt single-crystal electrodes with different surface
orientations in CO-saturated 0.1 M HClO4 . On all Pt(111)-, (110)-, and (100)-oriented surfaces, the
“low-potential CO oxidation current” disappeared during the potential cycles between 0.07 and
0.95 V in a CO-saturated solution. During the potential cycles, the surface structures changed.
On Pt(111)-oriented surfaces, the reconstruction proceeded only at steps [15]. On Pt(110), the
reconstruction was very large, probably changing the terrace structures too. On Pt(110)-oriented
Pt(431), the change in surface structure was larger than Pt(111)-oriented surfaces, but much smaller
than Pt(110), in spite of the largest numbers of steps and kinks, probably because of the less structural
irregularity and smaller surface energy compared with Pt(110)–(1 × 1). On Pt(100)-related surfaces,
the electro-oxidation of bulk CO proceeded at lower potentials than on other surfaces (Table 1).
Especially on Pt(100), the onset potential was as low as 0.29 V. The unstable Pt adatoms or clusters
were proposed to exist on the surface.
The potential-step measurements were carried out on Pt(111) in CO-saturated HClO4 . At 0.95 V,
the Pt atoms with low CNs were preferentially oxidized, which was then reduced to Pt(CO)4 2+ or
Pt(CO)2 upon the potential step to 0.60 V. Pt(CO)4 2+ was immobile at 0.60 V and inactive towards the
bulk CO electro-oxidation. At 0.05 V, only Pt(CO)2 existed as a single Pt species on the surface, which
were mobile and decreased its CN by the self-deposition at a stable site; the surface reactivity thus
decreased. On pure Pt electrodes, unfortunately, it seems improbable to maintain the “low-potential
CO oxidation current” after potential cycles in CO-saturated HClO4 . The active sites, allowing the
adsorption of water molecules through the CO adlayer [15], must be structurally maintained eventually
by a surface modification, such as by alloying.
The systematic investigation of the atomically-arranged defects on electrodes has started only
recently [15,16]. The combination of defects and adsorbates may create highly-active sites, but are
very small in number. As reported in this paper, those created sites could be unstable and changeable
too. New analytical technologies and theoretical backbones must be established and applied to this
new field.
Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/2/2/23/s1,
Figure S1. Schematic illustration of the adjustment of (111) and (100) facets by a laser beam reflection method.
48

Surfaces 2019, 2

Figure S2. Schematic illustration of a tree-compartment cell. The electrochemical measurements were carried out
with the hanging meniscus method.
Author Contributions: Conceptualization, J.I.; Methodology, H.U. and J.I.; Validation, H.U. and J.I.; Investigation,
K.A.; Resources, H.U. and J.I.; Writing—original draft preparation, J.I.; Writing—review and editing, H.U. and J.I.;
Visualization, K.A. and J.I.; Supervision, H.U. and J.I.; Project administration, J.I.; Funding acquisition, J.I.
Funding: This study was supported by Grant-in-Aid for Scientific Research (KAKENHI) and the New Energy
and Industrial Technology Development Organization (NEDO), Japan.
Acknowledgments: The authors thank Prof.
stepped-sample preparation

Nagakazu Furuya of the University of Yamanashi for the

Conflicts of Interest: The authors declare no conflict of interest.

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article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Spatially Resolved XPS Characterization of
Electrochemical Surfaces
Benedetto Bozzini 1 , Danjela Kuscer 2 , Matteo Amati 3 , Luca Gregoratti 3, *, Patrick Zeller 3 ,
Tsvetina Dobrovolska 4 and Ivan Krastev 4
1
2
3
4

*

Dipartimento di Ingegneria dell’Innovazione, Università del Salento, v. Monteroni, 73100 Lecce, Italy;
benedetto.bozzini@unisalento.it
Jožef Stefan Institute, Jamova cesta 39, 1000 Ljubljana, Slovenia; Danjela.Kuscer@ijs.si
Elettra - Sinctrotrone Trieste S.C.p.A. S.S. 14, km 163.5 in Area Science Park, 34149 Trieste-Basovizza, Italy;
matteo.amati@elettra.eu (M.A.); patrick.zeller@elettra.eu (P.Z.)
Institute of Physical Chemistry, Bulgarian Academy of Sciences, 1113 Sofia, Bulgaria;
tsvetina@ipc.bas.bg (T.D.); krastev@ipc.bas.bg (I.K.)
Correspondence: luca.gregoratti@elettra.eu

Received: 20 February 2019; Accepted: 9 April 2019; Published: 15 April 2019

Abstract: Synchrotron-based scanning photoelectron microscopy (SPEM) has opened unique
opportunities for exploiting processes occurring at surfaces and interfaces, which control the
properties of materials for electrochemical devices, where issues of chemical and morphological
complexity at microscopic length scales should be faced and understood. The present article aims
to demonstrate the present capabilities of SPEM to explore the surface composition of micro- and
nano-structured materials, focusing on cases relevant to electrochemical technologies. We report and
discuss a selection of recent results about three different systems, targeting hot topics in the fields
of electrochemical energy storage and electrochemical fabrication: (i) an in-depth analysis of Ag-In
electrodeposited alloys exhibiting dynamic pattern formation, (ii) the analysis of electrochemical
processes at the electrodes of a self-driven solid oxide fuel cell and (iii) an operando characterization
of a single-chamber solid oxide fuel cell. The last example has been performed at near-ambient
pressure conditions using a unique specially designed setup which extends the traditional capabilities
of scanning photoemission microscopes in the ultra-high and high-vacuum regimes to operating
conditions that are closer to realistic ones, contributing to overcome the so-called “pressure gap”.
Keywords: operando; near ambient pressure XPS; scanning photoelectron microscopy; solid oxide
fuel cells; surface science; electrodeposited alloys

1. Introduction
The ubiquitous and enabling role of electrochemical technologies in many different fields of
technology [1], but especially in energy storage [2], cannot be overemphasized. Electrochemical
processes and the devices in which they are implemented are highly complex systems in which
successful functioning is the result of the cooperation of different types of physics interacting on a range
of diverse space- and time-scales. The complex nature of electrochemical systems generally makes the
task of understanding and optimizing them an extremely challenging one. One of the main difficulties
underlying electrochemistry is the necessity of transferring charge between electrically and ionically
conducting phases in processes that are coupled with chemical reactions and phase-transformations.
Moreover, these complex processes occur mainly at the electrode–electrolyte interface and involve
thin layers of material. In these regions structural, chemical and electronic properties of the materials
may significantly deviate from the intrinsic bulk ones. An additional obstacle in the characterization
of the reactive interfaces is their position: they are often buried and not easily accessible without
Surfaces 2019, 2, 295–314; doi:10.3390/surfaces2020022

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Surfaces 2019, 2

disturbing the functionally relevant chemical and structural ambient. In the last decades, traditional
bulk-sensitive analysis tools, in particular, X-ray diffraction or fluorescence spectromicroscopy, largely
employed in the characterization of electrochemical systems, have been more extensively supported
by X-ray based surface sensitive techniques such as X-ray photoelectron spectroscopy (XPS) which can
maximize the information confined in thin layers [3]. XPS was initially conceived as a spectroscopy
technique necessarily requiring operation in ultra-high vacuum (UHV) or at least high-vacuum (HV)
ambient and it was developed and implemented with this approach for many decades. But, as for
other fields of science and technology, a state-of-the-art method for the investigation of electrochemical
systems should tackle the so-called “material gap” and “pressure gap”. The first one is linked to the
development of nanomaterials and nanotechnology and the corresponding need of understanding
processes at the nanoscale, while the latter aims to analyze an electrochemical reaction in the real
pressure working regime. The latter is a challenge for many vacuum-based characterization techniques.
In recent years a lot of efforts have been put in trying to overcome the two limitations for XPS, with the
successful outcome of developing novel approaches.
The fast progress in the development and implementation of photoelectron spectromicroscopy
techniques, i.e., adding spatial resolution and imaging capabilities to XPS, started in the last decade
of the 20th century with the construction of low emittance X-ray synchrotron machines, providing
tunable energy photon beams with very high brightness and variable polarization [4,5]. Thanks to the
high photon intensity provided by insertion devices it has become possible to add sub-micrometer
lateral resolution to XPS, realizing new instruments called photoelectron microscopes [6]. In X-ray
photoelectron microscopes high spatial resolution is achieved using either full-field or scanning
methods. In scanning instruments, called scanning photoelectron microscopes (SPEM), the photon
beam is demagnified to sub-micrometer dimensions using suitable photon optics. Modern SPEMs
can achieve spatial resolutions as low as 70–100 nm, allowing full material characterization at the
nanoscale [7,8]; despite the fact that full-field photoemission microscopes can achieve even better
lateral resolutions (20–30 nm), they impose severe constraints in the sample geometry and morphology
which, instead, do not affect the performance of scanning ones. This flexibility with respect to
sample requirements is particularly crucial for the operando characterization of electrochemical systems.
Since its construction in the mid of 90’s, the soft X-ray SPEM at the Elettra synchrotron research
center has investigated many electrochemical systems based on solid state electrolytes and electrodes
in the standard UHV or HV environment, often in operando conditions [9–11] demonstrating how
the material gap can be successfully overcome in the characterization of electrochemical systems.
With respect to the pressure gap problem in XPS, a fundamental contribution to overcome it was
achieved with the development of the differentially-pumped electron analyzers at the beginning of
this century; after almost 20 years of advances, XPS analysis at pressures as high as several tens of
millibar can be performed in gas and liquid environments not only at synchrotron beamlines but also
with conventional instruments [12,13].
In addition to standard near-ambient pressure XPS approaches (generally denoted by the acronym
APXPS) to the characterization of electrochemical systems, now largely diffused [14–16], it is worthwhile
to mention a couple of developments particularly adaptable to electrochemistry: (i) the implementation
of hard X-rays XPS, which allows spectroscopy investigation of deeper layers [17] such as buried
interfaces and (ii) the “dip and tip” method to study solid-liquid interfaces [18,19], available at few
synchrotron beamlines. In spite of their innovation all these approaches fail in combining APXPS
with submicron spatial resolution; this technological problem started to be solved as recently as in
the last decade thanks to the efforts of the team of the Escamicroscopy beamline at Elettra which
allowed the development and implementation of innovative technical solutions for SPEM, capable of
performing near-ambient pressure photoemission spectromicroscopy while keeping the typical SPEM
performance [20,21].
This manuscript is presenting three examples of electrochemical experiments performed with
SPEM in three different environmental conditions. In the first experiment, the characterization of an

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ex situ electrodeposited Ag-In alloy, exhibiting a dynamic pattern formation mechanism, has been
performed. In the second experiment the electrochemical processes at the electrodes of a self-driven
solid oxide fuel cell (SOFC) have been analyzed; in this case the operando condition was achieved in
HV pressure regimes, while the last case reports on the operando characterization of a single-chamber
solid oxide fuel cell (SC-SOFC) at near-ambient pressure conditions, using a newly developed cell.
2. Materials and Methods: the Soft X-rays Scanning Photoemission Microscope of Elettra
2.1. The Standard UHV—HV Setup
The standard configuration of the SPEM was originally designed to work in UHV conditions;
a schematic view of the focusing optics and photoelectron detection systems of the SPEM of the
Escamicroscopy beamline at Elettra is shown in Figure 1. The X-ray beam is focused on the sample
using a diffractive lens called zone plate (ZP), which allows to achieve a spot diameter of ~130 nm as
best performance. In the typical experiment the X-ray spot impinges normally to the sample even if
a polar rotation is possible. While samples are raster-scanned with respect to the X-ray microprobe,
the photoelectrons generated at the sample surface are collected and energy-selected by means of a
hemispherical electron analyzer (HEA), equipped with a 48-channels electron detector [22]. The takeoff
angle of photoelectrons is fixed to 30◦ . This configuration strongly enhances the surface sensitivity
of the SPEM because due to inelastic collisions only photoelectrons generated in the topmost layers
can reach the HEA. A SPEM can be operated in two modes: (i) imaging spectromicroscopy and
(ii) microspot spectroscopy. The imaging mode maps the lateral distribution of elements by collecting
photoelectrons within a selected kinetic energy window while scanning the specimen with respect to the
microprobe. By properly setting the energy range, the distribution of the chemical states of a selected
element can be measured, in addition to that of the bare elemental concentration. The microspot mode
is similar to conventional XPS, i.e., energy distribution curves are measured from the illuminated local
micro-spot area selected within an image.

ȱ
Figure 1. An illustrative sketch of the SPEM instrument. The X-ray beam is focused with a zone-plate
(ZP) and an order-sorting aperture (OSA), down to a spot of 130 nm and the sample is raster scanned in
front of it.

In this configuration, the samples are exposed to the common environment of the analysis chamber,
which hosts the focusing optics, the HEA and other ancillary equipment; the highest allowable pressure
in the case of experiments where samples need to be exposed to gases is about 1 × 10−5 mbar.

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2.2. Novel Solutions for Operando Near-ambient Pressure Approaches
The typical photon energies and ZP dimensions used in the SPEM imply distances between ZP
and samples of a few mm. Moreover, an additional pinhole, the order-sorting aperture (OSA), located
between the ZP and the sample, is necessary to select a single diffracted spot. These geometrical
constraints are the main reason why the technical solutions found to overcome the “pressure gap” for
XPS setups could not be transferred straightforwardly to SPEMs. At Elettra two special setups have
been recently developed, allowing experiments in near-ambient pressure conditions, preserving the best
achievable spatial resolution: (i) the Dynamic High Pressure (DHP) setup and (ii) the Near-Ambient
Pressure Cell (NAP-Cell).
The DHP system uses a highly collimated gas jet as a supersonic jet expansion into the SPEM
vacuum chamber [23]; gases are confined into a narrow beam and point to the sample surface by using
a pulsed jet instead of a continuous one. Similar setups have been already used in other research fields
to control the amount of injected gases [24,25]. The DHP set-up has been designed to fit into the SPEM
chamber: the gas jet is generated in the region around the sample by a thin needle connected to a pulsed
valve which produces a series of gas jet shots at a fixed repetition rate. Each shot gives rise to a short
burst of pressure in a small and confined volume near the sample and to an increase of the background
pressure in the SPEM chamber. By finely tuning (i) the shot duration, (ii) the sample-to-needle distance
and (iii) the DHP gas line pressure, the pressure at the sample can be controlled, while the injected
gas is efficiently diluted into the large chamber volume, so that the background pressure does not
exceed the HV limit in the analyzer. Figure 2 shows a typical single-shot pressure time-profile at the
sample surface for a pulsed valve aperture time of 3.2 ms and the corresponding SPEM background
pressure. The distance between the sample and the needle tip was fixed to 2 mm and the pressure of
the gas in the DHP gas line was 3.5 bar. As shown in Panel (a), in these conditions the sample surface
is experiencing a gas environment of ~10 mbar for a few ms. The corresponding change in the SPEM
pressure is reported in Panel (b). A valve aperture frequency of 0.35 Hz is needed to keep the SPEM
background pressure within the limit of ~1 × 10−5 mbar. More details can be found in ref [21].

ȱ
Figure 2. (a) Time-profile of a single pressure shot at the sample surface for a pulsed valve aperture
time of 3.2 ms. (b) Corresponding pressure change in the SPEM chamber.

To address experiments where a constant pressure value is needed, the NAP-Cell has been
developed. This device consists of a small vacuum-sealed cell-in which the sample is encapsulatedequipped with small pinholes as photon-in/photoelectron-out apertures [20]. A sectional view of this
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Surfaces 2019, 2

setup is shown in Figure 3. The gases needed in the experiment are delivered into the cell through a
flexible metal bellow which terminates the dosing line located inside the SPEM chamber. A gas-dosing
system allows regulation of the gas flow into the line at the desired flux/pressure. An encapsulated
heater allows to heat up the sample, at the back of which it is placed, to temperatures up to 900–950 K.
Additional electrical contacts are available for biasing samples in electrochemical setups. A pressure
gap between the SPEM vacuum chamber and the inner volume of the NAP-Cell is established by
the impedance induced by the small pinholes, which makes it possible to reach NAP conditions
(~0.1 mbar) inside the cell while maintaining HV condition outside it (10−5 mbar). It is worth noting
that, at variance with a standard NAP setup for spectroscopy, a differentially-pumped electron analyzer
is not needed.

ȱ
Figure 3. Working principle of the Near-Ambient Pressure Cell (NAP-Cell) technology for
SPEM. The reactive gases needed in the experiment are delivered into the cell from the bottom
aperture. Two separate holes are designed for the incoming X-rays and the outgoing photoelectrons.
An encapsulated heater and the electrical contacts are located inside the cell.

Regarding the microscopy capability, the entire cell is scanned with respect to the X-rays beam,
as for conventional UHV experiments. The “visible” area of samples hosted in the NAP-Cell is defined
by the size and orientation of the two holes and it is approximately a circle of 0.4 mm diameter.
A special design, characterized simply by one smaller pinhole, allows to reach up to 1 mbar, but at the
expense of a limitation in the field of view, down to 200 μm.
It is important to highlight that real devices such as catalytic or electrochemical systems operate at
pressure equal or higher than one atmosphere; therefore, it is the dream of each technique of analysis
to reach operando conditions similar to the realistic ones. The NAP-Cell developed at Elettra—in line
with most of the currently available NAP photoelectron systems around the world—may operate at
pressure levels which are still a few orders of magnitude lower than realistic ones; nevertheless these
conditions may activate chemical and physical processes which are pressure-dependent and which
do not occur at the typical UHV or HV pressure levels of standard XPS. An example can be found in
the last example of this manuscript: the partial conversion of the Sr signal from “lattice” to “surface”
found at NAP conditions could not be detected while operating at HV pressures. We are aware of
the fact that modern APXPS electron analyzers for spectroscopy can achieve pressures of a few tens
of mbar—which nevertheless, as hinted at above, does not essentially shrink the gap with respect to
practical operating conditions—but with a spatial resolution limited to hundreds of μm. The SPEM at
Elettra is the only available XPS microscope where a pressure value of 0.1 mbar can be achieved for
operando measurements.
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2.3. Sample Preparation
2.3.1. Ag-In Alloys
A spontaneously patterned Ag-In electrodeposit was obtained from an aqueous electrolyte of
composition: KAg(CN)2 (Umicore, Schwäbisch Gmünd, Germany) 30 mM, InCl3 (Alfa Aesar, Kandel,
Germany) 0.1 M, D(+)-glucose (Fluka) 0.1 M, KCN (Merck, Darmstadt, Germany) 0.5 M. The alloy
was grown onto a Cu foil cathode (2 × 1 cm2 ) that was electrochemically degreased and pickled in
H2 SO4 20 vol % prior to electrodeposition. To avoid Ag displacement-plating, the Cu cathode was
brought in contact with the electrolyte undercurrent. Alloy plating was carried out at room temperature
in a 2-electrode cell using a large-area (20 cm2 ) platinized-titanium anode. Pattern formation can
be controlled by appropriately adjusting: (i) the In3+ /Ag+ ratio in the plating bath, (ii) the current
density and (iii) the fluid dynamic conditions of the electrolyte. Figure 4 reports the experimental
conditions ensuring pattern formation in stagnant (diamonds) and stirred (crosses, magnetic stirring at
50 rpm, 100 mL cell) electrolytes. The alloy was grown by electrodeposition at 0.5 A/dm2 for 2.5 h,
yielding a coating thickness of ca. 40 μm. More experimental and theoretical details, as well as
comprehensive reference to relevant literature, can be found in ref. [26]. The sample analyzed in this
research (Figure 5) shows a double-spiral pattern with a spiral-core distance of ca. 0.5 mm and branch
spacing of 0.15 mm. The sample exhibits a 3-phase crystallographic structure, featuring α−Ag-rich
terminal solid solution (Figure 6; marked with red dots) and the Ag3 In and AgIn2 intermetallics,
as shown by the Powder Diffraction Files (PDF) Cards No: 29-0677 and No: 65-1552, respectively, in
keeping with the literature [27,28]. The same phase structure was found in the single-spiral sample
studied in [26]: the quantitative differences in phase composition and preferential orientations are
compatible with differences in electrodeposition current density.

ȱ
Figure 4. Electrodeposition conditions yielding pattern formation in Ag-In alloys: metal concentration
ratio in the electrolyte (Y-axis); current density (X-axis) and fluid-dynamic conditions of the bath
(1: stagnant, 2: stirred).

The photon energy employed was 652 eV, yielding an energy resolution of about 0.4 eV for
imaging. The intrinsic surface sensitivity of soft-X ray photoelectron spectroscopy, enhanced by the
grazing acceptance detection of SPEM (see Section 2.1 above), provides an effective probing depth
limited to 1 nm, with dominating contribution from the top two atomic layers. Details on the image
processing procedure are reported in the Supporting Information of ref. [26]. Ar+ bombardment
(1.5 × 10−5 mbar Ar, 1–1.5 kV, sample current of 0.8 μA) erosion was employed for depth-profiling
purposes; the ion beam was delivered at 45◦ with respect to the sample. Sequential sputtering periods
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were applied, corresponding to depths of 2, 4, 10, 20, 50, 65 and 90 nm, estimated according to standard
models for sputtering depth ([29] and references therein): for full methodological and quantitative
details, refer to ref. [26]. These sputtering steps effectively allow for monitoring of progressive
depth variations in the alloy pattern. Possible sputter-induced roughening of the original surface
topography [30] was monitored by AFM and found to be negligible and therefore to have no impact
on the measured variations in the composition. Moreover, the evident chemical changes taking place
in-depth demonstrate that intermixing and re-deposition of sputtered material are negligible.

Figure 5. The electrodeposited Ag-In alloy exhibiting a double-spiral pattern, subjected to depth-dependent
SPEM analysis. (a) The analyzed sample mounted on the SPEM sample-holder. (b) Optical-microscopy
image. (c) SPEM chemical-state map at the Ag 3d energy, showing the Ag intermetallic/Ag solid
solution ratio at a depth of ca. 50 nm.

Figure 6. X-ray diffractogram of the Ag-In electrodeposited sample analyzed by SPEM, exhibiting the
double-spiral pattern. The phase-structure features an α-Ag terminal solid solution and two Ag-In
intermetallics; the remaining reflexes correspond to Cu reflections from the substrate.

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2.3.2. Mn-Ni SOFC at High Vacuum
For this experiment, we have used planar, YSZ(100)-supported cells with Ni and Au/Mn layers:
the electrodes were fabricated by evaporating 100 nm thick Ni layers, and Au (70 nm)–Mn (20 nm)
bi-layers grown onto a Cr adhesion buffer (10 nm): the cell scheme is depicted in Figure 7 (for details
on fabrication, see, e.g., ref. [31]). Prior to measurements, the samples were subjected to Ar+ sputtering
to remove contaminant layers introduced during fabrication, followed by annealing at 650 ◦ C in
1 × 10−5 mbar O2 to restore the initial oxidation state of the materials. Electrochemical measurements
were run with a PAR potentiostat. The cell was short-circuited through a potentiostat working in a
ZRA (Zero-Resistance Ammetry) mode for simultaneous monitoring of the current response.

ȱ
Figure 7. Scheme of the electrolyte (YSZ)–supported the planar cell with Ni and Mn electrodes for
in situ electrochemical SPEM measurements. The images at the bottom are SPEM maps at the Ni 2p
and Mn 2p energies of the electrode-electrolyte interfaces.

To optimize the photoemission yield of the main core levels, a photon energy of 1000 eV was set.
As for all the experiments reported in this paper, a 5.0 purity grade of the gases was used.
2.3.3. LSM-NiO Single Chamber SOFC
The experiments on the single chamber SOFC were performed on (100)-oriented single crystal
YSZ (ZrO2 stabilized with 9.5 % Y2 O3 ) electrolyte on top of which a mixture of 50 wt % of LSM-YSZ
(La0.8 Sr0.2 )0.95 MnO3 and 50 wt % of (Y2 O3 )0.08 (ZrO2 )0.92 was used as cathode and a mixture of 60 wt %
NiO and 40 wt % of (Y2 O3 )0.08 (ZrO2 )0.92 (Ni–YSZ) as anode. The LSM-YSZ and Ni-YSZ powder
mixtures were blended with an organic vehicle and screen-printed on the electrolyte, forming two
symmetrical rectangles separated by a gap of ∼ 1.5 mm and sintered at 1200 ◦ C for 2 h. An Pt grid
was also screen printed onto the electrodes and subsequently fired at 1100 ◦ C for 30 min to provide a
more uniform distribution of potentials and more robust electrical contact. The SC-SOFC experiments
were carried out at 923 K with a CH4 /O2 gas mixture with a volume ratio of 2/1, in conformity with the
largest proportion of the available literature. A more detailed description of the sample preparation
can be found in ref [32].
To optimize the photoemission yield of the main core levels a photon energy of 750 eV was set.
In the results section of this paper, a collection of photoemission spectra and chemical maps will be

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shown; in most of the measurements, an overall energy resolution of 0.35 eV was set. The deconvolution
of the spectra was performed using Doniach–Sunijch functions convoluted with Gaussians, following
well-established procedures.
Electrochemical measurements have been carried out in a two-electrode configuration with a
Versastat potentiostat.
The planar configuration, with anode and cathode on the same plane, is a classical SC-SOFC
configuration also for practical applications (see, among others, the following reviews [33–35]).
The relative arrangement of the electrodes in practical systems is often of the interdigitated type, by
the main reason for this is the extent of the electrodes in larger devices and this has no impact on
microkinetics. One could argue about the optimization of the inter-electrode distance, but this aspect is
exclusively relevant if energetic optimization of the device is at stake. In electroanalytic terms, however,
the contribution of the ohmic drop of an unoptimized electrolyte gap can be very easily separated
from charge-transfer and mass-transport contribution by standard electrochemical measurements,
such as electrochemical impedance spectrometry. Therefore, with a straightforward ohmic correction,
all the electrokinetic information can be accurately retrieved. In particular, the link between electronic
conditions and surface chemistry can be established with utmost accuracy with our device.
3. Results
3.1. In-Depth Analysis of Ag-In Ex Situ Electrodeposited Alloys Exhibiting a Dynamic Pattern
Formation Mechanism
A characteristic aspect of electrodeposited alloy patterns is that these structures can be imaged
very precisely by visible-light microscopy, but exhibit a strikingly poor contrast in SEM, owing to
their extreme surface-confinement. As proved in ref. [26,36], SPEM, owing to its high chemical-state
sensitivity combined with surface sensitivity and lateral resolution, is the method of choice for
unravelling the morphochemical details of electrodeposited alloy patterns. In view of completing and
further validating the scenario set up in the two publications mentioned just above, in this study we
have explored the 3D space elemental and chemical-state distribution of an Ag-In sample, fabricated
under different conditions with respect to the one investigated in ref. [26]. Insight into the chemical
speciation of Ag and In had been gained in the detailed study reported in ref. [26], were we disclosed
that Ag is present in two chemical states: elemental (solid solution) and intermetallic, while In exhibits
an oxidized and an intermetallic form. This chemical-state scenario has been confirmed in the analysis
of the new sample and this information has been directly used for image processing of the Ag 3d and
In 3d SPEM maps, with the method expounded in the Supporting Information of ref [26]. Briefly, the
compositional inhomogeneity of the sample has been quantified by selecting the parts of the spectrum
recorded in correspondence of each pixel of the image, representing one of the two components of the
two key elements of the alloy. Figure 8 shows a representative selection of Ag 3d5/2 and In 3d5/2 SPEM
images of the electrodeposited double-spiral, obtained by eroding the sample to the indicated depths.
The maps depicted in Panels (a) and (c) were derived from the total intensity of the Ag and In 3d5/2
spectra and represent the spatial distribution of the Ag and In surface concentrations. Panels (b) and
(d) show the Ag and In chemical maps, obtained by selecting the spectral regions representative of
different chemical states. For Ag (Panel (b)) the contrast is due to the combination of intermetallicand solid solution-type metallic states. Regarding In (Panel (d)), instead, contrast is determined by
the distribution of In between Ag-In intermetallics and In3+ oxides/hydroxides. Coherently with our
findings of ref. [26], it is clear that elemental contrast (Panels (a), (c)) tends to fade out rapidly in depth,
while chemical-state contrast (Panels (b), (d)) is markedly more persisting. In addition, it is worth
noting that the Ag and In elemental (Panels (a), (c)) and chemical-state distributions (Panels (b), (d))
are anticorrelated, in keeping with the surface concentration ratio distribution, imaged in Panel (e).
Quantitative details, highlighting in-depth changes, can be more evidently pinpointed by averaging
over space, as shown in Figure 9. Our results confirm the alloy electrodeposition mechanism proposed
in [26] and further validates the mathematical DIB model of electrochemical pattern formation, recently
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developed in our group ([37,38] and references therein). More specifically, the present results provide
additional support of the model capability of predicting complex spiral-based morphologies. Briefly,
co-electrodeposition of In with Ag in amounts exceeding the solubility of the former element in the
latter, gives rise to the formation of Ag-In intermetallics following a two-step process. Ag+ and In3+
are reduced to the elemental state, forming an α -Ag terminal solid solution together with Ag-In
intermetallics, but simultaneously basic In3+ salts precipitate at the electrode-electrolyte interface.
The basic-salt film, in turn, can be electro-reduced forming fresh elemental In that can react with α-Ag
to form more intermetallic. This two-step process effectively explains why In3+ oxi-hydroxides can be
found at the alloy surface, but they are absent in the bulk of the layer. As far as pattern formation is
concerned, the DIB model provides a simple framework for the rationalization of this phenomenon,
in that it predicts that the coupling of electrodeposit morphology with adsorption of an electroactive
intermediate can give rise to the formation of dynamic patterns, among which double-spirals [39].
SPEM maps provide factual spectroscopic evidence of the presence of adsorbed In3+ electroactive
intermediates and of their organization in a typical pattern. In addition, a set of DIB-model parameters
can be identified, yielding a computed pattern with a remarkably good matching between the theoretical
map and the experimental pattern (Figure 10). The difference in the numerical values of the parameters
that correspond to the double-spiral pattern with respect to those found for the single spiral studied
in [26], is coherent with the different current-density values employed for the growth of the two alloys
(for details on the correspondence between model parameters and experimental variables, see [40]).

Figure 8. Ag 3d5/2 (a,b) and In 3d5/2 (c,d) SPEM surface concentration (a), (c) and chemical-state (b),
(d) maps, recorded at different depths. Topographical contributions have been removed from all maps.
The Ag/In ratio maps (e) outlines the compositional variations in depth.

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Surfaces 2019, 2

ȱ
Figure 9. Spatial averages, with respective standard deviations of compositional and chemical-state
observables from the maps of Figure 8. (a) Elemental amounts rationed to the background, from Panels
(a) and (c) of Figure 8; (b) intermetallic fractions, from Panels (b) and (d) of Figure 8; (c) Ag/In ratios
from Panel (e) of Figure 8.

ȱ
Figure 10. Comparison between optical-microscopy image (left, a) and computed (right, b) pattern
from the DIB model, obtained with the following parameter values: α = 0.5; γ = 0.2; k2 = 2.5; k3 = 1.5;
A1 = 10; A2 = 30; B = 22; C = 2.2.

3.2. Spectromicroscopy Analysis of Electrochemical Processes at the Electrodes of a Self-driven Solid Oxide
Fuel Cell
The present section contributes to the scantily covered field of in situ SPEM studies of self-driven
electrochemical systems, in which chemical energy is spontaneously converted into electrical energy.
In situ studies of self-driven fuel cells and batteries, though more difficult than those of externally-driven
systems such as symmetric cells or cells with a single electrode running the real reaction of the energetic
device, is notably important because the couplings between the anodic and cathodic phenomena are
known to play a crucial role for the efficiency and durability of devices for electrochemical energetics.
In the present study, we investigate a cell with a solid cathodic reagent, MnO2 , and an anodic gas, H2 ,
reacting at a Ni/NiO electrocatalyst. In particular, we compare operating conditions with active gas
pressures in the high vacuum and NAP ranges.
Sample annealing in 10−5 mbar of O2 at 923 K in the SPEM chamber, resulted in the oxidation of
the starting materials to NiO and MnO2 that, by exposure to different pressures of H2 , can transform
into Ni and MnO, generating different redox couples at the two electrodes, yielding in turn electrical
power, according to the scheme of Figure 11. The electrokinetics of the system was followed with
transient experiments consisting in switching from O2 10−5 mbar to H2 10−5 and 1 mbar, and recording

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time-dependent spectra and current time-series (Figure 12). In order to emphasize the current response,
the cell was short-circuited and the current was measured with a ZRA. Curve (A) in Figure 12 shows a
representative transient corresponding to switching from 10−5 mbar O2 to 10−5 mbar H2 . In keeping
with the thermodynamic analysis of ref. [41], H2 is oxidized to H2 O and MnO2 is reduced to MnO,
giving rise to the electrochemical reaction indicated in Figure 12, whereby O2− is transported through
the electrolyte. In correspondence, also Ni forms, enhancing the H2 oxidation rate, as witnessed by the
progressive rise in current, in keeping with the literature [42]. Curve (B) is obtained by switching from
10−5 mbar O2 to 1 mbar H2 . In these conditions, on the one hand, the electrochemical reaction is faster,
but, on the other hand, also the chemical reduction of MnO2 is anticipated, resulting in a progressive
decrease of the cell current.

Figure 11. Reaction scheme for the MnO2 /YSZ/NiO cell in H2 ambient.

ȱ
Figure 12. Current transients corresponding to O2 /H2 gas-switching experiments with a MnO2 |YSZ|NiO
cell at 923 K: (A) from 10−5 mbar O2 to 10−5 mbar H2 ; (B) from 10−5 mbar O2 to 1 mbar H2 . The image
is elaborated from Sci. Rep. 3 (2013) 2848, with permission.

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Panels (A) and (B) of Figure 13 report representative sequences of Ni 2p μXPS spectra acquired after
having switched the gas ambient from 10-5 mbar O2 to 10−5 and 1 mbar H2 , respectively. Replicated
experiments confirm in detail the observed kinetics. Ni in the initial state is fully oxidized and a
progressive reduction can be noticed, that is notably faster in the case of the lower-pressure experiment.
The reduction rates are summarized in Panel (C). Panel (D) reports the Mn 2p μXPS spectra acquired
for the initial and final states of the cathode at the two H2 pressures investigated: formation on
MnO can be clearly assessed from the characteristic satellite structure [43]. Coherently with the
recorded cell-currents and compatibly with the thermodynamic framework of ref. [41], at high pressure,
the electrochemical reaction with H2 oxidation at NiO and MnO2 reduction to MnO is faster than the
chemical Ni reduction and for this reason NiO is stabilized at short times. At low pressure instead,
the chemical reduction of Ni is favoured, resulting in a progressive enhancement of the electrocatalytic
performance of the anode.

ȱ
Figure 13. (A,B) Time-dependent Ni 2p μXPS spectra recorded at the indicated times after gas-switching
from 10−5 mbar O2 to: (A) 10−5 mbar H2 , (B) 1 mbar H2 . (C) Time dependence of the intensity of the
elemental Ni peak normalized on the maximum of the oxidized Ni peak. (D) Mn 2p μXPS spectra
measured at the beginning and at the end of the gas-switching transients.

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3.3. Operando Characterization of a Single-chamber Solid Oxide Fuel Cell at Near-ambient Pressure
3.3.1. Electrochemical Measurements
Since limited literature information is available on the performance of SC-SOFC at the pressures of
interest for APXPS, preliminarily to operando spectroscopy, we carried out calibration work in a range
of vacuum conditions spanning the interval including the actual operating conditions and slightly
higher ones, that allow the achievement of electrochemical responses that match the ones obtained at
conventional working pressures. In Figure 14 we report a representative selection of current transients
obtained by short-circuiting the cell in the reactive gas mixture: the actual operating condition during
SPEM. More information including pressure-dependent open-circuit potentials and electrochemical
impedance spectrometry can be found in ref. [26]. We have chosen to operate the SC-SOFC under
short-circuited conditions because, on the one hand, even though the current that can be drawn from the
device is rather low, it can be measured accurately and demonstrates the correct SOFC operation and,
on the other hand, since short-circuit represents the most aggressive condition, successful operation
in this condition is an a fortiori proof of the feasibility of the approach. These experiments consisted
in measuring the cell current with the ZRA method in CH4 /O2 = 2/1 mixtures of defined pressure.
One can observe that a sudden current surge is followed by a relaxation, the rate of which negatively
correlates with the pressure. After the transient, the current tends to a well-defined steady-state:
this allows to run NAP-SPEM experiments under perfectly controlled electrochemical conditions.

ȱ
Figure 14. Current-density transients for the short-circuited SC-SOFC, recorded with the Zero-Resistance
Ammetry (ZRA) technique in CH4 /O2 = 2/1 ambient at the indicated pressures at 924 K. (A) Experiments
carried out in a mock-up of the analysis chamber. Red arrows indicate pump-down transients. (B) Data
recorded during operando NAP-SPEM measurements.

3.3.2. Spatially Resolved Photoemission Measurements at Near Ambient Pressure
In order to prove the validity of our approach as robustly as possible, we have decided to carry out
operando XPS concentrating on the cathode of the NiO/YSZ/LSM SC-SOFC, since the chemical state of
the elements of this perovskite is far less sensitive to electrochemically-induced changes in comparison
with Ni [41]. In the SPEM, the cell temperature was set to 923 K for the whole measurement time.
To avoid any reduction of the electrode/electrolyte surface due to the combination of vacuum/high
temperature, in the stand-by periods between electrochemical measurements, a partial pressure of
5 × 10−2 mbar of O2 was introduced in the NAP-Cell. In these conditions, we have measured several
μXPS survey and high-resolution spectra as well as SPEM maps in different locations to probe the
chemical composition of the cathode and of the interface region and its morphology of pristine cells.
In addition to the main constituents of the cathodic/electrolyte material (La, Sr, Mn, Zr, Y, O) we also

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assessed the presence of two impurities: C and Si, the latter being typical for YSZ [42]. Figure 15
reports a representative selection of maps at the Sr 3d and La 4d energies (Panel (A)) and μXPS spectra
(Panel (B)) in the region across the interface between the LSM-YSZ patch (left) and the YSZ electrolyte
(right). These elemental maps show a clear topographic contrast in the cathode region, highlighting
micrometric and submicrometric granular morphology, combined with evident chemical contrast.
In addition, signal from Sr and La can be detected inside the electrolyte, proving diffusion of these
elements into the YSZ region. Differently from Sr and La, Mn remains strictly confined to the cathode
region. At the right of Panel (A), we show intensity profiles for Sr 3d and La 4d signal extracted
from the SPEM maps along the dotted lines and in Panel (B) we report μXPS spectra measured at
different positions along the same line, up to 240 μm from the electrode/electrolyte interface. The La
4d spectrum is partly overlapping the Si 2p one, that provides a convenient internal reference for
relative intensity variations of this perovskite component. The spatial dependence of the Sr and La
signal shows that the surface coverage of YSZ with Sr is constant in the investigated zone, while a
La gradient can be appreciated, highlighting a difference in the surface diffusion coefficients of these
two components of the cathodic material on the electrolyte. In this work, we concentrated on the
evolution of the cathode elements in SC-SOFC ambient and under electrochemical operating conditions.
In particular, we carried out a fine analysis of the La 4d, Zr 3d, O 1s and Mn 2p core levels, acquired on
different grains of the cathode, as well as in representative locations of the electrolyte, in the following
ambients: (i) 0.05 mbar O2 ; (ii) 0.1 mbar H2 and (iii) 0.1 mbar CH4 /O2 = 2/1. The electrochemical
conditions tested were: (a) OCP in all three ambients, (b) short-circuited cell in ambient (iii). Ambient
(iii) is the reactive gas mixture employed for self-driven SC-SOFC operation (for details, see [26,41]);
as hinted-at above, ambient (i) was used to preserve the original chemistry of the cell components in
stand-by conditions between electrochemical measurements and ambient (ii) was set for 1 h before
starting the SC-SOFC operation in order to activate the NiO anode [38]. Briefly, detailed analysis of
the La 4d, Zr 3d and Mn 2p core level spectra acquired on different grains of the cathode and on the
electrolyte under all investigated conditions, did not exhibit any effect either of the ambient or of the
electrochemical polarization. These spectra showed the typical binding energies and spectral shape
reported in the literature for similar materials [43–46] and the only differences were in the intensities,
as a result of the heterogeneous stoichiometry of the LSM-YSZ grains. Instead, the Sr 3d line exhibited
chemical-state sensitivity to the gas ambient: specifically, spectral changes have been found in H2 and
CH4 /O2 ambients, as detailed below. It is worth noting that the chemical state of Sr in SOFC materials
has been investigated by ex situ XPS in some detail in the literature [44,46–48]. Briefly, on the one hand,
Sr has been described in La-Sr perovskites with transition metals and, on the other hand, as a species
diffused on YSZ. The former state—characterized by an XPS spectrum with peaks at BEs between
131.4 and 132.5 eV—, corresponds to the perovskite lattice. The latter state—exhibiting peaks at BEs
between 132.8 and 133.8 eV—denotes the separation of a secondary oxidic phase. In the cases reported
in the literature, these two prototypical chemical states are typically found in combination within the
probed volume.
In our experiments carried out in H2 atmosphere, a slight, but well-defined spectral change in the
low BE tail of Sr spectra located on the cathode patch was found during the change from O2 to H2
atmosphere, while the same element, in the form expressed when diffused into the electrolyte region,
did not respond to the change (Figure 16, spectra a), b), d) and e)). This chemical change of Sr in LSM
cannot be reliably accounted for on the basis of the available literature, but can be justified in terms of
partial reduction.
As anticipated in ref. [26], the introduction of SC-SOFC CH4 /O2 2/1 gas feed in the NAP-Cell
brings about a notable, irreversible change in the Sr 3d spectrum measured at the cathode (Figure 16,
spectrum (c)), while no changes are found in Sr diffused onto the electrolyte patch (Figure 16,
spectrum (f)). The spectral change found in the cathode region occurs at OCP and it is not further
modified by applying electrochemical polarization within the range of self-driven SC-SOFC operation.
The μXPS spectrum recorded in the cathode region exhibits a strong surface contribution, while that

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measured on the electrolyte patch can be entirely accounted for by lattice Sr. The observed irreversible
surface segregation of Sr by XPS is in keeping with literature results reporting, on the one hand, the
formation of a SrO surface phase by crystallographic means [49] and, on the other hand, electrocatalytic
degradation due to blocking of the ORR active sites [50].

ȱ
Figure 15. (A) SPEM maps at the Sr 3d and La 4d energies and elemental concentration profiles,
recorded with a pristine cell in 0.05 mbar of O2 at 924 K. The imaged region is the interface between the
LSM-YSZ cathode (left) and the YSZ electrolyte (right). (B) Sr 3d and La 4d μXPS spectra measured at
the indicated distances from the cathode/electrolyte interface along the dotted line shown in Panel (A).

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Figure 16. Sr 3d μXPS spectra recorded on the cathode ((a), (b) and (c)) and electrolyte ((d), (e) and (f))
patches at 923 K. Spectra (a), (b), (d) and (e): comparison of OCP conditions in 0.05 mbar of O2 and
0.1 mbar H2 . Spectra (c) and (f): comparison between electrochemically inert 0.05 mbar O2 ambient
and reactive SC-SOFC mixture CH4 /O2 2/1 0.1 mbar.

4. Conclusions
Scanning photoelectron spectromicroscopy has undergone a fast development at the third-generation
synchrotron light sources, becoming a true microscopic tool for probing material properties, mainly
at their surface. In the field of electrochemistry this approach has been proved capable of following
morphological and chemical effects occurring at surfaces and interfaces. The present paper reviews
three different experiments showing how SPEM can overcome the “material” and “pressure” gaps in
the characterization of electrochemical surfaces. In the first experiment the fascinating spatiotemporal
patterns formed during an ex situ electrodeposition of an Ag-In alloy has been investigated. A depth
profile mapping of the distribution of the chemical states of Ag and In with high spatial and spectral
resolution has been performed providing a 3D frame of the patterns. The experimental data perfectly
match the mathematical predictions of the DIB model of electrochemical pattern formation.
The other two cases reported in the manuscript deal with the analysis of electrochemical processes
at the electrodes of solid oxide fuel cells. In the first case, an in situ study of a self-driven cell was
performed. The cell was formed by a solid cathodic material, MnO2 , and a mixture of Ni/NiO as
anode. The system was analyzed at high vacuum and near ambient environmental pressure conditions.
The main achievement is that, coherently with the recorded cell currents at high H2 pressure, its
oxidation at NiO and the corresponding MnO2 reduction to MnO is faster than the chemical Ni
reduction and for this reason, NiO is stabilized at short times. At low pressure instead, the chemical
reduction of Ni is favoured, resulting in a progressive enhancement of the electrocatalytic performance
of the anode.
In the last experiment, we have shown an operando analysis of a single-chamber solid oxide fuel
cell at near ambient pressure regime by using a unique specially designed setup which extends the
traditional capabilities of scanning photoemission microscopes in the ultra-high and high-vacuum
regimes to more realistic conditions of environmental pressure. The response of the cathodic site of a
NiO/YSZ/LSM SC-SOFC exposed to O2 , H2 and a mixture of CH4 and O2 showed a chemical-state
sensitivity of Sr: a well-defined spectral change in the low BE tail of Sr spectra located on the cathode
patch was found during the change from O2 to H2 atmosphere, while the same element, when diffused
into the electrolyte region, did not respond to the same change. This chemical change of Sr in LSM
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cannot be reliably accounted for on the basis of the available literature, but can be justified in terms
of partial reduction. Moreover, the CH4 /O2 2/1 gas feed brings about a notable, irreversible change
in the Sr 3d spectrum measured at the cathode, while no changes are found in Sr diffused onto the
electrolyte patch.
We believe the reported examples demonstrate the great potential of SPEM to gain insight into
electrochemical and chemical processes taking place in energy conversion and other systems, which
can guide the design of next-generation devices with improved performance.
Author Contributions: D.K. prepared the single-chamber SOFC test samples. All authors characterized the
samples, revised and edited the manuscript.
Funding: The research leading to these results has received funding from the European Community’s Seventh
Framework Programme (FP7/2007-2013) under grant agreement nº 312284 and the European Community’s
Horizon 2020 Framework Programme under grant agreement nº 730872. The author D.K. acknowledges the
financial support from the Slovenian Research Agency (research core funding No. P2-0105).
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Probing the Surface of Noble Metals
Electrochemically by Underpotential Deposition of
Transition Metals
Nolwenn Mayet, Karine Servat, K. Boniface Kokoh and Teko W. Napporn *
IC2MP, UMR 7285 CNRS, Université de Poitiers, 4 rue Michel Brunet B27 TSA 51106,
86073 Poitiers CEDEX 09, France; nolwenn.mayet@univ-poitiers.fr (N.M.);
karine.servat@univ-poitiers.fr (K.S.); boniface.kokoh@univ-poitiers.fr (K.B.K.)
* Correspondence: teko.napporn@univ-poitiers.fr
Received: 15 February 2019; Accepted: 2 April 2019; Published: 9 April 2019

Abstract: The advances in material science have led to the development of novel and various materials
as nanoparticles or thin films. Underpotential deposition (upd) of transition metals appears to be
a very sensitive method for probing the surfaces of noble metals, which is a parameter that has an
important effect on the activity in heterogeneous catalysis. Underpotential deposition as a surface
characterization tool permits researchers to precisely determine the crystallographic orientations
of nanoparticles or the real surface area of various surfaces. Among all the work dealing with upd,
this review focuses specifically on the main upd systems used to probe surfaces of noble metals
in electrocatalysis, from poly- and single-crystalline surfaces to nanoparticles. Cuupd is reported
as a tool to determine the active surface area of gold- and platinum-based bimetallic electrode
materials. Pbupd is the most used system to assess the crystallographic orientations on nanoparticles’
surface. In the case of platinum, Bi and Ge adsorptions are singled out for probing (1 1 1) and (1 0 0)
facets, respectively.
Keywords: underpotential deposition (upd); Au; Pt; Pd; nanoparticles; cyclic voltammetry;
electrocatalysis

1. Introduction
Surface of materials is a key parameter in heterogeneous catalysis. Moreover, in electrocatalysis,
it is the reaction site where the exchange of electrons occurs, thereby the structure of the materials or any
roughness affect their electrochemical response. Referring the current density to the geometric surface
area does not consider the surface structure of the material, however, the active sites at the surface of a
material depend on its structure, therefore, the characterization of material surface in electrocatalysis
becomes an important step in understanding the reactions that occur and evaluating their activity.
Several electrochemical methods are used for characterizing the material surface. The underpotential
deposition (upd) appears to be a powerful tool for probing the surface of electrocatalysts. The upd
is sensitive to local order, contrary to electron diffraction methods such as Low Energy Electron
Diffraction (LEED) or Reflexion High Energy Electron Diffraction (RHEED) which are more sensitive to
long-range order [1]. The upd can be carried out on polycrystalline surfaces, single-crystals, and recently
nanoparticles. The latter have different shapes and sizes that involve different surface structures
leading to a variety of activities observed in heterogeneous catalysis and in electrocatalysis. The
development of nanoscaled materials in electrochemistry implies the increase of the utilization of upd
for various applications, from surface probing to the synthesis of nanomaterials.

Surfaces 2019, 2, 257–276; doi:10.3390/surfaces2020020

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Surfaces 2019, 2

2. Generalities about upd
2.1. Principles and Thermodynamics of the upd Process
The upd of a metal is defined as the deposition of a metal onto a foreign metallic substrate in a
potential region where a new and pure phase does not form on the substrate.
Figure 1 shows a schematic representation of the voltammetric profile for upd and the overpotential
deposition (opd) of a metal on a foreign substrate. The upd occurs at higher potentials than the potential
of reduction (Eupd > E) because of the strong interactions between the deposited metal and the substrate
surface structure. In the upd region, the metal is deposited as a monolayer on the substrate through
self-limiting deposition. In the opd region, the metal is deposited as multilayers (bulk electrodeposition)
in the absence of self-limiting behavior, creating rough structures [2].

Figure 1. Schematic profile of a voltammetric curve showing the underpotential deposition (upd) and
overpotential deposition (opd) processes of a metal (in orange) on a foreign substrate (in black).

The reduction potential Eeq of the metal (M) associated with the reaction Mn+ + ne− → M is
determined by the Nernst equation:
◦

Eeq = E +

RT aMn+
,
ln
nF
aM

(1)

where Eeq is the Nernst potential of the reaction, E◦ is the standard potential of the reaction, R is the
universal gas constant, T is the temperature, n is the number of exchanged electrons in the reaction, F is
the Faraday constant, aMn+ is the activity of Mn+ ions, and aM is the activity of the condensed phase M
(aM = 1).
The difference between the upd onset potential (Eupd ) and the bulk electrodeposition potential (E)
is correlated to the work function of the deposited metal and the substrate [3]. The underpotential shift
(ΔU = Eupd − E) can be plotted against the difference in work function between bulk substrate and
bulk deposited metal ΔΦ [3]:
ΔU = αΔΦ, α = 0.5 V eV−1 .
(2)
This linear relation between underpotential shift and work function difference suggests that the
covalent part of the ad-atom-substrate bond does not differ appreciably to the bond strength between
the ad-atom and the surface of the same metal.
The thermodynamic aspects of upd has been extensively reported on both polycrystalline and
single-crystalline surfaces [4–6]. It shows clearly that the interaction between particle and substrate

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has a strong effect on the changes in work function and also the interaction between particles, which
leads to several energy states. The effect of partially charged upd species that leads to the absence of
discreteness of charge has been also reported [6]. The theoretical aspect of upd has been reported to
understand thermodynamics behind this phenomenon [7,8]. It is well known that crystallographic
orientation, defects (steps, kinks, grain boundaries), surface reconstruction [9], and alloying play a
crucial role in the initial steps of metal deposition. Figure 2 shows the three low Miller indices for a
faces-centered cubic (fcc) system. The plane (1 1 1) is the most compact arrangement by comparison to
the more opened (1 1 0) plane.

Figure 2. Representation of the three low Miller indices for a faces-centered cubic system.

2.2. A Method of Characterization: Cyclic Voltammetry
Different techniques can be employed to characterize upd layers [10]. Cyclic voltammetry is
the method most used for carrying out the upd of a metal over a substrate. It consists in linearly
applying a potential with a constant scan rate (dE/dt) on the electrode, between two fixed potential
limits. The current is monitored during the scans, the current-potential curves (voltammograms) show
the oxidation-reduction processes that occur at the surface. The upd curves display cathodic and
anodic peaks, for the deposition and dissolution of the deposited metal, respectively. The presence of
several distinct adsorption peaks on the voltammograms indicates that the formation of the monolayer
takes place at different energetic adsorption steps. The structure of the peak strongly depends on
the crystallographic orientation of the substrate and the defects. The peaks are not well defined for
polycrystalline electrodes, which present different crystallographic orientations and a high density of
defects (kinks, grain boundaries, steps, etc.)—by comparison with single-crystal electrodes.
2.3. The upd Metals
Various substrates were used to perform the upd of a metal. The studies began with polycrystalline
substrates [11,12], which have a mixed surface structure. To understand and explain the phenomenon,
well defined surfaces are proposed. Therefore, upd of a metal monolayer onto polycrystalline or
single-crystal substrates of a foreign metal has been studied intensively. These investigations included
various substrates such as Au, Pt, Ag, Cu, Rh, Ru, Pd, and many deposited ions like Ag+ [13,14],
Cu2+ [15–24], Cd2+ [25], Pb2+ [26–30], Tl+ [31–34], Cd2+ [35,36], Hg2+ [37], Sn2+ [38,39], Bi3+ [30,40–45],
Ge4+ [46–48], Sb2+ [49–51], and Ni2+ [52]. The charge corresponding to the upd permits researchers
to estimate the amount of deposited metal. It is assumed that the metal ions in the sub-monolayers
are completely discharged, even if partially charged deposits have been reported. In a large number
of studies, the effect of a foreign metal monolayer deposited by upd on a metallic substrate was
investigated in relation to electrocatalytic reactions [53,54]. Indeed, the upd of a monolayer of a foreign
metal on a substrate modifies its surface structure. This modification has the advantage of affecting
the adsorbed species on the surface and thus the electrochemical reaction. The metal deposited
is called ad-atom. It was used in electrosynthesis to modulate the fabrication of different reaction
products from the same compound [55,56]. The upd process also plays a crucial role in the synthesis of
nanoparticles. The layer obtained by upd can be the precursor for fabricating nanoparticles of noble
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metals by galvanic replacement [57]. In this short review, upd of transition elements on noble metals’
surfaces will be addressed. It includes how surface crystallographic orientations are probing and
describes the determination of electrochemically active surface area (ECSA).
3. The Mupd on Noble Metals: Au, Pt, and Pd
3.1. The Case of Au
3.1.1. Cu on Au
The underpotential deposition of Cu (Cuupd ) on low-index facets of Au appears to be the most
extensively studied system. A large number of studies concerns the Cuupd on gold single-crystal
Au(1 1 1) [16,18,20,23,58–65]. Investigations on the electrode surface permitted to provide evidences of
the Cuupd species at the strained gold surface [24,66]. It is difficult to calculate the upd shift of copper
ions using electrolytes containing different anions because of their effect. Indeed, anions provide a
great influence on the Cuupd on Au, leading to different surface structures and onset potentials of the
upd process [21,22,67,68]. Cuupd has been investigated on different gold surface orientations (low and
high index). Therefore, Cuupd on Au(1 0 0) has been scrutinized by scanning tunneling microscopy
(STM) in hydrochloric and sulfuric acids. A (1 × 1) copper ad-layer structure was reported in sulfuric
acid, while an incommensurate (2 × 1) one-dimensional structure was observed when hydrochloric
acid was added [59]. These anions are specifically adsorbate species because chloride anions are more
strongly bound on the gold surface than sulfate anions. Moreover, in the presence of bromide anions
on Au(1 1 1) and Pt(1 1 1), a study realized has demonstrated once again that the anions have an
influence on the upd process and the stability of the deposited ad-layer [17]. It also strongly depends
on the substrate because the different behaviors on Cuupd originate from geometric constraint imposed
by the lattice structure relative to the deposited ad-layer rather than from energetic considerations.
The studies of Cuupd on Au(1 1 1) in sulfuric acid media suggest the co-adsorption of 2/3 monolayer of
√
√
copper and 1/3 of sulfate (SO4 2− ) [15]. This gives a pseudomorphic ( 3 × 3) structure, followed by a
further deposition of 1/3 monolayer of copper and completion of a full epitaxial (1 × 1) monolayer of
√
√
copper which is still covered by a ( 3 × 3) layer of sulfate anions [18]. A recent study goes further into
the mechanism of Cuupd in the presence of sulfates [21,22]. Indeed, cyclic voltammetry investigations
were performed on Au(1 1 1) in both upd and opd regions. In the potential region where the Cuupd takes
place, the first peaks which correspond to the phase transition of sulfate anions on Au(1 1 1) substrate
were observed. During the cathodic potential sweep, the adsorption of 2/3 monolayer and then 1/3
monolayer of copper, both with the co-adsorption of sulfate anions, were respectively obtained. During
the anodic potential sweep, desorption of the formed monolayer occurs. When the CV is recorded in
both upd and opd regions at lower potentials, the process of adsorption/desorption of multilayer copper
√
√
is observed. According to these authors, the EC-STM images reveal the nucleation of the ( 3 × 7)
√
√
2−
sulfate structure. At −500 mV vs. Pt/PtO, the ( 3 × 3) structure of SO4 anions on 2/3 ML of copper
can be seen on HR-STM image. This structure is also observed on two-dimensional gold islands. These
islands are provided from the lifting of the (1 1 1) reconstruction. This study shows that different
√
√
structures are observed on terraces resulting from different local thickness of copper. A ( 3 × 3)-like
structure of sulfate is typical of an incomplete second layer of copper. Figure 3 shows a model (top
√
√
√
view) of the sulfate structure ( 3 × 3), and 2/3 of monolayer of Cu, and the sulfate structure ( 3 ×
√
7) on the first pseudomorphic monolayer of Cu deposited on Au (1 1 1).

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Surfaces 2019, 2

√
√
√
√
Figure 3. Model of the ( 3 × 3) sulfate structure, with 2/3 of Cu monolayer (left) and ( 3 × 7) sulfate
structure on the first pseudomorphic monolayer of Cu underpotentially deposited (right) on Au (1 1 1).

Cuupd studies were carried on stepped surfaces of gold with [(n−1)(1 1 1) × (1 1 0)] (n = 6, 7, 10)
structures [19]. Figure 4 shows the voltammograms for Cuupd on (3 3 2), (7 7 5), and (5 5 4) surfaces
of Au. These structures are made of (1 1 1) terraces and (1 1 0) steps. The voltammogram of Au(3 3
2) displays two broad and overlapping peaks in the first adsorption/desorption region (at 0.520 and
0.560 V vs. RHE). When the step density decreases (the width n of the terrace increases), the peak at
0.560 V vs. RHE is decreasing and leads to assign it to (1 1 0) steps. This potential is close to the main
upd peak on the Au(1 1 0) electrode (0.570 V vs. RHE). In the case of the peak at 0.520 V vs. RHE, the
√
√
comparison with copper upd on Au(1 1 1) indicates the formation of ( 3 × 3) R30◦ structure on the (1
1 1) terraces. It can be seen that this peak becomes better defined when the width of (1 1 1) terraces
increases. By deconvoluting these peaks with Gaussian functions, it is found that the charge calculated
for (1 1 0) steps on Au(3 3 2) is 17%, with respect to the total upd charge. By comparison, the relative
geometric area of the (1 1 0) steps domain covers 16.7 % of the surface, which is in good agreement
with the relative calculated charge. However, since Cuupd is a complex function of sulfate coverage, it
should be noticed that there is no significant change in the sulfate coverage in this potential range. It is
necessary to measure the charge of the co-adsorbed sulfate anions in order to perform a full quantitative
analysis of the copper upd charge on stepped surfaces. On the other hand, the voltammograms in
√
Figure 4 display peaks between 0.350 V and 0.400 V vs. RHE. These peaks are attributed to the ( 3 ×
√
◦
3) R30 → (1 × 1) phase transition on the (1 1 1) terraces. The splitting of this peak is also influenced
by the step density but can be a consequence of the reconstruction of stepped surfaces. This study
clearly shows that the steps have an influence on the energetic/kinetics of phase transition. It has been
suggested that this influence could be either electronic or structural (Smoluchowski effect).

Figure 4. Cyclic voltammograms of (a) Au(3 3 2); (b) Au(7 7 5); and (c) Au(5 5 4) in 0.05 mol L−1 H2 SO4
+ 1 mmol L−1 CuSO4 at 5 mVs−1 . Reprinted from reference [19]. Copyright (2004) with permission
from Elsevier.

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The Cuupd can be employed to quantify the surface area of nanoporous gold films [69]. The values
obtained by this method are in agreement with those obtained by integrating the reduction peak of
gold oxides. However, it should be noted that nanoporous gold films contain residual silver since
their preparation consists of the dealloying of bimetallic films. The presence of silver can influence the
Cuupd , since this process does not occur on silver.
Many papers reported the effect of metal upd on the electrocatalytic activity of a material. In some
cases, the underpotential deposited metal can inhibit or even totally deactivate the surface towards
electrocatalytic reactions. Cuupd has been used in order to elucidate the mechanism of hydrogen
evolution reaction on palladium-modified stepped gold surfaces [70]. The Cuupd allows us to block
(deactivate) the palladium step sites towards hydrogen evolution reaction, which is not favored in
the presence of copper. Figure 5 shows ball models for Au(h k l)/Pd surface, with steps blocked by
copper ad-atoms.

Figure 5. Ball models of Au(h k, l)/Pd (h = k, l =k−1) surface where the Pd steps sites are blocked by
Cu modified with 1 ML of Pd.

A technique named Dynamic Electro-Chemo-Mechanical Analysis (DECMA) has been developed,
which permits assessment of the strain effect on the adsorption process during Cuupd [24].
3.1.2. Pb on Au
Many studies of the surface structure of gold electrodes have been carried out through the upd of
lead (Pbupd ). These studies allow us to assign the different voltammetric peaks of different surfaces to
the deposition processes that occur on low index planes or terraces and steps. The influence of the
crystallographic orientation on the process has been largely discussed [28]. Indeed, Pbupd has been
reported on polycrystalline gold [71], low-index facets, on (1 0 0), (1 1 1), and vicinal facets [43,72],
as well as a large variety of stepped surfaces [73].
Lead and gold atoms have very different sizes (lead is about 20% larger than gold), which favor
the formation of incommensurate ad-layers. Indeed, the analysis of the structure of a lead monolayer
on Au(1 1 1) by X-ray diffraction shows a hexagonal structure [29]. Figure 6 shows a schematic
representation of the structure of a lead monolayer on Au(1 1 1).

Figure 6. Representation of the hexagonal monolayer of lead on Au(1 1 1).

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Surfaces 2019, 2

One of the interesting features of Pbupd on gold is the specific adsorption at different potentials
depending on the surface orientation, as illustrated on a polycrystalline gold electrode (Figure 7).
During the negative potential sweep (from high to low potentials), lead is deposited on (1 1 0) facets in
the potential range from 0.55 to 0.47 V vs. RHE. The deposition on (1 0 0) facets occurs at around 0.43 V
vs. RHE and at 0.38 V vs. RHE on (1 1 1). The peak profile (width, shift) depends on the domain width
or the presence of terraces. This feature permits assessing the coverage of lead on various facets of
gold and also the percentage of each facet on this surface [74].



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Figure 7. Cyclic voltammogram of Au polycrystalline in 0.1 mol L−1 NaOH + 1 mmol L−1 Pb(NO3 )2 ,
recorded at 20 mV s−1 and at 20 ◦ C.

The lead stripping from other gold surfaces with high-index facets (Figure 8) is investigated by
Hamelin et al. [27]. The evolution of the curves through a zone of stereographic triangle (in Figure 8) is
shown. Each stepped surface is also described as TLK (terrace, ledge, kink) notation. In Figure 8a
(zone (1 1 1)-(1 1 0)), the characteristic peak of lead stripping from (1 1 1) terraces decreases gradually,
while a broader peak increases due to the stripping from the steps. An inversion of the nature of the
step and terrace takes place at (3 3 1) facets. It should be taken into account that (1 1 0) is a stepped
surface, as is 2(1 1 1)-(1 1 1). In Figure 8b (zone (1 1 0)-(1 0 0)), the stripping peak of lead from (1 0 0)
terraces decreases progressively, as the facets are closer to (1 1 0). The (2 1 0) facet appears as a turning
point of the zone (1 1 0)-(1 0 0). In Figure 8c (zone (1 0 0)-(1 1 1)), the evolution of the stripping curves
are the same as the previous one, again with a turning point for the (3 1 1) facet.
Pbupd on single-crystals has also been investigated in alkaline media for surface characterization
applications. It was observed that the pH affects the position of the deposition/stripping peaks
because of the different chemical states of lead: Pb2+ in acid media and Pb(OH)3 − in alkaline media.
The charge values associated with the deposition of a monolayer of lead in 0.1 mol L−1 NaOH + 1
mmol L−1 Pb(NO3 )2 media are 444 μC cm−2 for (1 1 1), 340 μC cm−2 for (1 0 0) and 330 μC cm−2 for
(1 1 0) [75], respectively.
Figure 9 shows the cyclic voltammograms of three low-index Au single-crystals in alkaline
media [75]. Different behaviors were observed for each orientation. On Au(1 1 1), a single sharp
deposition peak (0.380 V vs. RHE) was observed, while its associated stripping peak occurs at 0.435 V
vs. RHE. In the case of Au(1 0 0) electrode, two peaks were associated with the Pbupd . The main peak
occurs at 0.430 V vs. RHE and the smaller one at 0.380 V vs. RHE. The two peaks associated with the
dissolution of lead are recorded at 0.385 and 0.475 V vs. RHE. For the Au(1 1 0) electrode, the peaks

77

Surfaces 2019, 2

associated with the deposition and stripping of lead are much broader and less defined than those of
the previous orientations. The process is more sluggish on an Au(1 1 0) surface.

Figure 8. Cyclic voltammograms for the stripping of Pbupd on Au(h k l) electrodes in 10 mmol
L−1 HClO4 + 1 mmol L−1 PbF2 . Au(h k l) surfaces are from the three main zones of the projected
stereographic triangle: (a) (1 1 1)-(1 1 0) zone, (b) (1 1 0)-(1 0 0) zone, and (c) (1 0 0)-(1 1 1) zone.
All the surfaces are identified with Miller index notation and with the step notation. Reprinted from
reference [27]. Copyright (1984) with permission from Elsevier.

78

Surfaces 2019, 2

Figure 9. Cyclic voltammograms of Au(1 0 0), Au(1 1 0), and Au(1 1 1) in 0.1 mol L−1 NaOH + 1
mmol L−1 Pb(NO3 )2 at 50 mV s−1 . Reprinted from reference [75]. Copyright (2004) with permission
from Elsevier.

Pbupd has been reported to characterize bimetallic surface like for Au-Ag, as the Pbupd is also
favorable on silver surfaces [35,76]. This technique can be used to quantify the surface coverage of
gold and silver in an Au-Ag catalyst, based on the dependence of the deposited lead charge density
on the nature of the underlying substrate. Indeed, the coverage of silver on the gold electrode affects
the Pbupd . Two peaks are observed during the stripping of adsorbed lead (Figure 10), the first one
at −0.02 V vs. SHE, which is clearly influenced by the silver coverage (θAg ). The calculations of the
charge densities corresponding to the Pbupd can be made from the cyclic voltammograms. A charge
density of 290 μC cm−2 was obtained for a bare gold electrode, while a value of 385 μC cm−2 was
calculated on Agupd -modified gold. The difference of charge densities for the two surfaces may be due
to the difference in the availability of sites of lead deposition or to the different adsorption energies [4].
The coverage of silver can be estimated with the following equation:
θAg QPb-Ag + θAu QPb-Au = Qnet ,

(3)

where QPb-Ag represents the charge density of Pbupd on silver (385 μC cm−2 ), QPb-Au is the charge
density of lead upd on gold (290 μC cm−2 ), Qnet is the total charge density (subtracted with background
current), and θAu and θAg are the gold and silver surface coverages, respectively (with θAu + θAg = 1).

79

Surfaces 2019, 2

Figure 10. Lead stripping voltammetric curves of Agupd -modified Au with different coverages (θAg ) in
10 mmol L−1 HClO4 + 0.2 mmol L−1 Pb(ClO4 )2 at 20 mV s−1 . Reprinted from reference [76]. Copyright
(2017) with permission from Elsevier.

3.1.3. Probing Gold Nanoparticles Surface by upd
Pbupd appears to be an efficient tool to characterize the surface of gold nanoparticles (AuNPs).
However, contrary to single-crystal surfaces, AuNPs exhibit generally different planes on their surface.
Moreover, synthesis of nanoparticles requires the use of shape-directing agents as surfactants. Pbupd
in alkaline media can be employed to verify the cleanliness of NPs’ surface and oxidize adsorbed
organics on their surface [75]. The surfactant species are oxidized at potential higher than 1.20 V vs.
RHE. The advantage of working in alkaline media is that PbO2 deposition takes place at potentials
lower than 1.10 V vs. RHE, which is lower than the onset potential for the formation of gold oxides.
Gold surface is protected by these species, avoiding its oxidation which can modify its surface, while
organic species are oxidized.
The Pbupd has been performed on AuNPs with different shapes and crystallographic
orientations [74,77–79]. Hebie et al. [74] have clearly shown the shape dependence of Pb
deposition/stripping on AuNPs. Indeed, Figure 11 shows cyclic voltammograms of Pbupd on different
AuNPs—nanorods (AuNRs), nanocubes (AuNCs), and nanospheres (AuNSs). The peak associated
with the lead stripping from (1 0 0) facets is clearly seen for AuNCs. It is well known that cubic NPs
are enclosed by (1 0 0) facets [80]. As the AuNCs studied in this case possess truncated edges, lead
stripping is also observed for (1 1 1) and (1 1 0) planes.

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Surfaces 2019, 2





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Figure 11. Cyclic voltammograms of different AuNPs (nanospheres AuNSs, nanorods AuNRs, and
nanocubes AuNCs) in 0.1 mol L−1 NaOH + 1 mmol L−1 Pb(NO3 )2 recorded at 20 mV s−1 and at 20 ◦ C.

A recent study dealt with the Pbupd on quasi-spherical and faceted AuNPs [81]. These
investigations have revealed a typical split stripping peak associated with Au(1 1 1) facets in the case of
faceted AuNPs, and a single peak in the case of spherical ones. The split peak results from a sluggish
kinetics of lead dissolution [82]. It appears that in the case of small terraces with more steps and defects,
this doublet stripping peak is not observed, contrary to the peak obtained with faceted NPs. The study
suggests that introducing a degree of roughness on (1 1 1) terraces by potential cycles can suppress the
splitting profile. The splitting of the peak associated with the desorption from (1 1 1) terraces can be
used as an indication of the surface crystallinity of polyfaceted AuNPs.
Thallium is also used for upd process on gold nanoparticles. It is known to be used for characterizing
AuNPs with defined shapes (nanocrystals) [83,84]. Tlupd has been employed to check the surface
cleanliness and the crystallinity of these nanocrystals before performing an oxygen reduction reaction.
3.2. Case of Pt
3.2.1. Hupd on Pt
The famous upd process widely studied on platinum is the upd of hydrogen (Hupd ) [85,86]. Hupd is
not strictly an upd process but the adsorption of hydrogen on the surface. The Hupd region acts as a
fingerprint for platinum surfaces. Indeed, the profile of the hydrogen adsorption/desorption region
obtained by cyclic voltammogram permits an indication of the facets, to assess the active surface area
and also to ensure the cleanliness of electrode and cell. Combined with the adsorption of organic
molecule, it permits determining the surface coverage. The underpotential deposition of metals can
give complementary information to quantify the crystallographic orientations [47]. Traditionally,
the Electrochemical Active Surface Area (ECSA) of platinum is determined by the charge corresponding
to the Hupd . Indeed, one platinum atom has the ability to adsorb one hydrogen atom. The charge
associated with Hupd indicates the number of platinum atoms at the surface. The density of charge
associated with the formation of a monolayer of hydrogen depends on the crystallographic orientations,
and the value of 210 μC cm−2 is the average value for polycrystalline platinum [87,88]. CO-stripping
is also used as a method to calculate the ECSA of platinum in acid media [89]. It is important to
notice that in acid media, the number of electrons per site during the adsorption of CO on platinum is
equal to 2.
The development of bimetallic surfaces for catalysis requests to determine ECSA values. In the
case of ruthenium-containing catalysts, Hupd is not suitable because of the overlap of the hydrogen

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Surfaces 2019, 2

and ruthenium oxidation currents [90]. Moreover, hydrogen can be absorbed in the Ru oxide lattice,
leading to more than one monolayer of hydrogen. CO-stripping on bimetallic Pt-Ru surface appears
to be suitable, however, the adsorption modes on both platinum and ruthenium needs to be deeply
elucidated. The mode of adsorption is an important parameter to calculate the charge.
3.2.2. Bi on Pt: A Probe for (1 1 1) Planes
It has been shown that bismuth is spontaneously adsorbed on a platinum surface when the
electrode is simply put in contact with a solution containing a salt of Bi(III) species [44]. When a linear
potential sweep is applied to the electrode, the oxidized and reduced forms of this compound stay
adsorbed at the platinum surface. For the planes (1 1 1), (1 0 0), and (1 1 0), the oxidation peaks are
located at 0.610 V vs. RHE, 0.825 V vs. RHE, and 0.915 V vs. RHE, respectively. Biupd appears as a tool
to highlight the presence of (1 1 1) facets, especially when the surface exhibits a combination of several
planes. It has been proposed that one atom of bismuth covers three atoms of platinum [47]:
Pt3 (1 1 1)-Bi + 2H2 O = Pt3 (1 1 1)-Bi(OH)2 + 2H+ + 2e− .

(4)

In the case of stepped surfaces, the redox peak associated with the adsorption of bismuth on
(1 1 1) terraces sites gives a charge which is proportional to the number of (1 1 1) terrace sites, without
interference from adsorbed bismuth to another site [47]. The adsorption of bismuth suppresses the
Hupd region, which means that H adsorption sites on platinum are fully blocked.
To characterize the (1 1 1) sites belonging to terrace domains, the charge density values of each
surface are obtained by integrating the voltammetric peaks. Figure 12 shows the plots of the charge
underneath the bismuth redox peak (qBi ) versus the calculated terrace charge [47]. The calculation
of the terrace charge depends on how bismuth is adsorbed on the step and terrace sites. Bismuth is
preferentially adsorbed on step sites, so the adsorption takes place on the terraces when all the step
sites have been covered [91]. A linear equation is obtained and can be used to calibrate the (1 1 1)
terrace sites:
(5)
qBi = (0.64 ± 0.02)qt(1 1 1)

Figure 12. Charge density values of the bismuth redox peak vs. the charge associated with the (1 1 1)
sites on the terraces, assuming that one electron is exchanged per site (: Pt(n,n,n−2), ◦: Pt(n+1,n−1,n−1)
electrodes). Reprinted from reference [47]. Copyright (2005) with permission from Elsevier.

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Surfaces 2019, 2

3.2.3. Ge on Pt: A Probe for (1 0 0) Planes
The spontaneous deposition of germanium (Geupd ) acts as a complementary tool to Biupd to
characterize platinum surfaces. It has been proposed that one atom of germanium covers four
atoms of platinum and the equation of the oxidation of adsorbed germanium on platinum can be
written as follows:
Pt4 (1 0 0)-Ge + H2 O = Pt4 (1 0 0)-GeO + 2H+ + 2e− .
(6)
As for Biupd , the oxidation of a spontaneously deposited layer of germanium on a platinum surface
will give information about the crystallographic orientations of the surface [46]. The Hupd is completely
inhibited when a full layer of germanium is deposited (Figure 13). Conversely, the Hupd is partially
inhibited when the surface is not fully covered by germanium. The sharp reversible peaks at 0.540 V vs.
RHE are associated with germanium ad-atoms on the (1 0 0) plane. During continuous cycling between
0.060 and 0.600 V vs. RHE, the voltammetric profile remains stable, meaning that the germanium
species are still adsorbed on the surface. Moreover, the pH value of the electrolyte does not affect the
peaks, suggesting that a germanium oxide is formed.

Figure 13. Voltammograms of a Pt(1 0 0) surface fully covered with Ge (continuous line) and Pt(1 0 0)
blank. Reprinted from Reference [46]. Copyright (1992) with permission from Elsevier.

The (1 1 0) plane does not show a reversible oxidation/reduction process but a quick desorption is
observed. For (1 1 1) planes, the oxidation peak is centered at 0.730 V vs. RHE. However, the oxidation
makes a partial dissolution of the germanium layer.
Germanium appears to be sensitive to (1 0 0) planes. Conversely to bismuth, germanium is not
irreversibly adsorbed on Pt and it is possible to recover the hydrogen region following the experiment.
As for the study related to (1 1 1) terraces toward bismuth adsorption, adsorbed germanium can
be used to characterize (1 0 0) terraces sites of stepped platinum surfaces [47,48].
A linear equation can also be obtained to calibrate terrace sites, as for bismuth on (1 1 1) terraces:
qGe = (0.56 ± 0.03)qt(1 0 0).

(7)

3.2.4. Cu on Pt: A Tool to Characterize Bimetallic Surfaces
As it was reported in this mini review, Cuupd is so far, one of the most studied systems on low-index
single-crystals as well as on stepped surfaces [92]. As for gold surfaces, Cuupd on platinum is strongly

83

Surfaces 2019, 2

dependent of the anions present in the electrolyte [17,93]. However, there are only few investigations
about its application as a tool for surface characterization. Cuupd is a suitable technique to calculate
electrochemical active surface area (ECSA) of Pt and alloys containing Pt. Indeed, Cuupd has been
proposed to characterize the bimetallic Pt-Ru surfaces. The integration of the copper stripping peak
area permits the estimation of ECSA. It has been shown that Cuupd is an accurate tool to estimate
ECSA of each platinum and ruthenium in Pt–Ru electrodes. The coverage of ruthenium can also
be determined by this method. Cuupd has also been reported for the calculation of ECSA of Pt-Pd
alloys [94]. Indeed, copper is a suitable ad-atom to deposit because its adsorption occurs at a potential
region where no other Faradaic reaction takes place on both platinum and palladium.
3.2.5. Probing the Surface of Pt Nanoparticles
Biupd and Geupd were also investigated on platinum nanoparticles to probe their surface [47,95–97].
In the case of nanoparticles containing (1 1 1) and (1 0 0) planes, the fraction of each domain is obtained
by the ratio of the Faradaic charge of germanium (Qf,Ge ) or bismuth (Qf,Bi ) with the charge associated
with the hydrogen desorption process (Qf,H ). The characterization of PtNPs prepared in different
conditions has been reported [98]. The use of Geupd and Biupd permitted determining the percentage of
(1 0 0) and (1 1 1) facets, respectively. The fraction of (1 1 0) sites has been estimated. Indeed, the peak
of Hupd associated with (1 1 0) facets decreases, while the difference from 100% of the sum of (1 1 1)
and (1 0 0) facets decreases.
3.3. Case of Pd
Cu on Pd
There are only a few studies about upd of metals on palladium surfaces [11,40]. This is due to
the difficulty for metal monolayers to be adsorbed on the double-layer region, which is narrow [99].
Palladium also has the ability to absorb hydrogen in its bulk, this process may mask the upd one.
It should be noted that there is no study about upd of metals to characterize palladium surfaces. For
example, Cu grows epitaxially with the Pd substrate.
Cuupd has been studied on polycrystalline and single-crystalline palladium electrodes [100].
This study has shown that, under adequate conditions, Cu adsorption can be restricted to a potential
range where no other Faradaic reaction takes place. Cuupd on palladium surfaces occurs at potentials
lower than the onset potential of oxides formation, and hydrogen absorption is inhibited by changing
the pH of the electrolyte. Figure 14 shows the cyclic voltammograms of upd/stripping on palladium
(1 0 0) and (1 1 1) single-crystal and polycrystalline palladium. It is clearly seen that the single-crystal
electrode present less complex voltammograms than polycrystalline palladium. On Pd(1 1 1), the upd
and desorption processes give a single sharp peak (A3 /D3 ) preceded by a small shoulder. On Pd(1 0 0),
two peaks are observed (A2 /D2 and A6 /D6 ) and a small shoulder noted A1 /D1 . The complexity of the
voltammogram obtained on polycrystalline palladium is due to the variety of crystal planes.

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Figure 14. Cyclic voltammograms of Pd single-crystal electrodes in 0.5 mol L−1 NaClO4 + 0.01 mol L−1
HClO4 + 1 mmol L−1 Cu2+ , at 10 mV s−1 : (a) Pd(1 0 0), (b) Pd(1 1 1), and (c) polycrystalline Pd.
Reprinted from reference [100]. Copyright (1988) with permission from Elsevier.

4. Conclusions
Investigations on upd of transition metals at noble metal surfaces have been increasing over these
last two decades. It appears to be a powerful electrochemical tool to probe and characterize the working
electrode surface. Indeed, upd techniques have been reported to calculate electrochemical active surface
area, especially in the case of bimetallic surface, and to probe the crystallographic orientations among
the surface of single-crystals and nanoparticles. Pbupd is mostly used to characterize the orientation
of gold nanoparticles and surfaces. In the case of platinum surfaces, the probe of crystallographic
orientation is made with bismuth and germanium adsorption. Cuupd is one of the most studied upd
systems that appear to be useful for characterizing the surface of bimetallic materials.
Author Contributions: All authors contributed equally to this paper.
Funding: The authors acknowledge the Region Nouvelle Aquitaine for the financial support.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Nitrogen-Doped Ordered Mesoporous Carbons
Supported Co3O4 Composite as a Bifunctional
Oxygen Electrode Catalyst
Jing Wang 1 , Shuwei Zhang 1 , Haihong Zhong 1,2 , Nicolas Alonso-Vante 2 , Dianqing Li 1 ,
Pinggui Tang 1 and Yongjun Feng 1, *
1

2

*

State Key Laboratory of Chemical Resource Engineering, Beijing Engineering Center for Hierarchical
Catalysts, Beijing University of Chemical Technology, No. 15 Beisanhuan East Road, Beijing 100029, China;
2017200993@buct.edu.cn (J.W.); 2016201069@buct.edu.cn (S.Z.); 2014200922@grad.buct.edu.cn (H.Z.);
lidq@mail.buct.edu.cn (D.L.); tangpg@mail.buct.edu.cn (P.T.)
IC2MP, UMR-CNRS 7285, University of Poitiers, F-86022 Poitiers CEDEX, France;
nicolas.alonso.vante@univ-poitiers.fr
Correspondence: yjfeng@mail.buct.edu.cn; Tel.: +86-10-6442-5385

Received: 20 February 2019; Accepted: 26 March 2019; Published: 29 March 2019

Abstract: It is increasingly useful to develop bifunctional catalysts for oxygen reduction and
oxygen evolution reaction (ORR and OER) for fuel cells, metal-air rechargeable batteries,
and unitized regenerative cells. Here, based on the excellent conductivity and stability of
ordered mesoporous carbons, and the best ORR and OER activity of Co3 O4 , the composite
Co3 O4 /N-HNMK-3 was designed and manufactured by means of a solvothermal method, using
ordered N-doped mesoporous carbon (N-HNMK-3) as substrate, and then the bifunctional
electrocatalytic performance corresponding to ORR, OER in alkaline media was carefully
investigated. The results showed that Co3 O4 /N-HNMK-3 composite, a non-precious metal centered
electrocatalyst, displayed excellent ORR performance (activity, selectivity, and stability) close
to that of commercial 20 wt.% Pt/C and a promising OER activity near 20 wt.% RuO2 /C.
The outstanding bifunctional activities of Co3 O4 /N-HNMK-3 was assessed with the lowest E
value of 0.86 V (EOER,10 mA cm−2 -EORR,−3 mA cm−2 ) with respect to the two commercial precious
metal-based electrocatalysts.
Keywords: Ordered mesoporous carbon; nitrogen doping; cobalt-based electrocatalyst; bifunctional
oxygen electrode; solvothermal method

1. Introduction
Hydrogen is deemed as the best way for renewable sources energy-storage to replace
petroleum-based energy resources because of its high calorific value, abundant reserve, and near
zero release. Yet, both the oxygen reduction reaction (ORR) and oxygen evolution reaction (OER)
individually play crucial roles in fuel cells for the practical application of hydrogen, and water
electrolysis for the high-efficiency production of high-purity hydrogen [1–4]. Recently, unitized
regenerative cells (URCs), integrating water electrolysis and fuel cell in one setup, have attracted
increasing attention due to potential applications in aviation, submarines and transportation as a
mobile power source. This issue remains key to exploring high performance and low cost. In terms
of electrocatalytic activity, Pt/C and Ru-/Ir oxides are generally considered to be the most advanced
electrochemical oxygen catalysts for the ORR and OER processes, respectively. However, high cost,
low methanol tolerance, and poor stability continue to hamper practical marketing development [5].
Therefore, it remains a big challenge to develop promising bifunctional oxygen electrodes for the
development of URCs.
Surfaces 2019, 2, 229–240; doi:10.3390/surfaces2020018

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Among non-precious metal electrocatalysts being developed such as cobalt-based chalcogenides,
CoSe2 , CoS, and Co3 O4 , show outstanding performance toward oxygen electrode reactions including
activity, selectivity, and durability in alkaline media [6–9]. However, cobalt-based chalcogenide
electrocatalysts have a low electrical conductivity and per se a negative effect on the performance of the
electrocatalyst. On the other hand, ordered mesoporous carbon as a conductive substrate has attracted
more attention because of large surface, high conductivity and long durability. [10–13]. In addition, the
doping heteroatom, such as nitrogen [14,15], phosphor [16,17], sulfur [18] and boron [19], can further
improve electrocatalytic performance towards oxygen electrode reactions due to new active species.
In a previous report, we developed N-doped ordered mesoporous carbons (CMK-3) with high ORR
activity, long stability, and outstanding methanol tolerance in 0.1 M KOH [20]. Therefore, it is possible
to explore high-performance, low-cost electrocatalysts using N-doped ordered mesoporous carbon as
a substrate to support cobalt-based chalcogenides.
Here, we manufactured the N-doped ordered mesoporous carbon substrate (N-HNMK-3) for
Co3 O4 composite (Co3 O4 /N-HNMK-3) by means of a solvothermal route using the home-made
ordered mesoporous carbon as carbon source, melamine as a nitrogen source, cobalt acetate as a cobalt
source, and urea as the precipitating agent. Figure 1 describes the corresponding synthesis process.
In addition, we carefully examined the bifunctional electrocatalytic performance of ORR and OER
processes in alkaline media using the rotating disc electrode technique. The compound, based on
non-precious metals, shows a promising electrocatalytic performance.

Figure 1. Schematic on preparation process of Co3 O4 /N-HNMK-3.

2. Materials and Methods
2.1. Chemicals
Cobalt acetate (Co(CH3 COO)2 ·6H2 O), urea (CO(NH2 )2 ), HNO3 (6M), melamine (C3 H6 N6 ), and
ethanol (C2 H5 OH) were A.R. grade and were directly used as received from Beijing Chemical Co. Ltd
(Beijing, China). CMK-3 was synthesized as reported previously [20].
2.2. Synthesis of Nitrogen-Doped Ordered Mesoporous Carbon (N-HNMK-3)
Nitrogen-doped ordered mesoporous carbon was fabricated by post-treated nitrogen doping
method using ordered mesoporous carbon (CMK-3, surface area, 1247.5 m2 /g; pore volume, 1.5 mL/g)
as carbon source, and melamine as nitrogen source. To obtain HNMK-3, CMK-3 (0.5002 g) was
dispersed in 6.0 M HNO3 (50.0 mL) under vigorous stirring at 70 ◦ C for 3 h, and the resulting
suspension was cleaned to neutral pH using deionized water. Thereafter, 0.5000 g HNMK-3 and
2.5000 g melamine were then mixed during grinding to form the homogeneous mixture. Finally, the
obtained mixture was calcined in a nitrogen atmosphere at 900 ◦ C for 4 h to produce nitrogen-doped
ordered mesoporous carbon (N-HNMK-3).

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2.3. Synthesis of Co3 O4 /N-HNMK-3 Electrocatalyst
Co3 O4 /N-HNMK-3 electrocatalyst was carefully generated via the solvothermal route using
N-HNMK-3 as conductive support; Co(CH3 COO)2 as cobalt source, and CO(NH2 )2 as the precipitating
agent. Typically, 0.0175 g of the as-prepared N-HNMK-3 was dispersed in ethanol (40.0 mL) by
ultrasound for 2h to form the suspension A. Simultaneously, 1.50 g Co(Ac)2 (6.0 mmol) was added in
deionized water (40.0 mL) and ultra-sonicated for 20 min to form solution B. Then suspension A and
solution B were mixed and ultrasonicated again to produce another suspension. At this point 0.72 g
CO(NH)2 was added under vigorous stirring at 80 ◦ C for 24 h to produce a suspension, which was
heated at 150 ◦ C in a Teflon-lined stainless-steel autoclave for 3 h. The cooled product was centrifuged
and washed with deionized water and ethanol three cycles, and dried at 60 ◦ C overnight. The collected
final product was Co3 O4 /N-HNMK-3. In addition, Co3 O4 , as a reference, was prepared in the same
procedure without N-HNMK-3.
2.4. Characterization
The materials phase was characterized by Rigaku UItima III XRD (Rigaku Corporation, Tokyo,
Japan) with Cu Kα radiation (λ = 0.154 nm) at 5 min−1 per 2θ. The morphologies were analyzed by
Zeiss Supra 55 SEM(Carl Zeiss Jena, Oberkochen, German), a JEOL JEM-2012 TEM (JEOL, Tokyo,
Japan), and HRTEM (JEOL, Tokyo, Japan) with a line resolution of 0.19 nm was used. Raman spectra
were examined via a Nanofinder 3.0 Raman spectrometer with a He-Ne laser beam of 532 nm. Binding
energies of chemical species were collected by XPS using VG ESCALAB 2201 XL spectrometer (Thermo
VG Scientific, West Sussex, England).
2.5. Electrochemical Measurements
All electrochemical tests were performed in a standard three-electrode system on the Chen Hua
electrochemical workstation (CH Instruments Ins., Shanghai, China) at 25 ◦ C. Here, a saturated calomel
electrode and a Pt-wire were employed as the reference and counter electrodes, respectively. In this
work, all potentials were referred to a reversible hydrogen electrode (RHE): ERHE = ESCE + 0.99 V
(0.1 M KOH); and ERHE = ESCE + 1.06 V (1.0 M KOH).
Prior to catalyst deposition, the GC disk was polished with γ-alumina (5A), and successively
ultra-sonically treated in water, and then ethanol for one min in each solvent. The ink was prepared
by mixing the catalyst powder (6.7 mg), Nafion (5 wt.%) (40 μL), ethanol (300 μL), deionized water
(1160 μL) and ultra-sonicated for 40 min. 4.2 μL of the ink was deposited onto the GC disk (4.0 mm,
diameter) with a catalyst mass loading of 150 μg cm−2 . The electrolyte was degassed with argon and
oxygen for 30 min before cyclic voltammetry (CV) tests. Also, two benchmarks: 20 wt.% Pt/C (Johnson
Matthey) and RuO2 /C were prepared as references with a corresponding loading of 60 μg cm−2 for Pt,
and RuO2 .
In the ORR test, the LSV curves were recorded in O2 -saturated 0.1 M KOH by scanning the disc
potential vs. RHE from 0.95 V to 0.15 V at 5 mVs−1 with the electrode rotated at 2500, 2025, 1600,1225,
900, 625, and 400 rpm, and analysis performed using the Koutecky-Levich (K-L) equation:

where:

1
1
1
1
1
= + = +
1
j
jk
jd
jk
Bω 2

(1)

B = 0.62nFC0 D02/3 v−1/6

(2)

jk = nFkC0

(3)

j, jk , jd are the corresponding measured, kinetic and diffusion limiting current densities, respectively [5];
C0 and D0 were the saturated concentration, and the diffusion coefficient of O2 (1.14 × 10−6 mol cm−3
and 1.73 × 10−5 cm s−1 ) in 0.1 M KOH, v was the kinetic viscosity coefficient (0.01 cm2 s−1 ).

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In the RRDE test, the HO2 − % and the number of electrons transferred(n) was calculated by the
following two equations:
IR /N
HO2− % = 200 ∗
(4)
ID + IR /N
n = 4∗

ID
ID + IR /N

(5)

where ID , IR and N is disk current, ring current and collection efficiency (0.424), respectively.
For the OER tests, measurements were carried out in O2 -saturated 1.0 M KOH with a rotating
speed of 1600 rpm at a scan rate of 5 mV s−1 from 1.0 V to 1.65V vs. RHE.
3. Results and Discussion
3.1. Structure, Morphology, and Chemical Composition
Figure 2 shows the powder PXRD patterns of N-HNMK-3, Co3 O4 , and Co3 O4 /N-HNMK-3.
Compared to CMK-3, the peak of N-HNMK-3 samples at 23.8◦ /2θ shifts to a lower angle with a higher
degree of graphitization, which could be the result of a change in electron distribution following the
incorporation of the heteroatom on the surface of carbon materials [21]. Based on the Bragg equation,
a series of Bragg reflection peaks in Co3 O4 and Co3 O4 /N-HNMK-3 samples correspond to (111), (220),
(311), (400), (511), and (440) crystal faces, that match well with the cubic spinel phase Co3 O4 (ICDD
PDF card No. 42-1467). This means that the urea-assisted solvothermal method is an available way to
prepare the Co3 O4 /N-HNMK-3 composite.

Figure 2. Powder XRD patterns of N-HNMK-3, Co3 O4 , and Co3 O4 /N-HNMK-3 composite.

Figure 3 displays the morphologies and microstructures of Co3 O4 and Co3 O4 /N-HNMK-3
samples. The Co3 O4 sample shows different rod-shape in size, Figure 3a. In comparison, the Co3 O4
particles, in Co3 O4 /N-HNMK-3 composite, are well-dispersed and supported on N-HNMK-3 substrate,
Figure 3b,c. In addition, the average particle size for Co3 O4 was 20–30 nm, Figure 3c. The HRTEM
image of Co3 O4 reveals three lattice distances such as 0.286, 0.244 and 0.202 nm, which individually
correspond to (220), (311) and (400) Co3 O4 crystal planes, Figure 3d.

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Surfaces 2019, 2

Figure 3. (a) The SEM images of Co3 O4 , (b) Co3 O4 /N-HNMK-3. (c) TEM and (d) HRTEM images of
Co3 O4 /N-HNMK-3.

In addition, Figure 4 shows the Raman spectra of N-HNMK-3, Co3 O4 , and Co3 O4 /N-HNMK-3
samples as well as the corresponding fitting profiles of N-HNMK-3 and Co3 O4 /N-HNMK-3. One
observes two characteristic bands for N-HNMK-3, and Co3 O4 /N-HNMK-3 centered at 1358 (D peak)
and 1590 (G peak) cm−1 . The former peak (D peak) refers to sp3 -like C atoms defect sites, whereas
the latter (G peak) refers to sp2 C=C stretching mode [22]. Remarkably, other peaks were detected at
192, 468, 516 and 677 cm−1 , assessing the presence of Co3 O4 [23]. The intensity ratio of two bands
(ID /IG ) illustrates the disordered degree of carbonization. A ratio of 2.96 for N-HNMK-3, and 2.78
for Co3 O4 /N-HNMK-3 suggest that the incorporation of Co3 O4 decreases the disordered degree,
increasing the conductivity of the substrate.

Figure 4. (a) Raman spectra of N-HNMK-3, Co3 O4 and Co3 O4 /N-HNMK-3. (b) The corresponding
Raman fitting profiles of N-HNMK-3 and Co3 O4 /N-HNMK-3.

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Surfaces 2019, 2

Surface elemental properties and electronic valence states were further characterized by XPS.
A survey spectrum of Co3 O4 /HNMK-3 displays the existence of C 1s, N 1s, O 1s, Co 2p signals at
284.5, 399.9, 530.5, 780.1 and 796.2 eV, respectively, as shown in Figure 5a. The corresponding atomic
percentage of every element of different samples, N-HNMK-3 has a similar nitrogen content with
Co3 O4 /HNMK-3, as summarized in Table 1. In the C1s high-resolution XPS spectrum, shown in
Figure 5b, the prominent peak centered at 284.6 eV corresponds to sp2 -hybridised carbon atoms in
graphitic structure. The series of fitted peaks at 285.7, 288.3, and 288.6 eV are attributed to C=N, O-C=O
and C-N, respectively. The XPS spectra of C=N and C-N further assess the presence of N-dopant atoms.
As to Co2p spectrum, Figure 5c there are two peaks centered at 779.7 and 795.1 eV attributed to Co
2p3/2 and 2p1/2 signals, with an energy separation of 15.4 eV, evidencing the presence of Co3 O4 in
the composites [24]. Two nitrogen species detected in the N 1s high resolution, shown in Figure 5d, at
~398.6 eV and 400.2 eV are assigned to pyridine and Pyrrole-N. Previous results already demonstrated
that pyridine N is an active site for the ORR [25]. The relative content of Pyrrole-N in Co3 O4 /HNMK-3
sample is 70.2%, which is higher than that of Pyridine-N (29.8%). It’s worth noting that N-HNMK-3
and Co3 O4 /HNMK-3 samples have similar N content.

Figure 5. (a) XPS survey scan of samples: N-HNMK-3, Co3 O4 and Co3 O4 /N-HNMK-3; XPS spectra of
(b) C 1s; (c) Co 2p; (d) N 1s of Co3 O4 /N-HNMK-3.
Table 1. The element amount (at. %) in N-HNMK-3, Co3 O4 and Co3 O4 /N-HNMK-3 samples.
Samples

C

N

N-HNMK-3
Co3 O4
Co3 O4 /N-HNMK-3

88.8

2.6

60.6

2.2

Co

O

Pyridinic-N

Pyrrolic-N

31.2

68.8

42.8
9.4

8.6
57.2
27.8

29.8

70.2

3.2. Electrocatalytic ORR Performance
The ORR performance of Co3 O4 /N-HNMK-3 composite was determined by RDE technique, and
compared to two commercial reference materials: Pt/C and RuO2 /C. Figure 6 displays the activity
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Surfaces 2019, 2

towards ORR in 0.1 M KOH electrolyte. Specifically, Figure 6a,b show the ORR activity difference
among five samples: N-HMNK-3, physical mixture of N-HMNK-3 and Co3 O4 , Co3 O4 /N-HMNK-3
composite, Pt/C, and RuO2 /C. Herein, it is observed that the Co3 O4 /N-HMNK-3 composite has a
much higher ORR activity compared to N-HMNK-3, the physical mixture, and RuO2 /C, and close
to Pt/C. Interestingly, the Co3 O4 /N-HNMK-3 composite shows much better ORR activity than the
mechanical mixture (Co3 O4 +N-HNMK-3). This difference in the ORR activity results from the possible
synergistic effect between two compositions, and from the high dispersion of Co3 O4 on the N-HNMK-3
substrate. In addition, cyclic voltammetry measurement, shown in Figure 6c, was conducted to gain
deeper insight into the ORR process. In the O2 -saturated electrolyte there is a cathodic peak centered
at 0.75 V, which is absent in the Ar-saturated electrolyte. The onset potential of Co3 O4 /N-HMNK was
detected at 0.90 V vs. RHE. The ORR hydrodynamics on Co3 O4 /N-HNMK-3 between 400 rpm to 2500
rpm is shown in Figure 6d with well-defined sigmoidal curves. The electron transfers number (n)
was estimated through the slopes of fitted K-L plots, see the inset in Figure 6d. Within the applied
electrode potential from 0.2 to 0.5 V, the estimated transfer number (n) per O2 molecule was ca. 3.94,
suggesting a 4-electron transfer pathway. Meanwhile, we used the RRDE to calculate the hydrogen
peroxide and electron transfer number in Figure S1 (See Supporting Information). The low hydrogen
peroxide below 15% and transfer number beyond ca. 3.75 in the range of 0.1–0.75 V suggested a close
4-electron transfer pathway.

Figure 6. (a) ORR polarization curves of N-HNMK-3, Co3 O4 +N-HNMK-3 (physical mixture),
Co3 O4 /N-HNMK-3, Pt/C, and RuO2 /C at 1600 rpm in O2 -saturated 0.1M KOH; (b) the limited
current density at 0.50 V/RHE and half-wave potential from (a); (c) CV curves of Co3 O4 /N-HNMK-3
in Ar- and O2 -saturated 0.1M KOH; (d) ORR polarization curves of Co3 O4 /N-HNMK-3 at various
rotation rates, the inset shows the corresponding K-L plot of Co3 O4 /N-HNMK-3.

In addition to activity, the selectivity (tolerance to small organics) and the stability are two other
key parameters for a high performance electrocatalyst. Figure 7a depicts the electrocatalytic selectivity
in the presence and the absence of 3.0 M methanol in 0.1 M KOH electrolyte. After the addition of 3 M
methanol, no significant change was observed in the polarization curves in Ar- and/or O2 -saturated

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Surfaces 2019, 2

electrolyte, see the inset in Figure 7a. In O2 -saturated electrolyte Co3 O4 /N-HNMK-3 showed a
complete tolerance to methanol. In contrast, as it is well established, Pt/C has no selectivity towards the
ORR in presence of methanol, since this material develops a mixed potential due to methanol oxidation
and oxygen reduction. Furthermore, the stability test, as demonstrated by chronoamperometry,
Figure 7b, was conducted at 0.70 V/RHE, and revealed that, after 18000 s, the ORR activity of
Co3 O4 /N-HNMK-3 dropped 18.2% against 23.9% on Pt/C. Meanwhile, CVs with multiple cycles in
oxygen- saturated solution for ORR were shown in Figure S2 (See Supporting Information), after the
accelerated aging test, the reduction potential only shifts 36 mV to a negative region for 1500 cycles.

Figure 7. (a) CV and LSV curves of Co3 O4 /N-HNMK-3 in O2 - saturated with or without 3.0M methanol
in 0.1 M KOH; (b) The durability test of Co3 O4 /N-HNMK-3 and 20 wt.% Pt/C at 0.70 V in O2 -saturated
0.1 M KOH at 1600 rpm.

3.3. Electrocatalytic OER Performance
The OER plays a crucial role in water electrolysis and metal-air batteries’ systems [26]. In
particular, high-performance OER electrocatalysts aim at reducing energy costs to produce high purity
hydrogen and promote the development of URCs. Figure 8 shows the activity of Co3 O4 /N-HNMK-3
compared to N-HNMK-3, physical mixture of Co3 O4 + N-HNMK-3, Pt/C, and RuO2 /C in 1.0 M
KOH. Apart from RuO2 /C, the Co3 O4 /N-HNMK-3 has the highest OER activity with a higher current
density of 19.56 mA cm−2 at 1.65 V, lower overpotential (365 mV) at the current density of 10 mA cm−2 ,
a value close to the best OER electrocatalyst RuO2 /C. The Tafel plots for the OER kinetics are shown
in Figure 8b. The slope of 93 mV dec−1 for Co3 O4 /N-HNMK-3 is higher than that of 74 mV dec−1
for RuO2 /C, and lower than that of the physical mixture. In addition, the OER stability test, via CV,
performed on Co3 O4 /N-HNMK-3 is shown in Figure 8c,d. No significant variation was observed after
1500 cycles, and to some extent, Co3 O4 /N-HNMK-3 showed a ca. 4.1% increase in the overpotential at
10 mA cm−2 after 1500 cycles. The results confirm that Co3 O4 /N-HNMK-3 is quite durable for the
OER in 1.0 M KOH electrolyte.

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Surfaces 2019, 2

Figure 8. (a) LSV curves of N-HNMK-3, Co3 O4 + N-HNMK-3 (physical mixture), Co3 O4 /N-HNMK-3,
20 wt.% Pt/C, and 20 wt.% RuO2 /C towards the OER in O2 -saturated 1.0 M KOH at 5 mV s−1 and
1600 rpm; (b) Tafel plots of each samples derived from curves in (a); (c) CVs of Co3 O4 /N-HNMK-3
from 1.05 V to 1.70 V at 200 mV s−1 in O2 -saturated 1.0 M KOH from 1st cycle to 1500th cycle; (d) LSV
curves of Co3 O4 /N-HNMK-3 from 1.30 V to 1.70 V at 5 mV s−1 , the inset presents the variation of η 10
at different CV cycles.

3.4. The Bifunctional ORR/OER Performance
The overvoltage between ORR and OER is an important parameter to evaluate the bifunctional
electrocatalytic activity of the material. The corresponding data are summarized in Table 2. The ΔE
value is defined as the difference in potential between the OER current density of 10.0 mA cm−2
and the ORR current density of −3.0 mA cm−2 . Among the listed catalysts [27–32] reported in the
literature, Co3 O4 /N-HNMK-3 shows the best and promising bifunctional activity with a minimum ΔE
value, which is attributed to the interaction developed between Co3 O4 species onto N-doped ordered
mesoporous carbons.

Δ

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Surfaces 2019, 2

Table 2. The oxygen bifunctional electrocatalytic performance of catalysts prepared in this work
compared with other bifunctional catalysts from the literature.
Catalysts

E vs. RHE
at j = −3 mA cm−2

E vs. RHE
at j = 10 mA cm−2

N-HNMK-3
Co3 O4 +N-HNMK-3
Co3 O4 /N-HNMK-3
20% Pt/C
20% RuO2 /C
Co3 O4 /Co2 MnO4
MnCo2 O4 /Nanocarbon
NiCo2 O4 /Graphene
CoFe2 O4 /Biocarbon
O-NiCoFe-LDH
Co(OH)2 CO3 /C

0.69
0.56
0.73
0.81
0.44
0.68
0.79
0.60
0.69
0.62
0.81

2.13
1.69
1.59
2.07
1.55
1.77
1.75
1.68
1.67
1.67
1.73

ΔE
(EOER@10 mA cm−2 )-(EORR@−3 mA cm−2 )
1.44
1.13
0.86
1.26
1.11
1.09
0.96
1.08
0.98
1.05
0.92

Refs.
This work
This work
This work
This work
This work
27
28
29
30
31
32

4. Conclusions
In this paper, a nitrogen-doped ordered mesoporous carbons supported Co3 O4 composite
(Co3 O4 /N-HNMK-3) was prepared through a solvothermal method. The Co3 O4 /N-HNMK-3
composite showed a much higher ORR performance, which was close to that of the commercial
Pt/C in 0.1 M KOH. An excellent OER performance near commercial RuO2 /C in 1.0 M KOH was also
obtained. Interestingly, Co3 O4 /N-HNMK-3 has outstanding bifunctional activity with a much lower
E value between EOER,10 mA cm−2 -EORR,−3 mA cm−2 as compared to Pt/C and RuO2 /C in alkaline
medium. Undoubtedly, the synthesis method developed here to obtain the oxygen electrode based on
Co3 O4 /N-HNMK-3, with non-precious metals will promote the practical development of fuel cells,
water electrolysis and unitized regenerative cells.
Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/2/2/18/s1,
Figure S1. (a) Ring (top) and disk (down) current density from RRDE measurements of Co3 O4 /N-HNMK-3
samples after annealing at different temperature in O2 -saturated 0.1 M KOH at 25 ◦ C with a sweep rate of
5 mV s−1 at a rotating speed of 1600 rpm; (b) Molar fraction of HO2 − formation and electron transfer number n
from rotating ring-disk electrode (RRDE) curves in (a). Figure S2. (a) CVs of Co3 O4 /N-HNMK-3 from 0.1 V to
1.0 V at 100 mV s−1 in O2 -saturated 0.1 M KOH from 1st cycle to 1500th cycle; (b) LSVs of Co3 O4 /N-HNMK-3
from 0.1 V to 1.0 V at 5 mV s−1 in O2 -saturated 0.1 M KOH from 1st cycle to 1500th cycle.
Author Contributions: Methodology, N.A.-V. and Y.F.; formal analysis, S.Z., G.Y. and H.Z.; investigation, J.W.;
writing—original draft preparation, J.W. and S.Z.; writing—review and editing, N.A.-V., D.L., P.T., and Y.F.;
supervision, Y.F.; project administration, Y.F.
Acknowledgments: This work is supported by the National Natural Science Foundation of China, and the
Fundamental Research Funds for the Central Universities (12060093063). H.H. Zhong specially thanks the
financial support from China Scholarship Council (201806880040).
Conflicts of Interest: The authors declare no conflict of interest.

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102

Article

An Electrochemical Route for Hot Alkaline
Blackening of Steel: A Nitrite Free Approach
Maximilian Eckl, Steve Zaubitzer, Carsten Köntje, Attila Farkas, Ludwig A. Kibler and
Timo Jacob *
Institute of Electrochemistry, Ulm University, Albert-Einstein-Allee 47, 89081 Ulm, Germany;
maximilian.eckl@uni-ulm.de (M.E.); steve.zaubitzer@uni-ulm.de (S.Z.); koentje@gmx.net (C.K.);
attila.farkas@uni-ulm.de (A.F.); ludwig.kibler@uni-ulm.de (L.A.K.)
* Correspondence: timo.jacob@uni-ulm.de; Tel.: +49-731-50-25401
Received: 23 February 2019; Accepted: 26 March 2019; Published: 29 March 2019

Abstract: Blackening belongs to the predominant technological processes in preserving steel surfaces
from corrosion by generating a protective magnetite overlayer. In place of the commonly used
dipping-procedure into nitrite-containing blackening baths at boiling temperatures that are far above
100 ◦ C, here we describe a more environmentally friendly electrochemical route that operates at
temperatures, even below 100 ◦ C. After an investigation of the electrochemical behavior of steel
samples in alkaline solutions at various temperatures, the customarily required bath temperature
of more than 130 ◦ C could be significantly lowered to about 80 ◦ C by applying a DC voltage that
leads to an electrode potential of 0.5−0.6 V vs. Pt. Thus, it was possible to eliminate the use of
hazardous sodium nitrite economically and in an optimum way. Electrochemical quantification of
the corrosion behavior of steel surfaces that were in contact with 0.1 M KCl solution was carried
out by linear sweep voltammetry and by Tafel slope analysis. When comparing these data, even
the corrosion rates of conventional blackened surfaces are of the same magnitude as a blank steel
surface. This proves that magnetite overlayers represent rather poor protective layers in the absence
of additional sealing. Moreover, cyclic voltammetry (CV), atomic force microscopy (AFM), scanning
electron microscopy (SEM) and auger electron spectroscopy (AES) characterized the electrochemically
blackened steel surfaces.
Keywords: Surface Modification; Blackening of Steel; Magnetite; Corrosion Protection;
Auger-Electron Spectroscopy

1. Introduction
The technology of blackening is of utmost importance for surface finishing and the refinement of
various types of steel. Chemical oxidation of an iron surface forms a tenebrous decorative passivation
layer, which, in particular, can inhibit corrosion processes [1]. The resulting black-bluish coloration
represents a convenient quality feature for steel components in a wide area of applications. The
traditional hot alkaline blackening process for steel parts, in the following simply called blackening,
involves five consecutive steps [2–4]: (1) degreasing and the removal of oils, (2) first water flushing,
(3) actual blackening, where the specimens are immersed into a boiling sodium hydroxide solution
that contains sodium nitrite as oxidizing agent and some other additives, (4) second water flushing
for the removal of blackening bath components, and (5) surface treatment of the blackened steel by a
water-resistant sealant. As a result, the blackening process leads to the oxidation of the surface iron
atoms by forming black iron oxide (magnetite Fe3 O4 ) [5–7]. In a similar manner to the well-known iron
phosphating process [8], corrosion protection by blackening is not just achieved by simple deposition
of an additional layer, but by a chemical transformation of the substrate’s surface itself, resulting in a
so-called conversion coating [2,9]. The blackened surfaces are usually sealed by a waterproof layer
Surfaces 2019, 2, 0; doi:10.3390/surfaces2020017

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Surfaces 2019, 2, 0

(e.g., oil, etc.) to retard the initial stages of corrosion at defects and pores that are formed during the
chemical oxidation process [5,6].
The formation of the magnetite surface layer is generally thought to involve the following two
steps [6,10,11]: First, iron surface atoms are oxidized to iron(II) hydroxide (Fe(OH)2 , see Equation
(1)) by nitrite (NO2– ), which acts as oxidizing reagent. Fe(OH)2 is then further oxidized to magnetite
(Equation (2)) via the Schikorr reaction [12,13], causing the black surface burning [3,5,6,14].
Fesurf + 2 OH− (aq) → Fe(OH)2,surf + 2 e−

(1)

3Fe(OH)2,surf → Fe3 O4,surf + H2(g) + 2 H2 O(l)

(2)

Although the Schikorr reaction is spontaneous, i.e., thermodynamically favored, even at room
temperature, the formation of magnetite is usually kinetically hindered at temperatures below
100 ◦ C [4].
The simplicity and robustness of the blackening procedure is highly esteemed in many industrial
sectors. However, hazardous and toxic sodium nitrite and the high energy costs represent major
drawbacks. Hence, the development of a more environmentally friendly and efficient process would
be highly desirable and promising for several branches. Attempts to form magnetite layers through
an electrochemical reaction have been reported earlier for steel in hot (70−98 ◦ C) ammonium nitrate
solution [14,15]. However, this solution is, also considered to be unsafe. Alternatively, a hydrothermal
synthesis of magnetite has been proposed [3,4,16], requiring an autoclave and likewise relatively high
temperatures (100−250 ◦ C), in comparison to the conventional process.
It is obvious to alternatively study the electrochemical route for the Schikorr-reaction (Equation
(3)), which means that electrochemical blackening might be feasible without the use of harmful
oxidants [5,6,17–20]:
3 Fe(OH)2,surf → Fe3 O4,surf + 2 H2 O(l) + 2 H+ (aq) + 2 e−

(3)

The anodic oxidation of steel has been known since the 1960s, where solutions of boiling 14.3 M
sodium hydroxide and high current densities (50−100 mA·cm−2 ) are used to fabricate a loose and
porous magnetite layer [18,19,21].
In this article, we demonstrate that blackening can easily be electrochemically achieved for steel
samples that were immersed into concentrated sodium hydroxide solution under potentiostatic control
at temperatures that are much lower than those used in the hitherto existing commercial process.
The surface passivation layer that is generated in the absence of a chemical oxidizing agent under
mild conditions by this new electrochemical reaction shows decorative properties comparable to the
ones that are achieved with the standard technology of the costly and harmful chemical process. The
quality and selected properties of the electrochemically blackened surfaces were further investigated
while using atomic force microscopy (AFM), scanning electron microscopy (SEM), auger electron
spectroscopy (AES), and electrochemical methods.
2. Materials and Methods
2.1. Materials
The steel discs of the standard S235JRG2C+C (geometric area: 2.83 cm2 ) were grinded before
each measurement with 320 SiC abrasive paper (Buehler, ITW Test & Measurement GmbH, Esslingen,
Germany), sequentially polished with 9 and 3 μm diamond suspension, 0.05 μm Al2 O3 suspension, and
finally successively rinsed with isopropanol (VWR, analytical grade) and water. Ultrapure water that
was obtained from a water purification system (Arium 611 UV; TOC ≤ 3 ppb, Sartorius AG, Göttingen,
Germany) was used as solvent in all experiments. Sodium hydroxide (NaOH, VWR, 99%), blackening
salt (DWE Brünofix GmbH, Rednitzhembach, Germany and CSP—Chemische Spezialprodukte Olaf

104

Surfaces 2019, 2, 0

Günther, Markkleeberg, Germany), and potassium chloride (KCl, suprapur®, Merck KGaA, Darmstadt,
Germany) were used without further purification.
2.2. Blackening
Both the conventional chemical and the electrochemical blackening were performed in the
same measurement cell, which could be used in a three-electrode setup. 30 g of blackening salt
or sodium hydroxide was dissolved in 30 mL water and then heated in an iron crucible to the
desired temperature between room temperature and 120 ◦ C. A platinum counter electrode, a platinum
quasi-reference electrode, and the steel sample were immersed into the solution. Following this process,
the samples were rinsed with ultrapure water and then dried in a nitrogen stream. A potentiostat
(Fritz-Haber-Institut, Berlin, Germany) was used for electrochemical blackening and for recording
cyclic voltammograms.
Immersion into a boiling solution (130 ◦ C) of blackening salt dissolved in water (concentration
1 kg·L−1 ) for 12 min chemically blackened steel samples. For electrochemical blackening, the steel
samples were immersed into the electrolyte (blackening salt or sodium hydroxide solution) for 12 min,
using the following experimental conditions:

•
•
•
•
•

Blackening solution at 80 ◦ C: E = 0.5 V vs. Pt
Blackening solution at 100 ◦ C: E = 0 V vs. Pt
23 M NaOH at 80 ◦ C: E = 0.6 V vs. Pt
23 M NaOH at 100 ◦ C: E = 0.1 V vs. Pt
23 M NaOH at 120 ◦ C: E = 0.1 V vs. Pt

At temperatures that were higher than 120 ◦ C, the electrolyte started boiling, in which
electrochemical measurements were no longer practicable. Therefore, the maximum temperature for
electrochemical measurements was set to 120 ◦ C. It should be noted that the formation of magnetite at
temperatures below 80 ◦ C by using more positive potentials (E > 0.6 V vs. Pt) was not successful.
2.3. Characterization of the Blackened Samples
Atomic force microscopy (Nanosurf FlexAFM, Nanosurf AG, Liestal, Switzerland) in tapping
mode using Tap190 GD-G cantilevers (Budgetsensors, Innovative Solutions Bulgaria Ltd., Sofia,
Bulgaria), scanning electron microscopy, and Auger-electron spectroscopy (Phi 660, Physical
Electronics, Chanhassen, MN, USA) characterized the blackened surfaces.
The electrochemical characterization and quantification of the corrosion behavior was studied
with a conventional three electrode setup, consisting of a working, a reference, and a platinum
counter electrode. All of the potentials were recorded against an Ag/AgCl reference electrode (Schott,
cKCl = 3 mol·L−1 ). The working electrode was contacted with the electrolyte by a hanging meniscus
configuration, resulting in a geometrical surface area of 0.785 cm2 . The electrolyte was a 0.1 M KCl
solution, which was purged with N2 for at least 1 h before each measurement. Electrochemical
characterization of blackened samples was done by cyclic voltammetry between −1.5 V to −0.2 V vs.
Ag/AgCl at a scan rate of 20 mV·s−1 . The corrosion protection was quantified by linear voltammetry
in a potential range between −0.1 and 0.1 V vs. EOC at a scan rate of 1 mV·s−1 . EOC is the open circuit
potential after 1 h immersion of the sample into the electrolyte. The corrosion potential Ecorr , Tafel
slope coefficient βc , βa , and corrosion current density jcorr (as indicators of the quality of the corrosion
protection) were obtained from the measured potential–current curve by Tafel slope analysis and
fitting to the Wagner–Traud equation. The polarization resistance (RP ) was obtained from the slope of
the polarization curve at EOC or by calculation while using the Stern–Geary equation.

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Surfaces 2019, 2, 0

3. Results and Discussion
3.1. Chemical Blackening
The steel samples were prepared according to the recipe that is described in DIN 50938 in order to
compare the electrochemical and the chemical hot alkaline blackening procedures [22]. The polished
steel slices were immersed into a boiling solution of blackening salt dissolved in water (concentration
1 kg·L−1 ) at 130 ◦ C for 12 min, yielding a uniform black-blackened surface. Auger electron spectroscopy
identified the iron oxide species (see Figure 1a). The resulting oxide species were determined by the
intensity ratio of the fine structure of the Fe(M2/3 VV) signals at 46 and 54 eV [23]. The relevant intensity
ratio of approx. 8 to 10 matches that of Fe3 O4 (Figure 1b). The magnetite surface appears relatively
smooth (profile roughness parameter Ra = 12 ± 1 nm) (see the SEM and AFM images in Figure 1b,c).
Besides some cracks, surface defects that have been formed during the blackening process are visible
in these images.
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Figure 1. (a) Auger electron spectrum of a steel sample blackened at 120 ◦ C in hot alkaline nitrite
solution, showing (b) the Fe(M2/3 VV) signals at 46 and 54 eV with the intensity ratio of 7:10 for
magnetite on the surface. (c) Scanning electron microscopy (SEM) image of the same sample (10 kV,
2000 magnification) with some defects in the blackened surface; (d) The atomic force microscopy (AFM)
topography image shows a 10 × 10 μm2 area of the steel sample, scaling-up a surface defect similar
seen in the SEM image.

3.2. Electrochemical Blackening
Besides conventional chemical blackening at variable temperature, the steel samples were
electrochemically blackened both in a nitrite containing blackening solution and in sodium hydroxide
solution that is free of nitrite. To determine the optimum temperature and the DC voltage for
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Surfaces 2019, 2, 0

electrochemical blackening, the electrochemical behavior of steel samples was studied at different
temperatures between room temperature and 120 ◦ C in blackening salt and 23 M sodium hydroxide
solution by cyclic voltammetry (Figure 2). The study of steel electrodes that are in contact with sodium
hydroxide solution was also performed until the boiling temperature of 120 ◦ C.
According to the Pourbaix diagram of iron [17,24] and the well-known electrochemical behavior
of iron in alkaline solution [11,25–30], the anodic peak A1 at −0.08 V vs. Pt for steel in 23 M NaOH
(Figure 2c,d) was identified as the oxidation of Fe to Fe(II) [28–30]. This is in agreement with the
reaction that is given in Equation (1). The peak A2 at around 0.1 V vs. Pt in blackening salt solution and
in the range of 0.06 to 0.16 V vs. Pt in sodium hydroxide corresponds to the oxidation of iron hydroxide
to magnetite that is in correspondence to the Schikorr reaction (Equation (3)) [11,27–29]. Magnetite is
finally oxidized to iron(III) oxide at peak A3 [11,27]. Furthermore, the two cathodic peaks of C2 and
C3 are observed for nitrite-free solutions (Figure 2c,d). They correspond to the related oxidation peaks
A2 and A3 [29]. Measurements that were obtained with blackening salt solution (Figure 2a,b) only
exhibit the anodic peaks A2, A3, and the hydrogen evolution regime. It is noteworthy that the current
and charge densities for sodium hydroxide solution are approximately twice as large as those that
were measured with blackening salt.

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(b)

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Figure 2. Cyclic voltammograms of steel in blackening salt solution (a) at various temperatures between
room temperature and 100 ◦ C, (b) at 120◦ C and (c) in 23 M sodium hydroxide solution at various
temperatures between room temperature and 100 ◦ C, and (d) at 120 ◦ C; scan rate 50 mV s−1 .

A black coloring of the surface already appeared after the first voltammetric cycle in boiling
blackening solution, while at room temperature, even after one hour, no visible changes in color
have been observed. Only a few signs of corrosion at grain boundaries were monitored in the AFM

107

Surfaces 2019, 2, 0

topography image (Figure 3). This can be explained by the slow kinetics of the Schikorr reaction,
as already mentioned above. The blackening of the surface becomes evident after increasing the
temperature to 80 ◦ C. For this reason, the temperature range between 80 and 120 ◦ C was examined in
detail, both with sodium hydroxide and blackening salt solution as electrolytes.

(a)

(b)

Figure 3. AFM topography images (25 × 25
of a steel slice cycled in (a) 23 M NaOH and (b)
blackening salt solution. No significant differences in the surface topography appear. Some corrosion
at grain boundaries and the presence of corrosion products resemble each other.
μm2 )

The open circuit potential (EOC ) that was established for steel in blackening salt solution after
12 min is −0.05 V vs. Pt, which is located around the onset of peak A2 (Figure 2b). For this reason, the
potential was set to 0 V vs. Pt for the electrochemical treatment with blackening solution at 100 ◦ C.
Applying a potential of 0.1 V vs. Pt in 23 M NaOH solution at 100–120 ◦ C (Figure 2c), which almost
corresponds to peak A2, yields a uniform black magnetite layer after 12 min. With the same potential
at 80 ◦ C in blackening salt and 23 M NaOH solution, insufficient blackening is obtained due to the
slow kinetics of the Schikorr reaction. It is obvious that the electrode potential and temperature are the
main parameters that control the quality of electrochemical blackening. Within the same process time
of 12 min, a more positive electrode potential of 0.5 V vs. Pt with blackening salt and 0.6 V vs. Pt with
sodium hydroxide solution leads to the formation of a compact Fe3 O4 overlayer, which Auger electron
spectroscopy subsequently characterized (Figure 5).
The potentials and the resulting current densities that were measured after 12 min (jfinal ) for the
two electrolytes at temperatures of 80, 100, and 130 ◦ C (boiling temperature of blackening salt solution)
are summarized in Table 1.
Table 1. Reaction conditions for electrochemical blackening of steel in blackening salt solution and in
23 M NaOH between 80 and 130 ◦ C, temperature T, electrode potential E, and final current density
jfinal . The profile roughness as obtained from AFM imaging is given in the last column. The use of
23 M NaOH solution at temperatures over 100 ◦ C increases the surface roughness by a factor of 5, in
comparison to the chemical hot alkaline blackening.
Electrolyte
Blackening salt solution 1 kg·L−1

NaOH (23 M)

T/◦ C

E/V vs. Pt

jfinal /mA cm−2

Ra /nm

80
100
130
80
100
120

0.5
0
−0.05 (EOC )
0.6
0.1
0.1

1.0
0.6
(0)
0.4
1.0
19.5

15 ± 4
38 ± 6
12 ± 1
30 ± 11
59 ± 7
61 ± 11

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Surfaces 2019, 2, 0

Electrochemical blackening at temperatures below 80 ◦ C by using more positive potentials was
not successful. Apparently, ferrate(VI) ions (FeO4 2− ) are formed at higher potentials, resulting in a
purple colored electrolyte [28,31,32]. UV-VIS spectroscopy (Figure 4) could confirm the electrochemical
formation of the ferrate(VI) ion, showing a band with an absorption maximum at 505 nm (green color,
complementary of purple) [32–34].

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Figure 4. UV-VIS spectrum of 23 M NaOH after electrochemical blackening at potentials E > 0.6 V vs. Pt.
The absorption maximum at approx. 500 nm indicates the formation and presence of ferrate(VI) ions.

In order to confirm the chemical composition of the iron oxide overlayers that were formed by
electrochemical blackening, the steel samples were examined by Auger electron spectroscopy. Figure 5
shows (exemplarily for all samples under study) the Auger electron spectrum of steel blackened in
NaOH at 80 ◦ C and 0.6 V vs. Pt. The intensity ratios for the Fe(M2/3 VV) signals at 46 and 54 eV
(Figure 5b) are similar to those that were obtained for the samples with classical chemical treatment
(Figure 1a,b). We conclude that the formed black oxide layer is indeed magnetite (Fe3 O4 ), which is in
agreement with the findings in literature [23].
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Surfaces 2019, 2, 0

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Figure 5. (a) Auger electron spectrum of steel blackened in NaOH at 80
and 0.6 V vs. Pt. (b) The
intensity ratio of the peaks at 46 and 54 eV indicates the formation of a magnetite overlayer (Fe3 O4 ),
(c,d) SEM images of a steel sample blackened in NaOH at 80 ◦ C and 0.6 V vs. Pt, indicating a corrosive
attack during the blackening process (10 kV, (c) 100x, and (d) 2000x magnification).

In addition to the Auger electron spectroscopy measurements, we also took SEM images, showing
sponge-like structures on the surfaces (Figure 5c,d). This indicates corrosive attack and surface
roughening by NaOH during the electrochemical blackening process. The electrochemical origin of
this corrosion process is a parallel reaction that is competitive with the one sketched in Equation (2). At
high pH values, iron(II)hydroxide reacts with excess hydroxide to soluble HFeO2 - (Equation (4) [17,26].
Fe(OH)2,surf + OH- (aq) → HFeO2 - (aq) + H2 O

(4)

AFM topography images of the blackened samples after different treatments were recorded
in order to quantify the surface roughness (Figure 6). Only minor changes could be observed for
blackening salt solution at 80 and 130 ◦ C (Figure 6a,b). This is in agreement with a relatively small
roughness factor of 12–15 nm (Table 1). In contrast, steel samples that were blackened in a solution
of sodium hydroxide at 80 ◦ C, and even more at 120 ◦ C, yield a rougher (Ra = 30–60 nm) and more
damaged surface (Figure 6c,d). These surface defects are triggered by the corrosive effect of the sodium
hydroxide solution and the high anodic current densities [31]. As expected, the higher the anodic
current density, the larger the roughness of the blackened surface. We would like to point out that, by
applying a potential in both blackening solutions at 100 ◦ C and 23 M NaOH, the roughness is at least
increased by a factor of two. This increase of roughening at higher temperature is possibly caused by
the enhanced formation of Fe(OH)2 (Figure 2) and the corrosive attack of NaOH (Equation (4)), in
competition with the Schikorr reaction.

(a)

(b)
Figure 6. Cont.
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Surfaces 2019, 2, 0

(c)

(d)

Figure 6. AFM images of steel samples electrochemically blackened in (a) blackening salt solution at
80 ◦ C and (b) 130 ◦ C, and (c) in NaOH at 80 ◦ C, and (d) 120 ◦ C. The images show a 25 × 25 μm2 area
of the individual samples.

3.3. Electrochemical Characterization
The electrochemical behavior of the blackened samples was studied by cyclic voltammetry in the
potential range between −1.5 and −0.2 V vs. Ag/AgCl, with 0.1 M potassium chloride solution at a
sweep rate of 20 mV·s−1 (see Figure 7). In the case of the chemically blackened sample, the positive
potential limit was expanded to 0 V vs. Ag/AgCl to study an additional anodic current peak at −0.27 V
vs. Ag/AgCl. Subsequently, a Tafel slope analysis was performed for the voltammetric data.

(a)

(b)

Figure 7. Cyclic voltammogram of blackened steel discs in 0.1 M KCl solution at (a) 80 ◦ C and (b)
120 ◦ C; black curve: blackening salt; red curve: NaOH; scan rate 20 mV s−1 .

The cyclic voltammograms for the various steel samples that were blackened at different
temperatures are all very similar in the potential range between −1.5 and −0.6 V vs. Ag/AgCl
(Figure 7). The surfaces that were blackened in NaOH start oxidizing to Fe(III) at −0.6 V vs. Ag/AgCl,
while this reaction requires higher potentials for samples that were treated in the blackening salt
solution. This can be explained by a more defect-free and smoother steel surface, which is apparent
from the images that were obtained by AFM.
The corrosion potentials (Ecorr ), corrosion current densities (jcorr ), and Tafel slope coefficients (βc
and βa ) for the various blackened samples (listed in Table 2) were obtained by Tafel slope analysis from
the polarization curves that are shown in Figure 8. In order to determine the quality of the blackened
surfaces and their corrosion behavior (qualitative salt spray test may give similar trends), the corrosion

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Surfaces 2019, 2, 0

rates (kcorr ) were calculated while using the Faraday laws. The polarization resistances (RP ) were
obtained by calculation using the Stern–Geary equation [35].
Rp =

βa |βc |
2.303( β a + | β c |) jcorr



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Figure 8. Potentiodynamic polarization curves for pristine and blackened steel surfaces in in 0.1 M KCl;
scan rate 1 mV s−1 . Electrochemical blackening was performed both in 23 M NaOH and blackening
salt solution.

A comparison of the corrosion potentials shows that the electrochemically-blackened samples
appear to be less noble than the classically blackened or untreated steel samples. Nevertheless, the
corrosion current densities and rates are similar to the classically blackened samples and about 70%
lower than for an untreated steel sample. The samples that were treated with sodium hydroxide
solution at 80 ◦ C have an approximately 40% lower corrosion current when compared to the blank
sample. The polarization resistance supports this trend, illustrating an improvement of the corrosion
resistance of the samples that were oxidized in blackening salt solution by a factor of approx. two as
compared to a blank surface, while the improvement of electrochemical oxidized samples in 23 M
NaOH is about 10–40%.
The rough and porous structure of the magnetite layer possibly caused the less noble behavior
and lower corrosion protection of the electrochemical-blackened samples (Figure 6c,d). In this case,
chloride ions could pass the oxide layer and attack the metal, forming a local cell beneath the oxide,
which reduces the corrosion potential.
Supporting this, the form of the Tafel plots in Figure 8 and the values of the anodic Tafel slope
coefficients (βa ) of steel samples chemically and electrochemically blackened in blackening salt solution
indicate protection of the surface by passivation due to the formed magnetite layer. In contrast, the
anodic Tafel slope coefficients (βa ) of steel samples that were electrochemically blackened in 23 M
NaOH seem to be similar to an untreated sample. This behavior suggests that, although the corrosion
rate decreased by a reduction of the active surface area, the porous structure of the magnetite layer
does not result in a surface protection by passivation. The data in Table 2 indicates that the chemically
and electrochemically formed magnetite overlayers represent rather poor protective layers in the
absence of additional sealing. Since the blackened surfaces are usually sealed, the rougher surface may
increase the adhesion strength of the sealant, providing even better corrosion resistance.

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Surfaces 2019, 2, 0

Table 2. Corrosion potential (Ecorr ), corrosion current (jcorr ), Tafel slope coefficient (|βc | and βa ),
polarization resistance (RP ) and corrosion rate (kcorr ) of steel samples blackened under different
conditions. Blank steel samples were used as reference.
Sample
Blank
Blackening salt
solution 80 ◦ C
Blackening salt
solution 130 ◦ C
NaOH 80 ◦ C
NaOH 120 ◦ C

Ecorr /V vs.
Ag/AgCl

jcorr /μA cm−2

|βc |/V dec−1

βa /V dec−1

RP /kΩ cm−2

kcorr /mm a−1

−0.67

7.1

0.07

0.12

2.7

0.082

−0.74

1.2

0.04

0.21

6.6

0.014

−0.57

2.0

0.03

0.18

5.6

0.023

−0.84
−0.75

4.4
2.8

0.04
0.03

0.12
0.12

3.0
3.7

0.051
0.033

4. Conclusions
The electrochemical behavior of steel in blackening salt solution and the concentrated sodium
hydroxide solution was investigated between room temperature and 120 ◦ C by cyclic voltammetry.
It is shown that the blackening of steel can be electrochemically achieved by applying a potential
in the range of 0 to 0.6 V vs. Pt at temperatures between 80 and 120 ◦ C. A further decrease of the
bath temperature to room temperature and a simultaneous increase of potential did not yield the
desired result due to the formation of Fe(VI) ions and the slow kinetics of the Schikorr reaction.
Auger electron spectroscopy confirmed the formation of magnetite on the surface during blackening.
AFM studies revealed surface roughening for the given reaction conditions. Electrochemical studies
for blackened steel in 0.1 M KCl that is based on cyclic voltammetry and corrosion measurements
revealed a similar electrochemical behavior and an equivalent corrosion protection for the following
experimental conditions:

•
•
•

23 M NaOH, 0.6 V vs. Pt and 80 ◦ C,
23 M NaOH, 0.1 V vs. Pt and 120 ◦ C, and
blackening solution, 0.5 V vs. Pt and 80 ◦ C.

Very similar surface properties are achieved in an efficient electrochemical blackening process at
lower temperature when compared to classical chemical blackening at 130 ◦ C with hazardous sodium
nitrite. The combination of electrochemical blackening and surface sealing may help to develop an
inexpensive, environmentally-friendly, and adequate method to obtain the aspired corrosion protection
of steel surfaces, together with a unique appearance.
Author Contributions: Conceptualization, M.E., S.Z. and L.A.K.; Data curation, C.K.; Formal analysis, M.E., S.Z.
and C.K.; Funding acquisition, T.J.; Investigation, C.K.; Methodology, A.F.; Resources, T.J.; Supervision, A.F.,
L.A.K. and T.J.; Validation, A.F.; Writing—original draft, M.E. and S.Z.; Writing—review & editing, L.A.K. and T.J.
Acknowledgments: The authors thankfully acknowledge Kneissler Brüniertechnik GmbH for supporting this
research study.
Conflicts of Interest: The authors declare no conflict of interest.

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2.
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Article

Cyclic Voltammetry Characterization of Au, Pd,
and AuPd Nanoparticles Supported on Different
Carbon Nanofibers
Anna Testolin 1 , Stefano Cattaneo 1 , Wu Wang 2 , Di Wang 2,3 , Valentina Pifferi 1, *, Laura Prati 1 ,
Luigi Falciola 1 and Alberto Villa 1, *
1

2
3

*

Dipartimento di Chimica, Università degli Studi di Milano, Via Golgi 19, 20133 Milano, Italy;
Anna.Testolin@unimi.it (A.T.); Stefano.Cattaneo2@unimi.it (S.C.); Laura.Prati@unimi.it (L.P.);
luigi.falciola@unimi.it (L.F.)
Institute of Nanotechnology, Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1,
D-76344 Eggenstein-Leopoldshafen, Germany; wu.wang@partner.kit.edu (W.W.); di.wang@kit.edu (D.W.)
Karlsruhe Nano Micro Facility, Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1,
D-76344 Eggenstein-Leopoldshafen, Germany
Correspondance: Valentina.Pifferi@unimi.it (V.P.); alberto.villa@unimi.it (A.V.);
Tel.: +39-025-031-4222 (V.P.); +39-025-031-4361 (A.V.)

Received: 18 January 2019; Accepted: 13 March 2019; Published: 19 March 2019

Abstract: Three types of carbon nanofibers (pyrolytically stripped carbon nanofibers (PS),
low-temperature heat treated carbon nanofibers (LHT), and high-temperature heat treated carbon
nanofibers (HHT)) with different graphitization degrees and surface chemistry have been used as
support for Au, Pd, or bimetallic AuPd alloy nanoparticles (NPs). The carbon supports have been
characterized using Raman, X-ray photoelectron spectroscopy (XPS), and cyclic voltammetry (CV).
Moreover, the morphology of the metal nanoparticles was investigated using transmission electron
microscopy (TEM) and CV. The different properties of the carbon-based supports (particularly the
graphitization degree) yield different electrochemical behaviors, in terms of potential window widths
and electrocatalytic effects. Comparing the electrochemical behavior of monometallic Au and Pd and
bimetallic AuPd, it is possible to observe the interaction of the two metals when alloyed. Moreover,
we demonstrate that carbon surface has a strong effect on the electrochemical stability of AuPd
nanoparticles. By tuning the Au-Pd nanoparticles’ morphology and modulating the surface chemistry
of the carbon support, it is possible to obtain materials characterized by novel electrochemical
properties. This aspect makes them good candidates to be conveniently applied in different fields.
Keywords: gold; palladium; bimetallic alloy; carbon nanofibers (CNFs); cyclic voltammetry (CV)

1. Introduction
Over the last decades, many studies have focused on metal nanoparticles, thanks to their excellent
catalytic, electrocatalytic, optical, and magnetic properties, that often differ from the ones of their
correspondent bulk metals [1–3]. There has also been interest in bimetallic nanosystems, because of the
possibility to create new materials characterized by specific and innovative properties due to synergistic
effects among precursors [4–7]. These effects reflect a general enhancement of the performances of the
final products with respect to their monometallic components. The possibility to have different kinds
of bimetallic samples (in terms of composition, structure, metal loading, morphology, etc.) has led to
widespread applications in a lot of fields such as engineering, electronics, catalysis, electrocatalysis,
and sensors [8–10]. Particularly in heterogeneous catalysis and electrocatalysis, metal nanoparticles
are often supported on specific substrates [11–17], which in turn possess different characteristics and
properties related to the amount of catalyst loaded on the solid material, or to the different nature
Surfaces 2019, 2, 205–215; doi:10.3390/surfaces2010016

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Surfaces 2019, 2

and functionalization of the substrate itself. However, in many cases, the effect of the support on the
properties of the metal nanoparticles and the possible synergistic effects between bimetallic composites
and supports are yet to be investigated in depth.
In the present work, we propose an integrated study of three types of carbon nanofibers (PS,
LHT, and HHT) modified with Au and/or Pd nanoparticles (AuNPs, PdNPs, and alloyed AuPdNPs).
AuPd systems are widely used in electrocatalysis due to their peculiar behavior compared to the
corresponding monometallic Au and Pd [18–23]. Besides surface characterization techniques, such
as Raman, TEM, and XPS, commonly involved in the study of these kind of materials [24,25], cyclic
voltammetry is here used as a supplementary and complementary characterization method [26–29].
The different properties of the supports, particularly in terms of graphitization degree, are reflected in
different electrochemical behaviors, in terms of potential window widths, electrocatalytic effects, and
preferential sites for anchoring the nanoparticles. Moreover, the results highlight how the bimetallic
samples behave differently with respect to the relative Au and Pd monometallic systems, underlying
the presence of an intimate contact between the two metals, which provides new materials, completely
different from the original ones. This aspect makes them good candidates to be conveniently applied
in many different applications.
2. Materials and Methods
2.1. Metal Nanoparticles/CNFs Preparation
Commercial carbon nanofibers (CNFs), PR24-PS (hereafter shortened as PS) from Applied Sciences,
inc. (average diameter of 100 ± 30 nm and a specific surface area of 45 m2 /g); CNFs PR24-LHT (LHT)
from Applied Sciences, inc. (average diameter of 100 ± 30 nm and a specific surface area of 43 m2 /g);
and CNFs PR24-HHT (HHT) from Applied Sciences, inc. (average diameter of 100 ± 30 nm and a
specific surface area of 41 m2 /g) were used as pristine carbon materials. NaAuCl4 *2H2 O (99%, Aldrich,
St. Louis, MO, USA) and Na2 PdCl4 (98 % Aldrich) were used as metal precursors, while NaBH4
(> 98%, Ventron) and polyvinyl alcohol (PVA, MW = 9000–10,000, 80 % hydrolyzed, Aldrich) were
used as reducing and protective agents, respectively.
AuNPs: 0.051 mol of solid NaAuCl4 *2H2 O and PVA solution (1% w/w, metal/PVA 1:1 w/w)
were added to 100 mL of Milli-Q H2 O. The solution was stirred for few minutes, after which NaBH4
(metal/NaBH4 = 1/4 mol/mol) solution was added under vigorous magnetic stirring. Within a few
minutes of sol generation, the sol was immobilized by adding the support (ca. 1 g) under vigorous
stirring. The slurry was filtered, and the catalyst washed thoroughly with distilled water and dried at
80 ◦ C for 4 h. The total nominal metal loading was 1 wt %.
PdNPs: 0.094 mol of Na2 PdCl4 and PVA solution (1% w/w, metal/PVA 1:1 w/w) were added to
100 mL of Milli-Q H2 O. The solution was stirred for few minutes, after which NaBH4 (metal/NaBH4
= 1/8 mol/mol) solution was added under vigorous magnetic stirring. Within a few minutes of sol
generation, the sol was immobilized by adding the support (ca. 1 g) under vigorous stirring. The slurry
was filtered, and the catalyst washed thoroughly with distilled water and dried at 80 ◦ C for 4 h.
The total nominal metal loading was 1 wt %.
AuPdNPs: 0.037 mol of solid NaAuCl4 *2H2 O, 0.025 mol of Na2 PdCl4 and PVA solution (1% w/w,
metal/PVA 1:1 w/w) were added to 100 mL of Milli-Q H2 O. The yellow solution was stirred for few
minutes, after which NaBH4 (metal/NaBH4 = 1/8 mol/mol) solution was added under vigorous
magnetic stirring. Within a few minutes of sol generation, the sol was immobilized by adding the
support (ca. 1 g) under vigorous stirring. The slurry was filtered, and the catalyst washed thoroughly
with distilled water and dried at 80 ◦ C for 4 h. The total nominal metal loading was 1 wt %.
2.2. Electrochemical Characterization
Cyclic voltammetric experiments were performed using an AutoLab PGStat128 (Metrohm
AutoLab, Utrecht, The Netherlands) equipped with the NOVA 2.1 Software (Metrohm AutoLab,

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Utrecht, The Netherlands). The experimental setting was composed of a conventional three electrodes
cell in which a saturated calomel electrode and a platinum wire were used as reference (RE) and
counter (CE) electrodes, respectively. A glassy carbon (GC) modified with carbon nanofibers (CNFs)
or with mono or bimetallic Au and Pd NPs supported on carbon nanofibers was used as working
electrode (WE). Before the modification, GC was polished with diamond powder (1 μm, Sigma Aldrich,
Milan, Italy) on a Struers DP-Nap cloth and washed with milli-Q water. 5 mg of the different CNF-NPs
powders were suspended in 1 mL EtOH (96 %, Sigma Aldrich) and 4.5 μL Nafion (5% solution in low
aliphatic alcohols, Aldrich, Milan, Italy). Then, 20 μL of the suspension was deposited on GC using an
automatic micropipette (Kartell, Noviglio, Milan, Italy) and allowed to dry for approximately 30 min.
Cyclic voltammetries were recorded in aqueous solution with 0.1 M H2 SO4 as supporting electrolyte.
The potential was scanned from −1 V (SCE) to +2 V (SCE), at a scan rate of 100 mV s−1 and a step
potential of 5 mV.
2.3. Morphological Characterization
Raman spectroscopy was performed with a Renishaw inVia Raman microscope for analysing
the graphitisation degree of the carbon nanofibers. Bare supports and fresh and used catalysts
were analysed. Typically, a sample of approximately 0.01 g was placed on a metal slide inside the
spectrometer. The powder was analysed under an IR class laser (514 nm) with a laser intensity of 50 %.
The sample was scanned at an attenuation time of 22 seconds, and 10 scans were carried out to give
a spectrum.
X-ray photoelectron spectroscopy (XPS) was performed on a Thermo Scientific K-alpha+
spectrometer (Thermo Fisher Scientific, Waltham, MA, USA). Samples were analysed using a
monochromatic Al x-ray source operating at 72 W (6 mA × 12 kV), with the signal averaged over
an oval-shaped area of approximately 600 × 400 microns. Data were recorded at pass energies of
150 eV for survey scans and 40 eV for high resolution scan with a 1eV and 0.1 eV step size, respectively.
Charge neutralization of the sample was achieved using a combination of both low energy electrons
and argon ions (less than 1 eV), which gave a C(1s) binding energy of 284.8 eV.
All data were analysed using CasaXPS (v2.3.17 PR1.1) using Scofield sensitivity factors and an
energy exponent of −0.6.
The Au, Pd specimens were characterized by transmission electron microscopy (TEM, FEI Titan
80−300, Thermo Fisher Scientific, Waltham, MA, USA). The samples were directly dispersed on copper
grids covered with holey carbon film. Morphology of the catalysts was characterized by high angle
annular dark-field (HAADF) scanning transmission electron microscopy (STEM), and its composition
information was acquired by EDAX S-UTW EDX detector in an FEI Titan 80−300 microscope operating
(Thermo Fisher Scientific, Waltham, MA, USA) at 300 kV. Analysis of STEM-EDX spectrum imaging
was carried out by using TEM image and analysis (TIA 4.7 SP3 version) software. Particle size of the
specimens was measured on HAADF-STEM images by using the ImageJ software (National Institutes
of Health, Bethesda, MD, USA) fitting the particles with ellipsoid shapes.
3. Results and Discussion
3.1. Morphological Characterization
Commercial carbon nanofibers (CNFs) thermally processed at different severity degrees [30,31],
were selected as supports. Raman spectroscopy has been used to investigate the graphitization degree
of CNFs treated at different temperatures. In particular, the ratio of the integral intensities of D and
G bands (ID /IG ), which is an index of the defectiveness of the graphite layers [32], was measured
(Table 1). G band is related to structurally-ordered graphite domains, whereas D band can be associated
with the turbostratic and/or disordered carbonaceous structures [32]. As expected, increasing the
temperature of the heat treatment results in a higher graphitization degree (ID /IG = 0.75, 0.60, 0.11 for
PS, LHT, HHT, respectively) (Table 1). X-ray photoelectron spectroscopy (XPS) has been used to confirm

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the different graphitization degrees. C1s XPS spectra show two main components: the first one at
284 eV, which can be attributed to sp2 -hybridised carbon species, and the second component at 285 eV
ascribable to the presence of sp3 -hybridised carbon species [33] (Figure 1 and Tables 1 and 2). Sp2 /sp3
increased with increasing annealing temperature (4.1, 7.0 and 16.8 for PS, LHT, HHT, respectively)
(Table 1). Accordingly, a higher amount of oxygen has been observed decreasing the graphitization
degree (C:O ratio of 88.3:11.7, 92.2:7.8, 97.3:2.7 for PS, LHT, HHT, respectively) (Table 2). Table 2
summarizes the oxygen species observed on the surface, their concentration, and the overall elemental
composition. Fitting the O1s peak, five contributions were observed (Figure 2). Binding Energy (BE)
of 530.2–530.5 eV corresponds to highly conjugated forms of carbonyl oxygen such as quinone [34].
The second contribution (531.6-531.9 eV) can be assigned to a carbon–oxygen double bond, whereas the
one at 533.0-533.5 eV corresponds to an ether-like single bond C-O-C and to carbon oxygen single-bonds
in hydroxyl groups C-O-H [35]. Oxygen species at 534.6–535.0 eV refers to the presence of carboxylic
groups COOH [35]. The last signal at 536.7–537.1 eV can be assigned to adsorbed water [36]. C=O,
C-O-H and C-O-C were the most abundant species present in all the CNFs (Table 2 and Figure 2).
Table 1. Atomic content of sp2 and sp3 carbon and ratio sp2 /sp3 from X-ray photoelectron spectroscopy
(XPS) and ID /IG ratio from Raman for the bare supports.
Catalyst

C sp2 (%)

C sp3 (%)

sp2 /sp3

ID /IG

PS
LHT
HHT

75.8
83.8
90.8

18.6
12.0
3.4

4.1
7.0
16.8

0.75
0.60
0.11

Figure 1. C1s XPS spectra of (a) pyrolytically stripped carbon nanofibers (PS), (b) low-temperature
heat treated carbon nanofibers (LHT), and (c) high-temperature heat treated carbon nanofibers (HHT).

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Table 2. XPS analysis of the carbon nanofibers (CNFs).
Sample

C=O

C-O-H

O1S
C-O-C

COOH

H2 O

sp2

sp3

C1s
C=O

C=C

Atomic
Ratio % C:O

PS

BE
Atom %

530.2
(10.2)

531.8
(40.7)

533.0
(33.1)

534.6
(9.8)

536.7
(6.0)

284.4
(75.8)

284.9
(18.6)

287.2
(3.7)

290.0
(1.8)

88.3:11.7

LHT

BE
Atom %

530.5
(11.6)

531.6
(34.2)

533.1
(40.3)

535.0
(8.8)

537.1
(5.1)

284.5
(83.8)

284.9
(12.0)

287.4
(2.6)

290.1
(1.6)

92.2:7.8

HHT

BE
Atom %

530.2
(13.9)

531.9
(46.2)

533.5
(24.0)

535.0
(10.1)

537.0
(5.8)

284.7
(90.8)

285.2
(5.4)

287.9
(1.6)

290.6
(2.2)

97.3:2.7

Figure 2. O1s XPS spectra of (a) PS, (b) LHT, and (c) HHT.

Monometallic Au and Pd, together with their corresponding bimetallic system (AuPd), were
prepared by sol immobilization method using polyvinyl alcohol (PVA) as protective agent and NaBH4
as reducing agent and immobilized on the three different CNFs. The monometallic and bimetallic
systems consist of nanoparticles with an average size of 3–4 nm [37,38]. To investigate the morphology
of the bimetallic systems, STEM-EDX spectrum imaging was performed on single nanoparticles to
map the interior metal compositions. Element maps of one AuPd particle are reported in Figure 3,
as example. Both Pd and Au signals were found in nanoparticles with a similar distribution, indicating
the formation of AuPd alloyed nanoparticles. Quantification of the integrated EDX spectra of this
nanoparticle indicates that the atomic ratio of Au and Pd is 64/36 close to the nominal 60/40 (Figure 3).
3.2. Electrochemical Characterization
As demonstrated by Raman spectra and XPS analysis (see Tables 1 and 2, Figures 1 and 2
in Section 3.1), the three types of carbon nanofibers differ in the ratio between the sp2 and sp3
conformation of carbon atoms, i.e., in the “graphitization” degree, an index of the order degree in
carbon-based materials [31,39,40]. These differences are reflected in different voltammetric behaviours,
especially in terms of potential window widths (the useful voltage range, between the upper and lower
limits where oxidation and reduction of the solvent/supporting electrolyte occur) and electro-catalytic
effects (Figure 4).

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Figure 3. The element maps obtained from Scanning Electron Microscopy- Energy Dispersive X-ray
Analysis (STEM)-EDX spectrum imaging in the area marked box of AuPd nanoparticles and the
integrated EDX spectra of this nanoparticle.

Figure 4. Cyclic voltammetry (CV) patterns of PS, LHT, and HHT nanofibers recorded in 0.1 M H2 SO4
at 100 mV s−1 . Zoom of the current density values in the inset.

The potential window width is in the order PS < LHT < HHT (Figure 4), both in the cathodic
and anodic side, pointing out the higher electrochemical stability and inertness of HHT (> LHT > PS)
towards oxygen evolution reaction (OER, in the anodic side) and hydrogen evolution reaction (HER,
in the cathodic side).
Moreover, PS nanofibers present oxidation and reduction peaks around +0.4 V (SCE), which are
less marked for LHT and basically absent in the case of HHT. Considering the surface morphology
of the investigated materials (Section 3.1), these peaks have been attributed to the presence of
oxygen functionalizations on the nanofibers [41,42], which are in higher concentrations in the case
of PS, lower for LHT, and almost absent in HHT (Table 2). It is known, in fact, that the lower the
graphitization degree, the higher the presence of surface impurities and defects [43]. The presence of
these functionalities can be responsible also for the different potential windows of the materials.
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In order to investigate if the various properties of the nanofibers reflect also in a different behavior
when they are used as support for the nanoparticles samples, we studied CNFs modified with
monometallic Au or Pd NPs and with bimetallic AuPd alloyed NPs.
Figure 5a reports the voltammograms of the Au-modified nanofibers. As previously observed in
Figure 4, the smallest potential window is related to PS system, followed by LHT and then HHT. In this
case, however, additional oxidation and reduction peaks related to the presence of gold nanoparticles
on the carbonaceous fibers are present (inset of Figure 5A). In particular, apart from the same peaks of
the bare nanofibers at +0.4 V (SCE), a well-defined signal (for LHT and HHT) at around +1.3 V (SCE)
is evident, whose peak-shape confirms the nanosized structure of the Au particles [44]. In the case of
PS, the signal is partially covered by the reaction of the background, resulting in a shoulder plateau.
Considering the oxidation peak potential position, a small catalytic effect for HHT (peak at +1.25 V vs
SCE, with respect to the value of +1.31 for LHT-supported sample) is observed, probably related to a
higher interaction between HHT nanofiber and AuNPs, due to the high graphitization degree of the
material which reflects in a favored electron transfer.

Figure 5. Cyclic voltammetry (CV) of PS, LHT, and HHT modified with AuNPs (a) or PdNPs
(b) recorded in H2 SO4 0.1 M at 100 mV s−1 . In the inset, zoom on the density current values.

The interesting behavior of the supported materials can be outlined considering the main
oxide-reduction peaks in the 0.5–0.8 V (SCE) region, to be compared with the peak position (at
about +0.8 V (SCE)) of analogous AuNPs system in the absence of the carbon nanofibers support [40].
In the present case, all the three different supported samples show the gold oxide reduction peaks
at lower potentials (and in the order PS < LHT < HHT). This behavior can be ascribed to a different
electrochemical stabilizing effect exhibited by the different carbonaceous fibers on the nanomaterials
that can be further and conveniently exploited in sensing and catalytic applications.
In the case of PdNPs (Figure 5b), the presence of a small oxidation peak around +1 V (SCE)
is evident, which is in line with the values of analogous colloidal PdNPs. In the case of LHT_Pd,
more intense signals at about 1.73 and 1.89 V (SCE) are present, probably due to the concomitant Pd
over-oxidation together with O2 evolution. In both HHT systems (HHT_Au and HHT_Pd), an evident
reduction peak at about −0.6 V (SCE) is present, whose nature is yet to be clarified (Figure S1).
The CV patterns of the bimetallic alloyed AuPd nanoparticles supported on the carbonaceous
materials (Figure 6) are more interesting. Also in this case, the broad oxidation peak related to the
reaction of the background is present in the PS fibers case, which covers the ones related to the metallic
nanoparticles. In the case of LHT, it can be noticed the presence of two oxidation peaks at about +1.25
and +1.74 V (SCE) related to the Pd and Au oxide formation, respectively, both shifted at higher anodic
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potentials with respect to their monometallic components [45]. This behavior can be attributed to
a stabilizing effect that the two metals exhibit when they are alloyed, as already evidenced in the
literature [7]. The single reduction signal at about −0.29 V (SCE) is probably related to the reduction
of the alloyed AuPd oxide, which occurs at a more cathodically shifted potential with respect to its
monometallic references, resulting in a great stabilization of the oxide, probably due to the synergistic
effects of the two metals in the alloy, and also due to a contribution by the carbonaceous support.
This unexpected potential shift of the main oxide-reduction peak (usually, on bare Pd, in this region
it is more likely to have proton absorption), has been observed in other homogeneous noble metals
alloys [46] and in other bimetallic Pd-based catalysts supported on carbon nanofibers [47].
For HHT, a small oxidation peak around +1.2 V (SCE) is present (see the inset of Figure 6), which
could be related to the formation of the palladium oxide, but no signal related to the oxidation of
AuNPs is evident. It seems that the AuNPs oxidation process occurs at higher potential with respect to
+2 V (SCE), where the signal of the OER takes place. Also, the reduction peak in this case is not evident
(probably extremely cathodically shifted and under the signal of the HER background), once more
evidencing the greater stabilizing effect of HHT on the alloyed NPs.

Figure 6. Cyclic voltammetric (CV) features of PS, LHT, and HHT modified with AuPdNPs recorded
in H2 SO4 0.1 M at 100 mV s−1 . Zoom of the density current values in the insets.

4. Conclusions
Three types of carbon nanofibers (PS, LHT, and HHT) characterized by a different degree
of graphitization (HHT > LHT > PS) have been studied by CV. The different properties of the
carbon-based supports yield different electrochemical behaviors, in terms of potential window widths
and electrocatalytic effects. Au, Pd, and bimetallic AuPd alloy nanoparticles have been prepared by sol
immobilization using PVA as a stabilizer and NaBH4 as a reducer, and deposited on the three carbon
nanofibers. The subsequent support of the metal NPs yields the different behavior of the latter in
comparison to the non-supported systems. A cathodic and anodic shift in the oxidation and reduction
peaks, related to the formation and disruption of the oxide, is evident, particularly in the case of
the use of the more graphitized HHT support. This effect can be attributed to a different interaction
between the metal systems and the carbon supports, which provides devices characterized by the
presence of the metal catalyst, but with unexpected large potential windows, to be conveniently used
in electrocatalysis and sensors applications.

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Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/2/1/16/s1,
Figure S1: Cyclic Voltammetric (CV) features of HHT nanofibers modified with Au, Pd and AuPdNPs recorded in
H2 SO4 0.1 M at 100 mV s−1 .
Author Contributions: V.P. and A.V. designed the experiments; A.T. performed the CV experiments and wrote
the article; A.V. carried out the XPS experiments and helped in the interpretation; W.W. and D.W. carried out the
TEM experiments and helped in the interpretation; V.P. and L.F. helped in the interpretation of the CV results; S.C.
prepared the materials; and L.P. and L.F. were involved in the writing and editing of the manuscript.
Acknowledgments: We acknowledge Karlsruhe Nano Micro Facility for using the TEM.
Conflicts of Interest: The authors declare no conflict of interest.

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Article

“Click” Chemistry on Gold Electrodes Modified
with Reduced Graphene Oxide by
Electrophoretic Deposition
Vladyslav Mishyn 1 , Patrik Aspermair 1,2 , Yann Leroux 3 , Henri Happy 1 , Wolfgang Knoll 2,4 ,
Rabah Boukherroub 1 and Sabine Szunerits 1, *
1

2
3
4

*

Univ. Lille, CNRS, Centrale Lille, ISEN, Univ. Valenciennes, UMR 8520—IEMN, F-59000 Lille, France;
vladyslav.mishyn@univ-lille.fr (V.M.); patrik.aspermair@univ-lille.fr or patrik.aspermair@cest.at (P.A.);
henri.happy@iemn.univ-lille.fr (H.H.); rabah.boukherrroub@univ-lille.fr (R.B.)
CEST Competence Center for Electrochemical Surface Technology, 3430 Tulln, Austria;
wolfgang.knoll@cest.at
Univ. Rennes, CNRS, ISCR—UMR 6226, Campus de Beaulieu, F-35000 Rennes, France;
yann.leroux@univ-rennes1.fr
AIT Austrian Institute of Technology, 3430 Tulln, Austria
Correspondence: sabine.szunerits@unvi-lille.fr; Tel.: +33-3-62-53-17-25

Received: 22 January 2019; Accepted: 8 March 2019; Published: 18 March 2019

Abstract: The coating of electrical interfaces with reduced graphene oxide (rGO) films and their
subsequent chemical modification are essential steps in the fabrication of graphene-based sensing
platforms. In this work, electrophoretic deposition (EPD) of graphene oxide at 2.5 V for 300 s followed
by vapor treatment were employed to coat gold electrodes uniformly with rGO. These interfaces
showed excellent electron transfer characteristics for redox mediators such as ferrocene methanol
and potassium ferrocyanide. Functional groups were integrated onto the Au/rGO electrodes by
the electro-reduction of an aryldiazonium salt, 4-((triisopropylsilyl)ethylenyl)benzenediazonium
tetrafluoroborate (TIPS-Eth-ArN) in our case. Chemical deprotection of the triisopropylsilyl function
resulted in propargyl-terminated Au/rGO electrodes to which azidomethylferrocene was chemically
linked using the Cu(I) catalyzed “click” chemistry.
Keywords: reduced graphene oxide; electrophoretic deposition; surface chemistry; click chemistry

1. Introduction
Accurate analysis of the presence of disease-specific biomarkers in biological fluids remains
of great importance in clinical settings [1] and electrochemical sensors can reach that goal by
converting a chemical or a biological response into a processable and quantifiable electrochemical
signal [2]. Graphene and its related derivatives have generated great expectations as a transducing
platform in biosensing, due to their good mechanical properties accompanied by biocompatibility,
electrical conductivity and fast charge transfer kinetics [3–6]. A mandatory step in the production of
biosensors is the modification of graphene-based materials with recognition elements. Covalent and
non-covalent strategies have been employed, including amide bond formation and π–π interactions,
among others [7–9], to integrate surface functionalities and ligands onto graphene-based transducers.
The development of these approaches depends on having robust graphene-coated interfaces at
hand. Next to drop-casting and spin-coating of reduced graphene oxide (rGO) suspensions onto
electrical interfaces, electrophoretic deposition (EPD) has been shown to be an effective technique
for manipulating graphene oxide (GO) suspensions with the aim of producing graphene-related
films [9–13]. The ability of EPD to be applied to different materials and to control the thickness of
the deposits has been well known for a decade [14,15]. EPD has gained increased interest as an
Surfaces 2019, 2, 193–204; doi:10.3390/surfaces2010015

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alternative processing technique for the deposition of various nanomaterials ranging from metal oxide
particles [16] to carbon nanotubes [17]. EPD is also relevant to the development of graphene-based
coatings in a cost-effective manner [13,18]. Its capacity to be used in more complex, integrated electrode
systems is an advantage over drop-casting and other less defined deposition techniques [10,11,19–24].
The group of Boccaccini added intensively to this field by deepening our understanding of EPD
through the investigation of GO-EPD kinetics as a function of deposition time and potential [12].
To design a powerful electrochemical sensor, working with highly reduced graphene oxide
nanosheets formed by EPD is required. Cathodic EPD would be the preferential approach, as it
allows the simultaneous deposition and reduction of GO to rGO [9]. The presence of carboxyl and
hydroxyl functions on GO results in an overall negatively charged material of about −41.3 ± 0.8 mV
for aqueous GO suspensions. Migration to the anode rather to the cathode occurs upon applying
a DC voltage [13]. The anodic EPD of GO results in GO with a low degree of reduction during the
deposition process. Thermal or chemical reductions [25] are necessary to restore the aromatic network
in order to obtain a material with good electron transfer properties. This is one of the reasons why GO
is often charged with a cationic polymer (e.g., polyethyleneimine) [8] or metallic cations (Ni2+ , Cu2+ ,
etc.) [21–23] to achieve a positively charged GO nanomaterial, which can be deposited by cathodic EPD.
The presence of polyethyleneimine (PEI) has been shown to be advantageous for the integration of
surface ligands and the formation of an immunosensor for the selective and sensitive electrochemical
detection of uropathogenic Escherichia coli [8], the detection of dopamine in meat [20] and Ni2+ for the
construction of non-enzymatic glucose sensors operating in a basic medium. However, the formation
of well-reduced rGO by EPD free of metal ions and other surface ligands remains a challenge.
In this study, we evaluate the effect of applied electrical current and applied voltage on the
electrochemical behavior of electrophoretically deposited rGO on gold thin film electrodes. It is shown
that the use of a voltage bias of 2.5 V for 5 min results in rGO thin films of good electrochemical behavior.
These films can also be submitted to further surface modification using diazonium electrochemistry
without altering their adhesion characteristics.
2. Materials and Methods
2.1. Materials
Potassium hexacyanoferrate(II) ([K4 Fe(CN)6 ]), hydrazine hydrate, phosphate buffer tablets
(PBS, 0.1 M), tetrabutylammonium fluoride (TBAF), ferrocenemethanol, copper(II) sulfate (CuSO4 ),
L-ascorbic acid, EDTA and N-butylhexafluorophosphate (NBu4 PF6 ) were purchased from
Sigma-Aldrich and used as received. Graphene oxide (GO) powder was purchased from Graphenea,
Spain. 4-((triisopropylsilyl)ethylenyl)benzenediazonium tetrafluoroborate (TIPS-Eth-ArN2 + ) was
synthesized as reported previously [26].
Azidomethylferrocene was synthesized according to Reference [27].
Au thin film electrodes were prepared by thermal evaporation of 5 nm of titanium and 40 nm of
gold onto cleaned glass slides.
2.2. Electrophoretic Deposition
Before electrophoretic deposition, the gold electrode was cleaned by UV/ozone for 5 min, rinsed
with acetone and water and dried under a nitrogen flow. The deposition took place in a two-electrode
system with a platinum foil (1 cm2 ) as the cathode and the cleaned gold surface as the anode (0.5 cm2 ).
The electrodes were placed in parallel to each other at a fixed distance of 1.5 cm. An aqueous GO
solution of 1 mg mL−1 was used for the EPD. Voltage biases of 1.25 V, 2.5 V, 5 V or 10 V were applied
using a potentiostat/galvanostat (Metrohm Autolab, Utrecht, The Netherlands) for 5 min. The modified
gold electrodes were slowly withdrawn manually from the solution and dried in a horizontal position
under ambient conditions for 1 h. After the deposition was complete, the modified electrodes were
placed into a Teflon autoclave (45 mL) and sealed with 1 mL of hydrazine hydrochloride. The autoclave

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was heated to 80 ◦ C and kept under a constant temperature for 4 h. The interfaces were then removed
and gently washed with water.
2.3. Surface Modification
2.3.1. Diazonium Chemistry
The electrografting of 4-((triisopropylsilyl)ethylenyl)benzenediazonium tetrafluoroborate (TIPSEth-ArN2 + ) (1 mM) in 0.1 M NBu4 PF6 in acetonitrile was performed using cyclic voltammetry with a
scan rate of 50 mV s−1 for five cycles between +0.60 V and −0.75 V vs. Ag/AgCl. The electrodes were
rinsed with copious amounts of acetonitrile and acetone and dried under a stream of argon.
2.3.2. “Click” Chemistry
Before “click” chemistry, the TIPS protection group was removed by the immersion of the
Au/rGO-TIPS surface into tetrabutylammonium fluoride (TBAF, 0.05 M in THF) for 20 min. The Au/
rGO interface was then immersed into an aqueous solution of CuSO4 (10 mM) and L-ascorbic acid
(20 mM) in the presence of azidomethylferrocene (0.83 mM in THF) and left for 1 h under an argon
atmosphere. The interface was then treated with an aqueous solution of EDTA for 10 min to chelate
any remaining Cu2+ residues and finally washed copiously with acetone and water and left to dry.
2.4. Surface Characterization Techniques
2.4.1. Scanning Electron Microscopy (SEM)
SEM images were obtained using an electron microscope ULTRA 55 (Zeiss, Paris, France)
equipped with a thermal field emission emitter and three different detectors (EsB detector with
filter grid, high efficiency in-lens SE detector and Everhart–Thornley Secondary Electron Detector).
2.4.2. X-Ray Photoelectron Spectroscopy (XPS)
X-ray photoelectron spectroscopy (XPS) was performed in a PHl 5000 VersaProbe-Scanning
ESCA Microprobe (ULVAC-PHI, Osaka, Japan) instrument at a base pressure below 5 × 10−9 mbar.
Core-level spectra were acquired at a pass energy of 23.5 eV with a 0.1 eV energy step. All spectra
were acquired with 90◦ between the X-ray source the and analyzer. After the subtraction of the
linear background, the core-level spectra were decomposed into their components with mixed
Gaussian−Lorentzian (30:70) shape lines using CasaXPS software. Quantification calculations were
conducted using sensitivity factors supplied by PHI.
2.4.3. Electrochemical Measurements
Electrochemical measurements were performed with a potentiostat/galvanostat (Metrohm
Autolab, Utrecht, The Netherlands). A conventional three-electrode configuration was employed using
a silver wire and a platinum mesh as the reference and auxiliary electrodes, respectively.
2.4.4. Micro-Raman Analysis
Micro-Raman spectroscopy measurements were performed on a LabRam HR Micro-Raman
system (Horiba Jobin Yvon, Palaiseau, France) combined with a 473 nm laser diode as the excitation
source. Visible light was focused by a 100× objective. The scattered light was collected by the same
objective in backscattering configuration, dispersed by a 1800 mm focal length monochromator and
detected by a CCD.
2.4.5. Atomic Force Microcopy (AFM)
Tapping mode AFM images in air and ambient temperature were recorded using a Bruker
Dimension 3100 AFM (Bruker, Champs-sur-Marne, France). The surfaces were imaged with a silicon
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cantilever (AppNano TM300, typical spring constant: 50 N/m) working at a frequency of 369 kHz.
Image treatment and root mean square (RMS) roughness Ra were obtained with WSXM software
(Bruker, Champs-sur-Marne, France). Surface roughness of the samples was measured by scanning
over a 5 × 5 μm area.
2.4.6. Profilometry
The thickness of the deposited films was determined by using an optical profilometer
(Zygo NewView 6000 Optical Profilometer with MetroPro software). This equipment uses non-contact,
three-dimensional scanning white light and optical phase-shifting interferometry, has vertical z-scan
measurements ranging from 0.1 nm to 15,000 μm and has capabilities of 1 nm height resolution with
step accuracies better than 0.75%. Images were taken with a 10× lens with a 14 mm field of view.
3. Results
3.1. Electrophoretic Deposition
The process for the formation of stable electrochemically active reduced graphene oxide (rGO)
thin films on gold thin film electrodes is schematically depicted in Figure 1A. It is based on a two-step
process in which, after EPD at 2.5 V for 5 min, the full reduction of GO to rGO is obtained by immersion
into hydrazine vapor for 4 h. The use of a potential of 2.5 V proved to be of high importance in the
process. Initial investigation revealed that the electrical current signature changes significantly at
voltages greater than 2.5 V (Figure 1B), where the electrical current passing through the interface
increases during the first 25 s before stabilizing. In contrast, at potential biases of 2.5 V and lower,
a decrease of the electrical current is first observed before its stabilization. Indeed, due to the insulating
character of GO, the electrical current should decrease during the deposition process as observed for
a low voltage bias. The current drop at 2.5 V is slower than that at 1.25 V. The reason for this is not
well understood, but could be due to a partial electrochemical reduction of GO under these conditions.
A similar behavior was observed by Diba et al. following the deposition of GO at 3 and 5 V [12].
The increase in current at elevated voltages indicates that next to the material deposition an
electrochemical reaction occurs, delaying surface passivation [12]. Visual inspection of the interfaces
(Figure 1C) shows clearly that the deposition occurred. However, a closer visual inspection of the
electrical interfaces (Figure 1D) reveals that the gold thin films were partially destroyed, most likely by
the gas evolution during water hydrolysis occurring in parallel at these voltage biases. This results in
the swelling of the GO deposit and eventual gold film rupture with poor film attachment.
Figure 2A shows SEM images of the different interfaces. While the films formed at 1.25 V display
granular like structures, the films deposited at 2.5 V exhibit the typical rGO-like wrinkle structures
with no evident local surface inhomogeneities and a complete and even coating of the gold thin film
interface. From the cross-section image (Figure 1A), no protrusions and hollow internal structures are
visible, indicating that at a lower potential local inhomogeneities associated with rGO formation are
avoided. The average surface roughness (Ra ), as determined by tapping-mode atomic force microscopy
(AFM) measurements, changed from Ra = 4.1 nm (Au) to Ra = 3.5 nm for Au/rGO (2.5 V deposition
potential).

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Surfaces 2019, 2

(A)

(B)

(C)

(D)
Figure 1. (A) Schematic illustration of the formation of an electrochemically active reduced graphene
oxide (rGO) thin film by anodic electrophoretic deposition (EPD) from an aqueous graphene oxide
suspensions (GO, 1 mg/mL) at a potential bias of 2.5 V for 5 min, followed by chemical treatment in
hydrazine vapor; cross-sectional SEM image of a gold interface coated with rGO using EPD at 2.5 V
for 5 min. (B) Change of current as a function of deposition time using different applied potentials:
concentration of GO (1 mg/mL) in water (pH 7.5), deposition interface: gold thin film electrode.
(C) Photographic images of gold thin film electrodes before (0.00 V) and after anodic EPD of an aqueous
solution of GO (1 mg/mL, pH 7.5) for 5 min at 1.25, 2.50, 5.00 and 10.00 V. (D) Photographs of the
different interfaces obtained (objective: 10×, numerical aperture: 0.9).

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Surfaces 2019, 2

To estimate the anodic efficiency of the EPD process, the gold interface was weighed before and
after rGO deposition and reduction. To convert the deposited weight (mrGO, 10 μg for 2.5 V for 5 min)
into a film thickness (drGO ), Equation (1) was used:
drGO (nm) =

mrGO×
A×ρ

(1)

where A is the surface area (0.8 cm2 in our case) and ρ is the density of rGO (2.09 g cm−3 ) [28]. A film
thickness of 59 nm was determined. This is in agreement with profilometry measurements where
a film thickness of about 55 nm was determined (Figure 2B). Samples prepared at 1.25 V for 5 min
had a deposited mass of 3 μg and an estimated rGO film thickness of 20 nm, below the acceptable
accuracy limit of profilometry to be validated. Increasing the potential to 5 V resulted in 120 nm
thick films. The Hamakar model [29] correlates the time-dependent (t = 300 s) amount of deposited
materials (mrGO in g) with the electrical field strength (E, 1.6 V m−1 in our case), the surface area of the
electrode (0.8 cm2 ), the electrophoretic mobility of GO (1.97 × 10−4 cm2 /(Vs)−1 ), the concentration of
the particle suspension (1 mg mL−1 , 0.001 g cm−3 ) and the anodic efficiency factor f (Equation (2)) [29].
mrGO = crGO × A × μ × E × t × f

(2)

In the case of f = 1, the amount of deposited rGO should be equal to mrGO = 47 μg. The determined
rGO amount, however, was only 10 μg, which implies an efficiency factor of f = 0.2.
As a voltage bias of 2.5 V seems to be the best condition for the EDP of rGO films onto the
gold electrodes, these interfaces were investigated in greater detail. The electrochemical behavior
of this interface using ferrocene-methanol redox couple in an aqueous solution is depicted in
Figure 2C. Whereas on a bare gold surface a fully reversible voltammogram was observed, an
irreversible voltammogram with a very small oxidation peak was observed on GO-coated gold
surfaces. To improve the electrochemical behavior of such coated gold surfaces, the interfaces were
further treated with hydrazine vapor, known for its strong reducing power. The cyclic voltammograms
of hydrazine treated surfaces were largely improved, showing a well-defined redox couple with
increased capacitance behavior as expected for rGO materials (Figure 2B, blue curve).
The chemical composition of the deposited graphene matrix was further evaluated using XPS
(Figure 2D). The high resolution C1s core level spectrum of the initial GO suspensions showed
contributions at 284.2 (C=C sp2 ), 285.0 (C–H/C–Csp3 ), 286.7 (C–O) 288.7 (C=O) and a small
contribution at 291.0 (O–C=O). The C1s XPS core spectra of GO-coated gold surfaces revealed similar
contributions with different intensities. In particular, the band at 284.2 eV due to C=C sp2 increased
compared to the GO solution, indicating the partial restoration of the sp2 network of the deposited
graphene material. The XPS of GO-coated gold showed additional bands at 285.0 and 286.7 eV with
a large band at 288.0 eV due to C=O. The hydrazine reduction of the GO–gold interfaces resulted
in a decrease of the epoxy/ether functions at 286.7 eV, in accordance with a partial reduction of GO
mostly likely due to the elimination of CO by a Kolb-like mechanism [12,13]. The ketone groups
remained preserved.
The Raman spectrum of GO-coated gold surfaces before and after hydrazine reduction (4 h)
is presented in Figure 2E. The increase of the D/G ratio from 0.86 to 0.97 after hydrazine chemical
reduction was observed.
Figure 3A summarizes the electrochemical behavior of a neutral redox species, ferrocene methanol,
on bare and GO-coated gold surfaces after the hydrazine chemical reduction. The electrochemically
active surface area of the different electrodes was determined by plotting the peak current as a function
of the square root of the scan rate (Figure 3B), according to Equation (3) [30]:
A = slope/(268600 × n3/2 × D1/2 × c)

132

(3)

Surfaces 2019, 2

where n is the number of electrons transferred in the redox event (n = 1), D is the diffusion coefficient of
ferrocene methanol (7.5 × 10−6 cm2 s−1 ) and c is the concentration of ferrocene methanol (1 10−6 mol
cm−3 ). Taking into account the experimentally determined slopes (9.265 × 10−5 AV1/2 s−1/2 for Au
and 9.980 × 10−5 AV1/2 s−1/2 for Au/rGO), the active surfaces of 0.126 cm2 (naked gold) and 0.136 cm2
(Au/rGO) were determined. By plotting the current density vs. potential (Figure 3B), a larger current
density was detected on the EPD coated gold interface due to the excellent electrochemical behavior of
the interface. This is in agreement with the deposition of an electrochemically active 3D rGO material.

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Figure 2. (A) SEM image of gold thin films before and after EPD from an aqueous graphene oxide
suspension (GO, 1 mg/mL) at different potential biases for 5 min. (B) Thickness of reduced graphene
oxide as a function of applied potential (t = 5 min). (C) Cyclic voltammograms recorded on gold
thin film electrodes (black), after EPD from an aqueous graphene oxide suspension (GO, 1 mg/mL)
at a potential bias of 2.5 V for 5 min (red); after further reduction with hydrazine (blue) using
ferrocenemethanol (1 mM)/PBS (0.1 M), scan rate = 100 mV s−1 . (D) C1s high-resolution spectra
of GO (black), EPD GO (red) and further reduced GO (blue). (E) Raman spectra of the EPD film formed
at 2.5 V before (red) and after hydrazine reduction (blue).

133

Surfaces 2019, 2

(A)

(B)

Figure 3. (A) Cyclic voltammograms recorded on gold thin film electrodes (black) and after EPD and
further reduction with hydrazine using ferrocene methanol (1 mM)/PBS (0.1 M), scan rate = 100 mV
s−1 . (B) Cyclic voltammograms recorded on gold thin film electrodes (black) and after EPD and further
reduction with hydrazine using ferrocene methanol (1 mM)/PBS (0.1 M), scan rate = 100 mV s−1 .

3.2. Surface Modification Using Diazonium Electrografting
To validate the stability of the formed interface, efficient covalent modification based on the
electroreduction of a triisopropylsilyl-protected ethynyl diazonium salt was performed (Figure 4A).
Covalent surface modification of graphene-based materials using the electroreduction of aryldiazonium
salts is a popular approach as it allows the introduction of different chemical groups. The major
drawback of this approach is the difficulty to control the extent of the reaction—notably, to limit the
reaction to the formation of a functional monolayer. The highly reactive nature of the formed aryl
radical results in the formation of disordered polyaryl multilayers, which in the case of a sensor might
limit the dynamic range of sensing. The introduction of bulky substitutions on the ArN2 + moieties
limits radical addition reactions and allows the formation of ultrathin functional layers [26,31,32].
On the basis of this concept, the precursor 4-((triisopropylsilyl)ethylenyl)benzenediazonium
tetrafluoroborate (TIPS-Eth-ArN2 + ) was used for the formation of an organic thin film on Au/rGO.
The layer was obtained by potential cycling in a solution containing TIPS-Eth-ArN2 + (Figure 4B).
As observed by cyclic voltammetry, the reduction peak of TIPS-Eth-ArN2 + at −0.2 V decreased
rapidly after five cycles, indicating the blocking of the Au/rGO electrode. Compared to glassy carbon
electrodes, the reduction peak shifted to more negative potentials [26]. The blocking properties of the
layer was further investigated by recording the cyclic voltammogram of the oxidation of ferrocene
methanol in water before and after the electrografting process (Figure 4C) and was found to be typical
of a totally blocked electrode.
To have access to the acetylene function, the modified Au/rGO-TIP surface was immersed
into tetrabutylammonium fluoride (TBAF, 0.05 M in THF) at room temperature. The success of the
deprotection step was evidenced by the recovery of the ferrocene oxidation signal, being almost
identical to that obtained on Au/rGO (not shown). The deprotected substrate was further treated by
“click” chemistry (Huisgen 1,3-dipolar cyclization) using azidomethylferrocene, as first reported by
Gooding et al. [33]. The success of the integration of ferrocene units onto the Au/rGO electrodes was
seen from the presence of the Fe2p component in the XPS survey spectrum (Figure 5A). In the absence
of copper catalyst, no click reaction took place and no Fe2p component could be recorded. Figure 5B
shows the cyclic voltammogram of the ferrocene-modified electrode examined in electrolytic ethanol
solution. A surface concentration of Γ = 2.5× 10−10 mol cm−2 of bound ferrocene groups was derived
from these measurements using Equation (4):
Γ = Q/nFA

(4)

where Q is the passed charge, n is the number of exchanged electrons (n = 1), F is the Faraday constant
and A is the electroactive surface of the electrode determined as 0.136 cm2 . Considering the ferrocene
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Surfaces 2019, 2

molecules as spheres with a diameter of 6.6 Å, the theoretical maximum coverage for an idealized
ferrocene monolayer can be estimated as Γ = 4.8 × 10−10 mol cm−2 [34]. In our case, about half of the
full coverage was achieved.

(A)

(B)

(C)

Figure 4. (A) Covalent modification of Au/rGO electrodes using a silyl-protected diazonium
salt followed by tetrabutylammonium fluoride (TBAF)-based deprotection and “click” reaction
using azidomethylferrocene as a model compound. (B) Electroreduction of 4-((triisopropylsilyl)
ethylenyl)benzenediazonium tetrafluoroborate (TIPS-Eth-ArN2 + ) on Au/rGO, scan rate 50 mV s−1 .
(C) Cyclic voltammograms of ferrocene methanol (1 mM)/0.1 M PBS of rGO/Au before (blue) and
after modification with TIPS-Eth-ArN2 + .

(A)

(B)

Figure 5. (A) High-resolution Fe2p X-ray photoelectron spectroscopy (XPS) core-level spectrum of
ferrocene modified Au/rGO interfaces. (B) Cyclic voltammograms in acetonitrile/ NBu4 PF6 0.1 M),
scan rate 100 mV s−1 .

135

Surfaces 2019, 2

4. Conclusions
In conclusion, by applying a low EPD voltage of 2.5 V, the deposition of partially reduced
graphene oxide film can be formed on gold thin film electrodes in a homogenous, and mechanically
and chemically stable manner. Further reduction of the films in hydrazine vapor results in reduced
graphene oxide-coated substrates showing excellent electron transfer characteristics. The films proved
to be of high robustness and could be further modified via the electroreduction of aryldiazonium
salts. Using a protection–deprotection approach allows the immobilization of a functional monolayer
which can be further functionalized by Cu(I)-catalyzed “click” chemistry. We have demonstrated that
a densely packed ferrocenenyl monolayer can be efficiently integrated on these rGO-modified surfaces.
Due to the versatility and mild conditions of the “click” chemistry reaction, a wide range of functional
groups can be immobilized on such surfaces. These results pave the way for the use of this technology
for the modification of more complex electrode configuration such as screen-printed or flexible
integrated electrode arrays and the integration of surface ligands for sensing-related applications.
Author Contributions: V.M.: electrophoretic deposition and electrochemistry; P.A.: click chemistry; Y.L.: surface
chemistry; H.H.: characterization; W.K.: editing of the draft; R.B.: characterization, editing of the draft; S.S.;
original draft preparation, supervision of work.
Acknowledgments: General financial support from the Centre National de la Recherche Scientifique (CNRS),
the University of Lille, the Hauts-de-France region and the EU through FLAG-ERA JTC2015 “Graphtivity”
is acknowledged.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

138

Article

Effects of the Interfacial Structure on the Methanol
Oxidation on Platinum Single Crystal Electrodes
Mohammad Ali Kamyabi 1,2 , Ricardo Martínez-Hincapié 2 , Juan M. Feliu 2, * and
Enrique Herrero 2, *
1
2

Department of Chemistry, Faculty of Science, University of Zanjan, Zanjan 45371-38791, Iran;
makamyabi@gmail.com
Instituto de Electroquímica, Universidad de Alicante, Apdo. 99, E-03080 Alicante, Spain;
ricuardomartin@gmail.com

Received: 11 February 2019; Accepted: 1 March 2019; Published: 12 March 2019

Abstract: Methanol oxidation has been studied on low index platinum single crystal electrodes
using methanol solutions with different pH (1–5) in the absence of specific adsorption. The goal
is to determine the role of the interfacial structure in the reaction. The comparison between the
voltammetric profiles obtained in the presence and absence of methanol indicates that methanol
oxidation is only taking place when the surface is partially covered by adsorbed OH. Thus, on the
Pt(111) electrode, the onset for the direct oxidation of methanol and the adsorption of OH coincide.
In this case, the adsorbed OH species are not a mere spectator, because the obtained results for the
reaction order for methanol and the proton concentrations indicate that OH adsorbed species are
involved in the reaction mechanism. On the other hand, the dehydrogenation step to yield adsorbed
CO on the Pt(100) surface coincides with the onset of OH adsorption on this electrode. It is proposed
that adsorbed OH collaborates in the dehydrogenation step during methanol oxidation, facilitating
either the adsorption of the methanol in the right configuration or the cleavage of the C—H bond.
Keywords: methanol oxidation; platinum single crystals; pH and concentration effects; adsorbed OH

1. Introduction
The methanol (MeOH) oxidation reaction has attracted significant attention in the last decades,
both from a fundamental and applied point of view [1–5]. From the fundamental approach, methanol
is a small organic molecule and the knowledge acquired by studying its oxidation mechanism can
be later transferred to more complex organic molecules. In fact, its mechanism can serve as a model
for the oxidation of more complex alcohols, such as ethanol or ethylene glycol. On the applied side,
methanol can be used as a fuel in fuel cells. When compared with hydrogen as fuel, methanol presents
several practical advantages: It is liquid at room temperature with a high volumetric energy density
and does not require the deployment of new distribution channels for its use. However, the anodic
over-potentials are significantly higher than those measured for the hydrogen oxidation reaction.
For this reason, the development of more effective electrocatalysis for the oxidation reaction is needed
if methanol is going to be employed as fuel. The intelligent design of more efficient electrocatalysts
requires a deep knowledge of the reaction mechanism.
Among all pure metals, platinum is the best electrocatalyst for the oxidation reaction. In the
complete oxidation, six electrons are exchanged according to:
CH3 OH + H2 O → CO2 + 6H+ + 6e

(1)

It is now generally accepted that the oxidation mechanism takes place following a dual path
mechanism [1]. In one of the paths, adsorbed CO is formed, which is considered a poisoning intermediate,
Surfaces 2019, 2, 177–192; doi:10.3390/surfaces2010014

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Surfaces 2019, 2

because it is strongly adsorbed on the surface and its oxidation requires high potentials. On the other
hand, the second path avoids the formation of CO and goes through an active intermediate [1]. In a
reaction in which six electrons are exchanged, several intermediates are involved, and as has been
proposed, the two paths may share some intermediates [6,7]. As happens in any reaction in which
adsorbed intermediates are involved, the reactivity is very dependent on the surface structure [8–14].
Moreover, not only the surface structure, that is, the local arrangement of the surface atoms, affects
reactivity, but also the interfacial structure and composition. First, the comparison between ultra-high
vacuum and electrochemical environments indicates that the bonding mode of the methanol molecule is
clearly affected by the presence of the interfacial water [10] because the proposed intermediates and their
binding modes are different. Second, the presence of specifically adsorbed anions on the electrode surface
inhibits the reactivity [11]. Thus, currents in the presence of sulfate or phosphate are significantly smaller
than those measured in perchloric acid solutions. Additionally, other factors that affect the interfacial
structure should be considered, such as, solution pH, and electrode charge. In all the cases, all these
parameters modify the interaction between the methanol molecule and the surface, which in turn controls
the reaction mechanism.
In order to have a deeper knowledge of the effects of interfacial on the methanol oxidation reaction,
experiments in solutions with pH values ranging between one and five have been carried out. Due to
the large impact of the specific adsorption of anions on the methanol oxidation currents [11], electrolyte
solutions should have an adequate buffering capacity without containing anions that can adsorb
strongly on the surface. For that purpose, mixtures of perchloric acid and fluoride have been used
because perchlorate and fluoride anions do not adsorb strongly on the electrode surface and the pKa of
HF/F- is 3.15 [15,16]. Furthermore, to better explore the relationship between the interfacial structure
and the reactivity, methanol concentrations between 10−4 and 10−1 M have been used. The low
concentration regime will allow detecting the modification of the adsorption behavior of the electrode
due to the presence of methanol and relating them with the oxidation mechanism and its selectivity.
2. Materials and Methods
Low index platinum single crystal electrodes, namely Pt(111), Pt(100), and Pt(110), were fabricated
from small single crystal beads (ca. 2.5 mm diameter). These beads were oriented, cut, and polished until
mirror finish using the procedure described by Clavilier [17,18]. Prior to any experiment, the electrodes
were cleaned by flame annealing, cooled down in a H2 /Ar atmosphere and protected with water
in equilibrium with this gas mixture [19]. The protected electrodes were immersed in the electrolyte
solution at a controlled potential (0.1 V) in meniscus configuration so that only the surface with the desired
orientation is in contact with the solution. For the Pt(111) and Pt(100), this procedure leads to the formation
of a well-ordered surface with the nominal orientation [20,21]. Due to the reconstruction phenomena
undergone by the Pt(110) surfaces, which is strongly affected by the treatment, the surface presented
here corresponds to the (1 × 2) or (1 × 1)/(1 × 2) structure with major (1 × 2) contribution [22,23]. Thus,
the observed behavior should correspond with that expected for the (1 × 2) surface, although contributions
from the (1 × 1) cannot be discarded. The cleanliness of the electrode surface and the solutions were
checked by voltammetry in the absence of methanol. The characteristic voltammetric profiles of the
different electrodes, which are stable upon cycling when clean solutions are used, were obtained.
Experiments were carried out at room temperature, 22 ◦ C, in classical two-compartment
electrochemical cells de-aerated by using Ar (N50, Air Liquide, Paris, France, in all gases used),
including a large platinum counter electrode and a Reversible Hydrogen (N50) Electrode (RHE)
as reference. The solutions were prepared using concentrated HClO4 (for analysis, Merck KGaA,
Darmstadt, Germany), NaF (99.99%, Suprapur, Merck KGaA, Darmstadt, Germany), KClO4 (Suprapur,
Merck KGaA, Darmstadt, Germany), and ultrapure water (Purelab Ultra, Elga, High Wycombe, United
Kingdom). Solutions with a pH between two and five were prepared using mixtures of NaF and HClO4
so that the buffering capacity is high enough to maintain the interfacial pH constant. Voltammograms

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were recorded using a potentiostat (EA161, eDAQ, Sydney, Australia) and a digital recorder (ED401,
eDAQ, Sydney, Australia) in a hanging meniscus configuration.
3. Results
3.1. Pt(111) Electrode
Figure 1 shows the voltammetric profiles of the Pt(111) electrode in 0.1 M HClO4 with two different
methanol concentrations: 10−2 and 10−4 M and the comparison with that recorded in the absence of
methanol for the same electrolyte solution. The comparison between the blank voltammogram of the
Pt(111) electrode and those obtained in the presence of methanol allows correlating the interfacial
structure with the reactivity. Moreover, relevant information regarding the oxidation mechanism can be
obtained from this comparison, depending on the examined region. As aforementioned, it is generally
accepted that methanol oxidation reaction occurs through a dual path mechanism, in which a path
involves the formation of adsorbed CO, a poisoning intermediate, and a second path going through
an active intermediate. Due to the high number of electrons exchanged in the process, both paths
should be interconnected and can share some intermediates leading to a complex mechanism [6,7].
The voltammetric profile in the supporting electrolyte can be divided into three regions, hydrogen
adsorption region below 0.4 V, the OH adsorption region above 0.6 V, and the double lager region
between 0.4 and 0.6 V. The analysis of the modifications of the profiles in these regions will allow
obtaining relevant data for the oxidation mechanism.
The profile measured for 10−2 M has the typical shape of those obtained at higher methanol
concentrations, with an effective onset potential of ca. 0.6 V [14,24], which coincides with the onset
of the OH adsorption process. The large currents measured for this concentration prevents detailed
analysis of the effects of the surface composition in the methanol oxidation in this region. To diminish
oxidation currents and to establish whether the OH adsorption process is affected by the reaction,
currents were recorded for 10−4 M methanol. For this concentration (Figure 1B), modifications of
the blank voltammetry recorded in the absence of methanol are only observed in the region where
OH is adsorbed. Thus, in the positive scan direction, a new peak appears at ca. 0.72 V, apparently
superimposed with the original OH adsorption signal, because some of the characteristic features of
this region, such as the spike at 0.80 V, are still visible. On the other hand, the profile in the negative
scan direction shows complex behavior, which is clearly a consequence of the addition of two different
signals: the negative currents associated with the reductive desorption of adsorbed OH and the
positive current of the methanol oxidation reaction. To determine whether OH adsorption process
is affected by methanol, the blank voltammetric profile obtained in the supporting electrolyte has
been subtracted from that measured for the 10−4 M methanol concentration (Figure 1C). As can be
seen, the resulting voltammetric profile has the same shape as that recorded with larger methanol
concentrations, and the only difference is the absolute values of the currents, which are proportional to
the methanol concentration. It should be stressed that a spike at 0.80 V in the blank voltammogram,
which has been associated to changes in the OH adsorption mode depending on the water structure,
is very sensitive to the actual composition of the solution [25]. The presence of traces of species in
concentrations well below 10−6 M alters the currents in this region. In this case, the subtraction of the
two profiles shows a bipolar peak in this region that is only due to the change in the composition of
the interface, but not related to the oxidative process of methanol. Then, the subtracted voltammetric
profile demonstrates that the OH adsorption process between 0.6 and 0.9 V is also occurring in the
presence of methanol and that the interfacial structure of the electrode is not significantly affected by
the presence of methanol.

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Figure 1. Voltammetric profiles of the Pt(111) electrode in (A) 0.1 M HClO4 + 10−2 M MeOH and
(B) 0.1 M HClO4 + 10−4 M MeOH. The red trace shows the profile in absence of MeOH. (C) Subtraction
of the profiles in panel B. Scan rate: 50 mV s−1 .

These results clearly show that significant oxidation currents are only recorded in the region where
OH is adsorbed, implying that the oxidation of methanol is taking place on a surface partially covered
by OH. Moreover, during the spike at 0.80 V, there is a change in the reactivity and the completion of
the OH layer at 0.85 V leads to the inhibition of the methanol oxidation reaction. Thus, according to
these results, adsorbed OH seems to be required for the oxidation of methanol, but also some free sites
where the methanol molecule could interact with the electrode.
The inspection of the hydrogen adsorption and double layer regions can be used to provide
information on the type of intermediates that are formed during the oxidation reaction. Probably, CO
is the most important intermediate, because its oxidation, which occurs through a well-characterized
Langmuir-Hishenlwood mechanism, requires the presence of adsorbed OH on the surface [26]. Thus,

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rates for the CO oxidation reaction on Pt are only significant above 0.7 V, and if CO is formed below this
potential value, adsorbed CO molecules are accumulated on the surface [12,14,27]. The accumulation
of CO on the surface blocks hydrogen adsorption and the hydrogen adsorption charge recorded below
0.4 V diminishes [28]. As can be seen, in 10−4 M methanol, the accumulation of CO is negligible
because the profile in the hydrogen region for this concentration overlaps with that recorded in the
absence of methanol. This fact implies that if CO is formed during the oxidation, it is oxidized at
the same rate that is produced and the accumulation during the time of the experiment is negligible.
For higher concentrations, there is a small diminution in the hydrogen adsorption charge, implying that
some CO has been accumulated on the electrode surface. However, the amount of CO accumulated
is very small. In fact, chronoamperometric measurements indicate that CO formation takes place
exclusively above 0.5 V and that the rate for this surface is the lowest of all the single crystal electrodes
studied [14].
Additionally, other possible intermediates may be formed and adsorb on the electrode surface
giving rise to characteristic signals. One of the possible intermediates in the oxidation of methanol is
formic acid. However, the observed behavior is not compatible with the formation of a measurable
amount of formic acid. In the acidic media, formic acid adsorbs on the platinum electrode as formate
in the bidentate configuration. The adsorption of formate modifies the voltammogram so that the
typical shape corresponding to OH adsorption above 0.6 V disappears and a new signal between 0.4
and 0.6 V can be observed [29]. Moreover, formic acid is readily oxidized on the Pt(111) electrode
above 0.35 V [30,31]. None of these new signals appear on the recorded voltammograms, even for
the highest recorded concentration (0.1 M). Thus, the measured profile between 0.4 and 0.6 V in the
presence of methanol matches that recorded in its absence, implying that formic acid is not being
formed at significant concentrations during the oxidation of methanol on this electrode surface in the
present conditions. It should be mentioned that a small fraction of formic acid has been detected by
HPLC analysis after long electrolysis time (ca. 15 min) and with high methanol concentrations [32].
Clearly, the electrolysis time and the solution concentration affect the product distribution. On the
other hand, the presence of methyl formate detected by mass spectroscopy [13,27] appears to be a
consequence of the attack of a methanol molecule in solution to some intermediate species adsorbed
on the surface and not as a result of the esterification reaction between solution methanol and formic
acid produced in the partial oxidation of methanol [33,34].
From all these results, it can be concluded that the absence of significant modifications in
the interfacial structure, in the presence of methanol (aside that relate to the small amount of CO
accumulated on the surface), indicates that the different adsorbed intermediates of the reaction are
not strongly adsorbed and their interaction with the surface must be weak. In fact, when anions are
strongly adsorbed, such as sulfate or phosphate added to the supporting electrolyte, the oxidation
currents diminish significantly, because methanol adsorbates cannot compete for the adsorption sites
with the strongly adsorbed sulfate or phosphate [14,24]. However, the presence of adsorbed OH seems
to be required for the oxidation. To further prove this relationship between adsorbed OH and methanol
oxidation, the reaction order for methanol was calculated. For this, the current densities measured at
0.65 and 0.75 V after subtraction of the current density at these potentials in the blank voltammogram
are plotted vs. the electrode potential in a double logarithmic graph (Figure 2). As can be seen, a linear
relationship is observed for 0.65 V with a slope of 0.53. For 0.70 V, there is a clear deviation from
linearity, especially for 0.1 M methanol, where currents are smaller than those predicted from the
behavior observed for low concentrations, and the measured slope is 0.62. The fractional reaction
orders and the dependence with the electrode potential and OH coverage clearly indicate that the
reaction mechanism is complex, and OH should be involved in the process.

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Figure 2. Double logarithm plot for the current density measured at 0.65 and 0.70 V in 0.1 M HClO4 vs.
MeOH concentration.

Additionally, the comparison with the observed behavior for formic acid will also provide some
additional insight. The onset of formic acid oxidation (in the absence of previously formed CO) matches
the onset of formate adsorption because the adsorption of formate in the favorable configuration is
required for the oxidation [35,36]. This adsorption process starts at a potential where hydrogen is
desorbed and the adsorption strength of formate is larger than that of OH because OH adsorption
is blocked in its presence. On the other hand, the results presented here indicate that the interaction
of methanol (and that of the intermediates except for CO) is weak. For this reason, it would be
expected that the interaction of methanol with the surface was more favorable in the double layer
region because there are no other adsorbed species on the surface. However, the currents in this
region are negligible and only where OH is adsorbed, oxidation currents are recorded, reinforcing
the relationship between methanol oxidation and adsorbed OH. Additionally, it should be mentioned
that the detection of soluble intermediate products increases when sulfuric acid is used as supporting
electrolyte [13,27,32]. The presence of adsorbed sulfate prevents the adsorption of OH and the possible
intermediates, and thus, the oxidation of methanol becomes more difficult and increasing amounts of
partial oxidation products are formed.
Regarding the role of adsorbed OH in this process, previous results with isotopically labeled
methanol and water indicate that the cleavage of the C-H bond is involved in the rate determining
step [24]. Then, adsorbed OH should be involved in positioning the methanol molecule in the
right configuration for the cleavage and/or in interacting with one of the hydrogen atoms of the
methanol molecule in such a way that the activation energy of this (these) step(s) is (are) diminished.
This interpretation agrees with the higher amount of partial oxidation products obtained when the
adsorption of OH is hindered by the presence of specifically adsorbed anions. This is not the only
role of OH in the oxidation mechanism because it is also required for the oxidation of CO in the
Langmuir-Hinshelwood mechanism [26,37,38].
Additional effects of the interfacial structure in the oxidation reaction can be obtained by changing
the solution pH. Two interfacial parameters are affected, the electrode surface charge and the water
structure. The equilibrium potential for the methanol oxidation reaction shifts 59 mV per pH unit,
as the RHE scale does. For this reason, it is expected that the methanol oxidation reaction occurs in the
same potential region in the RHE scale. However, as the solution pH increases, the electrode potentials
for the oxidation in a pH-independent potential scale, such as the Standard Hydrogen Scale (SHE)
scale, shift to more negative potential values. Since the potential of zero free charge of the Pt(111)
surface is constant and pH-independent (0.28 V vs. SHE [39]), as the pH increases, the electrode charge

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becomes more negative, and thus, affecting the interaction of the species with the surface. Additionally,
water structure also changes, altering the interaction of the molecules with water, which in turn can
influence the reactivity.
To study those effects, the methanol oxidation reaction was investigated in different pH solutions
up to a pH value of 5 (Figure 3). As can be seen, the qualitative behavior of the electrode is independent
of the solution pH, and the oxidation of methanol is only taking place in the region where OH is
adsorbed. Thus, at 10−4 M methanol, an oxidation wave is observed as superimposed with the
OH adsorption in the positive scan direction, whereas in the negative direction, the profile is more
complex, as happens for pH = 1. Additionally, the profile in the hydrogen region overlaps with the
measured in the absence of methanol, which implies that CO is not accumulated on this surface.
For higher concentrations, the general shape of the voltammogram in methanol containing solutions
do not change with the pH, showing a broad wave, whose onset potential coincides with that of
OH adsorption and currents in the negative scan direction, which are slightly smaller than those
recorded in the positive direction. However, a detailed analysis shows some modifications with pH,
as discussed below.
100

pH=3
-4
10 M MeOH

pH=5
-4
10 M MeOH

j/μAcm

-2

50
0
-50

2000
pH=3
-2
10 M MeOH

pH=5
-2
10 M MeOH

j/μAcm

-2

1500
1000
500
0
0.0

0.2

0.4

0.6

0.8

E vs. RHE/V

0.0

0.2

0.4

0.6

0.8

E vs. RHE/V

Figure 3. Voltammetric profiles of the Pt(111) electrode in pH = 3 (left column) and pH = 5 (right
column) with different concentrations of MeOH. The red trace shows the profile in absence of MeOH.
Scan rate: 50 mV s−1 .

The detailed analysis of the different effects require that the profiles at different pH values are
compared. Figure 4 shows the voltammetric profiles of the Pt(111) electrode in the presence and
absence of methanol for the different pH values. As can be seen, the changes in the blank voltammetric
profile are small, especially for pH values lower than four. The onset for the OH adsorption is the
same and the only major modifications are observed for the spike, which tends to disappear and as
the solution pH increases, in a clear evolution to the profile recorded in alkaline media [40]. Despite
these minimal changes, the oxidation of methanol is clearly affected by the pH. As shown in Figure 4B,
the onset of the oxidation displaces to more positive potentials in the RHE scale and the current
increases. Clearly, the changes in the interfacial properties are affecting this reaction. Additional
information on that can be obtained calculating the reaction order for protons. For an irreversible
reaction, reaction orders should be calculated at a constant potential value, in this case, in the SHE
scale. Thus, when plotting the measured current at 0.5 V vs. SHE as a function of the pH in a double
logarithm scale, a linear relationship is obtained in which the slope is ca. 0.45 (Figure 5). This value
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implies that the reaction order with respect to the proton concentration is –0.45. While the fractional
orders are difficult to explain with simple mechanisms, the straightforward interpretation would
be that protons are released before the rate-determining step. This observed behavior is in good
agreement with the isotopic effect observed [11], which also indicate that the release of protons is
involved in the rate determining step. An additional effect that should be taken into account when
analyzing this behavior is the effect of the surface electrode charge. Since the potential of zero free
charge is 0.28 V vs. SHE, as the pH increases, the region in which methanol oxidation takes place
moves from positive charge to neutral values. Thus, the interaction between the species involved in
the reaction (OH and methanol) and the surface is modified and affects the reaction rates. In this sense,
for the oxygen reduction reaction, it has been proposed that higher rates obtained in neutral values are
related to the surface charge values [41]. For neutral pH values, the charge of the surface is neutral in
the onset of the ORR. The observed behavior for methanol shows a very similar dependence.

Figure 4. Dependence of the voltammetric profile of the Pt(111) with pH (A) in absence of methanol
vs. RHE, (B) in 10−2 M MeOH vs. RHE and (C) in 10−2 M MeOH vs. SHE. For methanol oxidation,
only the positive scan direction is shown. Scan rate 50 mV s−1 .

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Figure 5. Current density for the oxidation of methanol at 0.5 V vs. SHE as a function of the solution pH.

3.2. Pt(100) and Pt(110) Electrodes
The major difference between the behavior of the Pt(111) electrode and the other two basal planes
is the presence/absence of a pure double layer region where no specific adsorption occurs. For the
Pt(110) and Pt(100) electrodes, OH is readily adsorbed at low potentials when hydrogen is being
desorbed [23,42–45], and thus there is no pure double layer region. Then, any molecule that interacts
with the surface should either compete with hydrogen or OH for the adsorption sites. Additionally,
the overlapping of the two regions prevents the determination of the potential of zero free charge.
For the Pt(100) electrode (Figure 6), hydrogen is adsorbed on the surface between 0.05 and 0.45 V
and the OH adsorption occurs in a broad wave between 0.45 and 0.7 V, just after the desorption
of hydrogen from the surface [43,45]. In the presence of 10−4 M methanol, the only significant
modification is the appearance of a peak at ca. 0.65 V in the positive scan direction, and a small
diminution in the hydrogen adsorption charge. Unlike the Pt(111) electrode, this peak is only visible
in the positive scan direction, and thus should be related to the oxidation of some CO molecules
that have been accumulated on the electrode surface at low potentials. However, the amount of CO
accumulated on the surface is very small and negligible currents are recorded in the negative scan
direction. As the methanol concentration increases, this CO oxidation peak increases in charge and
shifts to more positive potentials. In parallel, hydrogen adsorption charge below 0.45 V diminishes
due to the blockage of hydrogen adsorption sites by CO. Additionally, a second oxidation wave after
this peak CO emerges both in the positive and negative scan directions. This wave shows a hysteresis
between both scan directions, which is clearly visible for the solution containing 10−2 M methanol,
and should be related to the accumulation of CO in the positive scan direction at potentials below 0.6 V.
Thus, the presence of adsorbed CO in the positive scan direction hinders the oxidation of additional
molecules of methanol on the surface. After the oxidation of CO, methanol molecules can interact with
the surface and give additional oxidation current contributions. In the negative scan direction, any CO
formed during the reaction can be readily oxidized at potentials higher than 0.6 V, giving rise to the
hysteresis between both scan directions.

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Figure 6. Voltammetric profiles of the Pt(100) electrode in 0.1 M HClO4 with different MeOH
concentrations. Inset: magnification of the positive scan direction in the region between o and 0.6 V

The process of CO formation on the electrode surface can be followed for the 10−3 M methanol
concentration (Figure 7). As can be seen, the voltammogram for the first scan in the presence of 10−3 M
methanol overlaps with that recorded in the absence of methanol until 0.45 V, implying that no CO
has been accumulated on the surface up to this potential value. Coinciding with the onset of OH
adsorption, some additional currents are observed. These currents should correspond to the formation
of CO from methanol according to the reaction
CH3 OH → CO + 4H+ + 4e,

Figure 7. First and second voltammetric scan for the Pt(100) electrode in 0.1 M HClO4 + 10−3 M MeOH.
The blue trace shows the profile in absence of MeOH. Scan rate: 50 mV s−1 .

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In fact, the analysis of the transient currents shows that CO is formed between 0.45 and 0.6 V [14],
implying that the formation of CO is only taking place in the region where OH is adsorbed. In the
negative scan direction, some additional currents are observed in this region, which are also related
to the formation of CO, as revealed by the significant diminution of the hydrogen adsorption charge
observed below 0.4 V. After this first scan, the second scan does not show the currents at 0.45 V, because
CO is already adsorbed on the surface. It should be noted that the following scans are identical to the
second scan. For the 10−2 M methanol concentration, the process of CO formation is not as visible
as in 10−3 M (inset of Figure 6). While currents are recorded between 0.4 and 0.6 V, the peak is not
well developed. It should be highlighted that, for this concentration, the results are very dependent
on how the meniscus is formed. This fact implies that CO is being formed during the immersion of
the electrode and the meniscus formation. While the electrode is immersed at a controlled potential
(0.1 V), a few ms are required to establish the imposed potential. Due to the higher concentrations
and the concomitant higher reaction rates, CO is being formed during this process. As can be seen
in Figure 6 for this concentration, hydrogen desorption states have been displaced to lower potential
values and hydrogen desorption charge is smaller, which is the characteristic behavior of a Pt(100)
surface partially covered by CO [46]. Thus, a lower amount of CO can be formed in the positive scan,
resulting in a less defined peak for this process.
Regarding the behavior of the methanol oxidation reaction on the Pt(100) electrode in different
pH solutions (Figure 8), the general behavior is the same as that recorded in 0.1 M HClO4. In the
positive scan direction, the peak associated with the oxidation of the accumulated CO can be observed,
followed by the presence of a second oxidation wave, which is also visible in the negative scan
direction. In general, no significant variation of the currents is observed, and the measured currents
in the broad wave in the negative scan direction are almost independent of the pH. The observed
changes, especially in the shape of the peak associated with the oxidation of CO and the broad wave
are probably related to the changes in the OH adsorption mode. The voltammograms in absence of
methanol show a progressive evolution from the profile measured at pH = 1 (Figure 7) to that obtained
in alkaline solutions [47]. While the OH coverage is not significantly affected by the pH, the differences
in the voltammograms reflect that the energetics of the OH adsorption process is altered by pH [47].
If, as proposed, OH involved in the reaction, the changes in the energetics should imply changes
in the kinetics, that in this cases are translated in changes of the voltammetric shape. On the other
hand, a reaction order cannot be obtained using this wave due to experimental problems. In this case,
the raising part of the curve in the negative scan direction (that appearing between 0.6 and 0.7 V)
should be used because it is not affected by the presence of adsorbed CO or any other inhibition process
related to the formation of a compact OH layer. However, this region is very is very narrow and when
curves are transformed in the SHE scale, the regions for the different pH do not overlap, preventing
the accurate measurement of the current for the different pH values at a constant potential vs. SHE.
Nevertheless, the observed behavior is similar to that observed for oxygen reduction, where currents
are almost independent of the solution pH [41].

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Figure 8. Voltammetric profiles of the Pt(100) electrode in pH = 3 (left column) and pH = 5 (right
column) with different concentrations of MeOH. The red trace shows the profile in absence of MeOH.
Scan rate: 50 mV s−1 .

For the Pt(110) electrode, the voltammogram for the oxidation of methanol shows a wave with a
complex shape in the positive scan, whose onset is located at ca. 0.5 V in 10−4 M methanol, whereas,
in the negative scan direction, the currents are almost negligible (Figure 9). These small currents
cannot be assigned to the formation of adsorbed CO since the voltammetric profile in the low potential
region overlaps with that measured in the absence of methanol. It should be noted that the signals
appearing in the blank voltammogram between 0.06 and 0.30 V contain contributions both from
hydrogen and OH adsorption. Then, the low currents recorded in the negative scan direction should
be assigned to the oxidation of the surface, which is an irreversible process that occurs for this electrode
above 0.8 V. As the concentration increases, the current also increases. However, the peak current
difference between the positive and negative scan directions remains, due to the oxidation of the
surface. It should be stressed that the oxidation of methanol is taking place, also on this electrode,
on a surface partially covered by OH, reinforcing the connection between adsorbed OH and methanol
oxidation. On the other hand, the solution pH effect on this surface is very similar to that described for
the Pt(100) electrode. It should be noted that the voltammograms for this surface are very sensitive
to the actual structure of the electrode, likely related to the fraction of the surface having the (1 × 2)
reconstruction. This ratio depends, in turn, on the annealing conditions [23], so that the measured
currents for the methanol oxidation reaction for this electrode show large variability.

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Figure 9. Voltammetric profiles of the Pt(110) electrode in (A) 0.1 M HClO4 + 10−4 M MeOH and (B) 0.1
M HClO4 + 10−4 M MeOH. The red trace shows the profile in absence of MeOH. Scan rate: 50 mV s−1 .

4. Conclusions
The use of diluted methanol solutions with different pH has allowed establishing relationships
between the interfacial structure and the reactivity. First, the results indicate that methanol oxidation
is only taking place when the surface is partially covered by adsorbed OH. Thus, the dehydrogenation
step to yield adsorbed CO on the Pt(100) surface starts on the same electrode potential for the onset
of OH adsorption on this electrode. On the other hand, on the Pt(111) electrode, where CO is not
accumulated, the onset for the direct oxidation of methanol and the adsorption of OH coincide. In this
case, the adsorbed OH species are not a mere spectator, because the obtained results for the reaction
order for methanol and the proton indicate that OH adsorbed species are involved in the reaction
mechanism. Since previous results show that the rate determining step in the methanol oxidation
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reaction is a dehydrogenation step, probably adsorbed OH collaborates in this step, facilitating
either the adsorption of the methanol in the right configuration or the cleavage of the C—H bond
so that the activation energy is minimized. These results are important for the computational
modeling of the reaction because adsorbed OH should be incorporated in the model to reflect the real
interfacial structure.
Author Contributions: Conceptualization, J.M.F. and E.H.; methodology, E.H.; validation, M.A.K. and R.M.-H.;
formal analysis, E.H. and R.M.-H.; investigation, M.A.K. and R.M.-H.; writing—original draft preparation, E.H.;
writing—review and editing, M.A.K., R.M.-H., J.M.F. and E.H.; visualization, R.M.-H. and E.H.; supervision, J.M.F.
and E.H.; project administration, J.M.F. and E.H.; funding acquisition, J.M.F. and E.H.
Funding: This research was funded by the MINECO (Spain), grant number CTQ2016-76221-P.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Strategies to Hierarchical Porosity in Carbon
Nanofiber Webs for Electrochemical Applications
Svitlana Yarova, Deborah Jones, Frédéric Jaouen and Sara Cavaliere *
Institut Charles Gerhardt Montpellier, UMR CNRS 5253, Agrégats Interfaces et Matériaux pour l’Energie,
Université de Montpellier - ENSCM - CNRS, CEDEX 5, 34095 Montpellier, France;
svitlana.yarova@etu.umontpellier.fr (S.Y.); deborah.jones@umontpellier.fr (D.J.);
frederic.jaouen@umontpellier.fr (F.J.)
* Correspondence: sara.cavaliere@umontpellier.fr; Tel.: +33-46714-9098
Received: 4 February 2019; Accepted: 25 February 2019; Published: 5 March 2019

Abstract: Morphology and porosity are crucial aspects for designing electrodes with facile transport
of electrons, ions and matter, which is a key parameter for electrochemical energy storage and
conversion. Carbon nanofibers (CNFs) prepared by electrospinning are attractive for their high
aspect ratio, inter-fiber macroporosity and their use as self-standing electrodes. The present
work compares several strategies to induce intra-fiber micro-mesoporosity in self-standing CNF
webs prepared by electrospinning polyacrylonitrile (PAN). Two main strategies were investigated,
namely i) a templating method based on the addition of a porogen (polymethyl methacrylate,
polyvinylpyrrolidone, Nafion® or ZnCl2 ) in the electrospinning solution of PAN, or ii) the activation
in ammonia of previously formed CNF webs. The key result of this study is that open intra-fiber
porosity could be achieved only when the strategies i) and ii) were combined. When each approach
was applied separately, only closed intra-fiber porosity or no intra-fiber porosity was observed.
In contrast, when both strategies were used in combination all CNF webs showed high mass-specific
areas in the range of 325 to 1083 m2 ·g−1 . Selected webs were also characterized for their carbon
structure and electrical conductivity. The best compromise between high porosity and high electrical
conductivity was identified as the fibrous web electrospun from PAN and polyvinylpyrrolidone.
Keywords: carbon nanofiber; porous fiber; electrospinning; mesopore; micropore; porogen; ammonia
activation; surface area

1. Introduction
The performance of electrodes in electrochemical energy storage and conversion devices strongly
depends on two key factors: a) the density of active sites or active area where faradaic or capacitive
phenomena occur; and b) the mass- and charge-transport of species (gases, liquids, ions and electrons)
to and from the active sites across the porous electrode [1,2]. The design of porous and hollow
nanostructures is crucial to obtain high surface areas, nanoscale porous structures and a high density of
electrochemically accessible active sites. In particular, carbon is widely used as a support (e.g., for fuel
cells) or as an active material (e.g., in supercapacitors and in some batteries) due to its sufficiently high
electrical conductivity, low cost and reasonable stability in specific electrochemical potential windows.
Among carbon materials, carbon nanofibers (CNFs) are very promising for the fabrication of
electrochemical electrodes due to their high aspect ratio and the possibility to assemble CNFs in unique
three-dimensional web structures with very high macropore volume between the fibers. High porosity
within electrodes is critical to avoid a slow Knudsen diffusion regime in gas-diffusion electrodes
(as in fuel cells), or to avoid reactant/ion depletion (in supercapacitors, redox flow batteries, etc.).
A very efficient and upscalable approach to prepare a CNF network is via the electrospinning and
carbonization of a polymer solution [3,4]. Furthermore, CNF webs obtained via electrospinning can
Surfaces 2019, 2, 159–176; doi:10.3390/surfaces2010013

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often be designed to possess sufficient mechanical stability and through-plane electric conductivity to
be used as free-standing electrodes, as shown in recent studies on batteries [5,6], supercapacitors [7–9],
fuel cells [10–12], vanadium redox-flow [13] and Li–O2 batteries [14]. This technique may not only
simplify the electrode fabrication process, but also maximizes the macropore volume in the electrode
and avoids binders, otherwise necessary with dispersed CNFs or other carbon powder materials to
form the electrode. Such binders or additives often reduce the accessibility of active sites by ions
in the electrolyte, by electrons from the current collector and/or by reactants in the gas-phase or in
the electrolyte.
Electrospun CNFs are generally prepared by electrospinning a polymer precursor
(e.g., polyacrylonitrile (PAN), polyimide, polybenzimidazole or polyvinyl alcohol), followed by at least
one thermal treatment to convert the polymer fibrous network into conductive carbon [15]. The nature
and molecular weight of the polymer precursor, its concentration in the electrospun solution as well as
the thermal treatment conditions all can influence the morphological and physical properties of the
CNF web obtained. Typical CNF diameters obtained are in the range 100–500 nm, resulting in specific
surface area in the range of 4–40 m2 ·g−1 if the CNF does not possess internal porosity. In order to
improve the specific surface area of CNF-based electrodes obtained by electrospinning, and therefore
their electrochemical properties, it is crucial to develop strategies that result in CNF webs with internal
porosity inside the fibers. The other possible approach, namely reducing the CNF diameter to <10 nm
(to result in an expected specific area of ca 200 m2 ·g−1 ), is limited both in terms of the possible increase
of specific area and also by the limitations of the mechanical stability of a CNF web with such an
ultralow fiber diameter.
Two main strategies have been investigated previously to prepare porous carbon fibers by
electrospinning, one in which a porogen is introduced in the electrospinning solution and which
results in biphasic polymer-porogen fibers, and another in which the plain CNFs are subjected after
their synthesis to a reactive chemical (gas or solid). In the former case (labelled “pre-synthesis”),
the carbonization of the biphasic polymer fibers leads to the selective removal of the porogen during
the temperature ramp-up, resulting in CNFs having either a closed or open porosity. In the second
case (labelled “post-synthesis”), the selective etching of less-organized carbon domains in the CNFs
by the reactive chemical may result in the formation of open pores from the outer surface of CNFs
and the inward. However if the carbon structure of the CNF is homogeneous the reactive etching may
simply result in a thinning of the CNF, without the creation of internal porosity.
In the pre-synthetic approach, the electrospinning technique allows for the introduction of various
porogens in the polymer precursor solution, whereby the porogen acts as a hard or soft template during
the subsequent carbonization step. The most studied hard templates for preparing CNFs are preformed
nanoparticles or nanostructures (silica [8,16], nano-CaCO3 [17]) that generate porosity upon their
removal with chemical and/or thermal treatments. The hard template approach allows control of the
size and morphology of the pores and minimizes undesired chemical reactions between the template
and the carbon structure [18]. Hard templates can also form during the carbonization treatment
itself, as is the case when introducing metal salts (e.g., Co(NO3 )2 6H2 O, Ni(CH3 COO)2 , ZnCl2 ) in the
electrospun polymer solution. The transition metal salts form metallic or metal oxide particles (e.g., Co,
Ni, ZnO), which then serve as porogen, and in some cases, as a catalyst for enhanced graphitization
of the CNF formed from the polymer precursor carbonization [19–23]. Other templates including
Prussian blue analogues [24] and metal organic frameworks [25–27] have been used in conjunction
with polymer precursors, conferring both porosity and heteroatom doping to CNFs obtained by
electrospinning and carbonization. Sacrificial polymers or organic molecules such as polymethyl
methacrylate (PMMA) [6,14,28–30], polyvinylpyrrolidone (PVP) [31,32], poly(ethylene oxide) [33],
Nafion® [34], polysulfone [35], polystyrene [36], poly-L-lactic acid [37] and beta-cyclodextrin [38] can
also be spun together with the main carbon precursor (any polymer forming conductive carbon with
high yield during pyrolysis), allowing the formation of porous CNFs upon their removal by solvent
or thermal treatment. Different kinds of templates can further be combined to produce hierarchical

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porosity, which is crucial for the transport of different species (ions, gases, liquids) in the electrodes of
various electrochemical devices. For instance, the combination of PVP with Mg salts [22,39] gave rise to
hierarchical meso- and microporosity inside CNFs prepared by electrospinning. For the post-synthetic
approach, various chemical activations have been investigated to create porosity in preformed CNFs.
Immersion of CNFs in concentrated KOH solutions followed by a treatment at high temperature is
a recognized approach [40,41]. Other more direct approaches consist of re-pyrolyzing the CNFs in a
reactive atmosphere, such as steam [42,43] or ammonia [44].
While a number of studies on CNFs prepared by resorting to either a pre-synthetic approach
involving a given porogen and given carbon precursor, or a post-synthetic approach, few studies have
compared the merits of using different porogens in otherwise identical preparation conditions for CNFs,
and even less studies have reported on the combination of i) porogen-PAN electrospun solution and ii) a
post-synthetic treatment to further increase intra-fiber porosity. In this paper, we therefore investigated
the formation of intra-fiber porosity in self-standing CNF webs that were derived from electrospun
PAN, resorting either to pre-synthetic and/or post-synthetic approaches to create porosity. PAN was
selected for its relatively high melting point, high carbon yield and facility to be electrospun [15].
For the pre-synthetic approach, various hard (ZnCl2 ) and soft templating agents (PMMA, PVP, Nafion® )
were co-electrospun with PAN. For the post-synthetic approach, NH3 pyrolysis was applied to CNF
webs derived from electrospun porogen-PAN solutions. The resulting CNF webs were characterized
for their properties relevant to electrochemical applications, including not only mass-specific surface
area and porosity, but also morphology, structure, composition and electrical conductivity.
2. Materials and Methods
To prepare reference CNFs, PAN (Mw = 150,000 g·mol−1 , Sigma-Aldrich, Saint Louis, MO, USA)
was first dissolved in N,N-dimethylformamide (DMF, pure, Carlo Erba, Val de Reuil, France) for
12 h at 50 ◦ C (10 wt % PAN concentration), after which the solution was cooled to room temperature.
The polymer fibers were then electrospun at 20 ◦ C and collected on a drum rotating at 100 rpm
(Spraybase ® , Dublin, Ireland). The distance between the tip of the needle (22 gauge) and the collector
was 10 cm, and a voltage of 13 kV (Auto-Reversing High Voltage Power Supply Spellman CZE1000R,
West Sussex, United Kingdom) was applied to obtain a stable Taylor cone. The flow rate was kept as a
constant 1 mL·h−1 (syringe pump KDS 100 Legacy Syringe Pump, KD Scientific, Holliston, MA, USA).
The obtained PAN-based fibers were submitted to stabilization and carbonization to give rise to
carbon materials [45,46]. In particular, the electrospun PAN fibers were treated in air at 150 ◦ C for 2 h
with a heating rate of 2.5 ◦ C·min−1 , and then at 250 ◦ C for 3 h with a heating rate of 2.5 ◦ C·min−1 .
The stabilized nanofibers were finally carbonized at 1000 ◦ C (ramp rate 5 ◦ C·min−1 ) for 2 h under
flowing argon atmosphere. After 2 h at 1000 ◦ C, the heating was stopped and the sample cooled
down naturally to room temperature under flowing Ar. The obtained CNFs are labelled PAN10 -CNF,
the scalar 10 standing for the wt % PAN in DMF solution.
To prepare biphasic polymer fibers comprising PAN, the porogen was solubilized or dissolved
in a PAN/DMF solution. The investigated porogens are PMMA (Sigma-Aldrich, Mw = 15,000 or
120,000 g·mol−1 ), PVP (Sigma-Aldrich, Mw ~1,300,000 g·mol−1 , Steinheim, Germany), Nafion® (NR50,
Sigma-Aldrich, Saint Louis, MO, USA) and zinc chloride (ZnCl2 anhydrous, >98%, purchased from
Alfa Aesar, Kander, Germany). For PMMA and ZnCl2 , the PAN and/or porogen concentrations in
DMF were varied, since the introduction of the porogen had an obvious influence on rheological
properties of the solution and morphology of the electrospun polymer fiber web. Table 1 details the
investigated compositions of the electrospun solutions. Such biphasic polymer fibers were stabilized in
air and carbonized in flowing Ar in the exact same conditions as described above for pure-PAN based
fibers. The final materials are labelled as Porogenx -PANy -CNF, with x standing for the wt % of porogen
in the electrospun solution, y the wt % of PAN in the electrospun solution, and CNF indicating the
pyrolysis was performed in Ar (see labels in Table 1).

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Table 1. Composition of the polyacrylonitrile (PAN)/Porogen electrospun solutions and labels of the
resulting carbon nanofibers (CNFs) and NH3 -activated CNF webs (ACNFs). The DMF amount in the
electrospun solutions was (100-porogen wt %-PAN wt %).
Porogen (Mw )

Porogen wt %

PAN wt %

CNF Label
(after Carbonization in
Ar)

CNF Label
(after Carbonization in Ar
and NH3 Activation)

None

0

10

PAN10 -CNF

PAN10 -ACNF

PMMA (15,000)

2

8

15kPMMA2 -PAN8 -CNF

15kPMMA2 -PAN8 -ACNF

PMMA (120,000)

4

6

120kPMMA4 -PAN6 -CNF

-

PMMA (15,000)

2

8

15kPMMA2 -PAN8 -CNF

-

PMMA (120,000)

4

6

120kPMMA4 -PAN6 -CNF

-

PVP

5

10

PVP5 -PAN10 -CNF

PVP5 -PAN10 -ACNF

Nafion®

2

8

Nafion2 -PAN8 -CNF

Nafion2 -PAN8 -ACNF

ZnCl2

1

10

Zn1 -PAN10 -CNF

-

ZnCl2

3

10

Zn3 -PAN10 -CNF

-

ZnCl2

5

10

Zn5 -PAN10 -CNF

-

ZnCl2

7

10

Zn7 -PAN10 -CNF

Zn7 -PAN10 -ACNF

Chemical activation of selected CNFs was performed by applying a flash pyrolysis in ammonia
at 900 ◦ C. Unless otherwise indicated, the pyrolysis duration in ammonia was 15 min. The resulting
materials are identified by their label ending with ACNF, instead of CNF for non-activated material.
The morphology of the polymer fibers and/or CNFs was investigated by field emission-scanning
electron microscopy (FE-SEM) using a Hitachi S-4800 microscope (Hitachi Europe SAS, Velizy,
France). Data analysis and fiber diameter distribution were performed using an image processing
software Image J 1.48 v (U. S. National Institutes of Health, Bethesda, MD, USA). CNFs were
analyzed by transmission electron microscopy (TEM) using a JEOL 2200FS (Source: FEG) microscope
operating at 200 kV equipped with a CCD camera Gatan USC (16 MP) (Tokyo, Japan). For TEM
cross-sectional analysis, a microtome was used on resin-encapsulated sample and slices were deposited
on carbon-coated copper grids (Agar Scientific, Stansted, Essex, United Kingdom).
Surface area and porosimetry of the samples were analyzed with N2 physisorption with a Tristar
II Micromeritics instrument (Norcross, GA, USA) at 77 K. Prior to analysis, all samples were outgassed
overnight at 120 ◦ C under vacuum. The resulting isotherms being of type I according to the IUPAC
classification, BET plots were drawn below the relative pressure of 0.1 from the adsorption branches and
employed to evaluate the BET specific surface (SBET ). The alpha-plot method was utilized to determine
the mesoporous (Vmeso ) and microporous volumes (Vmicro ) as well as average pore diameter (dpore ).
C, H, N, O elemental analysis was performed with a Vario MICRO Element Analyzer
(Elementar Analysensysteme, Hanau, Germany).
Raman spectra were recorded on a LabRAM Aramis IR2 Horiba Jobin Yvon spectrometer
(Villeneuve d’Ascq, France) equipped with a He/Ne laser (λ = 633 nm) and a long work distance
objective ×50. The spectra were fitted with five bands using Origin software (OriginLab Corporation,
Northampton, UK).
The in-plane electrical conductivity of self-standing CNF webs was measured using a
2400 Keithley in a four-electrode configuration on a 5 mm × 40 mm × 0.05 mm carbon electrode strip
in a Fumatech MK3-L cell operated in the current range of 0–100 mA.
3. Results and Discussion
3.1. Characterization of Porogen-PAN Fibrous Webs After Carbonization in Ar
As a reference, the CNF web prepared only from PAN was characterized (labelled PAN10 -CNF).
The resulting CNF web was self-standing and flexible, as already reported by the authors [10].

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The morphology of PAN10 -CNF was investigated by SEM and TEM. The material was observed to be
randomly oriented cylindrical fibers (Figure 1a) with an average diameter of 200 nm (Figure 2a).
These fibers were smooth and dense, as demonstrated by the TEM cross-section in Figure 3a.
According to the latter image and to previous reports [47], the porosity and specific surface area
developed by such fibers is very low (ca 20 m2 ·g−1 ), mainly attributed to the outer surface area of
the fibers, with no internal porosity. The nearly flat N2 adsorption isotherm demonstrates the lack of
porosity inside the fibers (Figure 4).
In order to create internal porosity within fibers, a range of porogens were added to the precursor
PAN solution prior to electrospinning, as described in the experimental methods. The investigated
polymer templates (PMMA, PVP, Nafion® ) were selected for their relative stability and rigidity.
They are known to be stable up to 250 ◦ C in air, the conditions used during the stabilization step.
When heated to much higher temperatures in inert gas, they decompose and form mainly volatile
products, thereby acting as porogens, while PAN decomposes as well but is converted to carbon in
high yield.
The results obtained with PMMA are discussed first. PMMA of two different molecular weights
(15,000 and 120,000 g·mol−1 ) and in different ratios to PAN (1:4 and 2:3) was added in the electrospun
DMF-based solution. Those ratios were selected on the basis of a previous study on PMMA-PAN
composite fibers [36]. In all the cases, the addition of PMMA resulted in CNFs with slightly smaller
average diameter (compare Figure 1a,b) than for the reference PAN10 -CNF web (average diameter of
175 nm, see Figure 2b), but with otherwise practically identical features and properties (Figure S1).
Figure 1b depicts the sample 15kPMMA2 -PAN8 -CNF as an example. The decrease in fiber diameter
is most probably due to the lower concentration of PAN in the electrospun solution relative to the
reference PAN10 -CNF (8% vs. 10%), and was also expected due to the change in precursor viscosity
already observed upon addition of this polymer [6]. The developed porosity is clearly visible on
the cross-sectional TEM micrograph (Figure 3b), where pores with an average diameter of 10 nm
appeared within the fibers. The modification of the preparation parameters (higher molecular weight
for PMMA, or different PMMA:PAN ratio) did not have significant effects on the obtained morphology
and porosity (images not shown). The question as to whether the formed porosity is open or closed is
discussed later.
The results obtained with PVP porogen are now discussed. A PVP:PAN ratio of 1:2 was employed
(sample labelled PVP5 -PAN10 -CNF). In these conditions, carbon fibers with considerably greater
diameter compared to the reference PAN10 -CNF (average diameter of 750 nm) (Figures 1c and 2c)
were obtained. This can be attributed to a modification of the viscosity of the electrospun solution,
affected by the presence of PVP aggregates and consequent decrease in chain entanglement. The pores
generated inside the CNFs presented an average size around 3 nm (Figure 3c).

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Surfaces 2019, 2

ȱ
Figure 1. FE-SEM micrographs of: (a) PAN10 -CNF; (b) 15kPMMA2 -PAN8 -CNF; (c) PVP5 -PAN10 -CNF;
(d) Nafion2 -PAN8 -CNF; (e) Zn1 -PAN10 -CNF; (f) Zn3 -PAN10 -CNF; (g) Zn5 -PAN10 -CNF;
(h) Zn7 -PAN10 -CNF. The scale bar corresponds to 6 μm.

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Surfaces 2019, 2

ȱ
Figure 2. Histograms of the fiber diameter distribution for: (a) PAN10 -CNF; (b) 15kPMMA2 -PAN8 -CNF;
(c) PVP5 -PAN10 -CNF; (d) Nafion2 -PAN8 -CNF; (e) Zn1 -PAN10 -CNF; (f) Zn3 -PAN10 -CNF;
(g) Zn5 -PAN10 -CNF; (h) Zn7 -PAN10 -CNF.

The third polymer template investigated was Nafion® perfluorosulfonic acid. It has been
previously demonstrated to be effective in creating nanoscale porous domains inside carbon
nanofibers [34]. In this work, its addition to PAN (sample labelled Nafion2 -PAN8 -CNF) resulted
in the formation of thin carbon fibers (average fiber diameter of 180 nm) (Figures 1d and 2d). The CNF
webs also had an obviously a higher density of fibers (compare Figure 1a,d), which is beneficial for
improved conductivity and mechanical stability. The fibers have a well-developed internal porosity,
as shown by TEM analysis of the cross-sections, Figure 3d. A homogeneous distribution of pores with
a size of 4 nm can be observed within the CNF structure.

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Surfaces 2019, 2

ȱ
Figure 3. TEM micrographs of: (a) PAN10 -CNFs; (b) 15kPMMA2 -PAN8 -CNF; (c) PVP5 -PAN10 -CNF;
(d) Nafion2 -PAN8 -CNF; (e) Zn1 -PAN10 -CNF; (f) Zn3 -PAN10 -CNF; (g) Zn5 -PAN10 -CNF;
(h) Zn7 -PAN10 -CNF.

Finally, the inorganic template precursor ZnCl2 was investigated. The presence of ZnCl2 may
allow the formation of pores during the ensuing stabilization-carbonization steps in air and in
Ar with the following mechanism: during the stabilization step, hydrated chloride hydrolyzes
and forms an oxychloride, from which zinc oxide forms creating microporosity by etching carbon
atoms [23]. FE-SEM characterization of Figure 1e–h showed that increasing the ZnCl2 concentration
in the electrospun solution from 1 wt % to 7 wt % (samples Zn1 -PAN10 -CNF, Zn3 -PAN10 -CNF,
Zn5 -PAN10 -CNF, Zn7 -PAN10 -CNF) the diameter of the carbon fibers increased tremendously from
300 to 700 nm (Figure 2e–h). It is known that solution conductivity affects fiber uniformity. The high
conductivity of the PAN/ZnCl2 (52:1) solution was reported to lead to instability in the electrospinning
process with the formation of large fibers and bundles as a cotton-like 3D deposit [48]. However,
the reverse has also been observed: increasing the conductivity through the addition of salt may
produce finer, more uniform fibers, resulting in an increased elongational force exerted on the fiber
jet [49]. Indeed, other reports on the electrospinning of ZnCl2 /PAN solutions showed that the average
diameters of the obtained CNFs gradually decrease from 350 to 200 nm with increasing zinc chloride
content from 1 wt % to 5 wt % [23].

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Surfaces 2019, 2

Another particular feature of CNFs derived from electrospun ZnCl2 /PAN solutions is the
formation of progressively larger fiber bundles with increased Zn salt concentration (see Figure 1g,h
in particular). This phenomenon was already observed for PAN/polystyrene mixtures, where the
template not only acted as sacrificial decomposable phase, but also controlled the formation of these
architectures [13]. In the TEM micrographs of the cross-sections of Zn1 -PAN10 -CNF, Zn3 -PAN10 -CNF,
Zn5 -PAN10 -CNF, Zn7 -PAN10 -CNF samples (Figure 3e–h), no pores are visible inside the fibers.

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300$3$1$&1)
3933$1$&1)
3$1&1)




















5HODWLYHSUHVVXUHSS
Figure 4. Selected N2 adsorption-desorption isotherms of carbonized and NH3 -activated CNF webs
(ACNF). The carbonized web obtained from PAN-only is also shown as a reference (PAN10 -CNF).

Except for the latter Zn-PAN-CNF samples, all other CNFs derived from polymer/PAN
solutions containing the sacrificial polymers presented pores visible by TEM in the mesoscopic range,
in agreement with previous works [6,14,29,30,34]. The nitrogen adsorption/desorption isotherms
of all the investigated CNFs showed however very low SBET values (below 20 m2 ·g−1 , see * in
column 2 of Table 2), almost unchanged compared to the reference PAN-CNFs (Table 2). This result,
apparently in contrast with the micrographs of Figure 3 showing pores of different size in the
prepared CNFs and in contrast with previous reports using the same templates and similar synthesis
conditions [6,14,28–30,34], demonstrates that the obtained porosity is closed, and not accessible from
outside the CNFs. This observation may be due to the partial collapse of the percolating porous
network formed during the thermal removal of the template or porogen [34]. As a result, free-standing
porous fiber webs are obtained, with closed porosity and low surface area. The latter is assigned
only to the outer surface area of plain and smooth CNFs with diameters in the range 150–700 nm,
depending on the porogen used. The appealing approach of using a soft or hard template added
during the preparation step of electrospinning seems therefore ineffective in creating open porosity.
The conditions necessary to reach an open porosity are shown later in this study.
Graphitic character and the related electrical conductivity of the carbon based electrode materials
are relevant properties for electrochemical applications. To investigate them, Raman spectroscopy and
the four-electrode method were used to characterize the carbon nanofiber networks prepared in this
work. In particular, the modification of this properties upon the different steps of preparation was
studied. The Raman spectra of all CNFs prepared with different porogens (Figure S2) present two
intense and broad bands, the so-called D band at 1357 cm−1 ascribed to defects and disorder in the
graphitic structure and the so-called G band at 1560 cm−1 corresponding to the in-plane vibration of
sp2 -bonded carbon in graphite. The spectra were fitted with more contributions, namely the D4 band,
ascribed to sp3 -carbon (1180 cm−1 ), the D3 band (1500 cm−1 ) associated with an amorphous sp2 carbon
163

Surfaces 2019, 2

bonded in the graphitic phase and the D2 band (1580 cm−1 ) corresponding to graphitic lattices in the
structure (Figure S3) [50]. The relative intensities of the D and G bands (ID /IG ) as well as the relative
areas (AD /AG) are often used to estimate the degree of graphitization of carbon materials [39,51].
Lower ID /IG ratios indicate higher levels of crystalline sp2 -carbon [52,53]. The ID /IG and AD /AG
values obtained after deconvolution of the spectra of all CNFs derived from the carbonization of
porogen/PAN solutions are summarized in Table 3. It is evident that the addition of polymer or
inorganic salt porogen did not significantly influence the graphitization of the carbon fibers. Similar
ID /IG ratios were obtained for reference PAN10 -CNFs and the other CNF webs. Their values between
1.8 and 2.7 demonstrate that the fibers are composed by disordered carbon (due to the relatively low
carbonization temperature of 1000 ◦ C) with local graphite inclusions (turbostratic domains) already
evidenced for electrospun CNFs [24,39,47,54].
Table 2. Textural properties of carbonized (CNFs) and ammonia-activated (ACNFs) carbon nanofibers.
Fiber Precursor
PAN10
15kPMMA2 -PAN8
120kPMMA2 -PAN8
15kPMMA4 -PAN6
120kPMMA4 -PAN6
PVP5 -PAN10
Nafion2 -PAN8
Zn1 -PAN10
Zn3 -PAN10
Zn5 -PAN10
Zn7 -PAN10

CNFs

ACNFs

SBET , m2 ·g−1

SBET , m2 ·g−1

CBET

Vmeso , cm3 ·g−1

Vmicro , cm3 ·g−1

dpore , nm

20
35
n/a*
n/a*
n/a*
3
n/a*
n/a*
n/a*
n/a*
n/a*

n/a*
450
645
360
410
325
535
680
570
865
1083

3034
2410
2927
3324
2012
1993
1986
1707
1791
3558

0.0599
0.0877
0.0083
0.1166
0.0791
0.0721
0.0980

0.1243
0.1669
0.1006
0.1177
0.0941
0.1327
0.1814
0.1614
0.1982
0.3098

2.9
1.9
0.5
0.8
1.8
2.4
1.6
2.5
1.5
1.4

* SBET <3 m2 ·g−1

Table 3. Relative areas and intensities of D and G bands (ID /IG , AD /AG ) in the Raman spectra after
carbonization in Ar (CNF) and after a subsequent activation in ammonia (ACNF).
CNF

Fiber Precursor
PAN10
15kPMMA2 -PAN8
120kPMMA2 -PAN8
15kPMMA4 -PAN6
120kPMMA4 -PAN4
PVP5 -PAN10
Nafion2 -PAN8
Zn1 -PAN10
Zn3 -PAN10
Zn5 -PAN10
Zn7 -PAN10

ACNF

AD/G

ID/G

AD/G

ID/G

3.01
4.73
4.59
3.24
3.14
4.69
5.34
4.40
6.34
4.24
6.54

1.83
2.45
2.02
2.20
1.86
2.09
2.77
2.43
2.40
2.57
2.39

2.95
3.06
5.00
3.78
3.76
3.21
2.52
4.65
9.61
2.30
3.06

1.86
1.88
2.39
2.08
2.02
2.28
2.33
2.49
2.56
1.88
1.89

The electrical conductivities of the CNFs prepared by electrospinning, determined directly
on the self-standing CNF webs, are consistent with the partial graphitic character evidenced by
Raman spectroscopy. Except for the fibers prepared using zinc chloride as porogen precursor, which
show significantly lower conductivity, all CNF networks present similar conductivity values around
6–9 S·cm−1 (Table 4), which is in agreement with previous results obtained on PAN based electrospun
materials [10,47,55].

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Surfaces 2019, 2

Table 4. In-plane electrical conductivity and elemental content of all CNF webs and selected ACNF webs.
CNF

ACNF

Fiber Precursor

Conductivity
S·cm−1

C
wt %

N
wt %

O
Wt %

Conductivity
S·cm−1

C
wt %

N
wt %

O
wt %

PAN10
15kPMMA2 -PAN8
PVP5 -PAN10
Nafion2 -PAN8
Zn1 -PAN10
Zn3 -PAN10
Zn5 -PAN10
Zn7 -PAN10

6.6
8.8
9.7
7.0
1.1
0.8
3.3
0.8

66.8
75.8
76.2
74.3
70.3
71.5
61.2
68.1

5.6
5.5
2.2
5.5
5.2
4.6
6.8
2.5

17.5
11.1
10.9
16.0
14.6
17.1
20.9
19.2

3.8
1.9
4.8
2.4
0.3

78.4
77.8
78.8
70.3
71.3

7.0
5.8
7.0
5.5
6.1

7.8
6.7
8.2
9.6
11.3

The carbon fibers prepared from ZnCl2 -PAN10 fibers demonstrated very low conductivity,
between 3.3 and 0.8 S·cm−1 . The graphitic structure being the same for all the samples as indicated
by the Raman study, this decrease in conductivity from the pristine CNFs may be ascribed to
the microstructure of the ZnCl2 -PAN-CNFs in which, due to ZnCl2 removal, electron paths are
likely disrupted. Furthermore, its particular morphology with fibers forming large bundles can
affect the inter-fiber/inter-bundle connection (decreased number of connection points) and thus
the electron transport.
3.2. Characterization of Porogen-PAN Fibrous Webs After Carbonization and NH3 Activation
The post-treatment with ammonia at high temperature is a conventional approach to creating
porosity in carbon-based materials by etching, which can also introduce nitrogen functionalities on
the surface [56]. The mechanism for the formation of both pores and nitrogen groups is based on the
reaction of ammonia with carbon. This gasification reaction continuously occurs during NH3 pyrolysis
and react with surface carbon atoms forming volatile compounds such as HCN [57,58]. When the
reaction between ammonia and carbon occurs at different rates on the surface (due to carbon structure
heterogeneity), increased porosity is obtained. When the ammonia pyrolysis is stopped, nitrogen atoms
that were reacting with carbon surface atoms at that moment remain on the surface, with different
environments including amino, cyanide, pyrrolic, pyridinic and quaternary nitrogen groups [59,60].
We first subjected the reference electrospun carbon nanofibers PAN10 -CNF web to treatment
under pure flowing ammonia at 900 ◦ C [61]. The resulting activated nanofibers (ACNFs),
labelled PAN10-ACNFs, did not present significant modification of their average diameter (200 nm)
nor of their porosity, still not visible by TEM cross-sectional observations (Figure S4). As a consequence,
specific surface area determined from nitrogen adsorption/desorption isotherms remained low
(Table 2). Due to the low surface area obtained with the porogen/PAN template method after
carbonization in Ar, as described previously, and the low surface area obtained with ammonia
activation on pure PAN-derived CNFs, the combination of the two approaches was then investigated
as a last attempt to develop high specific area. All CNF webs that had been prepared by porogen/PAN
templating and carbonized in Ar (labelled CNF) were then further activated in flowing ammonia gas
in the same conditions as for PAN10 -ACNF.
N2 adsorption-desorption measurements were performed to characterize their porosity and
specific surface area (Table 2 and Figure 4). Figure 4 clearly shows that ACNF webs adsorbed a
significant amount of N2 , the three polymer porogens resulting in intermediate adsorbed volumes
and the ZnCl2 porogen resulting in the highest adsorbed volumes. The isotherms of the different
ACNF webs are of type I and/or type II according to IUPAC classification, which indicates that
fibers presented an overall microporous structure, as well as mesopores. With Nafion® as a porogen,
significant hysteresis of type H4 is observed during desorption, closing at P/P0 = 0.45. This is assigned
to mesopores with a bottleneck shape, which may be related to the existence of Nafion® polymer
aggregates in the electrospinning solution [62,63], while other polymers were fully dissolved. In the

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Surfaces 2019, 2

case of microporous structures, the BET energetic constant (CBET , related to the energy of adsorption
of the first layer of N2 adsorbate on the carbon surface) is expected to be higher than the usual value
expected for mesoporous carbon structure as seen in Table 2 [64]. The values of the specific surface
area, meso and micropore volume as well as average pore diameter of all ACNF webs are summarized
in Table 2. The specific surface area of all ACNF webs ranged from 325 to 1083 m2 ·g−1 , the maximum
value corresponding to ZnCl7 -PAN10 -ACNF, while the average pore size was comprised between
0.5 and 2.9 nm.
The carbon fibers derived from PMMA-PAN solutions presented surface areas between
360 and 645 m2 ·g−1 . The specific area increased by increasing the molecular weight of the sacrificial
polymer from 15,000 to 120,000 g·mol−1 and by decreasing the PMMA:PAN ratio from 2:3 to 1:4.
The carbon fibers derived from the PVP-PAN solution with a PVP:PAN ratio of 1:2 presented a surface
area of 325 m2 ·g−1 . This is comparable to the BET value obtained with PMMA in high ratio to
PAN (2:3), and may also be increased if the PVP:PAN ratio had been further optimized (lowered).
The carbon fibers derived from Nafion® -PAN solution with Nafion® :PAN ratio of 1:4 presented a
surface area of 535 m2 ·g−1 , also comparable to the SBET value obtained with PMMA:PAN ratio of 1:4
(450 m2 ·g−1 and 645 m2 ·g−1 , depending on PMMA molecular weight). The carbon fibers obtained
from ZnCl2 -PAN solutions showed a SBET value of 680 m2 ·g−1 already at a low ZnCl2 :PAN ratio of
1:10, further increasing with increasing amounts of ZnCl2 : PAN ratio from 3:10 to 7:10. Due to the high
value of 1083 m2 ·g−1 measured for the carbon fibers obtained using a ZnCl2 :PAN ratio of 7:10 and for
the regular ammonia activation duration of 15 min at 900 ◦ C, the effect of the duration of activation was
performed for this sample, at the same temperature of 900 ◦ C. It is noted that the ammonia pyrolysis
was performed in flash mode [65], (sample heated from room temperature to 900 ◦ C in the range
of 1–1.5 min), which allows very precise control of the pyrolysis duration down to 5 min. To stop
the pyrolysis, the quartz tube was immediately removed from the split hinge oven. The results are
depicted in Figure 5. The specific area was already high after only 5 min of ammonia activation, with
slightly increased values for increased durations up to 20 min. The optimal duration was considered
to be 15 min, corresponding to the maximum compromise between high developed area and high
mechanical resistance. A longer treatment of 20 min led to an increased fragility of the carbon web,
limiting its utilization as a self-standing electrode.










6%(7P J 





















'XUDWLRQRIDPPRQLDDFWLYDWLRQPLQ
Figure 5. Effect of the duration of ammonia activation at 900 ◦ C on the mass-specific surface area (SBET )
of Zn7 -PAN10 -CNF.

Raman spectra recorded on activated fibrous webs were very similar to those obtained with
carbonized samples (Figure S5). The calculated ID /IG ratios (Table 3) were practically unchanged upon
activation in ammonia, demonstrating that the degree of graphitization was not significantly impacted
by the etching treatment.

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Surfaces 2019, 2

As already mentioned, ammonia activation usually leads to (additional) nitrogen doping of
the carbon [56,59]. In order to evaluate the introduction of nitrogen sites in the prepared ACNF
webs, elemental analysis was performed before and after the NH3 treatment, for selected CNF webs
(Table 4). The nitrogen amount in the CNF derived from PAN only was around 5.6 wt % after
carbonization (PAN10 row in Table 4, CNF column), in agreement with previous reports and with the
carbonization temperature of 900 ◦ C. It is known that higher pyrolysis temperatures result in lower
nitrogen content [15]. For the reference PAN10 -CNF, the nitrogen amount slightly increased to 7.0%
after ammonia activation, while a decrease in O wt % was observed. The latter effect is explained
by selective etching of surface O-containing groups by ammonia. For PVP5 -PAN10 and Zn7 -PAN10 ,
an approximately three-fold increase in nitrogen amount was observed from CNF to ACNF webs,
accompanied by a reduction in O wt % [59]. All ACNF webs however had comparable N wt %
amounts, in the range of 6.1 to 7.0 wt %.
Nitrogen doping, in particular with ammonia (highly basic nitrogen groups), is interesting
for some electrochemical applications such as supercapacitors or oxygen-reducing sites (in alkaline
electrolytes in particular). The etching of some carbon masses during ammonia activation results in
carbon fibers with increased internal porosity and also possibly weakened inter-fiber connections,
which may impact the mechanical stability of self-standing fiber webs and also their electrical
conductivity. The electrical conductivity of selected carbon fiber webs (one selected for each type
of porogen) was halved after the ammonia activation (Table 4). For example, it dropped from
6.6 to 3.8 S·cm−1 for the reference web derived from PAN only, and from 9.7 to 4.8 S·cm−1 for
PVP5 -PAN10 -CNF and PVP5 -PAN10 -ACNF, respectively. Interestingly, the fact that the conductivity
decreased significantly even for PAN10 -ACNF while such fibers have no open porosity, indicates that
the decreased in-plane conductivity cannot be assigned to the formation of open pores in the fibers.
The decreased conductivity of carbon fiber webs after ammonia activation may be due to the nitrogen
doping itself (decreasing the intrinsic conductivity of carbon) or weakened electrical contact at the
nodes of the fibrous web (possibly due to etching of carbon at the nodes). As already mentioned
earlier for the series of Zn7 -PAN10 -CNF fibers subjected to various duration of activation in ammonia,
activation longer than 20 min still resulted in a self-standing web but with greatly increased fragility,
which is directly related to the quality of the nodes of the webs.
All these results lead to the conclusion that the open porosity in the PAN-derived CNFs resulted
from a two-step approach involving the combination of a template route with a chemical activation
step. The porogens incorporated in the electrospun PAN solution led first to carbon fibers with closed
pores after the carbonization in Ar. The closed pores could be opened and interconnected within the
fiber upon a post-treatment in ammonia atmosphere at high temperature. Other examples of etching in
combination with the use of porogens have been reported (e.g., water etching-assisted templating) [31].
The purpose of this work was the rationalization of an effective strategy to obtain CNF webs
with both high porosity between fibers and high specific area by introducing porosity inside the fibers
for application in electrochemical energy conversion and storage devices. Having this objective in
mind, it is also important to maintain the electrical conductivity of the self-standing CNF webs to
sufficiently high values. Typically, to be not be limited by the conduction of electrons across an electrode,
the electric conductivity must be about 100 times higher than the electrolyte conductivity. The results
indicate that the approach combining the introduction of templates and chemical activation led to
open micro- and mesoporosity inside the fibers, while the electric conductivity was approximately
halved during the ammonia activation step. This decrease in conductivity may be due to the removal
of a fraction of the carbon from the bulk of the CNFs during ammonia activation and/or the
fragilization of the interfibrous connection. In addition, when the electrical conductivity of self-standing
CNF ammonia-activated webs is plotted against their BET area, a negative linear correlation is
observed (Figure 6). The identification of such a correlation is useful to select the best compromise
between conductivity and BET area, which will depend on the exact electrochemical application
and range of typical current densities produced by an electrode (in the order of increasing current

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Surfaces 2019, 2







&RQGXFWLYLW\6FP 

density: batteries < fuel cells < supercapacitors). With current densities <10 mA·cm−2 , the reported
conductivities [66] are not a limiting factor in batteries, and thus any of the ACNF webs could be
selected. With current densities in the range of 1–2 A·cm−2 , fuel cells are more demanding in terms of
electron conductivity in the electrodes. By comparison, the effective proton conductivity in typical
active layer of PEMFC is only in the range of 1–2 S·m−1 . Therefore, if the electronic conductivity
through the active layer is >20 S·m−1 (0.2 S·cm−1 ), it should not be a factor limiting the porous electrode
performance [67–70]. All present ACNF webs could therefore be applicable for self-standing electrodes
in, e.g., PEMFCs. Lastly, in electrochemical supercapacitors the instantaneous current density can
reach extremely high values on the order of 10 A·cm−2 due to the non-faradaic process resulting in
higher electric conductivity requirements. Moreover, the high electric conductivity requirement is
accompanied by the high BET area requirement [71], which are shown in Figure 6 to be antagonistic.






















6%(7P J 
Figure 6. Electrical conductivity versus specific surface area of the ACNF fiber webs. •—PAN10 -ACNF;
§
—PVP5 -PAN10 -ACNF; —Nafion2 -PAN8 -ACNF; ْ—15kPMMA2 -PAN8 -ACNF; —Zn7 -PAN10 -ACNF.
§
§
§
§

4. Conclusions

•
•
•
•

The templating approach of PAN fibers with polymer or inorganic porogens (PMMA, Nafion® ,
PVP, ZnCl2 ) resulted in CNFs with closed porosity after carbonization in Ar.
Subsequent activation in ammonia at 900 ◦ C opened this porosity, resulting in specific surface
area in the range of 325–1083 m2 ·g−1 .
Ammonia activation systematically decreased the electric conductivity of the ACNF webs by a
factor of approximately two to three.
A negative linear correlation between electric conductivity of ammonia-activated ACNF webs
and their BET area is revealed.

Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/2/1/13/s1,
Figure S1. TEM micrographs of PAN CNFs prepared from PAN and PMMA after carbonization and activation;
Figure S2. Raman spectra of template based CNFs after carbonization; Figure S3. Example of deconvolution of a
Raman spectrum of CNFs (PAN10 -CNF) by using Gaussian and Lorentzian curves (Voigt); Figure S4. FE-SEM
(left side) and cross-section TEM (right side) micrographs of PAN10 -ACNF fibers; Figure S5. Raman spectra of
CNFs after activation.
Author Contributions: Conceptualization, S.C and F.J.; methodology, S.C., F.J. and D.J.; formal analysis, S.Y.;
investigation, S.Y.; resources, S.C., F.J. and D.J.; writing—original draft preparation, S.Y and S.C.; writing—review
and editing, F.J and D.J.; supervision, S.C and F.J.; project administration, F.J.; funding acquisition, F.J. and S.C.
Funding: The research leading to these results has received funding from the French National Research Agency
under the CAT2 CAT contract (ANR-16-CE05-0007). SC acknowledges the financial support from the European

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Surfaces 2019, 2

Research Council under the European Union’s Seventh Framework Programme (FP/2007-2013) / ERC Grant
Agreement n. 306682 and the French IUF.
Acknowledgments: Authors thank Nicolas Donzel for his help with nitrogen adsorption/desorption analysis
and electrical conductivity measurements.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Potential Dependent Structure and Stability of
Cu(111) in Neutral Phosphate Electrolyte
Yvonne Grunder 1 , Jack Beane 1 , Adam Kolodziej 2 , Christopher A. Lucas 1 and
Paramaconi Rodriguez 2, *
1

2

*

Department of Physics, University of Liverpool, Liverpool L69 7ZE, UK;
Yvonne.Grunder@liverpool.ac.uk (Y.G.); sgjbeane@student.liverpool.ac.uk (J.B.);
clucas@liverpool.ac.uk (C.A.L.)
School of Chemistry, University of Birmingham. Birmingham B15 2TT, UK; AXK595@student.bham.ac.uk
Correspondence: p.b.rodriguez@bham.ac.uk

Received: 30 January 2019; Accepted: 20 February 2019; Published: 24 February 2019

Abstract: Copper and copper oxide electrode surfaces are suitable for the electrochemical reduction
of CO2 and produce a range of products, with the product selectivity being strongly influenced by the
surface structure of the copper electrode. In this paper, we present in-situ surface X-ray diffraction
studies on Cu(111) electrodes in neutral phosphate buffered electrolyte solution. The underlying
mechanism of the phosphate adsorption and deprotonation of the (di)-hydrogen phosphate is
accompanied by a roughening of the copper surface. A change in morphology of the copper surface
induced by a roughening process caused by the formation of a mixed copper–oxygen layer could
also be observed. The stability of the Cu(111) surface and the change of morphology upon potential
cycling strongly depends on the preparation method and history of the electrode. The presence
of copper islands on the surface of the Cu(111) electrode leads to irreversible changes in surface
morphology via a 3D Cu growth mechanism.
Keywords: Cu(111); electrochemical interface; in-situ X-ray diffraction

1. Introduction
Compared to many metal catalysts, copper surfaces have been proven to electrochemically convert
CO2 to high value and energy-dense products, such as methane ethylene, formic acid, methanol and
ethanol amongst others. However, the efficiency and the selectivity are far from optimal and the
parameters controlling these factors are not-fully understood. Differences in reactivity and selectivity
have been variously ascribed to surface area, particle size, surface structure and roughness and the
electrolyte composition [1–6]. More recently, the effect of the role of the oxygen content of the copper
catalyst and the oxidation state of the copper on the electrocatalytic activity and selectivity have
illustrated the high complexity of the system. Oxide-derived Cu catalysts have shown high-selectivity
towards the formation of C2 products [7–10]. Such selectivity has been attributed to changes in
the surface structure, including roughness and defects, with active sites being generated during the
reduction pretreatment of the Cu-oxide catalyst. However, more recently, in addition to the structural
factors, the presence of Cu+ and residual subsurface oxygen has been suggested to affect the product
selectivity [5,11,12]. Mistry et al. demonstrated, via operando XAFS, the presence of Cu+ species and
subsurface oxygen during the carbon dioxide reduction reaction [5]. In another report, LeDuff et al. [7]
implemented a pulse sequence between reduction potentials where the carbon dioxide reduction
reaction (CO2RR) takes place (<−0.5 V) and a potential in the region between −0.2 and −0.35 V
vs. RHE where the co-adsorption of OH and other anions takes place [13]. They concluded that the
positive potential of the pulse has a significant effect in the catalytic activity and selectivity of the copper
single-crystal electrodes towards the CO2RR. This was associated to the adsorbed species, presumably
Surfaces 2019, 2, 145–158; doi:10.3390/surfaces2010012

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OH or anions, at the surface of the Cu single crystal electrodes which prevent any irreversible damage
or changes to the Cu surface structure [7].
In this paper, we present the in-situ characterization by surface X-ray diffraction of the Cu(111)
electrode at positive potentials with the aim to determine the composition and structure of the
electrochemical interface. The key to understanding the stability and reactivity of the Cu(111)
electrode is to control the surface morphology during the preparation of the clean surface. Depending
on the details of surface preparation, two different Cu(111) surfaces can be obtained; although
macroscopically rather rough, one surface is completely stable during potential cycling that involves
considerable modification of the surface structure. The other surface gives rise to X-ray scattering
features consistent with the presence of twinned Cu nano-islands and, although this surface exhibits
a similar potential-dependent restructuring, it is not stable during potential cycling. The direct
correlation of structural stability, only available via in-situ structural characterization, and reactivity is
vital to understanding structure–reactivity relationships, especially in studies of more active metal
electrodes [14,15].
2. Materials and Methods
The Cu(111) single crystal working electrode (MaTeck, miscut < 0.1◦ ) was prepared by
electropolishing for 10 s in 70 % orthophosphoric acid at 2 V against a high-surface area copper
mesh. The crystal was then rinsed in ultra-pure water, covered with the electrolyte solution and
transferred to the electrochemical cell. A copper wire was used as counter electrode and an Ag/AgCl
electrode was used for the reference electrode. All potentials are quoted against this reference electrode.
The experimental procedure for the surface X-ray diffraction experiments followed that of similar
studies reported previously [16,17]. Surface X-ray diffraction measurements were carried out on the
I07 beamline at Diamond Light Source (Harwell Science and Innovation Campus, Fermi Ave, Didcot
OX11 0DE, Oxford, UK) [18], with a monochromatic beam of 25 keV X-rays. Beam defining slits were
0.5 mm × 0.5 mm and the beam size at the sample was estimated to be 200 μm × 300 μm (vertical ×
horizontal). The sample was mounted on a (2 + 3) circle diffractometer [19]. A 2D Pilatus Dectris 1
M detector (DECTRIS Ltd., Baden-Daettwil, Switzerland) was used for the data acquisition. For the
measurement of crystal truncation rods, scans along the Qz (surface normal) direction at specific points
in reciprocal space were recorded. Background correction and standard instrumental corrections were
applied to the dataset to be able to model the intensity distribution [20,21]. Errors on the individual
data points were a combination of the statistical error and an estimated 10% systematic error. The
model of X-ray diffraction was obtained using a Python (Python Software Foundation, Beaverton, OR,
USA) program and the integrated lmfit leastsquare fitting routine, which returned the best values
for the relaxation, coverage and β-factor. The Cu(111) surface has a close packed structure with a
hexagonal unit cell, where the surface normal is along the (0, 0, L)hex direction and the surface plane
contains the ( H, 0, 0)hex and (0, K, 0)hex
by 60◦ . The Miller indices H, K,
√ vectors, which are separated
√
and L have units of a∗ = b∗ = 4π/ 3a NN and c∗ = 2π/ 6a NN , where the nearest neighbour for
copper is a NN = 2.556 Å.
3. Results
The Cu(111) surfaces were characterised in a phosphate buffered electrolyte by surface X-ray
diffraction and with electrochemical methods.
3.1. Electrochemical Characterization
The Cu(111) electrode was characterised in phosphate buffer solution of concentration equal
to 0.01 M, 0.05 M and 0.1 M at different pH values of 6 and 8 with cyclic voltammetry at different
scan rates (Figure 1A). In the anodic scan, a shoulder (A1) at approximately −0.65 V and a peak (A2)
approximately −0.5 V can be observed at both pHs; however, at lower pH, the second anodic peak is
more defined. At pH = 6, the anodic peak shifts towards higher potentials with decreasing phosphate
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concentration (Figure 1B). The total charge associated to the electrochemical process shows equal
charge in both the anodic and cathodic scan of 52–60 μC/cm2 , which corresponds to ~0.2 electrons per
surface atom. The charge values were found to be independent of the pH and phosphate concentration
(Figure 1C,D).

Figure 1. Cyclic voltammetry profiles of a Cu(111) electrode in a phosphate buffer solution at (A) pH = 6
and (B) pH = 8 with phosphate concentration equal to 0.01 M (red line), 0.05 M (blue line) and 0.1 M
(magenta line). Scan rate ϑ= 0.02 V s−1 . (C,D) Charge involved in the oxidative () and reductive ()
process at the Cu(111) electrode as a function of the phosphate buffer concentration for pH = 6 and
pH = 8 respectively.

3.2. Structural Characterization
The Cu(111) surface was investigated in phosphate buffer solution of concentration 0.1 M at pH = 8
with in-situ surface X-ray diffraction. Throughout the experiment, we noted the difference in the
quality/morphology of the starting surface, which resulted specifically in a difference in reversibility
of adsorption processes and stability. The differences in surface morphology are described in more
detail below. Prior to any prolonged potential cycling resulting in excessive time spent at potentials
where phosphate is adsorbed, all surfaces exhibited a rather similar change in structure as the potential
was cycled. Figure 2 shows the X-ray intensities measured at the reciprocal space positions (0, 0, 1.4)
and (0, 1, 0.5), shown as dotted and solid lines, respectively, and are representative of the potential
dependent behaviour. The non-specular (0, 1, 0.5) position is sensitive to any changes in the atomic
structure of the Cu(111) surface or to the adsorption of species into well-defined Cu adsorption sites,
whereas the specular (0, 0, 1.4) position is sensitive to any changes in the surface normal electron
density distribution including the electrolyte side of the interface. The results indicate that there is a
clear structural change at the interface, most likely involving restructuring of the Cu surface itself and
this exhibits considerable hysteresis consistent with significant rearrangement of surface atoms.

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Following the X-ray voltammetry (XRV) measurements, such as those shown in Figure 2, the
Cu(111) surface structures were characterised while changing the potential in steps of 0.1 V over the
range from −1.0 V to −0.5 V. For each potential, a rocking scan was measured at the (0, 1, 0.5) position
and the integrated intensity obtained from this measurement is also shown in Figure 2 as black squares.
It can be seen that these lie halfway in the hysteresis loop. The surface was characterised by detailed
crystal truncation rod (CTR) measurements at potentials in the order −1.0 V, −0.9 V, −0.8 V, −0.7 V,
−0.6 and −0.5 V, after which the potential was stepped back to −0.8 V to check the reproducibility and
stability of the surface structure. The exact potential history of the experiment is shown in Figure 3
together with the integrated intensity and the peak width at the (0, 1, 0.5) position. It can be seen that
the peak width did not change throughout the course of the experiment, indicating a stable surface,
i.e., there was no change in surface morphology or domain size (which is inversely proportional to the
width of the peak in the rocking scan). The integrated intensity changed with the applied potential
as expected from the results shown in Figure 2. This change was reversible, as indicated by the two
datasets recorded at −0.8 V, before and after the potential cycle, which gave the same structural
parameters within the experimental error.

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Figure 2. The potential dependence of the X-ray intensities measured at the non-specular, (0, 1, 0.5),
(solid line) and specular, (0, 0, 1.4), anti-Bragg positions (sweep rate = 2 mV/s). In addition, the intensity
obtained from measuring the integrated intensity at selected potentials under static conditions is also
shown (black squares).

The CTR data and the best fit to the data obtained through a least-square fitting procedure for
potentials of −0.5 V and −0.8 V are shown in Figure 4. The structural model used to fit the data
allowed variation in the coverage of the topmost copper layer, the relaxation of the two topmost copper
layers and also included a β-roughness factor which is typically used to model surface roughness in
CTR measurements [22]. For the data measured at positive potentials it was also necessary to include
an adsorbed oxygen layer with the oxygen atoms occupying fcc/hcp three-fold hollow Cu sites. The
oxygen coverage and distance to the topmost Cu layer was allowed to vary to obtain the best fit. The
obtained parameters for the best fits to the data at the different potentials are summarised in Figure 5.
At potentials positive of −0.7 V, the coverage of the topmost copper layer started decreasing and
the oxygen coverage increased, together with an increased surface roughness. The inclusion of the
oxygen layer in the structural model resulted in a decrease of the reduced χ2 (indicating the goodness

176

Surfaces 2019, 2

of fit in the least squares method) from 3.3 to 1.8 for the fit to the data recorded at −0.5 V. In-plane
scans were measured at all potentials at L = 0.4 along the high symmetry <1, 0, L>, <0, 1, L> and
<1, 1, L> directions but did not show any additional superstructure peaks. Thus, we conclude that
there is no specific ordering of the phosphate or any oxygenated species into a commensurate ordered
adsorbate or copper oxide layer, as such ordering would give rise to additional scattering features
that would be detected in the scans along the high symmetry directions. Previous works, in gas phase
and computational studies, suggest the formation of hexagonal or quasi-hexagonal structures on the
Cu(111) surface upon oxygen adsorption, structures which can be viewed as the initial layer of a
Cu2 O (111) film. The model proposed by Platzman et al. suggests a three-step oxidation mechanism
comprising the formation of a Cu2 O layer, followed by the formation of a metastable overlayer of
Cu(OH)2 and finally the transformation of this metastable overlayer phase into CuO layer [23]. Our
results in electrochemical environment do not show any of these ordered structures in the potential
range under study.

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Figure 3. The surface stability has been monitored throughout the course of the experiment. (a) The
full width half maximum (FWHM) and integrated intensity obtained from rocking scans at the (0, 1,
0.5) position are shown as function of the scan number during the experiment. The potential is also
shown in red. The peak width is stable, indicating a stable surface with a domain size of ~100 nm.
(b) The integrated intensity is changing reversibly with the applied potential. (c) The peak profiles of
the rocking scan and the fit of the peak with a Lorentzian line shape are shown for scan numbers 158
and 383 at a potential of −0.9 V and −0.6 V, respectively.

177

Surfaces 2019, 2

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Figure 4. CTR data measured at −0.8 V (blue symbols) and −0.5 V (black symbols) together with
the calculated fits to the data (solid lines) according to the structural model described in the text. At
−0.5 V, an overall decrease in intensity at the surface sensitive positions can be observed. This has been
modelled with an increased roughness and adsorption of an oxygenated species into the surface Cu
layer. The change in intensity induced by the potential is fully reversible.

As noted above, the morphology of the surface depends crucially on the surface preparation
and transfer into the electrochemical system. For the above case following the electropolishing of
the Cu crystal, the transfer process to the X-ray electrochemical cell was smooth and the working
electrode was contacted and held at −0.8 V as soon as the crystal was in the cell. Any exposure to
oxygen during the preparation or not applying a potential in the oxygen reduction region immediately
upon immersion into the cell resulted in a different surface morphology than for the surface described
above. Previous work in gas phase has shown that at room temperature the oxidation of Cu(111)
proceeds through the epitaxial growth of copper oxide islands [24,25]. It has been shown that, in
gas phase, the growth of oxides on Cu(111) depends on the oxygen pressure and temperature and it
can follows three possible processes: (i) growth from step edges; (ii) interrace growth from vacancy
islands; and (iii) growth of on-terrace oxide [24,25]. The CTRs in this case showed additional peaks at
(0, 1, 1), (0, 1, 4), (1, 0, 2) and (1, 0, 5) (Figure 6), which arose from stacking faults (twinning) induced
by the nucleation of copper atoms into hcp sites. The CTR data in Figure 6 were best modelled by
including additional Lorentzian peaks to represent the twinned Cu, rather than a smooth surface
with stacking faults, and indicate that the peaks arose due to nano-crystalline copper islands on the
surface. Due to the presence of the additional peaks resulting from the nano-crystalline copper at the
surface, the CTRs could be reproduced reliably with a very simple model including surface roughness
and Cu surface relaxation only, i.e., for this more complex surface, the fitting was not sensitive to the
inclusion of an oxygen layer. The surface roughness and relaxation rely on the overall intensity change

178

Surfaces 2019, 2

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across the whole range of L and the intensity distribution close to the Bragg peaks, respectively. We
were not able to reliably include any coverage of individual atomic layers or oxygen atoms, which
changed the intensity close to the anti-Bragg position as the intensity was dominated by the Lorentzian
peaks arising from the nano-crystalline copper. Although these parameters, which are summarized
in Figure 7, are only indicative, a similar trend in the potential dependent relaxation and roughness
was observed when stepping the potential slowly positively. Surprisingly, at −0.8 V, modelling of the
CTR data indicated a much lower surface roughness parameter, β, than the data shown in Figure 4.
We attribute this to the presence of the nano-clustered copper islands acting as nucleation sites for
copper ad-atoms, thus decreasing the overall surface roughness. The changes observed during the
potential steps were, however, not reversible. Upon stepping back to the negative potential, an increase
in the surface roughness, as indicated by the red symbols in Figure 7, could be observed. This increase
in surface roughness was accompanied by a decrease in surface domain size and an increase in the
size of the nano-crystalline copper domains present on the surface. This was shown by the rocking
scan measurements at the (0, 1, 0.5) position (Figure 8), which, as for the data shown in Figure 3,
also indicated the potential history of this sample. A gradual increase in the width of the rocking
scan, from 0.2◦ to 0.4◦ , was observed, corresponding to a domain size decreasing from ~100 nm to
~50 nm. Correspondingly, the height D of the nano-clusters could be estimated from the width of the
peaks observed in the CTR data in Figure 6, D = 2 Π/ΔQz , and was found to grow in the course of the
experiment from 20 to 40 nm. Thus, it can be concluded that this Cu surface—with Cu islands present
on the surface—was unstable during potential cycling.






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Figure 5. The structural parameters obtained from modelling the CTR data are plotted as a function of
potential: (a) the coverage of the surface Cu layer and the adsorbed oxygen layer; (b) the relaxation of
the top two copper layers; (c) the roughness as modelled through a β-roughness model; and (d) the
Cu-O layer separation.

179

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Figure 6. CTR data measured from a different Cu surface preparation at −0.8 V (blue symbols) and
−0.5 V (black symbols) together with the calculated fits to the data (solid lines) according to the
structural model described in the text. In comparison to the CTR data presented in Figure 4, additional
peaks arising from Cu nano-islands can be observed.

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Figure 7. The structural parameters obtained from modelling the CTR data in Figure 6 are plotted as
a function of potential: (a) the β-roughness factor; and (b) the relaxation of the two topmost atomic
Cu layers. The data obtained after having stepped the potential to the positive limit are shown in red
symbols. Although the potential-induced change in structure is similar, the roughness of the surface
has increased.

180

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Figure 8. The surface stability of the Cu sample, which showed additional peaks arising from the
presence of Cu nano-crystals, i.e., equivalent to the data shown in Figure 3: (a) the FWHM; and (b)
the integrated intensity obtained from rocking scans at the (0, 1, 0.5) position as a function of the scan
number during the experiment. The potential is also shown in red. The peak width is increasing
and the integrated intensity is not recovering after having held the potential positively, indicating a
change in surface morphology with increased roughness and smaller domain size during the course of
the experiment.

4. Discussion
In this paper, we present a structural investigation of the Cu(111) electrode surface in
phosphate-containing neutral electrolyte solution. We have found two different surface behaviours
depending on the initial morphology and defect density of the surface observed directly after
preparation and transfer into the X-ray electrochemical cell. In both cases, however, a clear structural
transition with potential was found at −0.6 V, which coincides with a current peak in the anodic
scan. The discussion of the structural results is based on the dataset presented in Figures 4–6, which
gave the most concise insight into the underlying surface processes. At potentials negative of the
observed transition, the coverage of the copper is constant and no potential dependent relaxation is
observed. An increase in roughness can, however, be observed for increasingly cathodic potentials.
The structural transition occurring at potentials >−0.7 V can be attributed to an increased surface
roughness combined with adsorption of an oxygenated species, followed by the formation of a mixed
copper-oxygen layer at positive potentials. This transition is accompanied by an inward relaxation of
the surface Cu layer and a slight outward relaxation of the second atomic Cu layer.
Cyclic voltammograms were recorded at two different pH values and different phosphate
concentrations to get an estimate of the charge transfer and possible coverage of an anionic layer
and/or oxygenated species. For all electrolyte solutions, an anodic peak with a shoulder could be
observed indicating a slow two-step adsorption process. The charge under the anodic and cathodic
current peaks in the cyclic voltammetry corresponded to 0.2 electron transferred per surface atom. The
181

Surfaces 2019, 2

in situ X-ray structural analysis of the Cu(111) surface at −0.5 V gave an oxygen coverage θO = 0.6.
It is known that phosphate species have a tetrahedral structure with a P-O bond length of the order
dP-O ~ 1.5–1.6 Å, thus the corresponding O-O distance is dO-O ~ 2.44–2.62 Å [26–28]. The atomic
distance of the copper atoms at the Cu(111) surface is 2.556 Å, which is similar to the O-O distance in
the phosphate, thus making the adsorption of phosphate or hydrogen phosphate with three oxygen
atoms sitting in the same adsorption site likely. This suggests that the adsorbate species at −0.5 V can
be either hydrogen phosphate or phosphate and not dihydrogen phosphate. The exact nature of the
phosphate species and the adsorption process is, however, not fully understood. From the structural
analysis at −0.6 V, it is clear that the difference between the data recorded at the two most positive
potentials is not just due to a change in coverage of the adsorbed oxygenated species. The change
in relaxation of the topmost copper layer and the change in the distance of the oxygen atoms to the
Cu surface indicate an intermediate structure at this potential. This is also evident from the XRV and
potentiostatic measurements presented in Figure 2.
To further elucidate the adsorption process, the peak positions in the cyclic voltammetry measured
in 0.1 M phosphate solution at pH = 6 and pH = 8 are shown as a function of the logarithm of the
scan rate in Figure 9. The cathodic peak and the first anodic peak (corresponding to the shoulder in
the cyclic voltammogram) did not show a scan rate dependence, indicating a fast adsorption process,
whereas the second anodic peaks showed a slope of approximately 0.12 V/log (scan rate/(V/s)),
independent of the pH of the electrolyte, indicating is a slow process. The rate determining step in
the adsorption process is thus associated with the second anodic peak. The surface X-ray diffraction
data show that at the corresponding potentials a substantial rearrangement of the surface atomic
structure took place. The X-ray voltammetry obtained in-situ shows a large hysteresis consistent with a
slow mass transfer process. Static CTR measurement indicate the incorporation of oxygen/phosphate
into the Cu surface resulting in mixed surface copper-oxide and increased surface roughness in this
potential region. Structural rearrangements of the surface Cu layer with similarly slow kinetics have
been observed in alkaline electrolytes [29,30]. In that case, the slow kinetics were associated with
formation of a single Cu(I) oxide layer and a decrease of the Cu atomic density by 30% [31].

Figure 9. Peak potentials obtained from cyclic voltammograms are shown as function of the log of the
scan rate: (A) pH = 6; and (B) pH = 8. The black and red symbols correspond to the first and second
anodic peaks, respectively. The green symbols correspond to the cathodic peak.

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With the acid dissociation constant pKa(H2 PO4 - /HPO4 2− ) = 7.2, at pH = 8 a proportion of 6%
H2 PO4 − and 94% HPO4 − is expected, thus mostly hydrogen phosphate. At pH = 6 a proportion of
87% H2 PO4 − and 13% HPO4 2− is expected in the electrolyte solution. On gold and platinum surfaces,
the adsorbed phosphate species has been found to be pH- and potential-dependent [32,33]. In contrast,
on copper surfaces, phosphate dihydrogen was found to adsorb for acidic electrolyte solutions (pH < 5)
but not in alkaline solution (pH > 9) [34]. Although we cannot totally exclude that dihydrogen is
adsorbing in the first step on the Cu surface, the fact that the two adsorption steps were observed at
pH = 6 and pH = 8 in the electrochemical data and did not show a difference in the kinetics of the
processes, suggest that the adsorption process occurred through a similar mechanism. In addition, a
more covalent bonding between the oxygen of the phosphate and copper was found, in comparison
to silver and gold [34], and this could lead to the atomic rearrangement and slow mass transport
observed in the X-ray measurements.
By combining the structural and electrochemical results, we propose the following adsorption
mechanism: The first peak observed in the cyclic voltammograms corresponds to adsorption of
hydrogen phosphate to the surface (either directly or through the deprotonation of dihydrogen
phosphate depending on the pH). This is supported by the structural data where a small amount of
oxygen was found to be specifically adsorbed into three-fold hollow sites on the Cu surface. The second
anodic peak is associated with the deprotonation of the hydrogen phosphate and the adsorption of the
phosphate anion which forms a strong bond with the Cu surface leading to an atomic rearrangement
involving the incorporation of three oxygen atoms from the phosphate into the surface Cu layer. This
process is totally reversible, as also confirmed by the electrochemical characterization which gives the
same charge under the peak in the cyclic voltammogram for the anodic and cathodic process.
The elucidation of the adsorption process helps in understanding the differences in stability of
the Cu(111) electrodes during adsorption process and the dependence on the details of Cu surface
preparation. The in-situ X-ray diffraction data enable three main properties of the surface contributing
to the morphology to be distinguished: (a) the domain size of the atomic terraces present at the
surface; (b) the presence of Cu nano-crystals at the surface; and (c) the overall atomic roughness,
described through the β-model and hereafter referred to as rugosity. We have shown that, even
though the Cu(111) surfaces exhibit identical terrace lengths, the rugosity of the surface at terrace
level might be different. The Cu(111) surface with large rugosity shows a reversible adsorption
process of phosphates and oxygen incorporation into the crystalline structure. On the other hand,
the Cu(111) surface with copper nano-crystals present on the surface showed degradation upon
phosphate adsorption/desorption driven by the potential cycling. Surprisingly the rugosity was larger
on the surface showing reversible adsorption, the reversible nature of the adsorption/desorption
being evident from the stable intensity at the anti-Bragg position of the CTRs. The oxygen induced
reconstruction of the Cu(111) surface at room temperature have shown the formation of a disordered
surface overlayer, with O and Cu atoms at different heights [35]. It has been proposed that the
formation of this disorder overlayer favours the oxygen diffusion on Cu (111) leading to fast nucleation
of a large number of oxide islands. We propose a similar process in electrochemical media through the
oxygens of the adsorbed phosphates.
As the adsorption of phosphate increases the rugosity of the surface, we associate the differences
in stability of the two surfaces to the presence of surface defects acting as nucleation sites for Cu
ad-atoms created during the phosphate incorporation into the surface Cu layer. Depending on the
surface morphology, these atoms can be incorporated into the nano-crystals present on the surface,
thus leading to further roughening and decrease in the surface domain size, i.e., the Cu surface is
unstable. On a Cu electrode surface with high rugosity, no single nucleation sites stand out and
the formation of nano-crystals is prevented. This is a more stable Cu electrode surface exhibiting a
reversible phosphate adsorption process.

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5. Conclusions
In this paper, we present electrochemical and in-situ surface X-ray diffraction measurements
of phosphate adsorption onto Cu(111) electrodes. The underlying mechanism of the phosphate
adsorption and deprotonation of the (di)-hydrogen phosphate was accompanied by a roughening of
the Cu surface. We report for the first time the roughening of the Cu surface through the formation of a
mixed copper–oxygen layer, where the oxygen from an adsorbed phosphate species was incorporated
into the surface Cu layer. The stability of the Cu(111) surface and the change of rugosity upon potential
cycling strongly depend on the preparation method and history of the electrode. It was shown that
the presence of Cu islands on the Cu(111) surface leads to a 3D nucleation and growth mechanism
which causes irreversible changes in surface morphology. The results demonstrate the importance
of controlling the surface preparation of copper catalysts, as this determines the stability of the
catalyst under operation conditions for the electrochemical CO2 reduction reaction. The incorporation
of oxygen into the metal surface from the adsorbed phosphate anion is a process that could also
be relevant to the study of similar oxoanions, such as sulphate, and their role in the stability of
electrocatalysts during oxidation processes.
Author Contributions: P.R. and Y.G. conceived the experiments. All authors contributed to the experimental
measurements. Y.G., J.B. and C.A.L. performed the analysis of the in-situ X-ray measurements. All authors
contributed to the discussion, writing and revision of the manuscript. All authors have given approval to the final
version of the manuscript.
Funding: This research received no external funding.
Acknowledgments: We thank the Diamond Light Source for supplying the beamtime necessary to conduct the
measurements presented in this paper. In particular we would like to thank our local contacts, Jonathan Rawle
and Hadeel Hussain, and the Diamond I07 team for their support during the experiment. AK acknowledges
the University of Birmingham for financial support through Ph.D. scholarships at the School of Chemistry.
P.R. acknowledges the University of Birmingham for financial support through the Birmingham fellowship
program. Y.G. acknowledges the Royal Society for funding through a Royal Society Research Fellowship.
Conflicts of Interest: The authors declare no conflict of interest. The funders had no role in the design of the
study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to
publish the results.

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Cu(0 0 1) in 0.1 M NaOH. J. Electroanal. Chem. 2003, 554–555, 113–125. [CrossRef]
Yaguchi, M.; Uchida, T.; Motobayashi, K.; Osawa, M. Speciation of Adsorbed Phosphate at Gold Electrodes:
A Combined Surface-Enhanced Infrared Absorption Spectroscopy and DFT Study. J. Phys. Chem. Lett. 2016,
7, 3097–3102. [CrossRef] [PubMed]
Gisbert, R.; García, G.; Koper, M.T.M. Adsorption of phosphate species on poly-oriented Pt and Pt(1 1 1)
electrodes over a wide range of pH. Electrochim. Acta 2010, 55, 7961–7968. [CrossRef]
Niaura, G.; Gaigalas, A.K.; Vilker, V.L. Surface-Enhanced Raman Spectroscopy of Phosphate Anions:
Adsorption on Silver, Gold, and Copper Electrodes. J. Phys. Chem. B 1997, 101, 9250–9262. [CrossRef]
Jensen, F.; Besenbacher, F.; Lægsgaard, E.; Stensgaard, I. Oxidation of Cu(111): two new oxygen induced
reconstructions. Surf. Sci. Lett. 1991, 259, L774–L780.
© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Interface Science Using Ambient Pressure Hard X-ray
Photoelectron Spectroscopy
Marco Favaro 1, *, Fatwa Firdaus Abdi 1 , Ethan Jon Crumlin 2 , Zhi Liu 2,3 , Roel van de Krol 1 and
David Edward Starr 1, *
1

2
3

*

Helmholtz-Zentrum Berlin für Materialien und Energie GmbH, Institute for Solar Fuels,
Hahn-Meitner-Platz 1, D-14109 Berlin, Germany; fatwa.abdi@helmholtz-berlin.de (F.F.A.);
roel.vandekrol@helmholtz-berlin.de (R.v.d.K.)
Advanced Light Source, Lawrence Berkeley National Laboratory, 1 Cyclotron Rd, Berkeley, CA 94720, USA;
ejcrumlin@lbl.gov (E.J.C.); liuzhi@shanghaitech.edu.cn (Z.L.)
School of Physical Science and Technology, ShanghaiTech University, Pudong, Shanghai 201210, China
Correspondence: marco.favaro@helmholtz-berlin.de (M.F.); david.starr@helmholtz-berlin.de (D.E.S.);
Tel.: +49-(0)30-8062-42835 (M.F.); +49-(0)30-8062-42320 (D.E.S.)

Received: 18 December 2018; Accepted: 24 January 2019; Published: 28 January 2019

Abstract: The development of novel in situ/operando spectroscopic tools has provided the
opportunity for a molecular level understanding of solid/liquid interfaces. Ambient pressure
photoelectron spectroscopy using hard X-rays is an excellent interface characterization tool, due to
its ability to interrogate simultaneously the chemical composition and built-in electrical potentials,
in situ. In this work, we briefly describe the “dip and pull” method, which is currently used as a
way to investigate in situ solid/liquid interfaces. By simulating photoelectron intensities from a
functionalized TiO2 surface buried by a nanometric-thin layer of water, we obtain the optimal photon
energy range that provides the greatest sensitivity to the interface. We also study the evolution of the
functionalized TiO2 surface chemical composition and correlated band-bending with a change in the
electrolyte pH from 7 to 14. Our results provide general information about the optimal experimental
conditions for characterizing the solid/liquid interface using the “dip and pull” method, and the
unique possibilities offered by this technique.
Keywords: in situ ambient pressure XPS; hard X rays; photoelectron simulations; solid/liquid
interface; TiO2 ; APTES

1. Introduction
Molecular-level processes occurring at solid/liquid interfaces are intriguing not only from a
fundamental physical-chemical perspective, but also from a practical perspective since they are an
essential part of (photo)electrochemical systems, which are key to future renewable energy storage
technologies. This provides a powerful driving force for the development of novel in situ/operando
characterization tools that can directly probe the interface [1–7]. For a complete understanding of the
solid/liquid interface, it is mandatory to investigate the properties of the solid bulk, the bulk liquid,
and the thin interface layer between the solid and the liquid phases [8]. This constitutes the most
important part of the solid/liquid junction where specific adsorption of ions, charge transfer dynamics
and electrical potential formation occur.
On the other hand, the measurement of the interface layer properties is a challenging task. First,
it requires operando or in situ experimental capabilities in order to capture the true makeup of
the interfacial layer under operating or realistic working conditions, respectively [1,8,9]. Secondly,
to characterize the true interfacial properties, one should use an experimental probe that limits the
perturbation of the junction itself. Therefore, several spectroscopic methods based on photon in/photon
Surfaces 2019, 2, 78–99; doi:10.3390/surfaces2010008

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out or photon in/electron out approaches have been applied to investigate solid/liquid interfaces. Among
them, synchrotron-based techniques such as surface X-ray diffraction, X-ray absorption/emission
and photoelectron spectroscopy, and “laboratory-based” techniques, such as infrared, Raman and
non-linear optical spectroscopies have been recently developed [10].
In this context, X-ray photoelectron spectroscopy (XPS) stands as an excellent characterization tool,
since it offers elemental and chemical sensitivity, simultaneously measuring local built-in electrical
potentials via the detection of rigid photoelectron kinetic energy shifts in both core and valence
levels [11–17]. Due to the high vapor pressure of many liquids of interest, differential pumping schemes
have to be used in photoelectron analyzers to minimize the elastic and inelastic scattering of electrons
in the gas phase above the liquid side of the junction [18–20]. In addition, small sample-to-analyzer
aperture working distances (WDs) must be used, for the same purpose [18–20]. A reasonable trade-off
needs to be found between limiting the electron scattering by the gas phase molecules and keeping
the pressure at the sample surface above 90–95% of the nominal pressure in the chamber. Usual
WDs (at which the analyzer focus is optimized) are about the diameter of the aperture itself [18–20].
Modern state-of-the-art electron analyzers are capable to operate at pressures of and above 30 mbar
(the vapor tension of water at room temperature) and at high photoelectron kinetic energies (KEs,
up to 12 keV) [11]. The extension of AP-XPS to high photon energies (i.e., high photoelectron KEs) has
two main advantages: the reduced photoelectrons/gas molecules scattering provides higher signal
intensity at higher gas pressure (i), whereas the high photoelectron KEs lead to increased photoelectron
inelastic mean free paths (IMFPs) in the liquid (ii). The latter enables the investigation of solid/liquid
junctions through liquid layers with thickness on the same order of the IMFP (ii).
Two additional challenges exist for the AP-XPS investigation of solid/liquid electrolyte interfaces.
The first is the preparation of stable liquid films thick enough to be representative of a realistic
solid/liquid interface, but thin enough to allow photoelectrons ejected from the interfacial region
to penetrate and emerge from the liquid on their path to the photoelectron analyzer. The second is
the tuning of the experimental conditions (detection geometry, photon energy etc.) to enhance the
photoelectron signal coming from the interface.
In this work, we briefly describe the so-called “dip and pull” method that can be used as a
way to investigate the solid/liquid interface in situ by coupling it with ambient pressure hard X-ray
photoelectron spectroscopy (AP-HAXPES) using X-ray energies between 2.0 and 10.0 keV [21].
By using a model system of a solid/liquid interface and numerical simulations, we estimate
the optimal photon energy range for the signal coming from the interface. The simulations have
been generated using the Simulation of Electron Spectra for Surface Analysis (SESSA) software [22].
Our model system is a TiO2 surface functionalized with 1 monolayer equivalent (1 MLE) of
3-aminopropyl triethoxysilane (APTES). Alkylsilanes adsorb at the surface of transition metal
oxides [23] as self-limiting monolayer systems, thereby constituting ideal interfaces. Furthermore,
APTES contains one nitrogen and silicon atom, whose corresponding 1s photoionization yield was
monitored as a function of the photon energy. The difference in binding energy (BE) between these two
core levels (N 1s BE ~ 400 eV, Si 1s BE ~ 1850 eV) will be used to highlight the different photoionization
yield trends with the photon energy.
In the last part of this work, we report on the evolution of the APTES-functionalized TiO2 surface
chemistry and correlated band-bending by systematically changing the pH of the aqueous electrolyte
from 7 to 14. This was performed at beamline 9.3.1 at the Advanced Light Source (ALS, Lawrence
Berkeley National Laboratory, Berkeley, U.S.A.) [11]. Our results provide general information about
the experimental conditions that allow for an optimal characterization of the solid/liquid interface
using the “dip and pull” method and AP-HAXPES, showing at the same time the capabilities offered
by this technique to address fundamental questions in energy materials and conversion research.

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2. Materials and Methods
2.1. Numerical Simulation Details
The simulations reported in this work have been generated using the Simulation of Electron
Spectra for Surface Analysis (SESSA) software (National Institute of Standards and Technology,
Gaithersburg, MD, USA.) [22].
The APTES thickness corresponding to 1 monolayer equivalent (MLE) was found as follows,
using SESSA simulations. First, 1 ML of TiO2 having a thickness of 3.57 Å (the relaxed Ti-Ti distance
along the rutile (110) direction [24]) was supported on a planar Cu substrate. Then, the APTES
layer thickness deposited atop the TiO2 film was systematically varied until the differential cross
section-normalized integrated photoelectron intensities of Ti 2p and Si 1s were identical. This leads
to an identical concentration of Si atoms with respect to the Ti atoms contained within 1 ML of TiO2
(for practical reasons chosen along the rutile (110) direction). Therefore, the thickness of APTES
corresponding to 1 MLE on rutile TiO2 (110) was found to be 12.7 Å.
Note that the differential cross sections (dσ/dΩ)nl (the probability for photoelectrons to be emitted
into the solid angle dΩ) were computed in SESSA using the first order nondipole approximation,
in which nondipole effects are treated in a perturbative manner. Within the Cooper parametrization,
the differential cross section (dσ/dΩ)nl for 100% linearly polarized radiation [25] can be expressed
as follows:
1
(dσ/dΩ)nl = (σnl /4π) × {[1 + βnl × (3cos2 α − 1)] + [sinαcosϕ × (δ + γ·cos2 α)]}
2

(1)

σnl is the total photoionization cross section (TPCS) for a core level described by a principal
and angular quantum number n and l, βnl is the corresponding asymmetry parameter describing the
anisotropy of the photoelectron angular distribution, α is the angle formed between the photoelectron
direction k and the polarization vector ε, ϕ is the angle formed between the radiation Poynting vector
S (parallel to the propagation direction) and the plane containing k and ε (see Figure 1 for the definition
of the aforementioned angles). Since the simulations have been performed considering radiation from
a synchrotron (horizontal photon beam) and bending magnet as source (electric field polarization
vector laying on the orbit plane, see Figure 1), the angle ϕ is equal to 0◦ for all α values (in our
simulations, α = 15◦ as implemented at BL 9.3.1, ALS). The term within the first square parentheses
represents the dipole contribution, whereas the second term is the parametrization of the first order
nondipole effects, which are non-negligible for photon energies above 2000 eV [25]. Within the
nonrelativistic single-electron approximation, the γ parameter represents the major correction term
corresponding to electric dipole-quadrupole interference. The magnetic dipole term, δ, can only be
present if core-relaxation occurs [25]. δ and γ, as well as σnl and βnl , are explicit functions of the photon
energy. Finally, the take-off angle θ (angle between the photoelectron direction k and the normal n to
the sample surface) was kept to 0◦ for all the simulations (normal emission detection, NE).
The atomic density (number of emitters per unit volume) of TiO2 and APTES used in the
simulations were equal to 9.568 × 1022 cm−3 and 9.522 × 1022 cm−3 , respectively (the APTES
stoichiometry used in the simulations was SiNO3 C9 H23 ). The energy gap between the valence and
conduction bands of TiO2 was set to 3.03 eV [26]. The energy gap between the HOMO and LUMO of
APTES was determined from the fluorescence spectrum of APTES, which exhibits an intense band
at about 438 nm [27]. The related HOMO-LUMO energy gap was therefore estimated to be equal to
2.83 eV.
To complete the model of the solid/liquid interface, a liquid water layer with a thickness of
20 nm was placed atop the sample surface, as schematically reported in Figure 1. This value has
been taken as a realistic representation of the liquid layer thickness typically obtained during “dip
and pull” experiments. The water atomic density and band gap used in the SESSA simulations were
1.003 × 1023 cm−3 and 6.9 eV [28], respectively.

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Figure 1. Schematization of the ambient pressure hard X-ray photoelectron spectroscopy (AP-HAXPES)
experimental geometry (implemented in Simulation of Electron Spectra for Surface Analysis (SESSA)
calculations) and 3-aminopropyl triethoxysilane (APTES)-functionalized TiO2 model system used in
this work.

2.2. Sample Preparation
The TiO2 precursor solution was made by mixing titanium-isopropoxide (Alfa-Aesar, Haverhill,
MA, USA.) with glacial acetic acid (Sigma-Aldrich, Saint Louis, MO, USA.) in anhydrous ethanol. 1 mL
of Ti-isopropoxide was added to 5 mL of anhydrous ethanol followed by adding varying amounts of
glacial acetic under vigorous stirring to give mole ratios ranging from 1:1 to 1:2, Ti-isopropoxide:acetic
acid. All solutions resulted in TiO2 films of similar quality. The precursor solution was spin-coated
onto FTO/glass substrates in four 100 μL aliquots at 2000 rpm for 120 s. Prior to spin coating the
substrates were sonicated in a 1% solution of Triton X-100 (Sigma-Aldrich, Saint Louis, MO, USA.)
in de-ionized water, DI water, and a 1:1 solution of acetone:absolute ethanol, followed by a final
sonication in DI water. The spin-coated films were sintered at 500 ◦ C for 12 h.
The TiO2 coated substrates were functionalized with APTES by placing the dried (300 ◦ C for
15 min) substrates into a solution of 0.25 mL APTES in 100 mL of anhydrous toluene under Ar(g) at
80 ◦ C for 5 h. Following reaction, the APTES/TiO2 coated substrates were dried at 70 ◦ C for 30 min
and packed into vacuum sealed bags in a N2 (g) filled glove box.
2.3. In Situ AP-HAXPES Measurements and Beamline 9.3.1 Experimental Details
The source for beamline 9.3.1 at the Advanced Light Source (ALS) of Lawrence Berkeley National
Laboratory is a bending magnet. A Si(111) double crystal monochromator (DCM) provides an energy
range between 2000 and 6000 eV (denoted as “hard” throughout the text to be consistent with the
acronym HAXPES). The minimum X-ray spot size on the sample is 0.7 mm (v) × 1.0 mm (h). All spectra
were taken with a photon energy of 4000 eV, at room temperature, and in NE (see Figure 1 for a
schematic of the experimental geometry). The pressure in the experimental chamber was kept at
the equilibrium vapor pressure of the aqueous electrolyte solution at room temperature. To limit
evaporation from the electrochemical cell a large volume (∼500 mL) of outgassed pure water was
introduced in the analysis chamber.

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The pass energy of the Scienta analyzer (R4000 HiPP-2, Scienta, Uppsala, Sweden) was set to 200 eV.
A step size of 100 meV and a dwell time of 300 ms was used. Under these conditions, the total resolution
(X-ray and electron analyzer) was equal to about 550 meV at 4.0 keV, determined by measuring the
spectral broadening of a gold Fermi edge taken on a clean polycrystalline surface. Binding energy (BE)
scale calibration was done by using the Au 4f7/2 (BE = 84.00 eV) photoelectron peak position taken on a
clean gold polycrystalline surface as reference values measured under all the experimental conditions.
During operation, the sample and the electron energy analyzer were commonly grounded. Spectral
fitting was carried out using a Doniach-Sŭnjić shape for the Au 4f photoelectron peaks, whereas
symmetric pseudo-Voigt functions (G/L ratio ranging from 85/15 to 75/25) were used to fit the Si 1s,
O 1s, Ti 2p and N 1s core levels (after Shirley background subtraction). During the fitting procedure,
the Shirley background was optimized together with the spectral components. Finally, chi-square (χ2 )
minimization was ensured by the use of a nonlinear least-squares routine, with increased stability over
simplex minimization.
2.4. “Dip and Pull” Method and in Situ Measurements
Prior to its introduction into the experimental chamber, the aqueous solutions at different pH
(see Table 1) were outgassed for at least 30 min at low pressure (between 10 and 20 Torr) in a dedicated
offline chamber. After degassing, the electrolyte solution was placed into the AP-HAXPES analysis
chamber, where the pressure was carefully decreased to just below equilibrium vapor pressure of the
electrolyte at room temperature (the chamber was pumped by a MVP 035-2 diaphragm pump, Pfeiffer
Vacuum, Aßlar, Germany). Then, the pumping via the diaphragm pump was stopped and the pressure
drifted up to the equilibrium vapor pressure of the used electrolyte solution at room temperature
(Table 1). At this point, the only active pumping was provided by the aperture of the HEA nozzle
that separates the high pressure analysis chamber from the differentially-pumped electrostatic lens
system (aperture diameter: 300 μm). To create the solid/liquid electrolyte interface, the sample was
immersed deeply into the aqueous solution. It was then slowly extracted from the liquid by raising the
sample manipulator at a constant rate (100 μm s−1 ). Following this procedure, a thin layer of aqueous
electrolyte film was formed on the sample surface. The latter was then positioned at the intersection of
the X-ray beam and the focal point of the hemispherical electron analyzer (HEA), thereby allowing
AP-HAXPES measurements of the solid/liquid interface.
Table 1. Concentration and pH values for the different aqueous solutions used in this study. The water
vapor partial pressure above pure water and KOH aqueous solutions at different concentration has
been determined using the Bridgeman-Aldrich [29] and the Balej equation [29], respectively.

Aqueous Solution

KOH Concentration
[M = mol L−1 ]

pH

H2 O Vapor Partial
Pressure at r.t. (= 298 K)
[mbar]

mQ H2 O
KOH 10−4 M
KOH 10−2 M
KOH 1.0 M

–
10−4
10−2
1.0

7
10
12
14

31.67
31.41
31.39
30.29

3. Results and Discussion
3.1. The Challenge of Preparing a Realistic Solid/Liquid Interface
The main challenge in the investigation of solid/liquid interfaces using a photon in/electron out
technique lies in the preparation of the interface itself. The electron detection implies that the electrons
ejected from the interfacial region need to penetrate through condensed phases whose thickness must
be less than two to three times the electron inelastic mean free path (IMFP). Although the use of hard
X-rays allows obtaining high photoelectron KEs (and therefore IMFPs), a photoelectron having a KE

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spanning from 2000 eV to 10000 eV is characterized by an IMFP on the order of tens of nanometers.
This poses severe limitations and constraints in the preparation of the solid/liquid interface.
Currently, two different preparation and investigation approaches are used, which differ from
each other for the side of the interface that is used for the X-ray incidence and electron detection [9].
Photoelectron spectroscopy can be conducted from the solid side of the solid/liquid interface
using few layer graphene membranes supporting the solid phase [30,31]. Although this technique is
appealing for the fact that gases or liquids can be flown through the system thereby providing facile
mass transport, it limits the investigation to thin solid films since the photo-emitted electrons must
travel through the solid phase/graphene membrane to reach the photoelectron analyzer. For instance,
a photoelectron travelling in Au with a KE of 500 and 5000 eV is characterized by an IMFP of about
0.8 and 4.2 nm, respectively. Using the definition of the information depth as three times the IMFP
(i.e., the depth in the material at which 95% of the ejected photoelectrons are inelastically scattered in
their path toward the surface), we find that the maximum probed thickness of Au achievable with soft
(hard) X-rays is equal to about 2 nm (13 nm).
A second approach, using X-ray incidence and electron detection from the liquid side, requires
the formation of a thin liquid layer atop the solid surface [9]. This approach has the advantage of
being able to investigate a broad range of solid materials of arbitrary thickness. This is particularly
important for photoelectrochemical (PEC) interfaces, due to the diffusion length of the excited
charge carriers (λCC ) that can span from tens to hundreds of nm. This poses a constraint on the
thickness (t) of the investigated semiconductor, where t ~ λCC . In addition, such approach is
applicable to fundamental investigations at interfaces such as the probing of the electrical potential
distribution simultaneously within the solid (i.e., band-bending) and the liquid side of the junction
(the double/diffuse layer) [13,17].
Similar to the previous discussion, the maximum thickness of the liquid layer must be three times
the photoelectron IMFP to allow the photoelectrons ejected from the interface to penetrate the liquid
film, to cross the liquid/gas interface and finally to travel through the gas phase on their path toward
the electron analyzer. Therefore, for photoelectrons traveling in water with a KE of 500 eV and 5000 eV,
the maximum water thickness must be equal to about 6 nm and 38 nm, respectively.
The preparation of liquid layers characterized by a thickness of the order of tens of nanometer
and that are stable for the duration of the measurements (often several hours) is not straightforward
and requires the development of novel techniques. As schematically reported in Figure 2, three
experimental procedures have been identified to obtain such “free-surface” liquid layers: the “emersion
technique” (recently called ‘’dip and pull”, Figure 2a), “the tilted sample” procedure (Figure 2b) and
the “offset droplet” method (Figure 2c).
The “dip and pull” method was developed from the early works on the extended liquid meniscus
of Bockris and Cahan [32], Siegbahn [33] and in particular by Hansen and Kolb [34–36]. It obtains
nanometer-thick liquid layers by partially extracting the sample from the liquid solution under
controlled conditions (i.e., equilibrium between the liquid and its vapor at a given temperature,
constant rate of pulling). This method will be discussed in detail in the next Section 3.1.1.
In “the tilted sample” technique developed by Kötz and coworkers [37], the liquid is contained in
a pocket in which the sample is immersed with a tilt angle. This tilt creates a region at the boundary
between the liquid free surface and the sample that can be positioned at the focal point of the analyzer
and irradiated by the X-ray beam. The authors were able to detect the photoelectron spectra of the
elements of the investigated ionic liquid (1-ethyl-3-methylimidazolium tetrafluoroborate) and of the Pt
sample simultaneously in one survey spectrum, using the Al Kα emission line as excitation source [37].
Similar to what has been described above, it is possible to create a stable evaporation/condensation
equilibrium by dosing water in the analysis chamber to match the water vapor pressure above the
liquid contained in the pocket, at the given temperature of the experiment.

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Figure 2. Schematization of the “dip and pull” (a), “tilted sample” (b) and “offset droplet” techniques
(c) that are successfully used to prepare free-surface liquid layers thin enough to probe the solid/liquid
interface using photoelectron spectroscopy.

Recently, Walton and coworkers developed instead an alternative way by using a fine capillary
(100 μm internal diameter) to inject and control a droplet deposited in situ on the sample surface [38].
Since an offset exists between the position of the capillary and the measurement spot, the authors
named this technique as the “offset droplet” method. The capillary is connected with an external HPLC
pump via a liquid feedthrough. By applying a fixed flow rate to the pump, the liquid can be directly
introduced into the analysis chamber and deposited on the sample surface. After an initial transient in
which the pressure in the analysis chamber will equilibrate to the liquid vapor tension, a steady-state
condition can be reached by choosing an appropriate flow to counterbalance the evaporation rate.
By rigidly moving the sample and the capillary, the droplet edge can be positioned at the intersection
of the photon beam and the analyzer focus, enabling therefore photoemission experiments at the
solid/liquid interface.
It is worth mentioning that the three aforementioned techniques do not allow a fine control of the
liquid layer thickness, although it is possible to “select” the suitable thickness for the experimental
purposes. With the “dip and pull” technique, it is possible to change the height of the measurement spot
by changing the positions of the sample and liquid container with respect to the X-ray beam-analyzer
focus intersection. The same procedure can be applied to the “tilted sample” and “offset droplet”
techniques, by changing the tilt angle and the liquid injection/suction rate, respectively. The described
procedure holds exclusively for wettable surfaces, where the contact angle ψ at the liquid meniscus is
smaller than 90◦ (indicating relatively intense interactions between the solid and the liquid). This poses
a severe restriction on the type of samples/surfaces accessible using these techniques, due to the fact
that no nanometric-thick continuous liquid layer can be formed on non-wettable surfaces (ψ ≥ 90◦ ).
A practical example of this drawback will be given in Section 3.3.
3.1.1. The “Dip and Pull” Method
Figure 3 reports a schematization of the “dip and pull” procedure. The sample is first placed in
front of the analyzer nozzle, aligning its bottom edge with the nozzle aperture (Figure 3a). The sample
is then lowered till its bottom edge enters in contact with the liquid free surface in the electrolyte
beaker (Figure 3b). By measuring these two positions using the sample manipulator encoder, it is
possible to get the value d, which is the distance between the nozzle aperture and the free surface
of the electrolyte (typically some millimeters). The sample is then immersed into the liquid by the
same length, plus an additional portion Δ (Figure 3c). The sample is then retracted from the liquid by
a distance d, as reported in Figure 3d. In this manner, the point P that was previously placed at the
liquid meniscus (liquid/gas interface) is now aligned with the nozzle aperture, thereby facilitating the
search of the optimal liquid layer thickness. By scanning the region around P (changing the height

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of the measurement spot as reported above) it is possible to find the suitable layer thickness for the
experiment (in the case of AP-HAXPES the film thickness can span from few to some tens of nm).
The liquid film thickness depends on a number of factors and their corresponding interplays:
on the wettability of the solid surface for a given liquid (i), the height of the measurement spot above
the free surface of the bulk liquid (ii), the presence of an eventual temperature gradient between the
two positions (iii), and the solvent evaporation rate (iv). To counterbalance water evaporation from
both the liquid layer and the container, two options can be used: either dosing water in the analysis
chamber through a valve connected with a heated external reservoir, or by introducing a second
(larger) water volume within the analysis chamber. Both solutions work at pressures close to the
water vapor pressure at room temperature (100% relative humidity, RH), thus generating a dynamic
equilibrium at the thin liquid layer where the evaporation and condensation rates are the same. Then,
by moving the sample (and the liquid container accordingly), the interface can be positioned at the
intersection of the incident X-ray beam and the focal point of the electron analyzer, thereby enabling in
situ photoemission measurements of the solid/liquid interface.

Figure 3. Schematization of the “dip and pull” method coupled with AP-HAXPES experiments. See
text for the detailed explanation of the experimental procedure. The pressure inside the analysis
chamber is generally close to the water vapor pressure at room temperature, thus enabling to work at
about 100% relative humidity.

The “dip and pull” method can be also utilized to perform in situ (photo)electrochemical and
photoemission measurements. Two additional electrodes can be mounted on the sample holder,
leading to an effective three-electrode electrochemical cell [11–17]. The bottom part of the electrodes
(Δ), being kept in the bulk electrolyte, provides the electrochemical continuity to apply a potential to
the thin electrolyte layer on the sample (working electrode) surface, thus enabling the investigation
of solid/liquid electrified interfaces [11–17]. A drawback of this method (and, in general, of the
“free-surface” liquid layer approach) is constituted by the mass transport limitation in the electrolyte
layer, since the latter is effectively static in the direction parallel to the solid surface (with the exception
of liquid flow due to eventual differences in the temperature between the electrolyte reservoir and
the measurement spot above the liquid meniscus, causing thermo-capillary or Bénard–Marangoni
convection). The mass transport limitations have been experimentally addressed in our recent
works [15,39]. We have demonstrated that the liquid layer undergoes instability for faradaic reactions
involving consumption of the electrolyte, such as during the oxygen evolution reaction (OER) in
1.0 M KOH aqueous solution [39]. Holding the potential at the Pt working electrode at +1.93 V vs.
RHE (reversible hydrogen electrode), we observed the loss of potential control within the liquid
layer in less than 2 h from the beginning of the experiment [39]. The loss of anodic polarization
was assessed by the deviation of the binding energy (BE) shift of the O 1s liquid phase water (LPW)
component and K 2p core levels from the applied OER potential, as a function of the observation
time [39]. The following mechanism is likely to occur: first, the hydroxyl anions are depleted from the
thin liquid layer due to the ongoing oxidation to molecular oxygen. Second, the consequent decrease
over time of the pH within the liquid layer leads to an increasing IR drop, which is responsible for
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the observed BE deviation from the applied potential. Interestingly, after the loss of potential control
occurred, we observed also an important decrease of the LPW signal mirrored by the K 2p intensity,
thereby indicating a progressive thinning of the liquid layer. To estimate the diffusion time scale of
the hydroxyl groups from the liquid meniscus through the liquid layer, we can use Fick’s first law of
diffusion (Equation (2)). The distance z of the liquid meniscus from the AP-HAXPES measurement
position is typically about 0.8 mm, and the diffusion coefficient D of hydroxyl anions in water at
infinite dilution at room temperature is equal to 5.30 × 10−5 cm2 s−1 [40].
J  = n/(Σ × Δt) = −D (dC/dz)

(2)

J  represents the diffusional flux parallel to the sample surface, which is the amount of hydroxyl
groups (n, in mol) flowing through the liquid layer cross-section (Σ) within the time Δt, whereas
(dC/dz) is the 1-dimensional concentration gradient within the liquid layer (parallel to the sample
surface) between the measurement spot and the liquid meniscus. First, we take into account a
liquid layer thickness of 20 nm, which is a typical value for the “dip and pull” technique, and
0.7 mm as the lateral dimension of the sample. The cross-section Σ of the electrolyte film is therefore
0.7 cm × 20 × 10−7 cm = 1.4 × 10−6 cm2 . Secondly, let us define a complete depletion of hydroxyl
within the liquid layer after a time Δt. This sets the concentration gradient (dC/dz) to be linear
between the measurement spot and liquid meniscus, where the OH− concentration (C) is nominally
equal to 1.0 M (1 mol L−1 ). Therefore, (dC/dz) = 1 mol L−1 /0.8 cm = 10−3 mol cm−3 /0.8 cm =
1.25 × 10−3 mol cm−4 . We can then determine the flux J of hydroxyl through the cross section Σ
per unit time: J  = n/(Σ × Δt) = C × Σ × z/Σ × Δt = C × z/Δt = 10−3 mol cm−3 × 0.8 cm/Δt =
8.0 × 10−4 mol cm−2 /Δt. Using Equation (2) we can calculate Δt, which is the time needed by the
hydroxyl groups to diffuse from the liquid meniscus to replenish the solution volume within the thin
electrolyte layer. Δt is found to be about 12 × 103 s (~3.3 h). This limitation in the ionic diffusion
rate is confirmed by a recent study of Shavorskiy et al., where the authors observe a significant IR
drop in the liquid film at the hematite/liquid electrolyte interface under PEC conditions (for applied
potentials above ~1.2 V vs. RHE) [17]. This mass transport limitation has also an important effect on
the achievable current densities within the liquid layer. In a recent study, we masked the bottom of a
Pt sample immersed in KOH 1.0 M aqueous electrolyte and compared to an unmasked electrode [15].
We determined that a bulk (unmasked) current density of about 1.0 mA cm−2 under OER conditions
corresponds to about 0.3 mA cm−2 in the emersed part of the Pt surface [15]. This observation was
confirmed by the polarization resistance (Rp ) for the masked and unmasked configurations, determined
using electrochemical impedance spectroscopy (EIS) measurements. In line with the current densities
results, the ratio between the two configuration Rp (unmasked to masked) was equal to 3.37 [15].
It is important to highlight that it was not possible to completely avoid the contribution arising
from the macroscopic liquid meniscus between the sample and the liquid free surface. Therefore,
the actual current density at the AP-HAXPES measuring spot might be even lower. As a general
remark, the interplay between the applied overpotential and the limitations in the mass transport
kinetics through the liquid layer plays a crucial role during “dip and pull” experiments, and needs
to be evaluated case by case (depending on the nature of the solid surface, the catalytic activity of
the (photo)electrocatalyst etc.). Moreover, the electrode potential at the measuring spot must be
corrected for the actual pH value (since an eventual concentration gradient leads to a pH gradient as
well within the liquid layer). This can be performed by calculating the electrolyte/water ratio using
the corresponding photoionization core levels (such as K 2p and O 1s for a KOH aqueous solution,
as reported in reference [15]).
The lowering of the reaction kinetics within the liquid layer due to the mass transport limitations
can be exploited to investigate the evolution of the interfacial properties on a time scale accessible
by AP-HAXPES (seconds). For instance, in a recent work we were able to monitor the light-induced
formation of a bismuth phosphate (BiPO4 ) layer atop a polycrystalline bismuth vanadate (BiVO4 )
surface, working at the half-cell open circuit potential and in a phosphate-containing electrolyte layer
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with a thickness ranging between 24 and 32 nm [41]. We found that the BiPO4 formation was reversible
upon restoring dark conditions. The BiPO4 formation and dissolution kinetics have been characterized
by fitting the temporal trend of the phosphate/water liquid phase (LPW) intensity ratio under visible
light and dark conditions, respectively (the spectral components have been determined from the O 1s
core level spectra acquired under light and dark conditions as a function of time). The retrieved time
constants of the BiPO4 formation and dissolution were found to be equal to 321 ± 61 s and 498 ± 89 s,
respectively [41].
Finally, a general challenge closely related to the use of the aforementioned methods concerns the
enhancement of the photoelectron signal coming from the narrow interfacial region atop the solid and
buried by the liquid side of the junction. In the next section, using a model solid/liquid junction and
numerical simulations, we will address this point and assess the optimal X-ray energy that allows the
enhancement of the interface photoelectron signal.
3.2. Optimization of the AP-HAXPES Experimental Conditions
The simulations have been generated using the SESSA software [22] (see “Materials and Methods”
section for the simulation details). Our model system is constituted by a TiO2 surface functionalized
with 1 MLE of APTES, characterized by a thickness of 12.7 Å. To complete our model of the solid/liquid
interface, a water layer with a thickness of 20 nm was placed atop the sample surface, as schematically
reported in Figure 1. This value has been taken as a realistic representation of the liquid layer thickness
typically obtained during “dip and pull” experiments and accessible by HAXPES [11–17].
The use of HAXPES in combination with AP experiments leads to several advantages:
1.

2.

3.

4.

5.

The high photon energy (and correspondingly the photoelectron KE) drastically decreases
the secondary electron emission cross sections. This has a beneficial effect in limiting the
radiation-induced damage suffered by the sample and the correlated radiolysis of water [42];
The inelastic scattering between photoelectrons and gas molecules decreases as a result of the
large energy difference existing between the X-ray energy and the rotovibrational/electronic
excitations and ionization thresholds in the gas molecules (typically falling within the UV-VIS
and soft X-ray regions, respectively, for most gases of interest such as O2 , N2 , H2 O, CO, CO2 ,
gaseous hydrocarbons and alcohols);
The elastic scattering between photoelectrons and gas molecules, which is responsible for
smearing out the photoelectron angular distribution, is less pronounced due to the generally
increasing forward focusing effect as the excitation energy increases [43];
To maximize the overall photoelectron yield in HAXPES, quasi-normal emission detection is
typically coupled with X-ray grazing incidence (α ≤ 5◦ ) [44–47]. This has the advantage of
minimizing the secondary (inelastic) electron background [46];
The high photoelectron KEs lead to increased photoelectron IMFPs in the liquid and thus enable
the investigation of solid/liquid junctions through liquid layers whose thickness is of the same
order of the photoelectron information depth.

On the other hand, the large probing depth reduces the relative contribution of the interface region
with respect to the overall detection (which includes the bulk liquid and solid signals). In addition,
the total photoionization cross sections (TPCSs) decrease with increasing photon energy, thereby
further reducing the signal coming from the interface (within our model, the signal coming from the
APTES overlayer whose thickness is comparable to typical electrical double layer thickness in solution).
Figure 4a,b clearly illustrate the scenario, reporting the photoelectron IMFP (right axes) and TPCS
(left axes) for Si 1s and N 1s core levels (from the APTES overlayer) as a function of photon energy.
The IMFP is calculated for photoelectrons traveling in pure water. The IMFP (Λe ) and TPCS (σ) trends
have been fitted using the following relations (Equations (3) and (4), respectively):
Λe = C × (hν)p
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(3)

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σ = σ0 × exp[−{(hν − hν0 )/τ}]

(4)

C, p, σ0 and τ are fitting parameters, whereas hν denotes the photon energy. Note that we have chosen
to report the TPCS instead of the differential PCS (DPCS) in order to provide a general discussion
about the photon energy dependency of the PCS, since the DPCS strongly depends on the specific
experimental geometry and utilized multipole expansion truncation (see “Materials and Methods”
section for details). The TPCSs have been determined by interpolation of the TPCS values reported by
Yeh and Lindau [48], whereas the electron IMFPs have been computed using the Tanuma–Powell–Penn
(TPP-2M) predictive equation [49]. Table 2 reports the determined p, τ and σ0 values for both analyzed
core levels, which will be used later to rationalize the observed behavior of the Si and N 1s core level
photoelectron intensity trend as a function of the photon energy.

Figure 4. Electron inelastic mean free path (IMFP, right axes) and total photoionization cross
section (TPCS, left axes) as a function of the photon energy, for Si 1s (a) and N 1s (b) core levels.
The photoelectron kinetic energy scale has been generated from the photon energy values, using
Einstein’s energy conservation law and a binding energy of 1845.0 eV and 400.0 eV for Si 1s and N 1s
core levels, respectively. The IMFPs are determined for photoelectrons travelling through water. Red
curves are fits according to Equations (3) and (4).
Table 2. p, τ, and σ0 fitting parameters describing the functional dependency of the electron inelastic
mean free path (IMFP) and total photoionization cross section (TPCS) on the photon energy, for Si 1s
and N 1s core levels. Note that the cross section unit is expressed in barn (1 barn = 1 · 10−24 cm−2 ).
Core level

Nominal Binding
Energy [eV]

p

τ [eV]

hν0 [eV]

σ0 [Mbarn]

Si 1s
N 1s

1845.0
400.0

0.843 ± 0.01
0.850 ± 0.01

1387 ± 18
622 ± 7

2500
1000

0.062 (at 2500 eV)
0.071 (at 1000 eV)

Therefore, it is necessary to find the optimal photon energy that enables to selectively enhance the
signal coming from the interfacial region. In addition, as Figure 4a,b and Table 2 show, this evaluation
needs to be performed for the different core level spectra involved in the investigation, since they are
characterized by different cross sections with different photon energy dependency.
Let us start by quantitatively defining the photoelectron intensity for a given core level
characterized by principal and angular quantum numbers n and l, respectively. The photoelectron
intensity Inl in vacuum, normalized by the incident X-ray flux integrated over the irradiated sample
volume and the analyzer acceptance solid angle over the surface, is given by Equation (5):


Inl =

0

∞

ρ(x, y, z) × (dσ/dΩ)nl × exp[−z/(Λe × cosθ)] dxdydz

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(5)

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ρ(x, y, z) is the number of emitters for unit volume (atomic density), z is the photoelectron path
in the material measured along the surface normal, Λe is the IMFP, (dσ/dΩ)nl is the DPCS, and θ is
the take-off angle formed between the photoelectron propagation direction and the surface normal.
We need now to solve the integral reported in Equation (5) for core levels belonging to the interface
(in our case, the APTES overlayer) by systematically changing the photon energy, in order to find the
energy value that maximizes Inl . This is given by the trade-off between the increase in the probing
depth and the DPCS (TPCS) lowering at increasing photon energies. To this end, we have performed
a series of SESSA simulations keeping the detection geometry fixed (α = 15◦ , θ = 0◦ ) and varying
the photon energy of the incoming radiation (considered as 100% linearly polarized in the orbit
(horizontal) plane) with a step of 50 eV. Note that the simulated detection geometry is the same to
the one experimentally available at BL 9.3.1, ALS. The results of these simulations are reported in
Figure 5a,b for the Si and N 1s core level photoelectron intensity (left axes), respectively.

Figure 5. Si 1s (a) and N 1s (b) photoelectron intensity evolution as a function of the photon energy, with
and without the X-ray window placed between the X-ray source and the sample. The two elements
belong to the APTES overlayer buried by 20 nm of water. For a better comparison, the reported
photoelectron intensities with no simulated absorber are normalized by the maximum value obtained
at 5390eV for the Si 1s core level. The right axes of Figure a,b report the X-ray transmission through a
Si3 N4 and polyimide (Kapton) window. For both materials, the simulated window thickness is 500 nm,
which is a common value for X-ray windows used during AP-HAXPES experiments. To simulate the
effect of the X-ray window on the photoelectron response, the X-ray transmission curves have been
convoluted with the photoelectron intensity trends obtained in absence of an X-ray absorber.

The intensity of the Si 1s core level spectrum (BE = 1850 eV) initially increases for photon energy
values between 2500 and 4500 eV (Figure 5a). This is due to the fact that the increasing photoelectron
KE leads to an increase of the IMFP, which balances the decay of TPCS (since the photon energy moves
away from the Si KLL threshold, centered between 1800 and 1900 eV). As the photon energy further
increases, the exponential decay of the TPCS starts to dominate over the other terms in Equation (5),
with the intensity curve leveling off and eventually reaching a maximum at a photon energy of 5390 eV
(the intensity curve has been fitted using a lognormal function in order to accurately characterized the
maximum). At this energy, the TPCS is about e−2.0 the value at 2500 eV (σ0 = 0.062 Mbarn), whereas
the probed depth (defined as the depth from which 95% of the emitted electrons are inelastically
scattered, i.e., 3·λe ) is about 40 nm (that is two times the thickness of the water layer). For higher
energies, the TPCS dominates and a monotone decrease of the photoelectron intensity can be observed
in line with the exponential decay of the TPCS. The intensity trend of the N 1s core level spectrum
as a function of the photon energy (Figure 5b) follows a similar qualitative trend, however shifted in

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photon energy (photoelectron KE) due to the lower energy threshold of the KLL transition (at around
400 eV). The maximum of the intensity curve is reached at a photon energy of about 2850 eV, for which
the TPCS is equal to about e−3.0 the value at 1000 eV (σ0 = 0.071 Mbarn). The information depth at
2890 eV is equal to about 25 nm.
We can therefore rationalize the observed dependencies drawing some general conclusions:
1.

2.
3.

For core levels whose electronic excitation energies (and therefore BEs) fall within the soft X-ray
region (below 1000 eV), the TPCS decreases rapidly as the photon energy increases within the
hard X-ray region (above 2000 eV). For core levels characterized by higher electronic excitation
energies (approaching or within the hard X-ray region), the corresponding TPCS decreases with
a lower rate with the increase of the photon energy;
The IMFP follows the same functional dependency with the photon energy for both type of
core levels;
The combination of the two aforementioned points lead to the following phenomenology: the
maximum of the photoelectron intensity of different core levels belonging to the interface region
lies at different photon energies. For soft X-ray core levels, the trade-off between relatively
high KEs and the fast decay of TPCS leads to probed depths at the curve maximum essentially
matching the electrolyte overlayer thickness. For hard X-ray core levels, instead, the slower
decay of the TPCS with the photon energy leads to information depth at the maximum of the
photoelectron intensity curve higher than the water overlayer thickness. In addition, the intensity
curve for hard X-ray core levels is characterized by a broader spectral range compared to that for
soft X-ray spectra. In our case, the full-width at half-maximum (FWHM) of the Si 1s intensity
curve is about 6500 eV, whereas the same for the N 1s core level is characterized by a FWHM of
about 4150 eV.

Furthermore, we simulated the effect of introducing an X-ray absorber between the X-ray source
and the sample, which is typically done in order to seal the X-ray source (beamline or anode source)
from the high pressure environment. Figure 5 reports the X-ray transmission (right axes) through
a 500 nm-thick Si3 N4 and Kapton window (the calculations of the X-ray transmission as a function
of the photon energy (1000 eV–10000 eV) and fixed incidence angle α at the window (α = 90◦ ) have
been performed using the database of the Center for X-ray Optics (CXRO) of the Lawrence Berkeley
National Laboratory (Berkeley, U.S.A) [50]). We can observe that working with hard X-rays (photon
energy above 2000 eV) and Kapton windows keeps the X-ray transmission close to unity even for
relatively thick windows (X-ray transmission above 95 %). On the other hand, for Si3 N4 windows
the Si KLL absorption edge between 1800 and 1900 eV decreases the transmission above 2000 eV
(by about 30%, 20%, and 10% at a photon energy of 2000, 2500, and 3100 eV, respectively). The Si
and N 1s photoelectron intensities have been then convoluted with the simulated X-ray transmission.
We can observe that the Si 1s intensity trend (Figure 5a) is basically not influenced by the presence
of the window (either Si3 N4 or Kapton), with an intensity decrease at the maximum (hv = 5390 eV)
of about 2% in the case of a Si3 N4 absorber. On the other hand, for the soft X-ray N 1s core level
(Figure 5b), the presence of the Si3 N4 window leads to an appreciable loss of the photoelectron
intensity at the curve maximum (by about 15 % compared to the simulation performed without
absorber). Moreover, the maximum of the intensity curve is blue-shifted (~250 eV) by the presence of
the Si3 N4 window compared to the curve obtained with no absorber. In line with the findings obtained
for the Si 1s core level intensity curve, the presence of the 500 nm-thick Kapton window does not
induce a shift of the N 1s core level intensity curve, with a loss of signal at the maximum less than 2%
compared to the calculation performed without absorber. We want to highlight that the presence of a
differentially-pumped aperture (pin-hole) instead of an X-ray window will lead to the same results
obtained for no X-ray absorber, since the former implies a simple photon energy-constant lowering of
the photon flux at the sample.

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Overall, this analysis elucidates the importance of tuning the photon energy for different core
levels belonging to the interface region buried by a nanometric-thick electrolyte layer, in order to
keep the optimal core level “brightness” and enhance the investigation of the solid/liquid interface.
Moreover, it shows that another important parameter to keep in mind for enhancing the photoelectron
intensity is the choice of the X-ray window, in particular for soft X-ray core levels.
3.3. Evolution of the Physical/Chemical Properties of the Solid/Liquid Interface as a Function of the
Electrolyte pH
In this section, we will show a practical example of the potentials offered by coupling the “dip and
pull” method with AP-HAXPES. We investigated a sol-gel spin-coated TiO2 surface functionalized
with ~1 MLE of APTES, using a fixed photon energy of 4000 eV. This value is close to the average
(4120 eV) of the two energies retrieved above to enhance the signal intensity from Si and N 1s core
levels within the APTES overlayer. In addition, the photon flux at the sample at photon energies of
2850 eV and 5390 eV at BL 9.3.1 at the Advanced Light Source was not optimal due to experimental
limitations of the beamline optics. We studied the solid/liquid interface formation by investigating the
surface in its pristine conditions (high vacuum, HV, about 10−6 mbar) (i), by exposing it to ~70% RH at
room temperature (hydrated conditions, HC) (ii), and finally after dipping the surface in pure water
and different aqueous electrolytes (iii), changing the pH from 7 to 14 (see also Table 1). The dipping
was performed in about 70% RH environment at room temperature (pwater ~ 21 mbar).
It is important to note that the experimental BE values measured in this work for light elements
such as O and N are slightly higher than those reported in the literature for similar systems. This might
be due to recoil effects when momentum is transferred from the ejected photoelectron to the emitting
atom. Recoil is present in all photoemission processes and its effects are non-negligible for high
photoelectron KEs and light elements [51,52]. At a photon energy of 4000 eV the calibration performed
on gold using the Au 4f7/2 core level signal (KE ~ 3916 eV) is essentially not influenced by the recoil
(the corresponding loss is about 10 meV). On the other hand, the energy loss for the O 1s (KE ~ 3470 eV)
core level spectrum is about 120 meV.
To monitor the formation of the liquid electrolyte layer on the surface, we acquired the O 1s core
level spectrum at all the aforementioned conditions, as reported in Figure 6a. At pristine conditions,
the O 1s spectrum exhibits an intense peak centered at a BE of about 530.7 eV, attributable to the
TiO2 lattice oxygen (O2− ) [53,54]. A minor component can be observed at about 533.1 eV, most likely
due to molecularly adsorbed water (H2 Oads. ) on Ti4+ sites [53] and on the APTES terminal –NH2
groups upon exposure to air [55]. It is in fact suggested by Meroni et al. in a recent work that APTES
chemisorption on TiO2 mainly occurs through the formation of one or two Si−O−Ti bonds involving
the Si headgroup [56], whereas adsorption via the terminal amino group seems instead considerably
more labile. Therefore, the Lewis base-character of the amine group might induce water adsorption at
the nitrogen atom through hydrogen bond formation with water molecules.
At HC and for the dipping experiments in the different aqueous electrolytes, we can observe
the presence of two new spectral features in the O 1s core level spectrum. The first, centered at a BE
between 533.6 and 534.0 eV is attributable to liquid phase water (LPW) due to the formation of the
liquid layer, whereas the second peak is associated to the gas phase water (GPW) (in the BE range
of 536.3 and 536.6 eV) [53]. The shift in BE of these two components is typically associated to work
function changes at the sample surface [57] due to the different explored conditions.
In addition, passing from the pristine conditions to HC and for the dipping experiments in the
different aqueous electrolytes, it is possible to observe the development of a third feature centered at
a BE of about 531.8 eV. The origin of such a spectral component, under the particular conditions of
the experiment, is not trivial and it can be due to a number of different causes. First, the observed
BE is typical of surface hydroxyl groups generated from the well-known dissociative adsorption of
water for pressures above ~ 10−3 mbar [53]. Second, oxygen-containing carbon compounds from
background contamination can also produce a peak at this BE, particularly after filling the chamber

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with water vapor for several minutes at relatively high pressures. The contamination originates most
likely from environmental CO2 and hydrocarbons containing COx groups. This is typically caused
by the displacement of molecules from the chamber walls upon exposure to water vapor at or above
the mbar pressure range or from molecules contained in the water vapor source itself [53]. Third,
the formation of non-volatile carbonates during the preparation of the alkaline electrolytes in ambient
conditions (CO2 + 2 KOH → K2 CO3 + H2 O) can lead to the “deposition” of carbonates on the sample
surface during the “dip and pull” procedure. The reported BE for carbonate groups (531.9 eV [58])
is also consistent with our findings. The “deposition” of carbonates at the highest investigated pH
might be the main cause of the important intensity increase of the spectral component observed when
passing from pH 12 to pH 14.
We want now to focus the reader’s attention on the photoelectron intensity evolution of the LPW
component during the experiment. Figure 6b,c report the LPW to O2− ratio and the corresponding
water layer thickness estimation, respectively. To do so, SESSA simulations have been performed
using the same sample structure reported in Figure 1. By changing the water overlayer thickness,
the simulated LPW/O2− ratio was adjusted to match the experimentally-determined value for each
condition. Passing from the first dipping in pure water to the third in KOH 10−2 M (pH 12),
the LPW/O2− ratio changes from about 0.3 to 0.5, which corresponds to an electrolyte layer thickness
between 3 and 4 nm. Within the experimental uncertainty, these are similar thicknesses found at HC,
after exposing the sample to about 21 mbar of water (~ 2.5 nm). This means that a thin layer of water
condensed on the surface at HC (~ 70% RH), and that the thickness of the liquid layer did not depend
on the successive dipping procedures in pure water (pH 7) and in KOH 10−4 and 10−2 M aqueous
solutions (pH 10 and 12, respectively). Differently, at pH 14 (KOH 1.0 M) the LPW/O2− ratio increases
up to ~2.5 (Figure 6b), which corresponds to an electrolyte layer thickness of about 13 nm (Figure 6c).

Figure 6. (a) O 1s core level spectra and corresponding multi-peak fitting for the explored experimental
conditions (the spectra are normalized by the intensity of the M–O2− component); (b): ratio between
the liquid phase water (LPW) and the lattice oxygen (O2− ) component determined via the multi-peak
fitting procedure; (c): corresponding estimation of the liquid layer thickness, performed using the
intensity attenuation of the O2− component. The inset reported in Figure c shows the BE negative
shift of the O2− component throughout the experiment, pointing to an upward band bending in TiO2
as consequence of the deprotonation of the surface –OH groups. The experiment was conducted at
room temperature.
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To find the reason for the observed phenomenology, we use the Henderson–Hasselbalch equation
to estimate the ratio between the unprotonated (–NH2 ) and protonated (–NH3 + ) amine groups as a
function of the electrolyte pH (–NH2 /–NH3 + = 10pH-pKa ). This procedure allows us to qualitatively
assess the net charge at the surface under the different explored conditions, thereby inferring about
the interaction between the functionalized charged surface and the liquid water. Using a value of the
APTES amine group acid dissociation constant (pKa) of 10.6 as reported by Notsu et al. [59], we find a
–NH2 /–NH3 + ratio of 2.5 · 10−4 , 0.25, 25, and 2500 at pH 7, 10, 12, and 14, respectively. This means that
for pH values below (above) the pKa, the amine group is present on the surface mainly in its protonated
(unprotonated) form. This is qualitatively demonstrated by the N 1s core level spectra reported in
Figure 7a for the pristine conditions, hydrated conditions and after dipping the sample in the KOH
1.0 M solution (pH 14). The spectrum acquired on the pristine surface exhibits a clear asymmetry
toward high BEs, and can be fitted using two spectral components centered at a BE of 400.4 and 401.7 eV
(with a FWHM of 1.8 and 2.4 eV, respectively). Taking into account a positive 140 meV shift of the BE
due to the recoil loss, the identified BEs for these components are in line with previous studies [56] and
can be associated with carbamate (–NHC(=O)OH) and –NH3 + moieties present at the APTES overlayer,
respectively. Interestingly, within the detection limit of the technique (about 1 at.%), we do not observe
the presence of unprotonated amine groups (which should generate a peak at a BE of 399.6 eV [56]).
This might be due to the fact that the –NH2 groups readily reacted with environmental CO2 to form
carbamates upon exposing the sample to ambient conditions, after the APTES functionalization of
the TiO2 surface [60]. Exposing the sample to ~ 70% RH at room temperature leads to an increase of
the –NH3 + surface concentration (in agreement with the Henderson–Hasselbalch equation for pH 7),
accompanied by an upward BE shift of both spectral components (by about 0.3 eV). After dipping the
sample in the KOH 1.0 M solution (pH 14), it is possible to observe the disappearance of the protonated
amine component as predicted by the Henderson–Hasselbalch equation. Under these conditions,
the N 1s core level peak could be fitted using only one spectral component, characterized by a FWHM
of 2.2 eV and centered at a BE of about 400.0 eV. This value is about 0.4 eV lower than the BE reported
above for the carbamate moieties present on the pristine surface (400.4 eV). We will give a plausible
explanation for such a shift later at the end of this section.
We can then conclude that for pH values below the –NH2 pKa (7 and 10), the presence of the
protonated amine groups, although positively charged, prevents the formation of a stable and relatively
thick water layer leading to the observed partial non-wetting behavior. Above the –NH2 pKa (pH 12
and 14), no protonated –NH3 + groups are essentially present on the surface, while the majority of the
free (unprotonated) amine groups are converted into carbamate moieties by the nucleophilic CO2 or
CO3 2− attack (keep in mind that the O 1s revealed an enhanced CO3 2− photoelectron signal at pH
14, most likely due to the high availability within the liquid layer of non-volatile carbonates formed
during the preparation of the solution). In this case, the absence of formal charges at the surface
might weaken its interaction with water. On the other hand, at pH 14 we observe a different behavior,
in which the surface shows a higher hydrophilicity character. It is reported that the surface hydroxy
groups on TiO2 possess acidic character [61] and can therefore easily donate H (Brønsted acid) in strong
alkaline conditions [61,62]. Therefore, the deprotonation reaction occurring at the solid/liquid interface
(–OH + H2 O → –O– + H3 O+ ) at pH 14 generates a surface negative charge that might be responsible
for the stabilization of a thicker liquid layer compared to those observed at lower pH values.
The development of a negative surface charge driven by the pH increase leads also to a second
effect. In Figure 6a it is possible to qualitatively observe a negative BE shift of the O2− component.
The inset reported in Figure 6c shows the O2− BE trend as a function of the different explored conditions.
The downward BE shift is equal to 0.4 eV passing from pristine conditions to the final dip in pH 14, and
it is attributable to an upward band-bending in TiO2 (formation of a depletion layer at the solid/liquid
interface). This is confirmed by the observation of the same negative BE shift when acquiring the Ti
2p core level spectrum, as reported in Figure 7b for the same experimental conditions. Passing from
pristine to pH 14 conditions, the BE measured on the Ti 2p3/2 spectrum is equal to 459.3 and 458.9 eV,

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respectively. The observed negative 0.4 eV shift is thus in agreement with the previous findings. It is
worth noting that the photoelectrons generated from the ionization of the O 1s and Ti 2p3/2 core levels
have similar KEs (the difference being about 70 eV), which leads to the same information depth at
a photon energy of 4000 eV (~17 nm). This means that the information carried by the BE shift is
generated by the spectral summation within the depletion layer, convoluted with the exponential
decay of the photoelectron intensity. In addition, we want to highlight that no appreciable FWHM
change is observed on the lattice oxygen spectral component or in the Ti 2p3/2 spectrum, passing from
pristine conditions to pH 14. This might be due to the difference between the information depth (using
O 1s and Ti 2p3/2 core levels) and the thickness of the space charge region (SCR). Using a concentration
of donors of 1 × 1018 cm−3 for intrinsic n-doped TiO2 [63] and a band bending potential of 0.4 V as
determined experimentally (assuming flat bands at the pristine conditions), the space charge region
thickness can be estimated to be about 45 nm [64].
We can make use of the adsorption of the APTES molecules at the TiO2 top-most layer to probe
the electrical potential value at the upper part of the band-bending potential distribution, by using a
core level spectrum from an element belonging to the APTES overlayer (e.g., Si and N 1s). In the case
of the Si 1s spectrum (Figure 7c), the negative BE shift between the pristine and the pH 14 conditions
is considerably larger than that observed for O 1s and Ti 2p3/2 , being equal to ~1.1 eV. This value
represents therefore the maximum band bending in the depletion layer of TiO2 .
On the other hand, as already introduced above, for the N 1s core level spectrum we observe a
negative shift of the carbamate group BE of about 0.4 eV, passing from pristine to pH 14 conditions.
This could be generated by two causes. First, the nitrogen atom is not “directly” adsorbed at the TiO2
top-most layer, differently than the silicon atom, but it is screened by three alkyl units from the latter.
Therefore, the presence of four single covalent bonds (Si–C, C–C, C–C and C–N) between the silicon
and the nitrogen atom might lower the potential experienced in the latter. Second, as schematically
reported in Figure 7d, the nitrogen atom is dangling from the surface into the liquid layer, thereby
experiencing the electrical potential drop in solution. The Debye length k−1 for a monovalent ion
in an aqueous solution at room temperature is related to the concentration C of the electrolyte by
k−1 = 3.04 × C−1/2 (C is expressed in mol L−1 , or M, and k in Å). This means that for a 1.0 M solution,
the Debye length is equal to about 0.3 nm. Defining the thickness of the Gouy–Chapman diffuse layer
as three times the Debye length (i.e., at which distance the potential at the sample surface is ∼95%
screened in solution), we find that such a thickness is equal to 0.9 nm. We performed a molecular
dynamics simulation of the adsorption of one APTES molecule on a rutile TiO2 (110) surface through
the formation of one Ti–O–Si bond (Figure 7d). To further confirm this, we used an UFF force field
and a convergence cut off of 10−6 . The simulated distance between the nitrogen and the ideal plane
containing the-top most layer oxygen atoms is equal to about 0.7 nm, which is similar to the thickness
of the diffuse layer calculated above. This means that the nitrogen atom might experience the local
potential at the slip plane (that is, at the ideal plane between the diffuse layer and the bulk solution).
The two described effects might then account for the different BE shift measured on the N 1s core level
spectra with respect to the value measured using the Si 1s core level.

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Figure 7. N 1s (a), Ti 2p (b) and Si 1s (c) core level spectra acquired at the pristine conditions, hydrated
conditions (HC) and after dipping the sample in the KOH 1.0 M solution (pH 14). (d) Relaxed structure
obtained from the molecular dynamics simulation of the adsorption of one APTES molecule on a rutile
TiO2 (110) surface through the formation of one Ti–O–Si bond (white: titanium; light grey: hydrogen;
dark grey: carbon; light green: silicon; red: oxygen; blue: nitrogen).

4. Conclusions
In conclusion, we have described the “dip and pull” procedure for the generation of
nanometric-thick layers of liquid electrolyte on solid surfaces. The prepared solid/liquid interfaces
can be then investigated using AP-XPS. The use of photon energies within the hard X-ray range (above
2000 eV) allows to probe solid surfaces buried by liquid layers as thick as tens of nanometers, enabling
the extension of the parameter space and the possibility to study a wide variety of systems, electrolytes,
and electrolyte concentrations.
We discussed the advantages and current drawbacks of using the static “dip and pull” method and
related techniques. The main limitation of such techniques is the mass transport in the direction parallel
to the interface, which sets restrictions on the current densities achievable in the liquid layer. On the
other hand, the lowering of the reaction kinetics within the liquid layer opens up the possibility of
performing time-resolved studies on time scales accessible by AP-HAXPES. In addition, this technique
is appealing for the investigation of fundamental properties of the solid/liquid interface, such as
electrical double layer structure and its modulation by electrical potential and illumination transients,
specific ion adsorptions and charge transfers across the interface. This is demonstrated by the recent
development of AP-XPS and AP-HAXPES instruments utilizing the “dip and pull” method, at the
Advanced Light Source (U.S.A.), Paul Scherrer Institute (Switzerland), MAXIV (Sweden), and BESSY
II (Helmholtz-Zentrum Berlin, Germany).
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By using a model system of a solid/liquid organic/inorganic hybrid interface and numerical
simulations, we retrieved the optimal photon energy range that enhances the signal coming from the
interface when coupling the “dip and pull” method with AP-HAXPES. This is an important parameter
that needs to be tuned for each investigated system, due to the fact that the use of hard X-rays to
induce photoemission implies relatively low photoelectron intensities due to the low cross sections.
Finally, we presented a practical application of the technique, investigating the
APTES-functionalized TiO2 polycrystalline surface in KOH aqueous electrolytes at different
concentrations. Our experiments show that it is possible to monitor in situ the physical/chemical
changes of the solid/liquid interface induced by the variation of the electrolyte pH. Our findings open
up the possibility to optimize the solid/liquid interface in order to study simultaneously the potential
profile in semiconductors (band-bending) and in the liquid side of the junction (double + diffuse layer),
using local molecular probes directly functionalizing the semiconductor surface.
We think that the detailed molecular level comprehension of the solid/liquid interface
properties will advance materials and process research in many key-technological fields, such as
(photo)electrocatalysis, corrosion protection, water remediation, and CO2 capture and conversion.
Author Contributions: Conceptualization, M.F. and D.E.S.; Formal analysis, M.F.; Investigation, M.F., F.F.A. and
D.E.S.; Writing—original draft, M.F.; Writing—review & editing, M.F., F.F.A., E.J.C., Z.L., R.v.d.K. and D.E.S.
Funding: This work was supported by the German Federal Ministry of Education and Research (BMBF project
“JointLab—Grundlagen elektrochemischer Phasengrenzen”, GEP, #13XP5023C). This research was partially
conducted at the Advanced Light Source of the Lawrence Berkeley National Laboratory, a DOE Office of Science
User Facility under contract no. DE-AC02-05CH11231. Z.L. was supported by the National Natural Science
Foundation of China (11227902).
Conflicts of Interest: The authors declare no conflict of interest.

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Article

Oxygen Reduction Reaction on Polycrystalline
Platinum: On the Activity Enhancing Effect of
Polyvinylidene Difluoride
Alessandro Zana 1 , Gustav K. H. Wiberg 1,2 and Matthias Arenz 1, *
1
2

*

Department of Chemistry and Biochemistry, University of Bern, 3006 Bern, Switzerland;
alessandro.zana@dcb.unibe.ch (A.Z.); gustav.wiberg@gmail.com (G.K.H.W.)
Department of Physical Sciences, Harold Washington College, Chicago, IL 60601, USA
Correspondence: matthias.arenz@dcb.unibe.ch; Tel. +41-31631-5384

Received: 20 December 2018; Accepted: 21 January 2019; Published: 24 January 2019

Abstract: There have been several reports concerning the performance improving properties of
additives, such as polyvinylidene difluoride (PVDF), to the membrane or electrocatalyst layer of
proton exchange membrane fuel cells (PEMFC). However, it is not clear if the observed performance
enhancement is due to kinetic, mass transport, or anion blocking effects of the PVDF. In a previous
investigation using a thin-film rotating disk electrode (RDE) approach (of decreased complexity as
compared to membrane electrode assembly (MEA) tests), a performance increase for the oxygen
reduction reaction (ORR) could be confirmed. However, even in RDE measurements, reactant mass
transport in the catalyst layer cannot be neglected. Therefore, in the present study, the influence of
PVDF is re-examined by coating polycrystalline bulk Pt electrodes by PVDF and measuring ORR
activity. The results on polycrystalline bulk Pt indicate that the effects of PVDF on the reaction kinetics
and anion adsorption are limited, and that the observed performance increase on high surface area
Pt/C most likely is due to an erroneous estimation of the electrochemical active surface area (ECSA)
from CO stripping and Hupd .
Keywords: ORR; Platinum; PVDF; PEMFC

1. Introduction
Efficient energy converters, such as the proton exchange membrane fuel cell (PEMFC), are integral
parts in an environmentally friendly hydrogen economy. Although the PEMFC technology has been
around for several decades, it still suffers from relatively high costs [1]. With the aim of making the
fuel cell technology more economical, significant effort has been invested in finding cheaper catalysts
and fabrication techniques, but also in optimizing properties of the solid membrane electrolyte [2–6].
There are several reports stating that the properties of the solid membrane electrolyte can be improved
by including different additives. For example, a performance enhancement of PEMFCs by including
additives such as polyvinylidene difluoride (PVDF) into the membrane has been reported [7].
If such studies are performed in single cell membrane electrode assembly (MEA) tests, it is
very difficult to analyze the origin of observed performance increases. This is due to the complexity
of the MEA system and the influence of multiple structural aspects on its performance. Therefore,
the effect of PVDF additions to the membrane is not yet fully clarified. In PEMFC research, a thin-film
rotating disk electrode (RDE) approach is often applied to reduce the system complexity and to study
electrocatalyst layers in a more fundamental fashion [8]. The previous work of our research group
used such an approach to mix PVDF into the ink of a high surface area catalyst. The electrochemical
properties of catalyst layers prepared by casting the ink on a glassy carbon disk were compared with
respect to PVDF free films [9]. The specific oxygen reduction reaction (ORR) activity established by
Surfaces 2019, 2, 69–76; doi:10.3390/surfaces2010007

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RDE in phosphoric acid electrolytes was found to improve when the catalyst layer contained PVDF.
Further work, however, showed that even in a RDE setup, the apparent (measured) specific ORR
activity is extremely sensitive to the film quality, which in turn depends on ink properties, such as
pH [10]. Thus, the observed activity increase can be due to intrinsic changes, such as oxygen solubility
and suppression of anion adsorption, or it might be due to improved mesoscopic film properties
enhancing the local reactant mass transport. To investigate the possible catalytic enhancement due to
PVDF further, in the presented work we therefore investigated thin PVDF films on a flat surface of
polycrystalline bulk Pt. This system excludes any influence of catalyst film properties, and therefore
focusses on the interface between catalyst and electrolyte.
2. Experiment
2.1. Sample Preparation
To prepare PVDF films onto a Pt surface of an RDE electrode, two different amounts of PVDF
(Polysciences Inc. Warrington, PA, USA) were dissolved in Tetrahydrofuran (THF) to obtain two
solutions containing 0.1 mM and 0.1 M of PVDF, respectively. The electrode tip for the PVDF coating
was a polycrystalline Pt disk inserted in a cylinder of polytetrafluoroethylene (PTFE), in such a manner
that only a circular area of the platinum is exposed. The Pt tip was sonicated in acetone, ethanol,
isopropanol, and MilliQ water, repeating each step three times. Following the sonication, the Pt tip
was dip coated into one of the two PVDF solutions for 30 seconds and pulled out of the electrolyte
with the surface facing down. The PDVF coated Pt electrode tip was left to dry in air in a closed glass
beaker. After drying, the mirror-like shine of the bare Pt surface was lost and the surface exhibited
a pale milky color. In the following, PVDF 1 will refer to the polycrystalline Pt surface dip coated in
the solution containing 0.1 mM PVDF, and PVDF 2 will refer to the Pt surface dip coated in the PVDF
solution containing 0.1 M PVDF. After the drying process, the coated Pt electrode was introduced in
an Ar-saturated 0.05 M H2 SO4 electrolyte solution for electrochemical characterization.
2.2. Electrochemical Characterization
The electrochemical characterization was performed by RDE according to the procedure described
in [11,12]. In brief, the electrochemical measurements were conducted in a three- compartment Teflon
cell, using an RDE system (Radiometer Analytical, Europe Hach Lange GmbH Headquarter Willstätterstr.
11, D-40549 Düsseldorf, Germany). A saturated calomel electrode (SCE) and a Pt mesh have been used
as reference and counter electrodes, respectively. The working electrode (WE) was a polycrystalline
Pt disk (Ø = 5 mm, A = 0.196 cm2 ). All the potentials in the paper are expressed with respect to
the reversible hydrogen electrode (RHE) potential experimentally determined for each measurement
series. The 0.05 M H2 SO4 aqueous electrolyte solution was prepared using Millipore Milli Q water
(> 18.2 MΩcm, TOC < 5ppb) and 96% H2 SO4 Suprapur (Sigma-Aldrich). The oxygen reduction rate
and specific activity (SA) are typically calculated from the positive-going polarization curves recorded
in oxygen saturated electrolytes at a scan rate of 50 mVs−1 and 1600 rpm rotation rate. In all cases,
a background subtraction was performed to remove any contributions from capacitive and surface
oxidation processes according to [13]. The iR-drop was measured and compensated according to the
methodology used in [13,14]. In brief, the solution resistance was recorded online with the potentiostat
(ECi-200 Nordic Electrochemistry Aps c/o Skelvangsvej 43, DK-8920 Randers NV, Denmark) by
superimposing a 5 kHz, 5 mV AC signal, and compensated for by an analogue positive feedback
scheme. The effective solution resistance was around 2Ω [15]. The electrochemical active surface area
(ECSA) was determined by Hupd integration and CO stripping experiments described in [12,13]. In all
the measurements the potential was cycled in the potential window between 0.05 and 1.05 VRHE .

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3. Results and Discussion
3.1. Surface Characterization in Ar Saturated Electrolyte
In Figure 1, the cyclic voltammograms (CVs) of the bare Pt surface and the two different PVDF
films on polycrystalline Pt recorded in Ar saturated electrolyte solution are shown. The two different
dip coatings were made from solutions containing 0.1 M and 0.1 mM of PVDF solution, respectively.
The voltammetric profiles of polycrystalline Pt are generally divided into three electrode potential
regions: (i) the hydrogen under potential deposition (Hupd ) region (0.05–0.3 VRHE ), (ii) the double
layer region (0.3–0.7 VRHE ), and (iii) the potential region of OHad and Oad formation and reduction
(>0.7 VRHE ) [16]. The effect of PVDF can be scrutinized by comparing the different CVs in the respective
potential regions. Starting with the Hupd region, two characteristic peaks of polycrystalline Pt in
contact with a H2 SO4 electrolyte solution are present for all samples. However, the Hupd currents
decrease according to the amount of deposited PVDF (concentration of PVDF in the dipping solution).
The formation and reduction oxygenated species exhibit the same trend, i.e., the intensity of the peaks
decreases as more PVDF is deposited. In contrast, the double layer potential region remains similar for
all the samples. Only a small increase in the pseudo-capacity is discernible for the Pt surfaces covered
by PVDF, which might be interpreted by a minor surface roughening or by a small increase in the
concentration of water dipoles close to the electrode surface. However, at the same time, the presence
of PVDF clearly reduces the amount of Hupd and OHad species adsorbed onto the Pt surface. Further,
the analysis of the Hupd peaks is a common method to determine the ECSA used for normalizing the
measured activity of a Pt electrode. Integration of the Hupd peaks and the resulting charge, Q Hupd ,
can be related to the charge of a freshly flame annealed polycrystalline Pt surface, which has a Q Hupd
per area unit 210 μC cm−2 [17]. From this analysis, we can determine the loss in ECSA due to the
PVDF coverage to 27 % and 43 % for PVDF 1 and PVDF 2, respectively.

Figure 1. Cyclic voltammograms recorded in Ar saturated 0.05 M H2 SO4 electrolyte solution for
bare and polyvinylidene difluoride (PVDF) modified Pt. The scan rate was 100 mVs−1 and the
measurements were recorded at room temperature. PVDF 1 stands for the film prepared by dip coating
the Pt electrode into 0.1 mM PVDF solution, PVDF 2 for the film prepared by dip coating into 0.1 M
PVDF solution. The roughness factors (Rf), calculated by Hupd for bare Pt, PVDF1, and PVDF2, are:
1.65, 1.20, 0.94. This corresponds to an ECSA loss with respect to the bare polycrystalline Pt electrode
of 27 % and 43 % for PVDF 1 and PVDF 2, respectively.

3.2. CO-Stripping Measurements
CO-stripping is an alternative technique to determine the ECSA [11]. A monolayer of CO is
formed on Pt from a CO saturated electrolyte solution, keeping the electrode at a low fixed potential
(typically around 50 mVRHE ). Thereafter, the dissolved CO is displaced from the electrolyte solution by
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saturating it with Ar, and the CO adlayer is oxidized in a positive-going voltage scan. As CO oxidation
is a two electron process, the surface area ACO , relating to the CO oxidation charge (QCO ), is calculated
by [17]:
QCO
(1)
ACO =
2 × 196 μC cm−2
In Figure 2, we show the CO stripping profiles for the bare surface and the two PVDF coated
polycrystalline Pt surfaces. At the start of the stripping sweep starting from 0.05 VRHE , the CO covered
Pt electrodes show no features, such as the Hupd , until the characteristic CO oxidation peak at around
0.8VRHE is observed [18]. The presence of the PVDF coating clearly reduces the amount of stripping
charge and at the same time the oxidation onset, as well as the peak position of the maximum being
shifted towards higher potentials. Calculating the ECSA loss from QCO , we determine losses of 17 %
and 64%, respectively (with respect to the roughness factor of the bare Pt surface). This means that the
ECSA loss calculated by Hupd and CO-oxidation led to comparable values, as seen above.

Figure 2. CO stripping curves recorded for polycrystalline Pt and the two PVDF modified surfaces
at room temperature in Ar saturated 0.05 M H2 SO4 electrolyte solution. The scan rate was 50 mVs−1 .
PVDF 1 stands for the film prepared by dip coating the Pt electrode into 0.1 mM PVDF solution, PVDF
2 for the film prepared by dip coating into 0.1 M PVDF solution. The roughness factors (Rf) calculated
by CO stripping for the bare Pt surface, PVDF1 and PVDF2 are: 1.40, 1.17, 0.51. This corresponds to
an ECSA loss with respect to the bare polycrystalline Pt electrode of 17 % and 64 % for PVDF 1 and
PVDF 2, respectively.

3.3. Oxygen Reduction
In order to determine the ORR activity, polarization curves of the bare and PVDF coated Pt
surfaces were recorded in an oxygen saturated 0.05 M H2 SO4 electrolyte solution, applying different
rotations rates. The resulting background corrected data and their analysis are summarized in Figure 3.
In Figure 3a, the ORR polarization curves of the three samples are separated into positive- and
negative-going curves and compared with each other. Based on the current profiles, two distinct
potential regions can be observed: the kinetic potential region at an electrode potential E > 0.8 VRHE ,
and the hydro dynamic mass transport limited potential region at E < 0.6 VRHE . The latter is easily
identified by its strong rotation rate dependency (not shown). The influence of the thick PVDF coating
(PVDF 2) can clearly be seen in the mass transport limited potential region, as well as in the kinetic
region by a significant decrease in current. Furthermore, all curves exhibit a hysteresis, i.e., a difference
in currents between the positive- and negative-going potential scan is observed. This hysteresis is
a typical feature in RDE measurements and can be interpreted due to the fact that the steady state
coverage of oxygenated species is not instantaneously achieved on Pt during potential sweeps when
coming from the reduced or oxidized surface (lower or higher potentials). Thus, the observed ORR
rate is artificially enhanced [19,20] in the positive-going polarization curves, whereas it may be too low
in the negative-going polarization curves. This is explicitly shown in Figure 3b in a logarithmic Tafel
plot of the kinetic current densities normalized to the geometric electrode area, i.e., JK . Note that the
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curves have been corrected for diffusion effects by the use of the Koutecký-Levich equation. In the
positive-going potential sweeps, significantly higher kinetic ORR rates are determined than in the
negative going potential sweeps. However, in the Tafel plot, the curves of the different kinetic currents
form almost parallel lines, indicating that the same ORR reaction pathway is valid for all investigated
surfaces. Furthermore, while larger amounts of PVDF inhibit the observed kinetic ORR current,
the kinetic ORR currents on the bare Pt surface and PVDF 1 almost overlap. This means that although
the ECSA on PVDF 1 is reduced, and the kinetic current is not affected.

Figure 3. (a) Oxygen reduction reaction (ORR) polarization curves recorded for bare polycrystalline Pt
and PVDF modified Pt surfaces. Dotted and solid lines indicate the positive- and negative-going
potential sweeps, respectively. The scan rate was 50 mVs−1 and the rotation rate 1600 rpm.
The electrolyte was 0.05 M H2 SO4 aqueous solution. (b) Tafel plot of the kinetic current.
(c) Koutecky-Levich plot of the currents at 0.35 VRHE .

As we consider the electrode during the negative-going potential sweep to be closer to the steady
state condition, we chose the negative going potential sweeps for further analysis. In this analysis,
we studied the hydrodynamic response of the system by changing of the convection (rotation rate of
1
the RDE). In Figure 3c, the inverse of the measured current density is plotted versus ω 2 , in a so-called
Koutecký-Levich plot based on:
−1
−1
Jm
= Jk−1 + JDL
= Jk−1 + ( BL (ω ))−1

(2)

where Jm is the measured current density, Jk the kinetic current density, JDL the diffusion limited
current density, BL the Levich B factor, and ω the rotation rate.
It can be seen that the data points in the Koutecký-Levich plot form parallel lines, which is a
clear indication that the hydrodynamic properties and the number of electrons transferred per oxygen
molecule are not affected by the PVDF [21]. Thus, we conclude that the PVDF coating does not favor
the formation of H2 O2 . The negative shift of the polarization curves of PVDF 2 can instead be related
to a reduced interfacial current. The reduced interfacial current can be the result of a 3D PVDF film
that reduces the diffusion to the electrode, or that a large part of the surface is blocked for the reaction
by the PVDF, or a combination of both [22].
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So far, we have only analyzed the kinetic current densities normalized to the geometric area, as
is often done when working with bulk metal electrodes. To obtain catalytic turnover rates, or the
equivalent in electrocatalysis, the (surface area) specific activity (SA), the kinetic current needs to be
normalized to the ECSA of the catalyst. The determination of the ECSA of a catalyst is an essential task
that can be challenging, in particular when working with catalysts for which no defined procedure
is yet established [23]. Even though this is not the case for Pt based catalysts, where Hupd and CO
stripping are standard procedures for the ECSA determination, the ECSA determination is not trivial.
The procedure implies that the catalytic reactions with H+ and CO and the ORR probe the same type
of active site on the catalyst surface. In addition, their accuracy is limited. In particular, the charge
related to the Hupd process is sensitive to the lower potential limit of the CV, where Hupd overlaps
with the hydrogen evolution reaction. Furthermore, it has to be mentioned that PVDF retains a dipole
moment that could interact with polar molecules or protons, and thus alter the adsorption process on
a PVDF coated surface. On the contrary, O2 does not retain any dipole moment, and low amounts of
PVDF are not expected to affect the ORR process. Deviations in the ECSA between determination by
Hupd and CO stripping, as shown here, are therefore not surprising, and can be considered within the
accuracy of the method (usually around 10–15 %).
In Figure 4, comparing the JK with the normalized SA (normalized with the CO stripping charge)
for Pt, PVDF 1, and PVDF 2, a complex behavior becomes apparent. In the positive-going potential
scan (which is in general analyzed in RDE measurements), JK decreases with the PVDF content.
However, this decrease is “overcompensated” by the loss in ECSA. This leads to a somewhat higher SA,
especially for PVDF 2. However, in the negative-going potential scan—which resembles most steady
state conditions—an increase in SA is observed for PVDF 1, but almost no increase in SA is observed
for PVDF 2. Taking all results into account, these findings indicate that the ORR kinetics are neither
significantly inhibited nor enhanced by the PVDF films. Improvement in SA seems mostly related to
the fact that the surface area determination is “more affected” by the PVDF than the ORR itself, which
might be related to the fact that for the ECSA determination, different probe molecules are used.

Figure 4. Summary of results. Specific activity (SA) values at 0.9 VRHE for bare and PVDF modified
Pt in the positive and negative potential scan, respectively. Relative (normalized to bare Pt) current
density JK ; relative ECSA based on Hupd and CO stripping charge; relative limiting current density Jlim
for bare and PVDF modified Pt.

It was not an arbitrary choice to use H2 SO4 as a supporting electrolyte. H2 SO4 is known
to be a “strongly adsorbing” electrolyte, and decreased ORR kinetics in H2 SO4 as compared to a
“non-absorbing” electrolyte, i.e. HClO4 , are observed. This is attributed to the site blocking nature of
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the HSO4 − and SO4 2− anions. In literature reporting a positive influence of PVDF on the ORR activity,
it was speculated that PVDF might positively influence (inhibit) anion absorption. However, in the
presented work we never observed an increase in Jk for a PVDF coated Pt surface, and thus conclude
that there is no positive influence of PVDF on anion absorption, as a diminished anion absorption
would result in a decreased overpotential of the ORR. The activity enhancing effect of PVDF that has
been previously observed for high surface area catalysts (i.e. Pt nanoparticles supported on a carbon
support (Pt/C)), therefore, is most likely due to different reasons—a combination of reduced active
surface area and improved film properties that has been seen to strongly influence the apparent ORR
activity of Pt/C in RDE measurements [10].
Lastly, it is seen that the diffusion limited current Jlim is affected by the amount of PVDF used to
cast the Pt surface. This could be the result of two different processes: a change in the reaction pathway
of the ORR, or the formation of a 3D diffusion barrier. As the Levich analysis shows a straight line
with constant B values (Fig. 3) for all the samples, a change in reaction pathway has to be discarded.
We therefore conclude that the main reason for the reduction of Jk is the formation of a permeable 3D
layer, which partially blocks the active catalyst surface.
4. Conclusions
We have examined ORR on polycrystalline Pt with different amounts of PVDF deposited.
The experimental results indicate that PVDF forms a 3D film which disturbs the mass transport
of oxygen to the Pt surface. The PVDF coating also exhibits an active site-blocking effect, as both
the ECSA and the apparent kinetic activity are reduced for thick PVDF layers. However, for a thin
PVDF layer, the kinetic region overlaps with the one measured for bare Pt, even though the ECSA is
decreased. Thus, the SA exhibits a slight increase. However, this increase in SA is associated with the
loss in ECSA rather than a real kinetic enhancement effect. Depending on the scan direction, the same
can be observed for thick PVDF films. Here, the kinetic current is reduced, but this reduction is
overcompensated by the ECSA loss. Thus, we conclude that PVDF does not affect the ORR kinetics in
0.05 M H2 SO4 . The reduction in ECSA with even low PVDF coverage might be related to a difference
between the interaction of the probe molecules (H+ and CO) with the PVDF covered surface and the
reactant. Thus, PVDF used in small amounts inhibits the adsorption of CO and H+ without strongly
blocking the ORR kinetics. This leads to the “artificial” increase in SA. Therefore, we can conclude that
the addition of PVDF to Pt bulk and Pt-based catalysts does not affect the kinetics of ORR. Previously
reported performance enhancing properties might be due the reduced ECSA, as well as other factors,
like improved film quality.
Author Contributions: Conceptualization, A.Z. and G.K.H.W.; Data curation, A.Z.; Formal analysis, A.Z. and
G.K.H.W.; Supervision, M.A.; Writing—original draft, A.Z. and G.K.H.W.; Writing—review & editing, M.A.
Conflicts of Interest: The authors declare no conflict of interest.

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Article

Surface Investigation on Electrochemically Deposited
Lead on Gold
Alicja Szczepanska 1,2 , Gary Wan 1,2 , Mattia Cattelan 3, *, Neil A. Fox 3 and Natasa Vasiljevic 1,2
1
2
3

*

School of Physics, HH Wills Physics Laboratory, University of Bristol, Tyndall Avenue, Bristol BS8 1TL, UK;
alicja.szczepanska@bristol.ac.uk (A.S.); gary.wan@bristol.ac.uk (G.W.); N.Vasiljevic@bristol.ac.uk (N.V.)
Bristol Centre for Functional Nanomaterials, University of Bristol, Tyndall Avenue, Bristol BS8 1TL, UK
School of Chemistry, Cantocks Close, Bristol BS8 1TS, UK; Neil.Fox@bristol.ac.uk
Correspondence: mattia.cattelan@bristol.ac.uk; Tel.: +44-117-394004

Received: 17 December 2018; Accepted: 14 January 2019; Published: 17 January 2019

Abstract: Electrodeposition of Pb on Au has been of interest for the variety of surface phenomena
such as the UnderPotential Deposition (UPD) and surface alloying. Here, we examined the interface
between the electrodeposited Pb film on Au, using surface sensitive techniques such as X-ray
Photoelectron Spectroscopy (XPS), Ultraviolet Photoelectron Spectroscopy (UPS), Energy-Filtered
Photoemission Electron Microscopy (EF-PEEM) and Work Function (WF) mapping. The initially
electrodeposited Pb overlayer (~4 ML equivalent thickness) was transferred from the electrochemical
cell to the UHV system. The deposited Pb layer was subjected to Argon sputtering cycles to remove
oxide formed due to air exposure and gradually thinned down to a monolayer level. Surface science
acquisitions showed the existence of a mixed oxide/metallic Pb overlayer at the monolayer level that
transformed to a metallic Pb upon high temperature annealing (380 ◦ C for 1 h) and measured changes
of the electronic interaction that can be explained by Pb/Au surface alloy formation. The results show
the electronic interaction between metallic Pb and Au is different from the interaction of Au with the
PbO and Pb/PbO mixed layer; the oxide interface is less strained so the surface stress driven mixing
between Au is not favored. The work illustrates applications of highly surface sensitive methods in
the characterization of the surface alloy systems that can be extended to other complex and ultrathin
mixed-metallic systems (designed or spontaneously formed).
Keywords: Lead OPD; surface alloy; XPS; UPS; EF-PEEM

1. Introduction
Electrodeposition of Pb on Au is a well-known system characterized by Stranski-Krastanov
growth [1]. The electrodeposition of Pb on Au starts by the formation of the epitaxial monolayer
positive from the equilibrium potential for bulk deposition known as Underpotential Deposition
(UPD), followed by the nucleation of 3D islands in the OverPotential Deposition (OPD) region [2,3].
The UPD of Pb Monolayer (ML) on Au (111) has been considered as one of the model UPD systems
and has been widely studied in the past few decades with a range of electrochemical and surface science
techniques [2,4–11]. An incommensurate, electrocompressed closely packed hexagonal structure of
a complete Pb UPD ML has been analyzed by a range of techniques including Scanning Tunneling
Microscopy (STM), Atomic Force Microscopy (AFM), ex situ Low Energy Electron Diffraction (LEED)
and in situ x-ray diffraction [8,12–16].
The UPD of Pb on Au has been used in a wide range of applications including measurements
of surface area/roughness, surface structure characterization, chemical composition measurements,
electrocatalysis, sensing as well as a mediation of the epitaxial thin film growth. For example, Pb UPD
on Au has been used for the determination of surface areas, surface structures and crystallographic
orientation of nanostructures such as gold nanorods [17], controlled-shape nanoparticles [17,18] and
Surfaces 2019, 2, 56–68; doi:10.3390/surfaces2010006

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nanoporous gold electrodes [19]. The Au surfaces modified by a Pb UPD layer have shown to
enhance the electrocatalytic activity and selectivity of oxygen reduction reactions [5,20,21], methanol
oxidation [22] and glucose oxidation reactions [23,24]. Moreover, the Pb UPD layers have been
successfully employed as mediators and manipulators of thin film growth kinetics in electrochemical
methods such as the Surfactant-Mediated Growth (SMG) [25,26] and Defect Mediated Growth
(DMG) [27,28] of epitaxial films of Ag, Co and Cu. In the recently developed Surface Limited Redox
Replacement (SLRR) method, the UPD monolayer of Pb has been exploited as a sacrificial layer for
the deposition of 2D epitaxial noble metal films, such as Pt, Ag, Au [26,29–36], nano-alloys [37–40]
and nanoporous fuel-cell electrodes [41]. The SLRR method led to the development of a special
class of highly active Pt-ML electrocatalysts for oxygen reduction reaction [42] and has become a
tool for the design of model systems for fundamental electrocatalytic studies of structure-activity
relationship [38] and catalytic stability [43]. The most recent advances include different ideas and
concepts for a so-called electroless (e-less) deposition of Pb monolayer that were integrated into a
deposition method called “electroless Atomic Layer Deposition” (e-less ALD) [44].
Pb on Au system is characterized by a large atomic size mismatch of 21%. Thus, it is an ideal
model for studies of the surface-stress evolution and the related phenomena. This is an immiscible
system, i.e., Pb and Au do not form bulk alloys. However, surface alloying has been reported in
both electrochemical [6,7,45–48] and ultrahigh vacuum [49,50] environments. Several early studies
argued that there was no evidence of surface alloy formation in this system [2,8,14,15]. However, later
investigations, by in situ STM and surface stress measurements, showed the surface transformations
that can be associated with the surface alloy formation/dissolution [6,7,46–48]. The in situ STM studies
showed roughening of the initially flat Au substrate (by vacancy and ad-islands formation) after a Pb
UPD layer dissolution [6,7,45,48]. The compressive stress evolution during monolayer formation of
the magnitude of 1.2 N/m [48] featured a very narrow tensile drop within the main Pb UPD peak,
corresponding to the formation of a high coverage Pb layer [47,48].
While most electrodeposition studies have been focused on Pb UPD, the bulk or so-called OPD
for this system has not been studied so extensively [48]. The morphology of Pb OPD on Au studied
by a combination of in situ STM [48] and ex situ AFM [3], have shown dependence on the deposition
potential and specific ion adsorption (i.e. nature of ions in the deposition solution [3]). In the ionic
liquid solution, the nucleation process of Pb leads to the formation of triangular-shaped vacancies on
the Au (111), followed by the creation of large, well-defined 3D crystals of Pb [51]. The compressive
stress evolution during 3D island creation as a result from isolated Pb island nucleation and growth
has been studied in combination with the in situ Oblique-Incidence Reflectivity Difference (OIRD) [48].
The most intriguing was the observation of the surface alloy formation during OPD deposition that
seems to be surface stress related and localized around the edge areas of 3D Pb crystals [48].
To better understand the phenomenon of surface alloy formation of Pb ML on Au, use of
state-of-the-art characterization methods is required. In our work, we present results from X-ray
Photoelectron Spectroscopy (XPS), Ultraviolet Photoelectron Spectroscopy (UPS), Energy-Filtered
Photoemission Electron Microscopy (EF-PEEM) and Work Function (WF) mapping. XPS is a technique
generally used to analyze the chemical composition of both bulk samples and surface alloys [52–55].
The UPS is a method that is highly surface sensitive and has been used to characterize a broad range of
surface alloys, such as Au/Ni, Co/Cu(001) [52], Pt-bimetallic systems such as Ptx M (M=Ni, Co, Fe,
Ti, V) [53,56,57]. EF-PEEM and WF mapping are methods that can provide information about surface
orientation and segregation [55,58,59]; also WF provides the information about surface activity [54].
These techniques are particularly useful due to their high surface sensitivity, ability to detect electronic
changes from chemical environments at the Pb/Au interface and imaging capabilities at the nanoscale.
Studying Pb UPD on Au surface alloying is difficult as the UPD layer stability without the
potential control can be compromised by exposure to air/oxygen [60]. Also, the bulk electrodeposited
thicker Pb films inevitably oxidize in the air (i.e. forming of oxides). Therefore, the approach used
in our work to study Pb ML level films was based on the deposition of bulk (OPD deposition) of

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equivalent 4 ML thickness films. Upon transfer from the electrochemical cell, PbO was removed
gradually from the surface by cycles of Ar-sputtering and gentle annealing until the layer was thinned
down to the desired ML level. Using a combination of XPS, UPS, EF-PEEM and WF mapping we then
explored the structure of the layer.
2. Materials and Methods
Electrochemical measurements were performed in a standard three-electrode cell with a BioLogic
VSP potentiostat (Bio-Logic Science Instruments SAS, Seyssinet-Pariset, France) and controlled by
built-in software.
A solution containing 1 mM Pb(ClO4 )2 and 100mM HClO4 was prepared by mixing high purity
chemicals (PbCO3 , 99.999%, Alfa Aesar, Heysham, UK and 70% HClO4 , 99.999%, Sigma Aldrich,
Steinheim, Germany). Before the experiment, the solution was deaerated with nitrogen gas (Oxygen
Free Nitrogen, CAS: 7727-37-9, BOC, Manchester, UK) for 30 min. The oxygen free environment was
maintained during the experiment with nitrogen flow above the solution.
Pb and Pt wires were used as pseudo Reference (RE) and Counter (CE) Electrodes respectively.
Prior to experiments, the electrodes were prepared by cleaning in HNO3 (≥65%, Sigma Aldrich,
Steinheim, Germany) for 10 s, rinsed with Milli-Q water (Millipore/Merck, Darmstadt, Germany) and
dried with nitrogen. The Pt wire was then flame annealed by a butane torch. All the potentials will be
reported with respect to the Pb pseudo RE (0.64 V vs Standard Hydrogen Electrode)
A vacuum evaporated Au thin film (250 nm) deposited on a glass slide (with a 4 nm Ti adhesion
layer) was used as a Working Electrode (WE). Prior to experiments, the Au sample was cleaned with
H2 SO4 (95.0–97.0%, Sigma Aldrich, Steinheim, Germany) for 10 min, rinsed with Milli-Q water and
dried with nitrogen. To obtain the Au (111) orientation, the sample was flame annealed for 2 min using
a butane torch and rinsed with Milli-Q water.
The sample was then immersed in the deaerated solution with the half of its exposed surface
area above the solution. The part of the substrate remained deposit-free, for the sake of comparison
between the Pb covered and bare substrate. The deposition of Pb film was done by Linear Sweep
Voltammetry (LSV) with a scan rate of 10 mV/s starting from the bare Au surface at a potential of
0.9 V to −0.1 V. The LSV potential range covered the Pb UPD ML deposition (between 0.9 V and 0.0 V)
followed by the OPD growth (from 0.0V to −0.1 V). The potential of −0.1 V was then held constant for
20 s allowing for further (excess) overpotential Pb deposition. The sample was then removed from the
cell, rinsed with Milli-Q water and dried with nitrogen gas. The equivalent thickness of the deposited
Pb film of 4 ML was calculated by integration of the deposition current normalized by a charge of
305 μC/cm2 corresponding to an ‘ideal’ ML of Pb (111).
The prepared sample was immediately transferred into the Bristol NanoESCA Facility’s Ultra
High Vacuum (UHV) chambers and heated overnight at 150 ◦ C to clean the surface from physisorbed
contaminants without desorbing and altering the deposited Pb. Other surface cleaning procedures
involved 0.5 kV argon sputtering (at 45◦ to the sample normal) for 5 min and 2 min 30 s (for a total of
7 min 30 s), and a final 1 h anneal at 380 ◦ C. All the surface analysis techniques were performed on both
the Au and Pb deposited sides of the sample before proceeding to the next procedure. Throughout the
experiments, the sample was kept in UHV chambers.
High resolution XPS was carried out using a monochromatic Al kα (hν = 1486.7 eV) excitation
source, with the analyzer at 45◦ to the sample normal. A Pass Energy (PE) of 6 eV was used for detailed
photoemission lines, giving an overall energy resolution of about 300 meV obtained by fitting the
Fermi edge of the polycrystalline Au sample using a sigmoid function; PE of 50 eV was used for
surveys. Angle resolved XPS acquisitions were performed using a non-monochromatic Al kα source
and a PE of 50 eV.
The NanoESCA II EF-PEEM was used for WF mapping employing a non-monochromatic Hg
source (hν ≈ 5.2 eV) with an overall energy resolution of about 140 meV and lateral resolution of about
150 nm. UPS was performed with monochromatic He-I (hν = 21.2 eV) and He-II (hν = 40.8 eV) with an

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overall energy resolution of about 140 meV. EF-PEEM was also used for chemical imaging on Pb 5d
photoemission lines using a monochromated He-II (hν = 40.8 eV) source; a lateral resolution of about
500 nm and energy resolution of 0.9 eV was used to improve the signal-to-noise ratio.
3. Results
3.1. Core-Level Spectroscopy
XPS is a surface sensitive technique that provides insight on the elemental composition and
chemical environment of a sample. The presence of Pb on the Au surface was identified and quantified
using XPS shown in Figure 1.

Figure 1. XPS surveys of the Pb deposited on Au after 150 ◦ C annealing (black), after 5 min of sputtering
(blue), after 7 min 30 s of sputtering (green) and after 380 ◦ C annealing (red). The spectra are shifted
vertically for clarity.

XPS was carried out on the Pb deposited side of the sample after each heating and sputtering
treatment stage, the XPS surveys are shown in Figure 1. After the first 150 ◦ C annealing step, the XP
spectrum showed the presence of Pb and Au, as well as carbon and oxygen from air contaminants and
oxidation. 5 min sputtering was used to gently clean the surface [59,61], successfully removing the
carbon contamination and significantly reducing the oxygen signal as well. However, the Pb quantity
strongly decreased (see Table 1) and the remaining oxygen was predicted to be from oxidized Pb.
Further argon sputtering for an extra 2 min 30 s did not appear to preferentially remove Pb oxide over
metallic Pb, while annealing at 380 ◦ C succeeded in removing the oxygen.
Table 1. Tabulated key results obtained from XPS after each surface treatment step.
Treatment

Pb 4f 7/2 Metal
BE

Pb 4f 7/2 Oxide
BE

Area Pb Oxide/Area
Pb Metal

O/Pb
(at. %)

Overlayer
Coverage

150 ◦ C anneal.
5 min sputter.
7.5 min sputter.
380 ◦ C anneal.

136.9 eV
136.9 eV
136.9 eV
136.8 eV

138.3 eV
138.1 eV
138.0 eV
/

12.0
4.5
3.5
/

1.7
0.9
0.8
/

11.4 Å (3.7 ML)
5.5 Å (1.8 ML)
3.0 Å (1.0 ML)
2.5 Å (0.9 ML)

To investigate the presence of Pb oxide, detailed scans of the Pb 4f and O 1s photoemission lines
were acquired, which are shown in Figure 2. After the initial 150 ◦ C annealing step, the fitted Pb 4f
peaks showed contributions from higher and lower binding energy (BE) doublets, which are from
oxidized and metallic Pb, respectively. The higher BE of the oxide doublets was due to the higher
electronegativity of oxygen with respect to Pb [62]. The O 1s peak also showed two components
after the first annealing step, where the lower BE component was expected to be from Pb oxide, and
the higher component to be mainly from oxygen on carbon contaminants [62,63]. This was further

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confirmed when 5 min of argon sputtering removed carbon, and significantly decreased the higher BE
component of the O 1s peak.

Figure 2. Detailed XP spectra of the Pb 4f (left) and O 1s (right) emission lines of the Pb deposited
on Au after different treatments, normalized by Au 4f areas. The acquired data are represented by
the dotted lines while the peak deconvolutions are drawn as thin lines. The black, blue, green and
red dotted lines represent the data after each treatment procedure, the same as in Figure 1. For Pb
4f the grey, purple and orange lines represent the fitted peaks for Pb metallic, Pb oxide and total fit
respectively. For O 1s, the light blue, purple and orange lines represent the fitted peaks for oxygen
bonded with carbon, with Pb and total fit respectively.

Table 1 lists the fitted BEs for the Pb-metal and Pb-oxide Pb 4f 7/2 peaks, and the stoichiometry of
Pb and O which were calculated by considering the Pb 4f area of oxidized Pb and the overall peak of
O; the peak areas ratios were normalized by their sensitivity factors.
The stoichiometry of O and Pb were not reliable for the first measurement after the 150 ◦ C
annealing due to carbon and oxygen surface contamination. But after sputtering, the carbon
contamination was removed and the O/Pb ratio after 5 and 7 min 30 s was found to be relatively
constant at 0.9 and 0.8 respectively, indicating that the layer after sputtering was mainly composed
of PbO. The BE of the Pb 4f 7/2 oxide decreased significantly with the sputtering procedure from
138.3 eV to 138.0 eV, probably due to the formation of suboxides and damage to the surface due to Ar
implantation (see also WF histograms discussion below).
Table 1 also reports the Pb overlayer estimated thicknesses. For the thickest layers of Pb (after
150 ◦ C annealing and after 5 min sputtering), the coverages have been determined using angle resolved
XPS as conducted in Ref. [64] approximating a homogeneous PbO layer. The coverage estimation
by angle resolved acquisitions on much thinner layers (after 7 min 30 s of sputtering and 380 ◦ C
annealing) is not reliable due to the low signal of the overlayers and their complex morphology. Their
coverages have therefore been extrapolated from the Au 4f /Pb 4f area ratio, calibrated with the
overlayer thickness of the first two acquisitions. The monolayer coverages were obtained using the
lattice parameter of PbO [65] after annealing at 150 ◦ C, 5 and 7 min 30 s of sputtering and using the
distance between atoms in bulk Pb (111) after the 380 ◦ C anneal.
Comparing the ratio of Pb oxide to metallic in the Pb 4f photoemission lines after the 150 ◦ C
anneal, we observed a ratio of 12.0 which then dropped to 4.5 and 3.5 after the 5 min and 7 min 30 s of
sputtering, respectively (see Table 1). The significant ratio difference after the first sputtering confirms
that not all the Pb layer was homogenously oxidized. This was indirect evidence of Stranski- Krastanov
growth, where the Pb islands, formed after the first monolayer, strongly oxidized when exposed to air
due to their high surface areas and probably maintained only a very small core of metallic Pb. The first
mild sputtering removed about half of the Pb layer (see Table 1), and the remaining layer was still

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partially oxidized but with the presence of more metallic Pb. Interestingly the oxide/metal ratio did
not change drastically upon a further 2 min 30 s sputtering. This proves, along with a similar O/Pb
stoichiometry, that the layer compositions at 1.8 ML and 1 ML of PbO are similar and that the oxides
do not form a uniform passivating layer on the top of the Pb metal. On the contrary, the oxidation
from air exposure created oxide particles all over the surface and therefore affected even the 1st Pb ML
in contact with Au [16,66], see also EF-PEEM discussion below.
The final annealing was then performed at 380 ◦ C to remove the Pb oxide without also desorbing
the metal Pb [67]. From Figure 2, the Pb oxide and oxygen emission lines disappeared after the
380 ◦ C anneal, but more interestingly, the metallic Pb peak increased significantly. This indicates
that the annealing did not simply desorb Pb oxide from the Au surface, but rather removed some
of the oxygen from the Pb oxide, probably facilitated by the Pb/Au interface. The calculated layer
coverage after 7 min 30 s was estimated to be about 1 ML, this value is preserved after the 380 ◦ C
anneal (see Table 1), confirming the transformation of the mixed oxide-metal layer to an only metallic
homogeneous overlayer. This latter finding is confirmed by WF maps, and EF-PEEM acquisition
reported below.
3.2. Valence Band Photoemission and Photoemission Microscopy
UPS (using He-I and He-II) was also performed on the sample after every annealing and sputtering
procedure (Figure 3), allowing for a much higher surface sensitive characterization due to the lower
photoelectron energies with respect to the XPS [68]. He-II (hν = 40.8 eV) is even more surface sensitive
than He-I and allows for emission from the Pb 5d levels [69].

Figure 3. UP spectra of Pb on Au after the different treatments, using (a) He-I and (b,c) He-II excitation.
(c) shows a magnified, normalized and background subtracted region of the Au 5d bands after 7 min
30 s sputtering and annealing at 380 ◦ C. UPS of the clean Au side of the sample is also plotted as a
reference on all graphs. The spectra are indicated in black after 150 ◦ C annealing, blue after 5 min of
sputtering, green after 7 min 30 s of sputtering, red after 380 ◦ C annealing and gold for the Au reference.

After the first annealing step (black lines), both UP spectra showed indistinct Au 5d bands, located
between 2 and 7 eV [70,71], but showed two broad features at about at 4.5 and 9.8 eV, as seen in
Figure 3a,b. Two distinct photoemission lines appeared at around 19.5 eV and 22.3 eV BE for the He-II
spectrum of the sample. These correspond to the Pb 5d lines and have distinctly higher BEs than the
18.0 eV and 20.6 eV of metallic Pb [69] confirming the presence of Pb oxide in agreement with the XPS
findings. The significantly higher oxide/metal ratio visible in Figure 3b with respect to the XPS was
due to the much higher surface sensitivity of He-II UPS than XPS.
The 5 min of sputtering significantly sharpened the Au 5d band peaks in both the He-I and He-II
UP spectra, while the Pb oxide peaks decreased and the metallic Pb peaks grew in intensity, indicating
that metallic Pb was initially buried under the Pb oxide and carbon contamination. Further sputtering
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Surfaces 2019, 2

of 2 min 30 s reduced both the metallic and oxide peak areas. As found with XPS, the oxide/metal
ratio visible in the Pb 5d region showed that between the 5 min and 7 min 30 s sputtering the layer
compositions were very similar. Also, the presence of both oxide and metal on the surface after 7 min
30 s of sputtering showed the presence of oxide in contact with Au, further indicating that the metal in
contact with Au oxidized during air exposure. Also confirmed by XPS, the final annealing at 380 ◦ C
completely removed the O from Pb and increased the metallic Pb signal at the oxide’s expense.
Figure 3c shows an enlarged view of the He-II UP spectra of the sample after the 7 min 30 s
sputtering and 380 ◦ C annealing along with undeposited Au region on the sample. The spectra
were normalized to the Fermi edge and subtracted of a Shirley background for easier comparison.
For the sample after 380 ◦ C annealing, a clear shift in the Au 5d bands to higher BEs and shrinkage
were observed, confirming electronic interactions between metallic Pb and Au. On the other hand,
comparing the Au 5d bands after 7 min 30 s sputtering, the shift and shrinkage were minimal and
possibly caused by the overlapping of the broad oxide features at 4.5 eV. Interestingly He-I acquisitions
for the 380 ◦ C annealed sample only showed a minimal shift in BE because the Au signal also originated
from deeper atom layers, He-II was necessary for the extreme surface sensitivity sampling, almost
exclusively at the Pb/Au interface.
To visualize the surface structure at the nanoscale, the low energy emission cutoffs from EF-PEEM
images were used to generate WF maps by performing an error function fitting at each pixel in the
field of view. Figure 4 shows the WF maps of the same position of the sample after 5 min of sputtering,
where the Pb signal was clear and without surface contamination, and after the final 380 ◦ C annealing
stage. The WF was roughly 3.95 eV and 4.15 eV respectively, in agreement with the literature where Pb
oxide has a lower WF than metallic Pb [72,73], considering also that the high extractor voltage used in
the PEEM reduces WF due to the Schottky effect [58].

Figure 4. WF maps of the Pb on Au after the 5 min sputtering (left) and 380 ◦ C annealing (center) steps.
On the (right), a normalized EF-PEEM image at 19.6 eV BE for Pb after 5 min of sputtering. The arrows
are a guide to the eye highlighting some features in the image. The images were taken in the same area,
with a field of view of 17.2 μm.

The WF map after 5 min sputtering showed a surface dominated by circular domains of about
300 nm in diameter, with a few stripes of low WF. EF-PEEM and WF maps obtained on the undeposited
Au side (not reported here) also showed similar domains and stripes, indicating that they came from
the substrate structure and that Pb OPD did not induce any changes in the general surface morphology
as observed with the selected lateral resolution (150 nm).
EF-PEEM was performed using the He-II excitation on the Pb 5d5/2 emission line, providing a
direct mapping of oxidized Pb at 19.6 eV. After 5 min sputtering, although very weak, some stripes
of higher intensity can be seen on the chemical map (indicated by red arrows). These features are in
the same position as the low WF stripes indicating that the low WF was due to Pb oxide. Notably, the
chemical map shows also that the oxide was not only localized on the stripes, but it was distributed on
all the sample’s surface. The poor lateral resolution (500 nm) and the low signal made it difficult to see
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Surfaces 2019, 2

its exact position, but it could be roughly related to the zone with WF < 4.1 eV in the WF map after
5 min of sputtering reported in Figure 4.
The WF contrast on the Pb side disappeared after annealing to 380 ◦ C suggesting that the
stripes were initially covered with Pb oxide but reduced to Pb after the annealing, confirming the
transformation of part of the oxide into the metal as found in the XPS analysis. This finding is further
confirmed by the fact that the chemical map at 18.0 eV (Pb 5d5/2 metallic) after annealing at 380 ◦ C
(not reported here) was homogeneous. Within our lateral resolution (150 nm), we could notice that the
annealing process did not change the surface morphology.
Histograms of all the WF maps were plotted in Figure 5 to show the distribution of WF across the
surface after each treatment step.

Figure 5. WF distributions after each treatment step plotted as a histogram after 150 ◦ C annealing
(black), after 5 min of sputtering (blue), after 7 min 30 s of sputtering (green), after 380 ◦ C annealing
(red) and for the Au reference (golden). All the distributions were taken from the same area other than
the Au reference, which was done on the undeposited half of the sample.

The WF varied throughout each step, where the sample at the first step (black) was carbon
contaminated, then (blue) was cleaned, but further sputtering (green) increased the damage to the
sample surface leading to surface roughening and broadening of the WF histogram; this broadening
can also be related to the formation of suboxides as discussed above in the XPS section. With the final
annealing (red), only metallic Pb remained on the surface with a very narrow WF distribution. The WF
of the Pb film remained lower than WF of the polycrystalline Au [74] throughout each step, indicating
that the OPD fully covered Au and the treatments did not expose the underlying Au, confirming the
XPS coverage calculation in Table 1.
4. Discussion
With the combination of multiple surface sensitive techniques, we were able to determine the
general structure of the Pb deposited on Au by OPD, its oxidation from exposure to atmosphere,
the transformation of the surface with sputtering and annealing, and then the changes and shifts in the
electronic bands from interaction between Au and Pb.
With XPS, we were able to determine the Pb oxidation species after sputtering to be PbO.
The amount of deposited Pb was also quantified and found to be as expected from the OPD charge.
Contaminants from the solution used in the electrochemical deposition were also not detected. XPS and
UPS also gave an indirect insight into the nanostructure of the materials, confirming Stranski-Krastanov
growth and that air oxidation reached the 1st Pb strained ML.

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Moreover, we observed that the Pb layer thickness after 7 min 30 s sputtering and the 380 ◦ C
of annealing did not change. This might be an indication of type-I alloying, which agrees with the
in situ STM results published by our group [48]. In type-I alloys, the surface adatoms are found to
be diffusing on the surface and are highly mobile, due to the kinetics of alloying and repulsion [75].
Indeed, we have found that the Pb was able to migrate over to the clean Au side of the sample (over
2 mm migration) after annealing to 380 ◦ C.
One particularly interesting finding is the interaction between Au and the metallic Pb from the
UPS in Figure 3c (red curve). The shift to higher BE in the Au 5d bands could be a result of differential
relaxation, exchange and correlation effects [76] due to the presence of Pb in the Au lattice [77]. Shifts
in the 5d band have also been reported for Au alloys [70] and were explained by charge transfer and
electron d-band repulsion where the band narrows [78]. Due to the alloying, the Au 5d peaks narrowed,
and the spin-orbit splitting decreased and overall shifted to higher BEs, as observed with Au alloyed
with Cd [77]. Comparing our findings to these reports we can assume we formed a Pb/Au surface
alloy after the annealing at 380 ◦ C.
From our UPS investigation we also showed that, in the case of partially oxidized Pb, the interaction
with Au was negligible. The interfacial strain at the interface due to the Pb oxidation was lower
than in the case of Pb metal. The Pb oxide affected the possibility of Pb metal to form a surface alloy
surrounding it with oxide (see Figure 4) and eliminated the stress at the interface, which can explain
why the Pb metal, present after only sputtering treatments, did not interact with Au. The high diffusion
rate of Pb has been observed during annealing at high temperature. It is possible that the Pb-Au bond
is much weaker at elevated temperatures compared to that at room temperature and the fast diffusion
of Pb on the Au surface suggests almost ‘lattice gas’ behavior. In the presence of the metallic Pb, the
system relaxes by minimizing the elastic energy through surface alloying. This is typical for immiscible
and highly strained systems such as ours [79,80].
The high temperature annealing, removing the oxide and transforming it into metal, is the only
treatment after which the metal enters in direct contact with the Au surface and forms a surface alloy.
On the other hand, this treatment allows faster diffusion of Pb during which the alloy formation is not
favorable. In future papers, we will explore these and similar materials at different temperatures to
investigate the interplay between the surface alloying and surface interaction (ad-atom diffusion).
Author Contributions: Data curation, M.C.; Funding acquisition, N.A.F.; Investigation, A.S., G.W. and M.C.;
Methodology, A.S., N.V., G.W. and M.C.; Supervision, M.C., N.A.F. and N.V.; Validation, M.C. and N.V.;
Visualization, M.C.; Writing—original draft, A.S., G.W. and M.C.; Writing—review & editing, A.S., M.C., N.A.F.
and N.V.
Funding: This research was partially funded by EPSRC Strategic Equipment, Grant number EP/K035746/1
and EP/M000605/1; Ph.D. studentships are funded through BCFN: EPSRC (EP/L016648/1), and Renewtec
Technologies, Al Hamad Group.
Acknowledgments: The authors acknowledge the Bristol NanoESCA Facility (EPSRC Strategic Equipment Grant
EP/K035746/1 and EP/M000605/1). A.S. and G.W. acknowledge Ph.D. studentships funded through BCFN:
EPSRC (EP/L016648/1), and Renewtec Technologies, Al Hamad Group, respectively.
Conflicts of Interest: The authors declare no conflict of interest.

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Article

Co3O4 Nanopetals on Si as Photoanodes for the
Oxidation of Organics
Leonardo Girardi 1 , Luca Bardini 1 , Niccolò Michieli 2 , Boris Kalinic 2 , Chiara Maurizio 2 , Gian
Andrea Rizzi 1, * and Giovanni Mattei 2
1
2

*

Department of Chemical Sciences, University of Padova, via Marzolo 1, 35121 Padova, Italy;
leonardo.girardi@phd.unipd.it (L.G.); luca.bardini@unipd.it (L.B.)
Physics and Astronomy Department, University of Padova, via Marzolo 6, 35121 Padova, Italy;
niccolotomaso.michieli@unipd.it (N.M.); boris.kalinic@unipd.it (B.K.); chiara.maurizio@unipd.it (C.M.);
giovanni.mattei@unipd.it (G.M.)
Correspondence: gianandrea.rizzi@unipd.it; Tel.: +39-049-827-5722

Received: 13 November 2018; Accepted: 9 January 2019; Published: 11 January 2019

Abstract: Cobalt oxide nanopetals were grown on silicon electrodes by heat-treating metallic cobalt
films deposited by DC magnetron sputtering. We show that cobalt oxide, with this peculiar
nanostructure, is active towards the photo-electrochemical oxidation of water as well as of organic
molecules, and that its electrochemical properties are directly linked to the structure of its surface.
The formation of Co3 O4 nanopetals, induced by oxidizing annealing at 300 ◦ C, considerably improves
the performance of the material with respect to simple cobalt oxide films. Photocurrent measurements
and electrochemical impedance are used to explain the behavior of the different structures and to
highlight their potential application in water remediation technologies.
Keywords: cobalt oxide; water oxidation; photo-electrochemistry; hydroxyl radical; electro-oxidation

1. Introduction
The photoelectrochemical oxidation of water is a major topic in modern chemical engineering;
key technologies, such as the production of hydrogen and oxygen as renewable fuels [1] and the
quantification [2] and abatement of organic pollutants [3,4] in the field of water remediation, in fact,
depend on it. The cornerstone for the development of scalable Photoelectrochemical Water Oxidation
(PEC-WO) processes is to find cost-effective, active, and durable semiconducting photo-electrodes able
to sustain oxygen evolution reaction (OER) for long periods of time. To this end, silicon is a preferred
material, as it possesses a suitable band gap and is readily available. The amount of oxygen produced
by Si photoelectrodes, however, is usually small. More importantly, the performances of silicon
photoanodes decay rapidly due to corrosion. To overcome these problems, Si-based water-splitting
electrodes require the use of a thin (210 nm) protective coating; the thickness of the coating should
be precisely controlled, as a layer which is too thick is reported to hamper the performance of the
underlying Si [5,6]. The metals commonly used for this purpose are Ru [7], Ni [8], NiOx [9], Ir [10],
Cu [11], CuOx [12], MnOx [13], Co, and CoOx [14].
Cobalt oxide, in particular, has been shown to be effective as a protective layer in a variety
of conditions [15,16], as it is particularly stable in the strongly basic conditions employed for the
electro-oxidation of water. The use of Co3 O4 thin films as passivating layers was also explored
by Yang et al. [17], who obtained a stable photocurrent as high as 17 mA/cm−2 in 1 M potassium
hydroxide (KOH; pH = 13.6). On the other hand, thanks to its absorption in the visible range,
cobalt oxide has also been used as a visible light sensitizer to increase the activity of wide band-gap
photoanodes at longer wavelengths [18–20].

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In this work, we show that cobalt oxide is an active participant in the electrochemical oxidation of
water under visible light irradiation, and that it can be effectively used as a stand-alone photoanode.
The thermal oxidation of a Co film, deposited by Direct Current (DC) magnetron sputtering on n-Si(100),
forms p-type Co3 O4 layers. These systems are electrodes with a stable anodic photocurrent in both
neutral Na2 SO4 and NaOH solutions. Moreover, we demonstrate that the controlled nano-structuring
of the cobalt oxide structure produces significant improvements in photoelectrode performance.
We further apply these findings to the oxidation of various organic molecules, and show that the
photocurrent increases proportionally with the organic content of the solution [21,22], indicating that
cobalt oxide photoanodes can be a promising material for applications in water remediation, as well
as a greener, more efficient alternative for the determination of chemical oxygen demand (COD), with
respect to commonly used dichromate methods [23]. In particular, we will show that different thermal
treatments of the Co layer give rise to different nanostructures characterized by different electrical and
sensing properties.
2. Materials and Methods
2.1. Sample Preparation
The substrates used were n-type Si(100) wafer, preliminarily washed in H2 SO4 :H2 O2 at 90 ◦ C.
The native silicon oxide passivation layer, whose typical thickness is about 2 nm [24], was not
removed prior to deposition. To this respect, it is worth mentioning that there are several examples
in the literature of Si-based photoanodes where the presence of the native silica layer was found to
improve the photo-electrochemical performances; some authors suggested that the silica layer may act
as an adhesive between the substrate and the deposited film [6]. 50 nm-thick Co films were deposited
on the Si substrates by DC magnetron sputtering, with a deposition rate of 1 Å/s. The deposited films
were annealed in O2 flux at different temperature values, as detailed in Table 1.
Table 1. Details of the post-deposition annealing performed in pure oxygen flux.
Sample Name

Treatment Conditions

PETAL
HYBRID
NOPETAL

4 h at 300 ◦ C
4 h at 300 ◦ C + 1 h at 450 ◦ C
1 h at 450 ◦ C

2.2. Structural and Morphological Characterizations
The thickness of the deposited films was measured by Atomic Force Microscopy (AFM)
using a NT-MDT PRO Solver microscope (NT-MDT, Moscow, Russia).
Scanning Electron Microscopy (SEM) measurements were carried out with a Zeiss Sigma HDF
Field-Emission Scanning Electron Microscopy (FE-SEM), equipped with InLens, Secondary Electrons
(SE), and Backscattered Electrons (BSE) detectors for the imaging (Zeiss, Jena, Germany).
Grazing Incidence X-ray Diffraction (GIXRD) spectra were collected with an X’Pert Pro
diffractometer (incidence angle = 1 deg) using the Cu Kα radiation. The GIXRD analysis was performed
by using MAUD software (University of Trento, Italy) [25].
Surface composition was determined by using X-ray Photoelectron Spectroscopy (XPS)
measurements, performed on a custom-built Ultra High Vacuum (UHV) chamber (base pressure
= 5 × 10−10 mbar) equipped with a non-monochromatized, double-anode X-ray source (Omicron
DAR-400, Scienta-Omicron GmbH, Uppsala, Sweden), a hemispherical electron analyzer (Omicron
EA-125, Scienta-Omicron GmbH, Uppsala, Sweden), and a 5-channeltron detector assembly.
The electron analyzer had an acceptance angle of ±4◦ , and the diameter of the analyzed area was
3 mm. The spectra were acquired with Mg Kα radiation.

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2.3. Electrochemical Measurements
Electrochemical measurements to test the activity were performed either in 0.1 M sodium
sulphate or in 0.1 M sodium hydroxide in milliQ water. The measurements were made with Autolab
PGSTAT204 potentiostat (Metrohm, Utrecht, The Netherlands) in a Teflon PhotoElectroChemical
(PEC) cell. A Pt coil wire and Ag/AgCl electrode were used as the counter electrode and reference
electrode, respectively. The samples were mounted outside the cell and kept in position by an O-ring
seal. The electrical contact was obtained by a metal tip which was firmly pressed onto the back
of the Si wafer by a spring. All samples were illuminated from the front. Photoelectrochemical
experiments were carried out with a white LED with an intensity of about 100 mW/cm2 . The light
intensity was measured using a photodiode at the same distance from the light source as the sample;
the light was passing through the electrolyte solution, and the quartz window mounted on the cell.
The light intensity was controlled by the optical bench (Metrohm-Autolab) coupled to the Autolab
PGSTAT204 instrument (Metrohm, Utrecht, The Netherlands). All Linear Sweep Voltammetry (LSV)
measures have been obtained with a scan rate of 5 mV/s. Electrochemical Impedance Spectroscopy
(EIS) data were obtained under illumination, and in the dark condition, the voltage amplitude for EIS
measurements was ±10 mV with the frequency range set from 105 Hz to 100 Hz, performing 50 points
with logarithmic distribution.
3. Results and Discussion
3.1. Structure and Activity
In Figure 1a–c, the FE-SEM top-view images of the three electrodes are shown. The heating process
of the Co metallic film at 300 ◦ C for 4 h (sample labeled PETAL) induces the growth of nanopetals
from the layer underneath (Figure 1a). The petals’ surface, ≈100–200 nm wide, is preferentially normal
to the original film surface, and the petal thickness is limited to about 10 nm. These structures are
similar to others which have previously been observed [26]. The petal formation is dependent on the
annealing temperature. Indeed, a similar treatment performed at 450 ◦ C (sample labeled NOPETAL,
Figure 1c) induces the formation of a rough surface characterized by relatively large grains and with no
nanopetals. If the PETAL sample is further treated at 450 ◦ C (sample labeled HYBRID), the nanopetals
are less visible (Figure 1b), and the overall surface morphology is intermediate between that of sample
PETAL and of sample NOPETAL. The GIXRD patterns of the three samples (Figure 1d) are all very
similar and show the signature of Co3 O4 nanocrystals, whose average size is about 25 nm for the
sample PETAL, 40 nm for the HYBRID one, and 50 nm for the NOPETAL one.
The different surface morphology has a direct impact on the electrochemical properties of the
material towards the oxidation of water. A sketch of the electronic band structure of the prepared
electrodes is shown in Figure 1e, where the n-Si substrate and the Co3 O4 layer (p-type semiconductor,
as shown below) form a “tunneling” p-n junction with an interposed SiOx dielectric barrier [27].
An approximate position of the Valence Band (VB) and Conduction Band (CB) edges of the materials,
before the junction is formed, can be easily derived from Pearson absolute electronegativity values [28],
where the band-gap values are the ones reported in the literature for Si(100), native SiO2 [29], and
Co3 O4 [30]. After Fermi-level alignment, the approximate position of the CB and VB edges can be
obtained for the sample PETAL from the Mott-Schottky (MS) plots (see SI Figure S2, curve acquired
at 2000 Hz) using the formulas [31–33]: Ecb (V) = V f b + 0.1 for Si(100) and Evb (V) ≈ V f b for Co3 O4 ,
where V f b is the flat-band potential obtained from the MS plots (V f b = 0.95 V vs. RHE for Si and
V f b = 1.66 V vs. RHE for Co3 O4 at 2000 Hz). Different values of V f b can be obtained in the case of
samples NOPETAL and HYBRID due their different nanostructuration [34]. The sketch of Figure 1e
indicates that, after the junction has formed, there is an electron migration from Si to Co3 O4 that shifts
the VB edge of Co3 O4 slightly above the VB of Si(100).

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Figure 1. (a–c) FE-SEM top-view images of the samples PETAL (a), HYBRID (b), and NOPETAL (c).
(d) Corresponding Grazing Incidence X-ray Diffraction (GIXRD) patterns. The Co3 O4 reflections are
marked (ICSD-63164). In the dashed region, a spurious diffraction signal from the Si substrate is visible
(the X-ray penetration depth is about 80 nm). (e) Sketch of the electronic energy bands of the prepared
electrodes. The theoretical positions of edges CB and VB, obtained from Pearson’s electronegativity
values, are in red. The approximate positions of CB and VB edges, after the junction is formed, obtained
from Mott-Schottky plots (SI. Figure S2) are indicated by green bars.

Electrochemical measurements were carried out either in 0.1 M NaOH or in Na2 SO4 , the former
being the ideal condition for water oxidation, the latter being more relevant for sensing and water
remediation. In fact, in NaOH electrolytes, the difference in the photocurrent values after the
addition of small volumes of different analytes (EtOH, glucose, etc.) is less evident, since it is
overshadowed by the already high photocurrent value due to hydroxyls oxidation. As shown in
Figure 2a, the samples HYBRID and NOPETAL show similar current outputs with respect to water
oxidation, despite the different surface morphology. The current for the sample PETAL, on the other
hand, is at least three times higher in sodium sulphate, and almost two times higher in alkaline
conditions, with a concomitant shift of the onset potential (SI Figure S1).

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Figure 2. (a) LSV in 0.1 M Na2 SO4 and 0.1 M NaOH of the three samples in dark and illuminated
conditions; (b) charge values at 1V vs. Ag/AgCl, from chronoamperometries (pH = 6), as a function of
theoretical Chemical Oxygen Demand (COD) of the sample PETAL with different organic molecules
in a 0.1 M sodium sulphate solution; (c) LSV with chopped illumination of the sample PETAL in
sodium sulphate before (orange) and after (blue) the addition of glucose. The behavior of the bare
silicon electrode is also shown for reference (black curve); (d) photocurrent dependence of sample
PETAL at 1 V vs. Ag/AgCl on the concentration of glucose.

For the sample PETAL, if a source of carbon like EtOH, MeOH, or glucose is added to the sodium
sulphate solution, a clear enhancement of the photocurrent is seen (Figure 2c), indicating that the
mineralization process is the decomposition of organic substances and can be described with the
reaction [2]: Cy Hm Oj Nk Xq + (2y − j)H2 O → yCO2 + q X− + kNH3 + (4y − 2j + m − 3k)H+ + (4y − 2j
+ m −3k − q)e− .
The process is partly due to electron capture by holes on the Co3 O4 surface (the other reaction
that leads to the mineralization to CO2 involves the reaction with hydroxyl radicals generated during
the oxidation of water). The electrode response to chopped illumination is characterized by ”spike
and overshoot” photocurrent transients. We attribute the photocurrent spikes of Figure 2c (blue and
orange curves) to a high degree of electron and hole recombination on the surface, and to the formation
of surface states (trap-states). This is a known behavior which is described, for instance, in [35,36].
When the light is turned on, holes generated in the space-charge region are swept rapidly towards
the semiconductor electrolyte junction. Due to the slow kinetics of the four-hole oxidation of water
to molecular oxygen, the concentration of holes increases considerably at the interface, until the rate
of arrival of holes is balanced in the steady state by the rates of charge transfer and recombination.
Since surface recombination leads to a flux of electrons towards the surface, the resulting photocurrent
transient is the sum of the hole and electron contributions. A careful analysis of Figure 2c also shows
that as soon as a hole scavenger such as glucose is added, the cathodic spikes are strongly attenuated.
It is worth reminding that, quite often, the COD value is obtained by measuring the total charge
due to the photocurrent after the degradation to CO2 is complete (see Figure 3a), and its value is
expressed as mg O2 /L, according to the equation [2]:

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Surfaces 2019, 2

COD(O2 mg/L) =

nC
Q
× 32000 =
× 32000 = K × Q
4
4FV

(1)

where n is the number of electrons, C is the concentration of substances, Q is the charge passed, V is
the volume of the solution, and F is the Faraday constant. We report in Figure 2b the total charge due
to the mineralization process for five different substances, as a function of theoretical COD (mg O2 /L).
This graph suggests that the charge is directly correlated to the theoretical COD values—in other
words, the higher the theoretical COD value, the higher the charge.
To better highlight the proportionality between the photocurrent/charge and organic content,
as well as to verify the concentration range to obtain a linear response, further tests were conducted
with glucose as a proof-of-concept carbon source, using the PETAL sample as an electrode. Since the
current arising from the oxidation of water is lower in sodium sulfate, the current generated by the
oxidation process of the organic molecule can be better highlighted. For this reason, the quantification
of different concentrations of glucose was reported in a sodium sulfate solution (Figure 2d).

Figure 3. (a) Photocurrent at 0.95 V vs. Ag/AgCl of the layer PETAL over time in a 3 mM phthalate, 0.1
M sodium sulphate solution (pH = 6); (b,c) fitted X-ray photoelectron spectroscopy (XPS) spectra [37]
of the Co 2p3/2 orbitals taken before (b) and after (c) 14 h of phthalate oxidation.

All samples show a linear response with glucose concentration, with the sample PETAL showing
the widest linear range of 0.4–4.2 mM (see SI, Table S1), a sensitivity of 7.5 ± 0.1 μA/cm2 mM−1 , a limit
of detection (LOD) of 2.3 ± 0.2 μM, and a limit of quantification (LOQ) of 7.7 ± 0.6 μM (The limit
of detection (LOD) and limit of quantification (LOQ) were calculated from the standard deviation
(σ) of the linear part of LSV measurements (Figure 2a) and from the slope (b) of the calibration
curve (Figure 2d). For LOD, we considered Δj = 3σ as the minimum value of Δj to detect glucose
(LOD = 3σ/b), while for LOQ, we considered it to be Δj = 10σ (LOD = 10σ/b). The error was calculated
using the propagation formula). Sensitivity and linear range are comparable with some of the values
found in the literature for metal oxide nanostructures [38].
Considering a previous report which demonstrates that a passivating film with a thickness of
10 nm is enough to significantly hamper the photoelectro-catalytic activity of silicon [6], we can safely
state that the surface of the Co3 O4 nanostructured layer (whose average thickness after annealing
is about 80 nm) is the only active species in the present scenario. This observation is supported by
Mott-Schottky plots (see SI, Figure S2), which show a positive slope (n-type behavior of the silicon
wafer) in the dark, but a negative slope (p-type behavior of Co3 O4 ) under illumination, showing that
cobalt oxide dominates the photoelectrochemical behavior of the electrode. In this case, the completely
different shape of the MS plots under illumination is probably due to inhomogeneities of the electrical
field and formation of surface states in the layer where the photocurrent is built up, as recently
described in a paper by Kirchartz et al. [39].
To further probe the performance of our nanostructured cobalt oxide layers, phthalate was
chosen as the next target molecule. Phthalates form a class of extremely common plasticizers and

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Surfaces 2019, 2

pollutants, which, due to their large diffusion, cause serious concerns for human health [40]. A
long-term test showed that after more than 10 h, the photocurrent had gradually decreased to
the value of the clean solution, containing only Na2 SO4 . In this case, it was possible to calculate
the value of COD according to Equation 1, and the proportionality constant K was found to be
1.594 × 10−4 O2 mg/L × C, a value similar to the ones found in the literature [2], indicating that
the mineralization process is mainly due to interaction with hydroxyl radicals. Moreover, X-ray
photoelectron spectra (XPS), taken on an as-prepared electrode and on the electrode after the long-term
test, showed no significant changes in the 2p peaks of cobalt; in particular, satellites which can be
related to Co(II) oxides/hydroxides are absent [41] (Figure 3b), highlighting the excellent stability of
the cobalt oxide surface even after prolonged operation.
3.2. Mechanistic Analysis
To fully characterize the effect of the surface structure of Co3 O4 on its electrochemical properties,
glucose was chosen for further analyses. In the literature, the oxidation of glucose is said to be caused
by the presence of Co3+ and Co4+ surface states according to the following reaction sequence [42,43]:
Co3 O4 + OH− + H2 O ↔ 3CoOOH + e− (0.3 V vs. Ag/AgCl)
−

−

(2)

CoOOH + OH ↔ CoO2 + H2 O + e (0.6 V vs. Ag/AgCl)

(3)

2CoO2 + C6 H12 O6 (glucose) → 2CoOOH + C6 H10 O6 (gluconolactone)

(4)

In the present case, cyclic voltammetries, acquired in NaOH electrolyte, as in the case of
Co3 O4 nanofibers [44], (CVs, SI Figure S3) never clearly showed the presence of oxidation peaks
at about 0.3 V and 0.6 V vs. Ag/AgCl. We attribute this behavior to a much lower surface area
of our samples compared with the examples found in the literature, where Co3 O4 nano-structrures
characterized by a very high surface area were used [45]. Moreover, it is likely that a different
mechanism was involved in our case, since an enhancement in the anodic (photo)current appeared
only when the sample surface was illuminated.
The effect of light is likely a combination of the population of the conduction band of the silicon
substrate, which allows the current to tunnel through the Co3 O4 /SiO2 /Si junction, as well as the
creation of active cobalt species on the surface of the electrode. The electrochemical impedance spectra
(EIS) reported in Figure 4 and performed in the dark and under illumination with a white LED in
Na2 SO4 electrolyte, also highlight the presence of two different contributions to the electrochemical
properties of the electrodes. EIS experiments were performed at two different overpotentials in order
to investigate the behavior of the polarized electrode both with (1 V) and without (0.1 V) the presence
of electrochemical reactions.
In the dark, significant differences were already apparent between the NOPETAL and the PETAL
layers. At low bias (100 mV vs. Ag/AgCl), the PETAL layer has two contributions of comparable
magnitude, one in the low frequency range and one in the high frequency one. The NOPETAL
layer shows a major contribution at low frequency, which largely overshadows the high-frequency
semi-circle. At low bias, the same behavior is present even when the sample is illuminated. For both
samples PETAL and HYBRID, on the other hand, there is a stark decrease in the low-frequency
contribution when the LED is turned on. This information can be taken as an indication that,
for samples PETAL and HYBRID, the lower-frequency part of the spectrum is dominated by the
contribution of surface states. For sample PETAL in particular, a description with three time constants
can also be proposed; in this scenario, the second time constant (roughly in the range 10 kHz–150 Hz)
disappears when the sample is illuminated. Such an interpretation would suggest the second
contribution being due to the space charge layer.
At a bias of 1 V in the dark, the low-frequency contribution remains dominant for both
NOPETAL and HYBRID layers; it is significant that, with the increase in bias, the impedance of
the high-frequency component for the sample NOPETAL increases. Conversely, in the case of the
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Surfaces 2019, 2

PETAL layer, the high-frequency contribution remains relatively unvaried, while there is an increase in
the mid/low frequency semi-circle, a behavior consistent with the application of a bias at a blocking
electrode-solution interface. These results indicate that the structuring of the Co3 O4 layer in sample
PETAL is also beneficial for the electrical connection with the underlying silicon substrate.

Figure 4. Comparison of Nyquist impedance spectra for the PETAL , NOPETAL , and HYBRID layers.
(a–c) Spectra taken at 0.1 V; (d–f) spectra taken at 1 V. In all figures, spectra recorded in the dark are
presented in blue, while spectra recorded under illumination are presented in orange. All spectra have
been acquired in 0.1 M Na2 SO4 electrolyte. Solid lines were obtained by fitting the EIS spectra with the
EC (equivalent circuit) represented in the upper part of each panel. CPE indicates a constant phase
element, Rs indicates the solution resistance. The inset of Figure (d) shows an enlargement of the fitting
results at low impedance values.

When the electrodes are illuminated at 1 V, Nyquist plots indicate a similar overall decrease in
impedance across the entire frequency range for all samples. For samples NOPETAL and HYBRID ,
a significant impedance response is also present at lower frequencies; a closer inspection of the behavior
of the phase angle reveals significant differences in the behavior of the three surface structures in this
region (Figure 5). While the PETAL layer shows near-ideal resistive behavior across the frequency
spectrum up to 100 Hz, a residual RC contribution at low frequency remains for the NOPETAL and
HYBRID layers. We also note that, upon illumination, the phase behavior, as a function of the frequency,
is intermediate between the ones of the HYBRID and PETAL layers, corresponding to the fact that the
surface morphology is also in between the two (see Figure 1a–c).The overall picture suggests that the
formation of surface states (i.e., CoOOH or CoO2 ) is linked to the presence of the petal-like structures.
The time constant visible in Figure 5 with light on, at around 50 Hz, can then be linked to the charge
transfer resistance between the electrode and the species in the solution. For sample NOPETAL , there
is significant resistance probably given by the fact that the reaction mechanism largely does not directly
involve the formation of surface states on the cobalt oxide surfaces. At the opposite extreme, sample

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Surfaces 2019, 2

PETAL does not have significant contributions in this region due to the direct involvement of Co3 O4
surface states in the reaction, which form a low impedance pathway for electrons. Sample HYBRID lies
in between; as can be seen in Figure 1, some petal-like structures can be spotted in sample HYBRID,
but most of them were destroyed by the thermal treatment.

Figure 5. Phase angle spectra recorded at 1 V in 0.1 M sodium sulphate for samples PETAL (red),
NOPETAL (blue), and HYBRID (green).

This qualitative description, implying that the existence of surface states is linked to the presence
of petal-like structures, is confirmed by fitting the impedance data by an equivalent circuit (EC),
containing, in the case of sample PETAL, three RQ elements corresponding to the Co3 O4 interface
with a SiOx /Si, Co3 O4 space charge region (characterized by the presence of surface states) and the
double layer at the interface with the electrolyte [46]. In the case of samples HYBRID and NOPETAL,
the EC consists in only two RQ (Q = CPE) elements, relatable to the Co3 O4 /SiOx /Si interface and
Helmholtz layer, confirming the fact that the contribution to the overall process of surface states in
these samples is negligible. A table containing the calculated resistances and CPE values, in the three
cases, is reported in the Supplementary Materials (Table S2).
The significantly different behavior of the sample PETAL is then to be assigned to a particularly
favorable interplay of surface area and the creation of active surface states, showing that
nanostructuring is not merely morphological, but has important implications on the electronic structure
of the Co3 O4 at the interface with the solution. It should be noted that one further element which
could contribute to the improved performance of the PETAL sample is the change in the wettability
of the surface, which can play an important role when gas-evolving reactions are involved [47].
Further experiments are needed to clarify this point. The possibility of combining the effect of surface
structuring, as described herein, with the creation of oxygen vacancies through surface doping [48]
makes cobalt oxide photoelectrodes a promising new substrate for the engineering of simple, stable,
and effective photoanodes.
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4. Conclusions
Nanostructured cobalt oxide layers obtained by oxidation of a 50 nm Co film deposited on 100 n-Si
by DC-magnetron sputtering have been produced. The choice of the annealing parameters allowed
for the production of a Co3 O4 layer decorated with nanopetals. This system can form a versatile and
durable photoanode, and has been tested as a sensor for different water-soluble organic molecules
in a sodium sulphate solution, obtaining, for glucose, a sensitivity of 7.5 ± 0.1 mA cm−2 M−1 .
This nanostructured electrode can also be effective for water remediation. Indeed, we have shown
that under positive bias and visible light illumination, the phthalate gets completely mineralized
after about 6 h of electrochemical work. Moreover, the composition of the Co3 O4 nanopetal layer
after repeated sensing experiments and electrochemical work under the application of illumination
and bias is virtually identical to that of the as-grown electrode, as shown by XPS analysis. Different
layer morphologies can be obtained by specific thermal treatments which remarkably influence the
electrical properties of these electrodes, as proven by EIS measurements and, consequently, their
sensing properties. We believe that this nanostructured electrode, formed by Co3 O4 nanopetals on
Si, can be quite easily engineered to produce sensing devices or larger-area electrodes for water
remediation applications.
Supplementary Materials: The following are available at http://www.mdpi.com/2571-9637/2/1/4/s1,
Figure S1: Linear sweep voltammetry (5 mV/s) under illumination of the PETAL sample in different electrolyte
solutions; Figure S2: Mott-Schottky plots of the prepared samples in the dark and under illumination in Na2 SO4 ;
Figure S3: Cyclic voltammetry performed on sample PETALS, in 0.1 M NaOH and 5 mM glucose with scan rate
of 40 mV/s in the dark and with illumination; Table S1: Sensing in Sodium Sulphate 0.1 M at 1 V vs. Ag/AgCl;
Table S2:Equivalent Circuit parameters.
Author Contributions: Investigation, L.G., L.B.; investigation, writing–original draft preparation, G.A.R.;
conceptualization, investigation, writing–review and editing: N.M., B.K., C.M.; writing–review and editing: G.M.
Funding: This research was partially funded by Università degli Studi di Padova, Physics and Astronomy
Department, grant number BIRD178923/17. Authors gratefully acknowledge the Italian Minister of University
(MIUR) for financial support to the SMARTNESS (Solar driven chemistry: new materials for photo- and
electrocatalysis) financed thorough the PRIN 2015K7FZLH.
Acknowledgments: We thank Alessandro Michieletto for helpful discussions.
Conflicts of Interest: The authors declare no conflict of interest. The founding sponsors had no role in the design
of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, and in the
decision to publish the results.

Abbreviations
The following abbreviations are used in this manuscript:
UHV
FE-SEM
GIXRD
VB
CB
EIS
PEC
COD
LSV
XPS

Ultra High Vacuum
Field Emission- Scanning Electron Microscopy
Grazing Incidence X-Ray Diffraction
Valence Band
Conduction Band
Electrochemical Impedance Spectroscopy
PhotoElectroChemical
Chemical Oxygen Demand
Linear Sweep Voltammetry
X-Ray Photoelectron Spectroscopy

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Surfaces 2019, 2

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article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Potential Driven Non-Reactive Phase Transitions of
Ordered Porphyrin Molecules on Iodine-Modified
Au(100): An Electrochemical Scanning Tunneling
Microscopy (EC-STM) Study
Tomasz Kosmala 1,2, *, Matías Blanco 1 , Gaetano Granozzi 1 and Klaus Wandelt 2,3, *
1
2
3

*

Department of Chemical Sciences and INSTM Unit, University of Padova, Via F. Marzolo 1,
35131 Padova, Italy; matias.blancofernandez@unipd.it (M.B.); gaetano.granozzi@unipd.it (G.G.)
Institute of Physical and Theoretical Chemistry, Bonn University, Wegeler-Strasse 12,
D-53115 Bonn, Germany
Institute of Experimental Physics, Plaza Maxa Borna 9, 50-204 Wroclaw, Poland
Correspondence: tomasz.kosmala@unipd.it (T.K.); k.wandelt@uni-bonn.de (K.W.)

Received: 3 July 2018; Accepted: 24 July 2018; Published: 25 July 2018

Abstract: The modelling of long-range ordered nanostructures is still a major issue for the scientific
community. In this work, the self-assembly of redox-active tetra(N-methyl-4-pyridyl)-porphyrin
cations (H2 TMPyP) on an iodine-modified Au(100) electrode surface has been studied by means of
Cyclic Voltammetry (CV) and in-situ Electrochemical Scanning Tunneling Microscopy (EC-STM) with
submolecular resolution. While the CV measurements enable conclusions about the charge state of
the organic species, in particular, the potentio-dynamic in situ STM results provide new insights into
the self-assembly phenomena at the solid-liquid interface. In this work, we concentrate on the regime
of positive electrode potentials in which the adsorbed molecules are not reduced yet. In this potential
regime, the spontaneous adsorption of the H2 TMPyP molecules on the anion precovered surface
yields the formation of up to five different potential-dependent long-range ordered porphyrin phases.
Potentio-dynamic STM measurements, as a function of the applied electrode potential, show that the
existing ordered phases are the result of a combination of van der Waals and electrostatic interactions.
Keywords: porphyrins; self-assembly; surface nanostructures; in situ EC-STM; metal-electrolyte
interface; potential-dependent structures; combined non-covalent control

1. Introduction
One of the major challenges in the development of supramolecular nanoarchitectures is to find
effective and flexible methods in order to create nanostructures with long range order [1,2]. In recent
years, supramolecular self-assembly attracted much attention by the scientific community: as a
well-known bottom-up method, molecular self-assembly on surfaces is a simple and fast tool to
build nanoscale structures, which can be readily tuned by the on-purpose synthesis of appropriate
molecular building blocks, and, thereby, the design of the noncovalent interactions between the
molecules and the substrate and the molecules themselves [3–7]. Moreover, to direct the ordering
process of the supramolecular nanoarchitectures, one can take advantage of well-defined single crystal
surfaces, which can serve as electronic and/or geometric templates due to the operation of specific
adsorbate-substrate interactions. In this case, the resulting supramolecular structures may be dictated
by the symmetry and periodicity of the substrate surface. In contrast to the assembly of supramolecular
architectures in three-dimensional (3D) bulk environments, e.g., solutions, where assembly is
governed only by specific and directional intermolecular interactions, a complex interplay between
adsorbate-adsorbate and adsorbate-substrate interactions controls the two-dimensional (2D)-phase
Surfaces 2018, 1, 12–28; doi:10.3390/surfaces2010003

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Surfaces 2018, 1

formation behavior when surfaces come into play [8–10]. Furthermore, if these molecular
building-blocks are deposited electrochemically in the form of ions, the electrochemical (EC) potential
becomes a very useful control parameter, which allows for influencing the self-assembly process [11–13].
In the end, the capability of many organic molecules to spontaneously self-assemble into molecular
monolayers on suitable solid substrates may lead to functional surfaces with broad applications in
electronic devices with nanometer dimensions, “green energy” catalysts, or novel light-stimulated
sensors [14,15].
Among other examples, porphyrins, due to the wide range of potential applications of this
class of molecules, e.g., in cancer therapy [16,17], catalysis [18,19], and sensing [20,21], have become
prototypical systems for the fabrication and design of supramolecular nanoarchitectures and the
characterization of their formation mechanism [22,23]. As an example from the point of applications,
tetra (N-methyl-4-pyridyl)-porphyrin p-tolylsulfonate molecules (H2 TMPyP, Figure 1) were proven to
be good sensors for the detection of benzene and heavy metal ions (Hg2+ , Pb2+ , Cd2+ ) in water [24].
From self assembling studies in ultra-high-vacuum (UHV), it is known that adsorbed porphyrins on
bare metal surfaces lie flat due to their large conjugated π-electron systems, which have the tendency
to maximize the π-bonding with the surface [25]. In the EC environment, however, it is expected
that electrostatic interactions of the porphyrin cations with the charged metallic electrode surface will
also have a strong influence. Moreover, when compared to an adsorbed porphyrin layer in UHV,
the electrode surface in contact with an electrolyte may be modified by the presence of specifically
adsorbed anions from the supporting electrolyte. The influence of these preadsorbed anions on the
deposition of organic cations depends obviously very much on their charge state. In the present
work the surface is modified by iodide anions, which are known to be almost uncharged, so that,
the iodine-porphyrin interactions are mainly dispersive in nature, i.e., of van der Waals type [26–28].
Actually, self-assembly of H2 TMPyP has been already investigated on various iodine modified metallic
surfaces, such as Au(111) [26,27], Pt(100), Ag(111) [29,30], Cu(111), and Cu(100) [11,28]. To the best
of our knowledge, studies on the self-assembly of porphyrins on the iodine precovered Au(100)
surface have not been reported yet. Here, we present the results of combined Cyclic Voltammetry
(CV) and in situ Scanning Tunneling Microscopy (STM) investigations. While the CV measurements
enable conclusions about the charge state of the adsorbed molecules, in particular, potentio-dynamic
STM measurements provide insight into their structural self-assembly on the surface. In this work,
we concentrate on the “non-reactive” potential regime, i.e., the potential regime in which the molecules
retain their oxidation state upon adsorption. Already in this restricted regime at positive potentials up
to five different potential-dependent long-range ordered phases of the H2 TMPyP molecules are found
on the iodine-modified Au(100) electrode. The different observed phases are thus not a consequence of
the variation of the redox-state of the molecules themselves, but only of the electrode potential (and
the concomitant change in iodine coverage). The potential induced modification of the non-covalent
interactions, namely van der Waals and electrostatic forces between the assembled molecules and
the substrate, is the only driving force for the observed phase transitions between the five porphyrin
structures in their constant oxidized form ([H2 TMPyP(0)]+4 ).

Figure 1. H2 TMPyP molecular structure.

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2. Materials and Methods
The results that are presented in this work were carried out using a home build in-situ EC
scanning tunneling microscope at constant current mode described in detail by Wilms et al. [31]. All of
the electrolytes were prepared by using deionized water from a Millipore® -Pure (Merck, Burlington,
MA, USA) water system (with a specific resistance of 18 MΩ·cm and a residual amount of organic
impurities in the ppb regime) and purged with suprapure argon gas for several hours before use.
Chemicals used (purchased from Sigma-Aldrich, Saint Louis, MO, USA) have the highest commercially
available quality level and were used without further purification. The STM tips were electrochemically
etched from a 0.25 mm tungsten wire in 2 M KOH solution and subsequently rinsed with high purity
water, dried, and coated by passing the tip through a drop of hot glue.
The Au(100) electrode (MaTeck Company, Juelich, Germany) was annealed before each
measurement in order to clean and smooth the surface. For this purpose, the sample was located on a
Ceran® plate and was annealed by a butane-oxygen gas flame for three minutes up to faint red glow
(600–700 ◦ C). Subsequently, the crystal was cooled down to room temperature (about 15 min) in an
atmosphere of argon. Then, it was heated again to a red glow and cooled for six minutes. To prevent
the re-contamination of the surface after the cleaning procedure, the crystal was covered by a protective
drop of Millipore® -water and immediately mounted into the EC cell of the EC-STM.
Cyclic Voltammetry (CV) and STM measurements were first performed in the pure supporting
aqueous 5 mM H2 SO4 + 1 mM KI electrolyte. In this electrolyte,
the iodine-modified Au(100)-surface
√
√ 
shows several phase transitions, starting from a
2 × p 2 at negative potentials, over a
√
√
√ 
√ 
2 × 2 2 to a pseudo-hex-rot structure at positive potentials. Both the
2 × p 2 and the
pseudo-hex-rot structure exhibit the phenomenon of reversible electro-compression [32–34]. For the
adsorption experiments the supporting electrolyte in the cell was replaced by a solution containing
sulfate, iodide, and H2 TMPyP porphyrin molecules (5 mM H2 SO4 + 1 mM KI + 0.01 mM H2 TMPyP).
The potential of the gold electrode was controlled with respect to a Pt/PtI quasi-reference electrode,
whose potential is related to that of the reversible hydrogen electrode (RHE) by the relation:
RHE = Pt/PtI + 580 mV.
The base vectors that were employed for the description of the ordered structures of the various
in the Results and Discussion section follow the common code:
→

•

a x for the base vector of the Au(100) lattice (x = 1,2);

•

bx and bx for the base vectors of the iodine lattice (x = 1,2); and,

•

c x , d x , ex , f x and gx for the base vectors of the porphyrin phases PI –PV (x = 1,2).

→

→

→ → → →

→

3. Results and Discussion
3.1. Cyclic Voltammetry and H2 TMPyP Structures
The structural investigations of the organic layers that are presented in this work have exclusively
been carried out under “non-reactive” conditions, i.e., in the regime where no redox-processes of
the porphyrin molecules take place, as revealed by the survey cyclic voltammogram (CV) of the
overall system shown in Figure 2a (see the red indicated region in Figure 2a). The first surface
reduction/re-oxidation step of the molecules is indicated by the peak couple Pred /Pox , as verified by
the following EC-STM measurements. Also, the adsorption/desorption of the iodine anions is marked
in the CV traces. As can be seen, the iodine desorption and the hydrogen evolution reaction (HER) are
retarded due to the presence of the molecules. Hence, we focus and place emphasis exclusively on the
structures of the [H2 TMPyP(0)]4+ adlayer and their changes merely driven by changes of the electrode
potential on the positive side of Pred /Pox .
It is well known that iodide anions adsorb specifically on the Au(100) surface, and even low
coverages stabilize the unreconstructed Au(100)(1 × 1) structure of the substrate. At high coverages
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the iodide anions form highly ordered layers, whose structure is potential dependent in terms of
electro-compression or electro-expansion [32–34]. However, within the electrode potential range
examined here, iodide forms a commensurate pseudo-hex-rot iodine structure on Au(100) [32–34],
and our results
indicate a pseudo-hex-rot iodine structure that is rotated by about 2.6◦ with respect to the
√
substrate 2 direction (Figure 2b), and consequently exhibits a periodic variation of the iodine binding
sites, resulting in the observed height, i.e., brightness, modulation. A ball model of this phase acquired
at an electrode potential of −150 mV vs. Pt/PtI is shown in Figure 2c, and it includes the unit cell.

Figure 2. (a) Cyclic voltammograms of the Au(100) surface in 5 mM H2 SO4 + 1 mM KI electrolyte
solution (black curve) and in the H2 TMPyP containing electrolyte (blue curve, 5 mM H2 SO4 + 1 mM
KI + 0.05 mM H2 TMPyP), dE/dt = 10 mV·s−1 ; GDR = gold dissolution reaction, HER = hydrogen
evolution reaction, and the red bars indicate non-reactive regimes; (b) high-resolution in-situ Scanning
Tunneling Microscopy (STM) image of the pseudo-hex-rot iodine structure obtained from a Au(100)
surface in 5 mM H2 SO4 + 1 mM KI solution; image parameters: 13.73 nm × 13.73 nm, It = 1 nA,
√
Ub = −93 mV, E = −150 mV vs. Pt/PtI; (c) ball-model of the (4 × 221) iodine structure shown in (b).

Exposure of this iodine-modified Au(100) surface to the H2 TMPyP molecules containing 1 mM
KI + 5 mM H2 SO4 solution at potentials between the oxidative gold dissolution reaction (GDR)
and the first reduction step of the porphyrin species (Pred ) leads to spontaneous adsorption and the
subsequent lateral ordering of H2 TMPyP molecules at the surface. The solution contains actually
[H2 TMPyP]4+ cations, but for the sake of ease we simply refer to H2 TMPyP molecules throughout this
paper. The accompanying p-tolylsulfonate- as well as the SO4 2− -anions from the supporting electrolyte
adsorb weaker than iodide anions and therefore do not displace the latter [34].

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Figure 3a presents a highly ordered layer of H2 TMPyP at an electrode potential −200 mV vs.
Pt/PtI. Each bright dot represents one porphyrin molecule. Close-ups of this layer (Figure 3b–d) reveal
that the molecules are self-assembled into differently ordered domains of molecular rows covering
the atomically flat terraces. In addition, step edges (panel a) are oriented parallel to and decorated by
molecular rows (see white and yellow lines in Figure 3a). A careful examination of Figure 3a–d clearly
indicates the coexistence of five different porphyrin phases (PI , PII , PIII , PIV , and PV ). Within each phase,
translational domains (Figure 3b,c molecular rows of the domain highlighted by yellow solid lines)
and the respective domain boundaries between them (Figure 3b,c yellow dashed lines) are observed.

Figure 3. Electrochemical Scanning Tunneling Microscopy (EC-STM) images of self-assembled
H2 TMPyP molecules on an iodine-modified Au(100) electrode in the non-reactive regime. PI –PV refer to
structurally different phases. (a) 83.55 nm × 83.55 nm, It = 1 nA, Ub = −320 mV, E = −200 mV vs. Pt/PtI;
(b) 42.35 nm × 42.35 nm, It = 1 nA, Ub = −472 mV, E = −200 mV vs. Pt/PtI; (c) 56.47 nm × 56.47 nm,
It = 1 nA, Ub = −491 mV, E= −200 mV vs. Pt/PtI; (d) 56.47 nm × 56.47 nm, It = 1 nA, Ub = −303 mV,
E = −200 mV vs. Pt/PtI.

In order to shed light on these porphyrin phases, STM images with atomic resolution were
registered (Figure 4a–d). These images clearly reveal that the organic macrocycles are lying flat on the
substrate due to their large molecular π-electron system and the tendency to maximize the π-bonding
to the surface, as reported before [11,26–28]. The STM images show that each flat-lying porphyrin
molecule can be recognized as a square-shaped motif with the characteristic four additional lobes
placed at the four corners of each square (see inset and blue molecules in Figure 4a). The center-to-center
distance measured diagonally across one molecule between the corner spots is 1.2 ± 0.1 nm, which is
consistent with the distance between two diagonally located pyridinium units [26]. Furthermore,
a detailed analysis of images, like in Figure 4, yields the characteristic lattice parameters, like vectors,
rotation angles, molecules per unit cell, and surface coverage of the five different phases (PI –PV ).
These structural data are collected in Table 1.

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Surfaces 2018, 1

Figure 4. High-resolution EC-STM images of the ordered H2 TMPyP adlayer on an iodine-modified
Au(100) electrode in the non-reactive regime. (a) 14.12 nm × 14.12 nm, It = 1 nA, Ub = −472 mV,
E = −200 mV vs. Pt/PtI. Size of inset: 1.94 nm × 1.94 nm; (b) 14.15 nm × 14.15 nm, It = 1 nA,
Ub = −291 mV, E = −200 mV vs. Pt/PtI; (c) 14.12 nm × 14.12 nm, It = 1 nA, Ub = −472 mV,
E = −200 mV vs. Pt/PtI; (d) 14.12 nm × 14.12 nm, It = 1 nA, Ub = −303 mV, E= −200 mV vs. Pt/PtI.
Table 1. Structural data of the five different self-assembled, potential dependent H2 TMPyP structures
PI –PV on an iodine-modified Au(100) electrode surface as detected by in situ STM in the non-reactive
potential regime.
Parameter

PI

PII

PIII

PIV

PV

lattice constants
[±0.05 nm]

c1 = 1.61
c2 = 1.80

d1 = 1.94
d2 = 1.56

e1 = 1.61
e2 = 2.06

f1 = 1.94
f2 = 1.76

g1 = 1.46
g2 = 2.06

angle
[±1◦ ]

92◦

74◦

89◦

77◦

85◦

molecule rotation
[α ±1◦ ]

+30◦ vs. c1

−28◦ vs. d1

+32◦ vs. e1

+26◦ vs. f 1

+27◦ vs. g1

surface concentration
[molecules/cm2 ]

3.453 × 1013

3.437 × 1013

3.016 × 1013

3.014 × 1013

3.333 × 1013

→

→

→

→

→

Based on these experimental results of the symmetry and orientation of the H2 TMPyP molecules,
in Figure 5 we propose a schematic model of the ordered H2 TMPyP porphyrin molecules for each
phase. It is important to emphasize that the value of the surface coverage of phase PV lies in the middle
of those of the other phases. Furthermore, PV was observed rather rarely in comparison to phases
PI –PIV : it only appears when the adsorption process started at an electrode potential of ca. −200 mV
vs. Pt/PtI, and it vanishes altogether at those potential values where the phases PI –PIV were observed
within more extended regimes of the electrode potential (vide infra). This suggests that Pv is less stable
than PI –PIV and it is observable only under specific conditions in this very narrow electrode potential
window, e.g., after adsorption at −200 mV vs. Pt/PtI and within a narrow range of electrode potential.
As mentioned previously, within the potential window investigated here, the porphyrin molecules
do not undergo any redox process. Hence, the observed phase transitions are not triggered by a
248

Surfaces 2018, 1

variation of the chemical state of the porphyrin due to an electron transfer process; but it can only be
traced back to changes of the interactions between the porphyrin molecules and the substrate [11,35,36].

Figure 5. Schematic models of the five ordered H2 TMPyP molecular phases in the non-reactive regime
on I/Au(100), with the denoted vectors of the unit cells, the angles between them, and the angle of
rotation of individual molecules off the molecular row direction (see Table 1).

3.2. Electrode Potential Stability of H2 TMPyP Phases on I/Au(100)
In order to check the stability and existence of phases PI –PV as a function of potential, and whether
there is a competition (i.e., a difference in stability) between them expressed by their different
equilibrium coverages, we also performed potentio-dynamic STM studies. These potentio-dynamic
STM measurements were started at −200 mV vs. Pt/PtI (see Figure 2) with a highly ordered porphyrin
layer comprising the first four phases PI –PIV , as shown in Figure 6a). The following images in panels
b–i of Figure 6 are a selection out of a series of 170 successive images and are taken at the indicated
potentials. They are largely registered at the same surface position (indicated by the red arrow in
Figure 6a–i). However, due to thermal drift this marker slowly drifted out of the scanned area so
that the last three images 6j–l were recorded around a new point of the same area (marked with a
light blue arrow). From the starting point (Figure 6a reflects the very initial state, which was allowed
to equilibrated for 210 s before the series of 170 images was started with Figure 6b), the electrode
potential was first scanned in the cathodic direction. The areas covered by PI and PII are observed to
shrink while those of PIII and PIV start to grow. When the potential reaches −300 mV vs. Pt/PtI, within
less than 5 min, the phases PI and PII are finally completely replaced by PIII and PIV (see Figure 6c,d).
Moreover, new bright spots appear in the images (see white arrows in Figure 6c), which represent the
growth of a completely new ordered phase. A height profile measured along the short orange line
in Figure 6d and shown in the inset yields a height difference between the bright spots and the dark
region between them of around 0.3 nm. This value approximately matches the thickness of one flat
lying porphyrin molecule [37].
These observations suggest that, when approaching the reactive regime (below −300 mV vs.
Pt/PtI), the new bright spots represent adsorption of molecules in the second layer. This is supported
by lowering of the electrode potential even further to −300 mV and −400 mV, which leads to an
increase of the bi-layer coverage and the creation of new phases under these reactive conditions,
named PVI and PVII (Figure 6d–f). Both monolayer phases PIII and PIV phases are finally completely
replaced by the bi-layer when the potential reaches the maximum of peak Pred at −395 mV vs. Pt/PtI,

249

Surfaces 2018, 1

which represents the first reduction step of the porphyrin molecules. Thus, the appearance of the
bi-layer in the STM images in coincidence with the first reduction peak in the CV clearly marks the
cathodic limit of the non-reactive monolayer adsorption regime. The full analysis of the EC-STM
data under reactive conditions will be the subject of a forthcoming paper. Here, we concentrate on
changes on the surface upon sweeping the potential back into the positive direction. No changes where
noticed until the electrode potential reached again −300 mV vs. Pt/PtI. At this potential, re-entering
in the non-reactive regime, a disordering of the lateral structure at the domain boundaries as well as
desorption of second layer porphyrin molecules starts to take place (see yellow arrows in Figure 6g).
At the potential −250 mV vs. Pt/PtI (corresponding to the anodic peak Pox ) the second layer has
almost completely disappeared, and at a potential of ca. −200 mV vs. Pt/PtI, represented by Figure 6i,
phases PI and PII have reappeared on the surface. Changing the electrode potential even further in
anodic direction leads to growth of PI and PII at the expense of PIII and PIV (see Figure 6j,k), and finally
a complete replacement of the latter two phases. Moreover, at a potential of −100 mV vs. Pt/PtI the
GDR starts at the step-edges as revealed by the course of the step edge in Figure 6k. After reaching
the electrode potential of −50 mV vs. Pt/PtI where the GDR proceeds very quickly (see Figure 6l),
PII completely disappears from the surface, while PI is still clearly observable.

Figure 6. Cont.

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Surfaces 2018, 1

Figure 6. Potential induced phase transition of molecular H2 TMPyP adlayer on an I/Au(100) electrode
surface, STM series: 50.32 nm × 50.32 nm; (a) It = 1 nA, Ub = −291 mV; (b) It = 1 nA, Ub = −267 mV;
(c) It = 1 nA, Ub = −267 mV; (d) It = 1 nA, Ub = −267 mV; (e) It = 1 nA, Ub = −208 mV; (f) It = 1 nA,
Ub = −208 Mv; (g) It = 1 nA, Ub = −208 mV; (h) It = 1 nA, Ub = −208 mV; (i) It = 1 nA, Ub = −250 mV;
(j) It = 1 nA, Ub = −477 mV; (k) It = 1 nA, Ub = −461 mV; (l) It = 1 nA, Ub = −461 mV. For explanation
of the different colored arrows see the text.

Figure 7 represents a diagram that indicates the stable potential and coexistence regimes of the
five phases PI –PV . The above potentio-dynamic STM measurements reveal a competition between
these phases, i.e., their relative stability as a function of electrode potential, as manifested by their
different surface concentrations. Furthermore, from a correlation of the structural changes with the
relative coverage of each phase, we arrive at a strict relationship between the potential regime of
each phase and its surface concentration: according to Table 1, PI has the highest surface density
(3.453 × 1013 molecules/cm2 ) and its existence regime is at the most positive potentials in comparison
to all other phases. Next, PII with a lower density (3.437 × 1013 molecules/cm2 ) than PI occurs at lower
potentials. Finally, PIII and PIV with the lowest surface concentrations among all five monolayer phases
(of 3.016 × 1013 molecules/cm2 and 3.006 × 1013 molecules/cm2 , respectively) have their existence
regimes at the most negative potentials.
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Surfaces 2018, 1

Figure 7. Diagram of the stability regimes of the PI –PV monolayer phases of adsorbed H2 TMPyP on
an iodine-modified Au(100) electrode surface.

Therefore, in order to get more insights and to understand the relations between the applied
electrode potential and the different surface concentration of each porphyrin phase, a set of STM
images was taken, which provides a direct correlation between the structure of the different organic
phases and that of the underlying iodine and gold lattice.
3.3. Substrate-Adlayer Structure Correlations
3.3.1. Phases PI , PIII
Owing to the overlap of the existence regimes, it is possible to record a set of STM images that
comprise both phases PI and PIII (14.12 nm × 14.12 nm, Figure 8a,b and Figure 8d,e) within the same
surface area. It is well known that by changing the tunneling conditions (namely tunneling bias) a
resonant tunneling may be achieved [38,39]. At high bias voltage (e.g., Ub = −416 mV in Figure 8a,
and Ub = −472 mV in the lower half part of Figure 8d), almost all electrons flow via resonant tunneling
through quantized energy levels that are defined by the molecular orbitals or the band generated
by the periodic arrangement of the molecules. Thus, tunneling through these states emphasizes the
contribution of the organic molecules in the respective STM images. On the other hand, at lower bias
voltage (e.g., Ub = −9 mV in the upper-left part of Figure 8d and left part of Figure 8e), the normal
tunneling process dominates the tunneling current, where electrons tunnel through the potential well
without interacting with the localized discrete molecular levels. Most importantly, in some cases at an
intermediate bias voltage it is possible that both tunneling channels through the potential well and the
localized discrete levels contribute to the image. Thus, images that are registered at this intermediate
bias voltage comprise both types of features that can be associated either with the substrate or the
adsorbate (see e.g., Figure 8b at Ub = −13 mV, and its Reverse Fast Fourier Transform (RFFT) in
Figure 8c in which we can observe a signal from the porphyrin molecules as well as from the iodine
lattice; and the upper-right part of Figure 8d and right part of Figure 8e at Ub = −9 mV, with the RFFT
presented in Figure 8f performed on the right part of Figure 8e, in which contributions from gold and
the porphyrin lattice are observed). Such images enable a direct correlation between the structure of
the organic overlayer and that of the underlying iodine and the supporting gold lattice. Moreover,
the simultaneous observation of features from both the molecules on the one hand, and iodine or gold
underneath on the other hand, excludes that the observed bias dependence of the images is due to
a physical removal of the molecules from the substrate by the tip at low bias voltages (Figure 8b,e).

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Surfaces 2018, 1

This approach circumvents any problem with drift during image taking because both contributions,
that from the molecules and that from the substrate, are included in the same image.

Figure 8. Correlation of the structures of the PI and PIII monolayer phases of adsorbed H2 TMPyP
molecules with the underlying iodine and gold layer, (a) PI porphyrin layer at high tunneling bias,
14.12 nm × 14.12 nm, It = 1 nA, Ub = −416 mV, E = −200 mV vs. Pt/PtI; (b) Iodine layer underneath PI
at low tunneling bias, 14.12 nm × 14.12 nm, It = 1 nA, Ub = −13 mV, E = −200 mV vs. Pt/PtI;
(c) a Reverse Fast Fourier Transform (RFFT) of the STM image (b); (d) upper part: Iodine layer
underneath PIII at low tunneling bias: Ub = −9 mV, lower part: PIII porphyrin layer at high tunneling
bias: Ub = −472 mV, 14.12 nm × 14.12 nm, It = 1 nA, E = −200 mV vs. Pt/PtI; (e) left part: Iodine layer
underneath PI at low tunneling bias: Ub = −9 mV, right part: combination of the Porphyrin layer and
the substrate layer at low tunneling bias: Ub = −9 mV, 14.12 nm × 14.12 nm, It = 1 nA, E = −200 mV vs.
Pt/PtI; (f) a Reverse Fast Fourier Transform (IFFT) of the left part of STM image (e).

Therefore, after close analysis and superimposition of the signals that are described above from
the porphyrin-, iodine-, and Au-structure presented in Figure 8, it is possible to relate the structure of
the molecular phases PI and PIII to those of the iodine and the gold lattice underneath (see Figure 9).
The iodine layer under both molecular phases was found to be a pseudo-hex-rot phase, but when
considering that this iodine phase is electro-compressible (the iodine lattice becomes more compressed
with increasing electrode potential), a differently compressed pseudo-hex-rot iodine phase should exist,
and was actually found, under the two H2 TMPyP phases. With the set of images from Figure 8 in
hand, it is thus possible to devise structure models of both phases PI and PIII , and the corresponding
data are summarized in Table 2.
Starting with PI (Figure 9), we detect that in the direction of the vectors b1 every nineteenth and in
the direction of the vector b2 every seventh iodine atom in the respective iodine row occupies a site on
top of a gold atom. On the other hand, the molecular rows of the H2 TMPyP layer of PI are not aligned
with any high symmetry direction of the substrate, neither of the iodine lattice nor of the Au(100)
surface. However, when considering the symmetry axis of individual molecules that are centred atop
of an iodine atom, we observe that this axis is rotated by 30◦ off the row direction (see Table 1) and
coincides with close-packed iodine rows running in the [0 7 35] direction of the substrate, respectively.

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Surfaces 2018, 1

On the basis of this model the surface coverage of the porphyrin adlayer is calculated to be 0.0286
monolayers (ML) relative to the density of the gold layer.

Figure 9. Structure models of the phases PI and PIII on an iodine modified Au(100) surface.

In turn, for phase PIII , the iodine rows follow the direction of the vectors b1 and b2 instead,
which in these directions leads to periodic atop binding sites of an iodine atom on a gold atom every
twenty-first and forty-first iodine, respectively. The symmetry axis of individual molecules is rotated
by 32◦ off the row direction (see Table 1), which coincides with the [0 9 37] substrate direction in this
case. As a consequence, the iodine coverage under PIII (0.250 ML) is slightly lower than that under PI
(0.0286 ML). This lower iodine coverage is a consequence of the electro-expansion of the pseudo-hex-rot
phase with decreasing electrode potential (0.52 ML under PI vs. 0.50 ML under PIII , see Table 2).

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Surfaces 2018, 1

Table 2. Data correlating the structures of the phases PI , PIII and PIV of self-assembled H2 TMPyP
molecules on an iodine modified Au(100) electrode with those of the respective iodine layer and the
Au(100) substrate underneath.
Parameter
Au lattice
parameters

vectors
direction
lattice
constant

Iodine lattice
parameters

vectors
direction
lattice
constant
Θ*
Matrix with
respect to
Au(100)

Porphyrins
lattice
parameters

vectors
direction
Θ*
Matrix with
respect to
iodine
Matrix with
respect to
Au(100)

PI

PIII

PIV

a1 = [0 1 1]
a2 = [0 1 1]
a1 = 0.289 nm
a2 = 0.289 nm

⎛ →
⎝ b1
→

b2

→
c1
→
c2
→
c1
→
c2

b1 = [0 7 35]
b2 = [0 11 5]
b1 = 0.404 nm
b2 = 0.410 nm
0.52 ML
⎞


3 −2
⎠=
3
8

→



a1
a2

→

c1 = [0 1 3]
c2 = [0 17 5]
0.0286 ML ⎛
⎞
 
 →
3 −2 ⎝ b 1 ⎠
=
→
3
4
b2
 

→ 
2.5 −5
a1
=
→
5.5
3
a2

⎛ →

⎝ b1
→
b

2

→
e1
→
e2
→
e1
→
e2

b1 = [0 9 37]


b2 = 0 77 39
b 1 = 0.388 nm

b 2 = 0.439 nm
0.50 ML
⎞


11.5 −7
⎠=
19
58

→



a1
a2

→



e1 = 0 1 3
e2 = [0 19 7]
0.0250 ML ⎛
⎞
 
 →
3 −2 ⎝ b  1 ⎠
=
→
4
4
b 2
 

→ 
2.5 −5
a1
=
→
6.5
3
a2

⎛ →

⎝ b1
→
b

2

b1 = [0 9 37]


b2 = 0 77 39
b 1 = 0.388 nm

b 2 = 0.439 nm
0.50 ML
⎞


11.5 −7
⎠=
19
58

→



a1
a2

→

f 1 = [0 61 73]
f 2 = 0 77 39
0.0250 ML ⎛
⎞
→  
 →
5 2 ⎝ b 1 ⎠
f1
=
→
0 4
f2
b 2
⎛ → ⎞


→ 
33.5 −3
f1 ⎠
a1
⎝ 5→
=
→
9.5
29
a2
5f
2

* Surface coverage relative to the density of the gold layer.

Even though the determined differences in coverage of PI and PIII , and, even more so, for the
respective underlying iodine structures, may appear small, there is no doubt that they are real,
because they are obtained from the coexisting phases in the same image of the same measurement.
Thus, PI with a monolayer coverage of 0.0286 ML is situated on an iodine layer of coverage
0.52 ML, while PIII of coverage 0.0250 ML is found on an iodine underlayer of coverage 0.50 ML
at lower potentials. In both phases, the direction of closest-packed molecular rows do not coincide
with any high symmetry direction of neither the iodine nor the gold lattice underneath, but all of the
molecules are centred on an iodine atom, and a symmetry axis of the molecules is aligned with the
direction of a close-packed iodine row. Since both phases, PI and PIII , occur within the non-reactive
potential regime, and thus, retain their charge state, only differences in density and atomic arrangement
of the iodine underlayer can be taken responsible for the differences between PI and PIII, namely their
different lattice parameters, and, in particular, their different atomic density.
Two arguments can be put forward to explain the higher coverage of PI when compared to PIII .
Firstly, the denser iodine anion layer underneath PI (at higher potential) may correspond to a higher
density of negative charge, which attracts more molecular cations. This argument, however, may not
be valid because among the halides iodine binds most covalently, in particular the higher its coverage
(at higher potentials) is. Iodine layers on noble metals, like Au or Pt, were found to be hydrophobic
and almost uncharged [26]. The same conclusion was drawn from studies with iodine adsorption
on Cu(100) [40,41]. Secondly, the denser the iodine layer the more effectively it screens the positive
charge of the metal substrate, and thereby, reduces the electrostatic repulsion between the electrode
and the molecular cations, and, instead, relatively strengthens van-der-Waals interactions between
the molecules and the covalently bound iodine layer. Even though this dominance of van-der-Waals
interactions has been favoured in the literature, our results indicate a combination of van-der-Waals
and electrostatic forces. While the simultaneous increase of iodine- and H2 TMPyP-coverage with
increasing electrode potential supports the notion of dominant van-der-Waals bonding of the molecules
to the surface, the match of the symmetry direction of the individual molecules with the nearly
rectangular iodine lattice underneath points to a contribution of electrostatic interactions, namely
interlayer attraction between the positively charged N-methyl-4-pyridyl-ligands and the iodine lattice
underneath, and intralayer repulsion between these ligands of the neighbouring molecules. Therefore,
these latter electrostatic forces are actually responsible for the lateral order formation.

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Surfaces 2018, 1

3.3.2. Phases PIII , PIV
A similar comparative analysis can be done with phases PIII and PIV based on images, as shown
in Figure 10. According to Table 1, these two phases have essentially equal surface coverage and
they are stable in the same potential range, as revealed by the potentio-dynamic STM measurements
shown in Figure 6. In order to achieve a structural correlation between the molecular adlayer and
the substrate underneath, we performed again bias dependent STM measurements (Figure 10a–c).
As listed in Table 2, the pseudo-hex-rot structure of the iodine structure on the substrate is the same


under both phases, and is described by the vectors b1 and b2 with directions of [0 9 37] and 0 77 39 ,
respectively. In these directions, every twenty-first and forty-first iodine atom, respectively, occupies
an atop binding site on a gold atom, and the iodine coverage is 0.50 ML (see Table 2). While the
structural correlation of phase PIII with the iodine lattice underneath has already been described in
the previous section, molecular rows of phase PIV described by f 1 and f 2 with directions [0 61 73]


and 0 77 39 , respectively, run parallel to the iodine b2 direction, but, interestingly, have the same
orientation of the molecular symmetry axis as in PIII (coinciding with the [0 9 37] direction of close
packed iodine rows), as shown in Figure 11a. Therefore, if both unit cells of PIII (blue) and PIV (red)
are superimposed (as shown in Figure 11b) a new coincidence lattice (S1 ,S2 ) (green in Figure 11b)
becomes apparent.
 The vector
 S1 of this structure is aligned with the densely packed iodine rows
in the direction 0 9 37 . The new (green) unit cell contains four molecules from PIII and four
molecules from PIV , plus one shared additional molecule (the green atom can be either blue or red in
the model) every eleventh iodine atom underneath. The fact that both phases PIII and PIV have the
same coverage (0.025 ML), means that both phases adsorb with equal probability on the surface at this
potential regime.

Figure 10. Structure correlation between the H2 TMPyP layer of PIII and PIV and the underlying iodine
and gold layer, (a) Iodine layer underneath at low tunneling bias, 14.12 nm × 14.12 nm, It = 1 nA,
Ub = −17 mV, E = −200 mV vs. Pt/PtI; (b) upper part: Iodine layer underneath at low tunneling bias:
Ub = −17 mV, lower part: Porphyrin layer at high tunneling bias: Ub = −472 mV, 14.12 nm × 14.12 nm,
It = 1 nA, E = −200 mV vs. Pt/PtI; (c) Porphyrin layer at high tunneling bias, 14.12 nm × 14.12 nm,
It = 1 nA, Ub = −472 mV, E = −200 mV vs. Pt/PtI.

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Surfaces 2018, 1

Figure 11. Structural model of PIII and PIV on iodine modified Au(100) surface (a); structure model of
superstructure created by superimposing of Phase III and IV (b).

4. Conclusions
In this study, we have investigated in situ the self-assembly of H2 TMPyP molecules on an iodine
modified Au(100) electrode surface by EC-STM methods exclusively in the non-reactive regime, i.e., in
the potential regime in which the molecules retain their oxidized [H2 TMPyP(0)]4+ state throughout.
The exposure of the pseudo-hex-rot iodine modified Au(100) surface to the porphyrin containing solution
results in the spontaneous adsorption and lateral ordering of the molecules on the surface, with the
organic macrocycles lying flat on the substrate. Up to five different porphyrin phases (PI –PV ) were
detected on the surface of the substrate, whose stabilities, surface concentrations, and existence regimes
were demonstrated yet to be potential dependent. In fact, potentio-dynamic STM measurements
revealed that PI and PII , stable structures at more positive potentials, transform into phases PIII and
PIV by applying more negative electrode potential, while PV was found to be stable and detected in a
very narrow potential window only. Therefore, there is a competition between the different phases,
resulting in a strict relationship between their respective surface concentrations and their existence
regimes as a function of the applied electrode potential.
In addition, by applying adequate imaging conditions, i.e., bias voltages, STM images were
obtained, which enable a correlation between the structures of the porphyrin adlayer and of the
underlying iodine and gold lattices. The lattice parameters regarding the relative orientation of the
iodine underlayer and the organic overlayer for PI (at more positive potentials) on the one hand, and for
PIII and PIV (at more negative potentials) on the other hand were analysed. The results reveal that
the observed phase transitions are accompanied by a change of the underlying electro-compressible
iodine lattice, in particular, its density, as a function of the electrode potential. Since iodide anions
are largely discharged upon adsorption and essentially bound covalently as neutral iodine atoms,
it has been suggested in the literature [27,29,30] that porphyrin molecules on iodine modified metal
surfaces are predominantly controlled by van der Waals forces. However, our results about potential
dependent lattice coincidences and molecular orientations strongly support the notion that not only the
van-der-Waals forces are responsible for the self-assembly of the porphyrin molecules on the polarized
substrate, but that also electrostatic interactions between the molecules and the iodine modified metal
substrate, as well as between the molecules themselves, play a decisive role in the 2D ordering process.

257

Surfaces 2018, 1

Author Contributions: All authors have made a substantial contribution to the work, and approved it for publication.
Funding: This work has in part been funded by the collaborative research centre 624 of the German
Science Foundation.
Acknowledgments: M.B. gratefully acknowledges the Clarin CoFund Program postdoctoral fellowship
(ACA17-29), funded by Gobierno del Principado de Asturias and Marie Curie Actions (grant 600196).
Conflicts of Interest: The authors declare no conflict of interest.

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259

Article

Methanol Oxidation on Graphenic-Supported
Platinum Catalysts
Gladys Arteaga, Luis M. Rivera-Gavidia, Sthephanie J. Martínez, Rubén Rizo, Elena Pastor * and
Gonzalo García *
Área de Química Física, Departamento de Química, Facultad de Ciencias, Universidad de La Laguna (ULL),
Instituto de Materiales y Nanotecnología (IMN), 38200 La Laguna, Tenerife, Spain;
alu0100846713@ull.edu.es (G.A.); mlriviera@gmail.com (L.M.R.-G.);
sthephanie94@hotmail.com (S.J.M.); rjrizo57@gmail.com (R.R.)
* Correspondence: epastor@ull.edu.es (E.P.); ggarcia@ull.edu.es (G.G.);
Tel.: +34-922-318-071 (E.P.); +34-922-318-032 (G.G.)
Received: 19 November 2018; Accepted: 29 December 2018; Published: 6 January 2019

Abstract: Graphene oxide (GO), reduced graphene oxide by thermal treatment (rGO-TT),
nitrogen-modified rGO (N-rGO), and carbon Vulcan were synthesized and employed in the current
work as catalyst support for Pt nanoparticles, to study their properties and impact toward the
methanol oxidation reaction (MOR) in sulfuric acid medium. Several physicochemical techniques,
such as X-ray photoelectron spectroscopy (XPS), X-ray powder diffraction (XRD), Transmission
electron microscopy (TEM), energy-dispersive X-ray spectroscopy (EDX), Raman, and elemental
analysis were employed to characterize the novel materials, while potentiodynamic and potentiostatic
methods were used to study catalytic performance toward the methanol oxidation reaction in acidic
medium. The main results indicate a high influence of the support on the surface electronic state of
the catalyst, and consequently the catalytic performance toward the MOR is modified. Accordingly,
Pt/N-rGO and Pt/rGO-TT show the lowest and the highest catalytic performance toward the
MOR, respectively.
Keywords: electrocatalysis; catalysts; methanol oxidation reaction; graphene; DMFC; Pt

1. Introduction
The search for alternatives to the use of fossil fuels, and the study and development of new green
technologies for energy supply is an imperative necessity in order to fight global warming. Direct
methanol fuel cells (DMFC) are an example of these energy-conversion devices [1–8]. They are able
to supply energy from the electro-oxidation of a cheap resource like methanol, resulting not only in
the reduction of release of harmful emissions, but also in the supply of high energy densities at low
temperatures. The design of the catalysts to be used in DMFC systems is of outstanding importance.
Many studies have confirmed that platinum is the most active metal when employed as a catalyst for
the methanol oxidation reaction (MOR) [9–11].
However, the scarcity of this metal in the world and its high cost hinder the commercialization of
such systems. In order to reduce as much as possible the amount of the catalytic material, the synthesis
of Pt nanoparticles with high surface area/volume ratios, as well as the use of catalyst supports for
a good dispersion of the metallic nanoparticles, is essential [12–14]. Furthermore, a high electrical
conductivity of the catalyst support is the key factor for electrochemical reactions like the MOR. In this
sense, graphene has emerged as a new-generation catalyst support, because of its excellent electrical
conductivity and high surface area [15–18]. Many works have described chemical exfoliation methods
of graphite to obtain so-called graphene oxide (GO) as a catalyst support. However, the conductivity
of such materials is lower than the highly desirable two-dimensional graphene, because of the high
Surfaces 2019, 2, 2; doi:10.3390/surfaces2010002

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Surfaces 2019, 2, 2

amount of oxygen groups on the surface. To fulfill conductivity problems, a great diversity of physical
and chemical methods have been described in the literature for GO reduction to obtain reduced GO
(rGO), with higher conductivity than GO itself [15–18]. On the other hand, it has been demonstrated
that the surface chemistry of the catalytic support plays an essential role in the activity of the catalyst
toward the MOR [19–23]. For example, the modification of carbon materials with heteroatoms, like
nitrogen or sulfur, causes an electronic modulation of the carbonaceous lattice, resulting in a change in
the electrocatalytic properties toward the MOR and other reactions of interest [15,18,23–25].
In the present paper, the synthesis of reduced graphene oxide (rGO) by a thermal treatment
(rGO-TT) and nitrogen-modified rGO by the caffeine route (N-rGO) is reported, both of which are
employed as supports for Pt nanoparticles. A deep physicochemical characterization of the novel
catalysts was performed, and their catalytic performance toward the MOR was studied and compared
with that of Vulcan XC72R supported Pt catalyst (Pt/C).
2. Materials and Methods
Sulfuric acid (p.a.; Merck, Darmstadt, Germany), methanol (p.a.; Merck, Darmstadt, Germany),
potassium permanganate (>99.8%; Sigma-Aldrich, Saint Louis, MO, USA), caffeine (>99%; Sigma-Aldrich,
Saint Louis, MO, USA), Vulcan XC-72R (Cabot, Boston, MA, USA), graphite (>99.8%; Sigma-Aldrich,
Saint Louis, MO, USA), hydrogen peroxide (30% v/v; Foret, Barcelona, Spain), hydrogen (5% H2 /95%
N2 ; Air Liquide, Tenerife, Spain), (8 wt % H2 PtCl6 ·6H2 O, Sigma-Aldrich, Saint Louis, MO, USA), Nafion
solution (5 wt %; Sigma-Aldrich, Saint Louis, MO, USA), Ar (99.999%, Air Liquide, Tenerife, Spain), CO
(99.997%, Air Liquide, Tenerife, Spain), and water (18.2 M Ω·cm−1 ; Milli-Q, Millipore, Burlington, VT, USA)
were purchased and then used for the synthesis of the graphene-based materials and the preparation of
electrolyte solutions.
2.1. Synthesis of Graphene Oxide
Graphene oxide (GO) was prepared by following a modified Hummers method [15]. Briefly, 1 g
of the graphite powder was mixed with 30 mL of concentrated H2 SO4 cooled in an ice bath. Then,
3.5 g of KMnO4 are slowly added while being stirred and cooled continuously. After removal from
the ice bath, the mixture was diluted with Milli-Q water under stirring for 1 h at 35 ◦ C. Then, the
solution was heated up to 95–98 ◦ C over 30 min. Next, 200 mL of ultrapure water was gradually
introduced, followed by 1.25 mL of 30% v/v H2 O2 , and stirring was maintained for 30 min. Finally,
the dispersion was centrifuged with Milli-Q water, until a pH of 7 was achieved in the supernatant
liquid. The material was dried using an oven at 60 ◦ C overnight.
2.2. Synthesis of Reduced Graphene Oxide Materials
In order to produce N-rGO, an adequate amount (4 mmol) of reducing agent (caffeine) was
ultrasonically dispersed in Milli-Q water and mixed with a GO aqueous dispersion (0.015 g mL−1 ).
The final dispersion was placed into a Teflon-lined autoclave and heated at 160 ◦ C for 10 h. Then,
N-rGO was washed by centrifugation, employing Milli-Q water, and transferred to an oven at
60 ◦ C for 24 h.
The rGO-TT was obtained by reduction of GO through a heat treatment under a reductive atmosphere
(5% H2 /95% N2 at 100 mL/min) in a tubular furnace. The thermal treatment comprised a ramp
temperature of 5 ◦ C/min from room temperature, up to 200 ◦ C. This temperature was maintained
for 30 min; finally, a ramp temperature of 5 ◦ C/min was applied up to 450 ◦ C where it was preserved for
1 h.
2.3. Synthesis of Pt/C, Pt/N-rGO, and Pt/rGO Catalysts
Platinum nanoparticles supported on carbon black Vulcan XC-72R, N-rGO, and rGO-TT were
synthesized following the formic acid method [6]. Briefly, catalyst support was dispersed in 2 M
formic acid solution under sonication. Then, the mixture was heated at 80 ◦ C under stirring, and
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an appropriate amount of metal precursor (H2 PtCl6 ) was employed to obtain theoretical platinum
loading of 20 wt % in the catalyst. The suspension was filtered, washed with ultrapure water, and
dried at 60 ◦ C overnight.
2.4. Physicochemical Characterization
Powder XRD spectra were acquired from X’Pert PRO X-ray diffractometer (PANalytical, Madrid,
Spain) to determine the crystal structure. CuKα radiation (λ = 1.5405 Å) was employed, and 2θ data
were collected from 20◦ to 100◦ with a scanning rate of 0.04 s−1 . Crystalline phases were identified
by comparing the experimental diffraction patterns with the Joint Committee on Powder Diffraction
Standards (JCPDS).
Elemental analysis with an experimental error close to 0.04% was performed using an Elemental
Analyzer CNHS FLASH EA 1112 (Thermo Scientific, Madrid, Spain).
Raman spectra were collected employing a RENISHAW confocal Raman microscope, model
inVia (RENISHAW, Gloucestershire, United Kingdom), with a green laser (λ = 532 nm) in the 100 to
3200 cm−1 range.
Transmission electron microscopy (TEM) images were obtained from a HRTEM JEOL JEM 2100
operating (JEOL Ltd., Tokyo, Japan) at 200 eV. Obtained images were used to evaluate the morphology
and agglomeration degree, as well as to calculate the average Pt nanoparticle sizes.
Chemical composition of the catalysts was determined by energy dispersive X-ray analysis
(EDX, Oxford 6699 ATW, Oxford, Oxfordshire, UK), and X-ray photoelectron spectroscopy (XPS)
data were obtained with a SPECS GmbH customized system (SPECSTM , Berlin, Germany) for surface
analysis, equipped with a non-monochromatic X-ray source XR 50 and a hemispherical energy analyzer
(PHOIBOS 1509MCD). Photoelectrons were excited by using MgKα line (1253.6 eV) operating at
200 W/12 kV. Powder samples were attached onto Cu foil and were placed first in the pre-treatment
chamber at room temperature for several hours, before being transferred to the analysis chamber.
The XPS data were signal averaged for a large enough number of scans to get a good signal/noise
ratio, at increments of 0.1 eV and fixed pass energy of 25 eV. High-resolution spectra envelopes were
obtained by curve fitting synthetic peak components, using the software XPSpeak. The raw data
were used with no preliminary smoothing. Symmetric Gaussian—Lorentzian product functions were
employed to approximate the line shapes of the fitting components. Binding energies were calibrated
relative to the C 1s peak from the graphitic carbon at 284.6 eV.
2.5. Electrochemical Characterization
All electrochemical experiments were carried out at room temperature using a three-electrode
glass cell, a carbon rod as a counter electrode and a reversible hydrogen electrode (RHE) as a reference
electrode. All the potentials in this work were given against the RHE.
The working electrode consisted of a certain amount of the catalyst deposited as a thin layer over
a glassy carbon disc (0.196 cm2 ). The ink was prepared by mixing 2 mg of the catalyst powder, 15 μL
of Nafion solution, and 500 μL of Milli-Q water.
Electrochemical measurements were performed with a PC-controlled Autolab PGSTAT30
potentiostat-galvanostat (Metrohm, Herisau, Switzerland) in 0.5 M H2 SO4 electrolyte solution.
A total of 2 M CH3 OH + 0.5 M H2 SO4 solution was used to study the methanol oxidation reaction
(MOR). Ar was employed to deoxygenate all solutions, as well as CO for CO stripping experiments.
The last experiments were recorded at 0.02 V s−1 after bubbling CO through the cell for 10 min while
keeping the electrode at 0.10 V, followed by Ar purging for 30 min to completely remove the excess
CO. CO stripping voltammograms were recorded, by first scanning negatively until 0.05 V so that
the entire hydrogen region was probed, and then scanning positively up to 1.2 V. Methanol cyclic
voltammograms (CVs) were recorded between 0.05 and 1.0 V at a scan rate of 0.02 V s−1 , while
current transients were obtained by stepping the potential from 0.05 V to the final oxidation potential
(0.50 < Ef < 0.90 V). Tafel plots were calculated from stationary currents achieved at 300 s during the

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chronoamperometry experiments. Currents were normalized by the electroactive surface area (ESA)
calculated from CO stripping experiments, by using a conversion factor of 240 μC cm−2 (12.6 cm2 for
Pt/C, 8.2 cm2 for Pt/rGO-TT and 9.1 cm2 for Pt/N-rGO).
3. Results and Discussion
3.1. Physicochemical Characterization
Figure 1 displays the X-ray diffractograms for graphite, carbon Vulcan, GO, rGO-TT, N-rGO,
Pt/C, Pt/rGO-TT, and Pt/N-rGO. Graphite depicts a sharp diffraction peak at 24.6◦ , which is related to
the (002) plane, and a small diffraction peak at 54◦ associated to the (004) facet [15]. The GO formation
from the oxidation of graphite is detected by the broad diffraction peak at 10.7◦ , which is related
to the (101) diffraction plane and by the absence of the diffraction peaks associated to graphite [15].
Additionally, Table 1 shows the expansion of the C–C interplanar spacing from 0.34 nm (graphite) to
0.82 nm (GO), due to the growth of oxygen functional groups between the graphitic layers [15].

Figure 1. XRD patterns of graphite, graphene oxide (GO), Vulcan carbon (C), reduced graphene oxide
by thermal treatment (rGO-TT), nitrogen-modified reduced graphene oxide (N-rGO), Pt/C, Pt/rGO-TT,
and Pt/N-rGO.
Table 1. X-ray diffraction (XRD) parameters for all studied materials (dhkl = interplanar spacing).
Catalyst

Positionhkl 2θ/◦

dhkl /nm

Number of Layers

Pt Crystallite Size/nm

Graphite
GO
Vulcan
rGO-TT
N-rGO

26.5 C(002)
10.7 C(002)
24.7 C(002)
25.2 C(002)
25.3 C(002)

0.34 C(002)
0.82 C(002)
0.36 C(002)
0.35 C(002)
0.35 C(002)

121
11
6
12
9

-

Pt/C
Pt/rGO-TT
Pt/N-rGO

67.4 Pt(220)
67.5 Pt(220)
67.6 Pt(220)

0.23 Pt(111)
0.23 Pt(111)
0.23 Pt(111)

-

2.3
2.4
2.0

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On the other hand, Vulcan carbon, rGO-TT, and N-rGO reveal similar interplanar spacing (Table 1)
and X-ray patterns (Figure 1), which indicate the restoration of the graphite C–C lattice and the
elimination of oxygen functional groups between the graphitic layers in rGO materials. Also, a small
diffraction peak at ca. 45◦ associated to the (100) plane is observed at the three carbon supports.
The number of graphene layers that was estimated from the Debye–Scherrer equation and the Bragg
law is also included in Table 1. The success of GO, rGO-TT, and N-rGO formation can be easily
perceived by the diminution of the graphene layers in comparison to graphite. In this context,
a graphite behavior was suggested for materials with more than 12 graphene layers, and therefore,
a graphenic-like behavior is expected for all materials of the current work [26]. On the other hand, it is
also important to note that N-rGO reveals a small diffraction peak at 10.7◦ associated to GO, which
suggests a partial reduction of GO when caffeine is employed as a reducing agent [15].
Additionally, Figure 1 depicts X-ray diffractograms of Pt/C, Pt/rGO-TT, and Pt/N-rGO, in which
the typical face centered cubic (fcc) structure of Pt is discerned at 2θ = 40◦ , 46◦ , 68◦ , and 82◦ , associated
to the subsequent planes: (111), (200), (220), and (311) (JPCDS 00-004-0802). It is remarkable that the
absence of the diffraction peak is at 10.7◦ for Pt/N-rGO, which suggests a reduction of the oxygen
functional groups by the formic acid that is employed to reduce the metal precursor. In addition, it is
important to highlight the absence of Pt crystalline oxides at all catalysts. However, the presence of Pt
amorphous oxides should not be ruled out, since the oxygen content in graphenic-based materials is
high (see Table 2). The most important crystallographic parameters of Pt-based catalysts are depicted
in Table 1, in which similar interplanar spacing for all catalysts, smaller crystallite size for Pt/N-rGO
than Pt/C and Pt/rGO-TT, and a small but visible lattice contraction in graphenic-based catalysts
are discerned.
Table 2. Composition acquired by energy dispersive X-ray analysis (EDX) analysis and * elemental
analysis for all studied materials.
Composition (% Weight)
Material

C

O

N*

Graphite
Carbon
GO
rGO-TT
N-rGO
Pt/C
Pt/rGO-TT
Pt/N-rGO

80.0
96.5
48.0
75.3
73.0
96.4
86.9
78.3

20.0
3.5
52.0
24.7
21.0
3.6
13.1
15.7

6.0
6.0

Scanning electron spectroscopy (SEM) coupled with energy dispersive X-ray spectroscopy (EDX)
was employed to study the morphology, distribution, surface topography, and semi-quantitative bulk
composition of the catalysts. The amount of Pt incorporated into each sample was close to the nominal
value (19.0 ± 1.5 wt %). Figure 2 displays SEM images in the microscale for all materials synthesized.
Pt/C reveals a smooth and homogeneous surface, while a granular and porous structure is observed
at Pt/rGO-TT and Pt/N-rGO.

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Figure 2. Scanning electronic microscopy (SEM) images for Pt/C (left panel), Pt/rGO-TT
(middle panel), and Pt/N-rGO (right panel).

Table 2 discloses a strong increment of the oxygen content after GO production, which is reduced
in the first stage after rGO-TT and N-rGO formation, and in the second stage after Pt deposition.
The last stage is in agreement with the X-ray diffraction pattern of N-rGO, as the small peak associated
with oxide species disappears after the addition of formic acid. It is noteworthy that even after Pt
deposition, both graphenic catalysts show a higher amount of oxygen than carbon-supported Pt.
Raman spectroscopy (spectra are not shown) was employed to study the structural changes,
the disorder degree, and the crystallographic size in the basal plane (La ) of graphenic materials. With
this end, the following equation was employed [27,28]:
  −1


I
L a (nm) = 2.4 × 10−10 ·λ4l · D
IG

(1)

where λl is the excitation wavelength of the laser (532 nm), ID is the peak intensity at 1360 cm−1
(associated to sp3 carbon domains), and IG is the peak intensity at 1580 cm−1 (related to sp2 bonds
into the graphitic grid). The main results indicate an increment of the disorder (ID /IG ~ 0.9) compared
to graphite (ID /IG ~ 0.15), and the similar La ~ 20 nm for Vulcan carbon and all graphenic materials.
The latter is the expected result, since structural disorder by edges, finite size, and carbon hybridization
by heteroatom–C bond formation is usually reported in this type of materials [15,18].
Morphology, particle size, and agglomeration degree were studied by transmission electron
microscopy (TEM). Figure 3 shows TEM images of Pt/C, Pt/rGO-TT, and Pt/N-rGO catalysts.
Meanwhile, Figures S1–S3 depict high-resolution TEM images for all electrocatalysts employed in the
current work. Spherical Pt nanoparticles are homogeneously dispersed on carbon sheets; however,
the agglomeration degree increases with the amount of oxygen in the catalyst in the subsequent way:
Pt/C < Pt/rGO-TT < Pt/N-rGO. Therefore, Pt nanoparticles seem to nucleate and grow referentially at
surface oxygenated sites. Regarding the particle sizes, Figure 3 includes histograms for each catalyst,
and shows increments in the following way: Pt/ N-rGO (3.14 ± 0.50 nm) < Pt/C (3.28 ± 0.25 nm) <
Pt/rGO-TT (3.32 ± 0.43 nm). The last is in agreement with the trend observed for the crystallite size
values calculated by XRD, although the values are slightly lower than those achieved by TEM, which
is expected (i.e., crystallite size ≤ grain size ≤ particle size).
X-ray photoelectron spectroscopy (XPS) was carried out to identify and quantify the surface
composition and the chemical state of the elements (see Table 3). Deconvoluted spectra for Pt, C, O,
and N are presented in Figures 4 and 5. The Pt 4f signal (Figure 4) is resolved into three doublets in all
electrocatalysts. The most intense Pt 4f7/2 component at lower binding energy is attributed to metallic
Pt0 , while the second doublet is assigned to Pt2+ , and the third doublet at higher binding energy is
related to the higher oxidation state of Pt4+ . According to Table 3, there is a peak shift of metallic Pt
to a lower binding energy for graphenic-based supports. The latter indicates a charge transfer from
graphenic-based supports to platinum. However, Pt2+ species (i.e., PtO) increase to the detriment of
metallic Pt at Pt/N-rGO material.

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Figure 3. Transmission electron micrographs and particle size distribution histograms of the catalysts.
Table 3. X-ray photoelectron spectroscopy (XPS) analysis of the catalysts.
Elements

Pt4f

C1s

O1s

Catalysts

Pt/C

Pt/rGO-TT

Pt0

71.5
(62)

C–C

284.6
(81)

Pt oxides

530
(5)

Pt2+

72.4
(31)

C–N

-

O=C

531.1
(10)

Pt4+

73.5
(7)

C–OH

285.8
(14)

OH–C

532.8
(72)

C=O

287
(4)

O–C=O

534.1
(13)

O–C=O

289.1
(1)

H2 O

C–C

284.6
(68)

Pt oxides

Pt0

71.2
(63)

266

530
(3)

Surfaces 2019, 2, 2

Table 3. Cont.
Elements

Pt4f

C1s

O1s

Catalysts
Pt2+

72.3
(30)

C–N

-

O=C

531.2
(23)

Pt4+

74.3
(7)

C–OH

285.8
(15)

OH–C

532.9
(58)

C=O

287
(10)

O–C=O

533.9
(9)

O–C=O

288.7
(7)

H2 O

535.2
(7)

Pt/rGO-TT

Pt0

71.2
(49)

C–C

284.5
(52)

Pt oxides

530.1
(4)

Pt2+

72.1
(42)

C–N

285.3
(14)

O=C

531.2
(33)

Pt4+

74.0
(10)

C–OH

285.9
(12)

OH–C

532.7
(48)

Pyridinic

398.7
(29)

C=O

287.1
(12)

O–C=O

534
(12)

Pyrrolic

400.2
(49)

O–C=O

288.5
(10)

H2 O

535
(2)

Quaternary

401.2
(22)

290.0
(2)

-

-

Pt/N-rGO

π–π

N1s

-

-


Figure 4. Pt 4f (left panels) and C1s (right panels) XPS spectra of the catalysts.

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Surfaces 2019, 2, 2

The C1s signal provides detailed information about the presence of oxygen and nitrogen species
bonded to carbon atoms (Figure 4). As can be seen, the carbon signal is mainly composed by graphitic
carbon (C1), followed by different oxygenated species at higher binding energy values, such as C–OH
(C3), C=O (C4), and carboxylic bonds (C5) [29]. In addition to these signals, Pt/NrGO reveals an
additional component at ca. 285.3 (C2) ascribed to nitrogen species bonded to carbon atoms. Indeed,
the strong N1s signal (Figure 5) confirms that the last statement since the nitrogen species, such as
pyridinic, pyrrolic, and N-quaternary, are clearly identified. Also, a small but visible signal at ca.
289 eV ascribed to a π–π* shakeup satellite for Pt/Vulcan and Pt/NrGO catalysts is observed [30].
Interestingly, the O1s signal (Figure 5) reveals different surface oxygenated species, which strongly
depend on the substrate. In agreement with the C1s signal, C=O (O2:~531.05 ± 0.05 eV) and C–OH
(O3:~533.5 ± 0.5 eV) species are the main components for all materials [29]. Signals related to carboxyl
groups (O4) and platinum oxides (O1) appear in minor proportion, while intercalated water (O5)
is only perceived at graphenic materials. On the other hand, Pt/NrGO shows similar intensities
for C–OH and C=O species, and the most intense component is associated to C–O species. All XPS
information is summarized in Table 3.


Figure 5. O1s (left panels) and N1s (right panel) XPS spectra of the catalysts.

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Surfaces 2019, 2, 2

3.2. Electrochemical Characterization
CO stripping experiments were performed to (i) characterize the catalyst surface; (ii) achieve
catalytic activity towards CO oxidation, which is one of the most important catalyst poisons; and
(iii) acquire the electroactive surface area (ESA). Figure 6 shows a similar CO oxidation peak value
(COP = 0.83 V) for all catalysts. Noticeable, Pt/N-rGO reveals a broader CO oxidation peak, with an
anodic tail that ends at 1.2 V. Additionally, the onset potential for the CO oxidation reaction is 0.70 V
for Pt/rGO-TT, 0.72 V for Pt/C, and 0.75 V for Pt/N-rGO. These parameters indicate a higher CO
tolerance for Pt/rGO-TT than for Pt/C, and the even lower one for Pt/N-rGO. Also, an increase of
the double-layer capacitive current is observed at graphenic-based materials, especially relevant for
Pt/N-rGO. The latter results and the low catalytic activity towards the CO oxidation on Pt/N-rGO
may be ascribed to the high amount of surface-oxygenated species, i.e., PtO and PtO2 . Regarding the
ESA, the values increase in the following way: Pt/rGO-TT (8.2 cm2 ) < Pt/N-rGO (9.1 cm2 ) < Pt/C
(12.6 cm2 ). The last is in agreement with TEM analysis, in which high Pt agglomeration degree was
detected at graphenic-based catalysts.

Figure 6. CO stripping voltammograms recorded at Pt/C, Pt/rGO-TT, and Pt/N-rGO catalysts in
0.5 M H2 SO4 . Ead = 0.1 V; scan rate = 0.02 V s−1 .

The catalytic performance toward the methanol oxidation reaction (MOR) was evaluated by
potentiodynamic (Figure 7) and potentiostatic (Figure 8) experiments. Figure 7 shows the lowest
performance for Pt/N-rGO toward the MOR, which is in disagreement with what has been reported
by other authors [31,32]. This different catalytic performance may be ascribed to the elevated amount
of Pt surface oxide species (PtO and PtO2 observed by XPS analysis) by using the caffeine route for the
synthesis of the N-rGO, which inhibits the methanol adsorption step (see below).

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Surfaces 2019, 2, 2

Figure 7. Methanol oxidation recorded for Pt/C, Pt/rGO-TT, and Pt/N-rGO catalysts in 2 M CH3 OH +
0.5 M H2 SO4 . Scan rate = 0.02 V s−1 .

Figure 8. Current transients for methanol oxidation on Pt/C, Pt/rGO-TT, and Pt/N-rGO catalysts in
2 M CH3 OH + 0.5 M H2 SO4 . Ei = 0.05 V, 0.50 V ≤ Ef ≤ 0.90 V.
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Surfaces 2019, 2, 2

Additionally, similar voltammetric profiles with comparable onset potentials (~0.5 V) and
analogous anodic (EPa ~ 0.93 V) and cathodic (EPc ~ 0.76 V) peak potentials are observed for Pt/rGO-TT
and Pt/C. However, a close inspection of Figure 7 reveals higher anodic currents at Pt/rGO-TT than
Pt/C during the positive sweep potential. Indeed, the pre-peak at ca. 0.77 V reveals higher current
density at Pt/rGO-TT (2.4 mA cm−2 ) than at Pt/C (2.0 mA cm−2 ). The current peak during the anodic
sweep (jPa ) is 4.9 mA cm−2 at Pt/rGO-TT and 4.2 mA cm−2 at Pt/C, while the current peak during the
cathodic sweep (jPc ) is 5.4 mA cm−2 at Pt/rGO-TT and 5.6 mA cm−2 at Pt/C. The jPa /jPc ratio is usually
employed to estimate the tolerance of the catalyst to adsorbed reaction intermediates (mainly COad ).
Pt/C shows the lowest jPa /jPc value and appears as the catalyst with the highest tolerance, which is
not in agreement with the CO stripping experiments in Figure 6. However, a recent work states that
the jPc is not related to a residual intermediate of the MOR, and the jPa /jPc ratio is associated with the
oxophilicity degree of the catalyst [33]. Thus, higher jPa /jPc value indicates higher oxophilicity degree,
and consequently, Pt/rGO-TT appears to be the material with the highest oxophilicity. The last is in
agreement with the physicochemical characterization, as well as with the CO stripping experiments, in
which Pt/rGO-TT reveals the highest CO tolerance.
Figure 7 depicts current transients achieved from an initial potential of 0.05 V, in which methanol
does not oxidize, to a final potential (0.50 ≤ E ≤ 0.90 V) where the MOR takes place. In agreement
with potentiodynamic experiments, Pt/N-rGO develops the lowest performance toward the MOR in
all potential ranges studied. In the same way, Pt/rGO-TT and Pt/C reveal similar better performance
toward the alcohol oxidation. Interestingly, Pt/rGO-TT and Pt/C develops inflection points at short
times at the same potential range (0.65 ≤ E ≤ 0.75), where the onset potential for the CO removal
(Figure 5) and the pre-peak at the potentiodynamic experiments in presence of methanol is observed
(Figure 6). This result suggests a change of rate-determining step (RDS), which may be attributed to an
enhancement of the water dissociation reaction (i.e., OHad formation) necessary to remove CO from
the catalyst surface [34,35]. Additionally, it is observed that at E ≥ 0.85 V (anodic peak during the
positive sweep potential, Figure 6), the current decay is faster at Pt/rGO-TT and Pt/C, which is related
to the low performance of Pt oxide surface toward the MOR [36].
Further characterization of the catalytic activity of the synthesized materials was carried out by a
Tafel plot analysis. A Tafel slope is usually employed to determine the rate-determining step (RDS),
and therefore, the reaction mechanism may be inferred from Tafel slope values. Figure 9 depicts a
Tafel plot obtained from the current transients of methanol oxidation for the three catalysts (Figure 8)
at different final potentials (0.50 ≤ E ≤ 0.90 V) in an acidic medium. Pt/C shows a Tafel slope of
120 mV dec−1 , while Pt/rGO-TT and Pt/C reveal a Tafel slope of 75 mV dec−1 . In order to clarify the
operating reaction mechanism on the catalysts involved in the current work, the main steps for MOR
are briefly described.
First there is the electrosorption and dehydrogenation of methanol onto a suitable surface.
The complete electrosorption and dehydrogenation steps yield to adsorbed carbon monoxide (COad ),
otherwise formaldehyde and formic acid may be produced [37]:
H3 OH  CH3 OHad

(2)

CH3 OHad → CH2 OHad + H+ + e−
+

CH2 OHad → CHOHad + H + e
+

CH2 OHad → CH2 O + H + e

−

(4)

−

(5)

CH2 O  CH2 Oad

(6)

+

CH2 Oad → CHOad + H + e

−

+

CHOad + H2 O → HCOOH + H + e
CHOad → COad + H+ + e−
271

(3)

(7)
−

(8)
(9)

Surfaces 2019, 2, 2

Second, there is the electrooxidation of COad . Water dissociation on the catalyst surface (formation
of adsorbed OH) is needed to perform the COad oxidation to CO2 , and consequently liberate the catalyst
surface [31,38,39]:
H2 O  OHad + H+ + e−
(10)
COad + OHad → COOHad
COOHad → CO2 + H

+

+ e

−

(11)
(12)

dec−1

A Tafel slope of 120 mV
suggests that the rate-determining step is related to the first
electrochemical step, and therefore the rate of MOR at Pt/N-rGO is limited by the “electrosorption
and dehydrogenation step”, i.e., Reaction 3. The latter is in complete agreement with the low current
developed by the Pt/N-rGO materials during the MOR, and should be ascribed to a non-suitable
surface for methanol dehydrogenation, which is comprised of strongly bonded oxygen species
onto the surface, i.e., PtO and PtO2 [36]. On the other hand, a Tafel slope close to 60 mV dec−1
(Pt/C and Pt/rGO-TT) indicates a chemical reaction after an electrochemical reaction as the RDS,
i.e., the surface diffusion of COad toward the active site (reaction 11). Both materials revealed
similar catalytic performance toward the MOR, although a slight enhancement of the CO tolerance
at Pt/rGO-TT was perceived. The last seems to be related to electronic effects (charge transfer
from rGO-TT to Pt) and to a higher amount of C–OH species onto the graphenic-based catalyst
(bifunctional effect).

Figure 9. Tafel plots obtained at 300 s from Figure 8.

4. Conclusions
Graphene oxide (GO), reduced graphene oxide (rGO-TT), nitrogen-modified rGO (N-rGO),
and carbon-supported Pt materials using these supports were successfully synthesized and fully
characterized. Results have shown that the support conditions the state of oxidation of metal
nanoparticles at the surface of the catalysts, which control the activity of a specific electrochemical
reaction. Indeed, a small addition of nitrogen into the catalyst support induces a change of the
rate-determining step from a chemical reaction after an electrochemical reaction to the first electron.
Thus, Pt/N-rGO displays the worst performance toward the methanol oxidation reaction, while

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Pt/rGO-TT not only enhances the catalytic activity toward the alcohol oxidation, but also toward the
most important catalyst poison.
Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/2/1/2/s1,
Figure S1: TEM image for Pt/C electrocatalyst, Figure S2. TEM image for Pt/rGO-TT electrocatalyst, Figure S3.
TEM image for Pt/N-rGO electrocatalyst.
Author Contributions: Data curation, G.A. and S.J.M.; Formal analysis, L.M.R.-G.; Writing – original draft, R.R.;
Writing – review & editing, E.P. and G.G.
Acknowledgments: The Spanish Ministry of Economy and Competitiveness (MINECO) has supported this work
under project ENE2017-83976-C2-2-R (co-funded by FEDER). L.M.R. thanks the ACIISI for his PhD fellowships.
G.G. acknowledges the Viera y Clavijo program (ACIISI and ULL) for financial support.
Conflicts of Interest: The authors declare no conflict of interest.

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Article

Electrochemical Behavior of Pt–Ru Catalysts
Supported on Graphitized Ordered Mesoporous
Carbons toward CO and Methanol Oxidation
Juan Carlos Calderón Gómez 1, *, Verónica Celorrio 2,3 , Laura Calvillo 4 , David Sebastián 1 ,
Rafael Moliner 1 and María Jesús Lázaro Elorri 1, *
1
2
3
4

*

Instituto de Carboquímica (CSIC), Miguel Luesma Castán 4, 50018 Zaragoza, Spain;
dsebastian@icb.csic.es (D.S.); rmoliner@icb.csic.es (R.M.)
UK Catalysis Hub, Research Complex at Harwell, RAL, Oxford OX11 0FA, UK; v.celorrio@ucl.ac.uk
Kathleen Lonsdale Building, Department of Chemistry, University College London, Gordon Street,
London WC1H 0AJ, UK
Dipartimento di Scienze Chimiche, Università di Padova, Via Marzolo 1, 35131 Padova, Italy;
laura.calvillolamana@unipd.it
Correspondence: jccalderon@icb.csic.es (J.C.C.G.); mlazaro@icb.csic.es (M.J.L.E.);
Tel.: +34-976-733-977 (J.C.C.G. & M.J.L.E.)

Received: 29 November 2018; Accepted: 27 December 2018; Published: 3 January 2019

Abstract: In this work, graphitized ordered mesoporous carbons (gCMK-3) were employed as support
for Pt and Pt–Ru nanoparticles synthesized by different reduction methods. The catalysts displayed
metal contents and Pt:Ru atomic ratios close to 20 wt % and 1:1, respectively. A comparison of
the physical parameters of Pt and Pt–Ru catalysts demonstrated that Ru enters into the Pt crystal
structure, with well-dispersed nanoparticles on the carbon support. The Pt catalysts exhibited similar
surface oxide composition, whereas a variable content of surface Pt and Ru oxides was found for the
Pt–Ru catalysts. As expected, the Pt–Ru catalysts showed low CO oxidation onset and peak potentials,
which were attributed to the high relative abundances of both metallic Pt and Ru oxides. All the
studied catalysts exhibited higher maximum current densities than those observed for the commercial
Pt and Pt–Ru catalysts, although the current–time curves at 0.6 V vs. reversible hydrogen electrode
(RHE) demonstrated a slightly higher stationary current density in the case of the Pt/C commercial
catalyst compared with Pt nanoparticles supported on gCMK-3s. However, the stationary currents
obtained from the Pt–Ru/gCMK-3 catalysts surpassed those of the commercial Pt–Ru material,
suggesting the suitability of the prepared catalysts as anodes for these devices.
Keywords: Pt–Ru catalysts; ordered mesoporous carbons; graphitization; CO oxidation; methanol
oxidation; direct methanol fuel cells

1. Introduction
A quite interesting carbon material used as support for electrodes in direct methanol fuel cells
(DMFCs) is the ordered mesoporous carbon (OMC). This material is synthesized by nanocasting
using a silica template, which is impregnated with a carbon-based resin that, after carbonization and
silica removal, results in an OMC with a 3D interconnected porous structure, high surface area, and
rich content of mesopores [1], which promote the formation of uniformly distributed nanoparticles
on the carbon surface [2]. Moreover, it is possible to modify its surface chemistry by creating
surface functional groups that act as anchoring sites and enhance the electronic transfer between
the catalytic nanoparticles and the support. The challenge is to maintain the ordered structure during
the functionalization treatments in order to preserve the mesoporous structure as well as the high

Surfaces 2019, 2, 1–15; doi:10.3390/surfaces2010001

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Surfaces 2019, 2

surface area [3,4], which has been demonstrated to be crucial for the performance of OMC-supported
Pt nanoparticles [5,6].
Despite the mentioned outstanding properties of OMCs, their low electrical conductivity hinders
their use in fuel cells. The relatively low temperatures employed during the carbonization stage seem
to be the cause for the poor formation of extensive graphite layers [7,8]. To avoid this drawback,
graphitization of OMCs has been suggested by means of different procedures, such as the use of
aromatic and polyaromatic hydrocarbons as carbon precursors during the silica template impregnation,
as they can produce more graphitic domains after carbonization. Some of these organic molecules
could be benzene, naphthalene, anthracene, and pyrene [9]. A similar procedure indicated that Fe
phthalocyanines can also generate graphitic planes, but this methodology is expensive due to the high
cost of these molecules, even when low temperatures are used for carbonization [10]. Postsynthesis
heat treatments are accepted as a useful way to obtain graphitized OMCs, although part of the ordered
structure can be altered, broadening the pore size distribution [11]. Some of these treatments can
be performed at temperatures lower than 1000 ◦ C in the presence of transition metal salts [12–14]
or increasing the temperature until 1500 ◦ C in N2 or Ar atmosphere, avoiding the use of these
metal compounds.
Graphitized OMCs have been employed to support Pt nanoparticles, finding enhanced stationary
current densities for the oxidation of methanol compared with other Pt catalysts supported on carbon
nanofibers, carbon nanocoils, and carbon black [6]. Pd nanoparticles have also been supported on
these materials to test their viability as anode catalysts in formic acid fuel cells. The current densities
associated with the oxidation of formic acid on these catalysts were higher than those observed on
a Pd catalyst supported on carbon black [15]. Regarding their use as cathode supports, Rivera et al.
prepared Pt2 CrCo alloys supported on graphitized OMCs to assess the oxygen reduction reaction
(ORR) on them. The authors obtained materials with outstanding activity toward the ORR, even in
the presence of methanol, considering the crossover of this fuel between the anode and the cathode in
DMFCs [16].
In this work, the activity of Pt and Pt–Ru catalysts supported on different graphitized ordered
mesoporous carbon (gCMK-3) was studied in order to determine their behavior toward the CO
and methanol electro-oxidation. It is important to highlight that the activity of Pt–Ru catalysts
supported on graphitized ordered mesoporous carbons toward the mentioned molecules have not
been reported in the literature, whereas the state of the art related to the behavior and performance of Pt
catalysts supported on graphitized ordered mesoporous carbons toward the methanol electrochemical
oxidation is in an incipient state. The CMK-3 was prepared using different TEOS/P123 mass ratios and
subsequently treated at 1500 ◦ C in Ar atmosphere to induce the formation of graphitic planes [6,15,16].
The Pt and Pt–Ru catalysts were prepared employing different reducing agents to verify their influence
on the physical and electrochemical behavior of the catalysts. The gCMK-3-supported catalysts
displayed higher CO tolerance and improved methanol oxidation current densities than those observed
for commercial Pt and Pt–Ru catalysts supported on carbon black from E-TEK, which were used as
reference to evaluate the results displayed for these catalysts.
2. Materials and Methods
2.1. Synthesis of Graphitized Ordered Mesoporous Carbon (gCMK)
CMK-3 was prepared by an incipient wetness impregnation procedure as described in
Reference [15]. Briefly, the SBA-15 silica (Sigma-Aldrich, Madrid, Spain) was impregnated with
the carbon precursor (furan resin, Huttenes-Albertus, Düsseldorf, Germany). The volume of furan
resin used in the synthesis was equal to the silica porous volume (~1 cm3 g−1 ). After letting the material
dry, it was subjected to a thermal treatment in nitrogen atmosphere at 700 ◦ C for 2 h, and the silica was
then removed using HF (40%, Fluka, Bucharest, Romania). Two silica templates with different textural
properties were used for the synthesis. They were prepared by following the procedure described in

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Reference [17] and using a TEOS/P123 (Sigma-Aldrich, Madrid, Spain) mass ratio (R) equal to 2 and 8
(SBA-15/R2 and SBA-15/R8, respectively). The carbon supports derived from these silica materials
were labeled as CMK-3-R2 and CMK-3-R8 [17,18]. Finally, the CMK-3 materials were heat-treated at
1500 ◦ C in a graphite electrical furnace for 1 h under argon flow in order to obtain the graphitized
material (gCMK-3) [15].
2.2. Synthesis of gCMK-3-Supported Pt and Pt–Ru Catalysts
Pt and Pt–Ru catalysts supported on gCMK-3s were synthesized employing three different
procedures described below. In all methods, a solution of the metal precursor salts (H2 PtCl6 8% w/w
solution and RuCl3 , 99.999%, Sigma-Aldrich, Madrid, Spain) was prepared in water or ethylene glycol
(Sigma-Aldrich, Madrid, Spain) in order to obtain materials with a Pt:Ru atomic ratio of 1:1.
Method 1 (sodium borohydride method, BM): gCMK-3 was dispersed in water, and the
solution containing the metal precursors was then slowly added under vigorous stirring. Finally, a
26.4 mM sodium borohydride (Sigma-Aldrich, Madrid, Spain) solution was added dropwise under
sonication [3].
Method 2 (formic acid method, FAM): gCMK-3 was dispersed in a 2 M formic acid (Merck,
Darmstadt, Germany) solution and then heated at 80 ◦ C. Afterward, the metal precursor solution was
slowly added and kept at this temperature for 12 h [4].
Method 3 (ethylene glycol method, EG): the metal precursors and the carbon support were
dispersed in EG, and the pH was adjusted to 11 using a NaOH (Sigma-Aldrich, Madrid, Spain)
solution in EG. The dispersion was heated at 195 ◦ C for 2 h, quickly cooled down to room temperature,
and the pH adjusted to 1 with HCl (Merck, Darmstadt, Germany) [19].
Finally, all the materials were filtered, washed, and dried at 60 ◦ C. The catalyst were labeled as
Pt/gCMK-3-RX-Y or Pt–Ru/gCMK-3-RX-Y, with X = 2, 8 and Y = BM, FAM, EG.
2.3. Physicochemical Characterization of the Catalysts
Energy dispersive X-ray (EDX) was employed to determine the metal content of the synthesized
catalysts using a scanning electron microscope (SEM) Hitachi S-3400N (Hitachi, Düsseldorf, Germany)
coupled to a Röntec XFlash analyzer (RÖNTEC GmbH, Berlin, Germany). This instrument is equipped
with a Be window and a Si(Li) detector (Hitachi, Düsseldorf, Germany) operating at 15 keV.
X-ray diffraction (XRD) analyses were carried out to study the crystalline properties of the
catalysts. XRD patterns were obtained by means of a Bruker AXS D8 Advance diffractometer (Bruker,
Karlsruhe, Deutschland), with a θ-θ configuration and a Cu-Kα radiation (λ = 0.154 nm). The scans
were achieved between 0◦ and 100◦ for 2θ values. The (220) peak located at 2θ ≈ 70◦ and the Scherrer’s
equation were used to calculate the crystallite size [20].
X-ray photoelectron spectroscopy (XPS) analysis was performed in order to determine the
chemical state and the surface composition of Pt and Ru in the catalysts. An ESCAPlus Omicron
spectrometer (Omicron, Houston, TX, USA) equipped with a concentric hemispherical analyzer, seven
channeltron detectors, and an Al/Mg monochromated X-ray source (Kα = 1253.6 eV) was employed.
This equipment was operated at 15 kV and 15 mA. The XPS spectra were recorded at pressures lower
than 8 × 10−9 mbar. Data processing was performed with the CASAXPS® software (SurfaceSpectra,
Manchester, UK), fitting the experimental data with Gaussian–Lorentzian curves.
In order to determine the particle size and the dispersion of the nanoparticles on the gCMK-3s,
transmission electron microscopy (TEM) analyses were made using a 200 kV JEOL-2000 FXII
transmission electron microscope (JEOL, Peabody, MA, USA). The catalysts were dropped from
an ethanol suspension on a carbon grid. The images were taken with a MultiScan CCD Gatan 694
(Gatan, Pleasanton, CA, USA) camera and treated with the ImageJ® software (National Institutes for
Health, Bethesda, MD, USA).

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2.4. Electrochemical Characterization
A three-electrode cell connected to an AUTOLAB NS85630 modular equipment (Metrohm,
Herisau, Switzerland) was employed to assess the electrochemical activity of the synthesized materials.
A reversible hydrogen electrode (RHE) placed inside a Luggin capillary was used as the reference
electrode, and a glassy carbon bar was used as the counter electrode. A catalyst-modified glassy carbon
disk was used as the working electrode. The catalyst inks were prepared with 2.0 mg of catalyst, 15 mL
of Nafion® (5 wt %, Sigma-Aldrich, Madrid, Spain) and 500 mL of ultrapure water. Afterward, a 60-μL
aliquot was deposited and dried on the glassy carbon disk. All the aqueous solutions were prepared
with high resistivity deoxygenated 18.2 MΩ cm H2 O. 0.5 M H2 SO4 (95%–97%, Merck, Darmstadt,
Germany) was used as the supporting electrolyte. For the methanol electrochemical characterization,
2.0 M methanol (99%, Merck, Darmstadt, Germany) solutions in the supporting electrolyte were
employed. CO activity was studied by bubbling CO (99.999%, Air Liquide, Madrid, Spain) into the
electrochemical cell for 10 min at 0.2 V vs. RHE and 20 ◦ C in order to form a CO monolayer on the
deposited catalyst; then, nitrogen (MicroGeN2, GasLab, Fremont, CA, USA) was bubbled for 20 min to
remove the dissolved CO from the supporting electrolyte. Two potential scans at 0.020 V s−1 between
0.05 V and 1.0 V vs. RHE were applied. All the experiments were performed at 20 ◦ C, while all the
current densities presented in this paper have been normalized by the electroactive area calculated
from the charge associated with the oxidation of an adsorbed CO monolayer on Pt (420 μC cm−2 Pt ).
The results obtained from the physical and the electrochemical characterization were compared
against those exhibited by the commercial Pt and Pt–Ru catalysts supported on carbon black from
E-TEK (DeNora, Milan, Italy). These catalysts have a metal content close to 20 wt% and a Pt:Ru ratio
near 1:1.
3. Results
3.1. Physical Characterization
Table 1 summarizes the metal content of the different studied catalysts. In all the cases, metal
contents around 20 wt% were observed, agreeing with the nominal values. Pt:Ru atomic ratios were
close to 1:1, except in the case of the catalyst prepared by the EG method, which exhibited a higher Pt
atomic content (66%).
Figures 1 and 2 depict the XRD patterns for the prepared Pt and Pt–Ru catalysts, respectively.
All the catalysts presented the typical (110), (200), and (220) peaks for the face-centered cubic structure
of Pt located at 2θ of 39.75◦ , 46.31◦ , and 67.48◦ , as those observed for the commercial Pt/C and Pt–Ru/C.
On the other hand, the peaks for the Pt–Ru materials exhibited a shift toward higher 2θ values with
respect to those of the Pt catalysts (see Figure 2), which is attributed to the structure contraction due to
the incorporation of Ru into the fcc Pt crystalline structure [21–23]. This fact was corroborated by the
smaller lattice parameter calculated for the Pt–Ru materials compared with the Pt values (see Table 1).
Furthermore, a peak around 25◦ in the XRD patterns of the Pt and Pt–Ru commercial catalysts was
detected, which corresponds to the (002) reflection of the graphite basal planes present in the carbon
black acting as support for these materials.

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Table 1. Physical characterization of the catalysts supported on graphitized ordered mesoporous
carbons (gCMK-3s) and the commercial Pt/C E-TEK. Metal surface area: SA (m2 g−1 ) = 6 × 103 /ρd,
with d as the average crystallite size (nm) and ρ as the alloy density (g cm−3 ). ρ was determined by
considering that ρPt–Ru (g cm−3 ) = ρPt Xm Pt + ρRu Xm Ru , with ρPt = 21.4 g cm−3 and ρRu = 12.3 g cm−3
and Xm Pt and Xm Ru as the weight percentage of Pt and Ru, respectively.
Catalyst

Metal Content,
wt %

Pt:Ru Atomic
Ratio

20
22
16
24
23
20
26
22
24
20

—–
—–
—–
—–
—–
—–
53:47
67:33
55:45
45:55

Pt/gCMK-3-R2-BM
Pt/gCMK-3-R2-FAM
Pt/gCMK-3-R8-BM
Pt/gCMK-3-R8-EG
Pt/gCMK-3-R8-FAM
Pt/C E-TEK
Pt–Ru/gCMK-3-R2-FAM
Pt–Ru/gCMK-3-R8-EG
Pt–Ru/gCMK-3-R8-FAM
Pt–Ru/C E-TEK

Crystallite and Particle Size, nm
XRD

TEM

5.9
3.7
5.6
6.2
4.1
3.0
6.6
3.6
5.6
4.4

3.9 ± 1.3
3.2 ± 1.0
4.1 ± 1.4
5.9 ± 1.4
4.9 ± 1.6
—–
5.2 ± 2.3
6.7 ± 2.8
3.9 ± 1.3
—–

Lattice
Parameter, Å

Metal Surface
Area, m2 g−1

3.923
3.926
3.920
3.919
3.920
3.921
3.896
3.921
3.898
3.866

47
76
50
45
68
93
49
85
57
76







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Figure 1. XRD patterns of the Pt catalysts supported on gCMK-3s and the commercial Pt/C from E-TEK.

The crystallite sizes were calculated from the (220) reflection, and the values are reported in
Table 1 [21]. For the Pt catalysts, the smallest crystallite sizes were obtained for the FAM method and
the largest ones for the EG, whereas the opposite was observed for the Pt–Ru materials. This result is
in agreement with that reported by Calderón et al. for Pt–Ru catalysts supported on carbon nanofibers
and prepared by different synthesis routes [21]. In general, the addition of Ru led to an increase in
the crystallite size and thus the metal surface areas (SA) agreeing with the expression described in
Table 1. As expected, SA diminished with the increase in the crystallite size, with Pt/gCMK-3-R8-EG
displaying the lowest value in the group of the Pt/gCMK-3 catalysts. In the case of Pt–Ru materials,
Pt–Ru/gCMK-3-R2-FAM presented lower value in the SA as a consequence of the largest crystallite
diameter determined for this material.

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Surfaces 2019, 2










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Figure 2. XRD patterns of the Pt–Ru catalysts supported on gCMK-3s and the commercial Pt/C and
Pt–Ru/C from E-TEK.

Figure 3 shows the TEM images obtained for the Pt/gCMK-3 catalysts, which demonstrate
well-dispersed nanoparticles on the carbon support, although some agglomerates can be observed
in some cases, such as Pt/gCMK-3-R2-BM and Pt/gCMK-3-R2-EG. Pt/gCMK-3-R2-FAM showed
the smallest particle size, while Pt/gCMK-3-R8-EG exhibited the highest one, following the trend
observed for the crystallite size values determined by XRD. For the Pt–Ru catalysts (Figure 4), larger
particles were observed, especially for Pt–Ru/gCMK-3-R2-FAM and Pt–Ru/gCMK-3-R8-EG. These
materials displayed a wide distribution of diameters between 1 and 12 nm, although the catalyst
Pt–Ru/gCMK-R8-FAM showed a narrower distribution and the smallest particle size among the
synthesized Pt–Ru catalysts.
The difference in the magnitude between the crystallite and the particle diameters can be explained
as follows: In XRD, an average of the crystallite size is obtained as a consequence of the X-ray beam
passing through the crystallites, even if they are agglomerated. During the TEM analysis, a counting
and measuring process is performed, which considers dispersed and nonagglomerated crystalline and
amorphous particles. Therefore, smaller values are determined [24].

Figure 3. TEM images and particle size distributions of the synthesized Pt/gCMK-3 catalysts.

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Surfaces 2019, 2

Figure 4. TEM images and particle size distributions of the synthesized Pt–Ru/gCMK-3 catalysts.

XPS analysis was carried out to identify the chemical and electronic states of the metals present in
the catalysts as well as their relative surface abundance, and the results are reported in Table 2. From
the deconvolution of the binding energy spectra of Pt 4f for Pt/gCMK-3 and Pt–Ru/gCMK-3 catalysts,
three pairs of Pt peaks were observed, which correspond to the three oxidation states of this metal (see
Figure S1 in the Supplementary Information). The transitions generated in the 4f 7/2 and 4f 5/2 orbitals
of Pt (0) are described by the peaks close to 71 and 75 eV, while the second and third pair of peaks,
located at 73–76 and 75–79 eV, are assigned to Pt (II) and Pt (IV), respectively. In the case of the catalysts
Pt/gCMK-3, a similar surface composition of the nanoparticles was observed as all of them displayed
a Pt (0) relative surface abundance around 60%. The oxidized states Pt (II) and Pt (IV) exhibited similar
abundances between them (close to 20%), indicating that neither the synthesis routes nor the carbon
support affect the surface composition of these catalysts. However, in the case of the Pt–Ru/gCMK-3
catalysts (spectra depicted in Figure S2 in the Supplementary Information), several differences were
observed in the relative abundance of the Pt surface oxidation states, which are attributed to the
entrance of Ru into the Pt lattice that induces changes in the electronic structure and interactions of
this metal [25–27]. The catalysts supported on gCMK-3-R8 displayed a higher abundance of Pt (0),
a result that matched with the trend observed for the Pt catalysts supported on gCMK-3-R8, which
displayed a slightly higher abundance of the Pt (0) oxidation state compared with that observed for
the Pt nanoparticles supported on gCMK-R2. Regarding the catalyst Pt–Ru/gCMK-3-R2-FAM, it
is possible to conclude that, on this carbon material, the presence of Ru induced the formation of
nanoparticles with a high surface content of Pt oxides.

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Surfaces 2019, 2

Table 2. Electronic and composition parameters of the catalysts supported on gCMK-3s obtained from
the XPS data.
Catalyst

Chemical State

Relative Area, %

Binding Energy, eV

Metallic Pt
Pt2+
Pt4+

59.4
21.7
18.9

Pt 4f 7/2
71.4
72.6
75.3

Metallic Pt
Pt2+
Pt4+

60.9
20.6
18.5

Pt 4f 7/2
71.4
72.6
75.3

Metallic Pt
Pt2+
Pt4+

61.1
19.5
19.4

Pt 4f 7/2
71.4
72.6
75.3

Metallic Pt
Pt2+
Pt4+

61.7
20.1
18.2

Pt 4f 7/2
71.4
72.6
75.4

Metallic Pt
Pt2+
Pt4+

62.1
18.8
19.1

Pt 4f 7/2
71.5
72.7
75.5

Metallic Pt
Pt2+
Pt4+

28.7
17.4
53.9

Metallic Ru
Ru4+
Ru4+ hydrate

14.9
27.9
57.2

Metallic Pt
Pt2+
Pt4+

59.1
22.2
18.7

Metallic Ru
Ru4+
Ru4+ hydrate

43.6
18.6
37.8

Metallic Pt
Pt2+
Pt4+

63.5
19.2
17.28

Metallic Ru
Ru4+
Ru4+ hydrate

18.9
40.7
40.4

Pt/gCMK-3-R2-BM

Pt/gCMK-3-R2-FAM

Pt/gCMK-3-R8-BM

Pt/gCMK-3-R8-EG

Pt/gCMK-3-R8-FAM

Pt–Ru/gCMK-3-R2-FAM

Pt–Ru/gCMK-3-R8-EG

Pt–Ru/gCMK-3-R8-FAM

Pt 4f 7/2
71.4
72.6
74.2
Ru 3p 3/2
461.0
463.2
465.4
Pt 4f 7/2
71.4
72.7
75.7
Ru 3p 3/2
461.6
463.2
465.3
Pt 4f 7/2
71.6
72.9
75.4
Ru 3p 3/2
461.1
463.3
465.1

Table 2 also reports the results related to the XPS analysis of Ru and its signals. The studied
transitions correspond to the 3p orbitals located between 450 and 494 eV (see Figure S3 in the
Supplementary Information). A doublet attributed to 3p 1/2 and 3p 3/2 transitions from Ru (0)
around 461 and 483 eV was determined, besides two doublets around 463-485 and 465-487, assigned
to the transitions of Ru (IV) and hydrated Ru (IV), respectively. The deconvolution analysis suggests
an influence of the synthesis method on the Ru oxides surface abundance as the catalysts reduced

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Surfaces 2019, 2

with formic acid (FAM) exhibited higher Ru (IV) abundances than that observed for the catalyst
reduced with ethylene glycol. This means that the obtaining of different relative abundances for the
Ru oxidation states depends on the employed reducing agent, with formic acid leading a less complete
reduction of Ru and ethylene glycol promoting a major reduction of Ru.
3.2. CO Stripping
The activity of the catalysts toward the carbon monoxide electro-oxidation was evaluated, and the
results are depicted in Figures 5 and 6. Pt/gCMK-3-R8 samples exhibited lower onset potential (values
between 0.52–0.61-V) than those supported on gCMK-3-R2 (0.63–0.67 V), indicating that linearly
adsorbed CO is promoted on these materials due to the high content of metallic Pt in these catalysts,
which was detected by XPS (see Table 2) [28]. From this analysis, it is possible to conclude that CO
oxidation starts at low potentials on catalysts with a high metallic Pt abundance. Regarding the CO peak
potentials, all the Pt samples exhibited a main oxidation peak value around 0.80–0.83 V vs. RHE, similar
to that observed for the commercial Pt/C. However, Pt/gCMK-3-R8-EG and Pt/gCMK-3-R8-FAM
displayed an additional peak at 0.71 V vs. RHE, indicating a higher ability of the materials supported
on gCMK-3-R8 to oxidize the adsorbed CO and thus suggesting an improved tolerance of these
materials toward this intermediate formed during the methanol oxidation [21].

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Figure 5. Cyclic voltammograms for the COad oxidation in 0.5 M H2 SO4 aqueous solution on the
Pt/gCMK-3 catalysts. Experiments were performed at 0.020 V s−1 and 20 ◦ C. CO adsorption potential:
0.2 V vs. reversible hydrogen electrode (RHE).

As expected, the Pt–Ru catalysts oxidized CO at lower potentials than the Pt catalysts (Figure 6)
due to the presence of Ru, which helps to oxidize the CO adsorbed on Pt [21]. These values were
between 0.43 and 0.47 V compared with 0.50 V detected for the commercial Pt–Ru/C from E-TEK.
For the Pt–Ru samples, the main oxidation peak was attained at around 0.54–0.57 V vs. RHE,
slightly lower than the commercial one (0.60 V). Only the Pt–Ru/gCMK-3-R8-FAM sample showed an
additional peak at lower potential (0.49 V), which could be interpreted as enrichment of Pt and Ru
oxides on the surface of the catalyst as this catalyst displayed the highest metallic Pt abundance and
also a high presence of Ru oxides [21,29]. This argument is also valid for explaining the positive CO

284

Surfaces 2019, 2

onset and peak potential values observed for the catalyst Pt–Ru/gCMK-3-R8-EG, which also exhibited
the lowest Ru oxides abundance values (see Table 2).


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Figure 6. Cyclic voltammograms for the COad oxidation in 0.5 M H2 SO4 aqueous solution on the
Pt–Ru/gCMK-3 catalysts. Experiments were performed at 0.020 V s−1 and 20 ◦ C. CO adsorption
potential: 0.2 V vs. RHE.

These results suggest a positive effect of the use of gCMK-3 carbons on the CO oxidation reaction,
which could be attributed to its enhanced diffusion through the gCMK-3 ordered mesoporous structure
and a good electronic transfer between the nanoparticles and the support, as already proposed by
Calvillo et al. for a Pt catalyst supported on gCMK-3. The mentioned authors reported a significant
decrease in these values compared with a commercial Pt/C catalyst from E-TEK, a result explained by
the changes in the electronic structure of the metal due to its interaction with the carbon support [6].
On the other hand, the graphitization of the carbon supports did not improve the catalytic activity
toward the CO electrochemical oxidation as more positive potential values for the oxidation of this
adsorbate were observed compared with those reported for catalysts supported on non-graphitized
CMK materials [30]. At the same time, this result matches those found in Reference [6], which were
attributed to a lack of surface oxygen groups, which play a key role in the oxidation of this adsorbate.
3.3. Methanol Oxidation
Figure 7 shows the behavior of the Pt catalysts supported on gCMK-3s toward the methanol
oxidation. The catalysts supported on gCMKs developed significantly higher current densities than the
commercial Pt/C. Moreover, the magnitude of the current densities were higher than those reported
by Salgado et al. for the oxidation of this fuel on non-graphitized CMK-3-supported Pt and Pt–Ru
catalysts [30]. Our results match with those reported in Reference [6], which can be explained by
the enhanced electrical conductivity of gCMK-3, which promotes the obtaining of increased current
densities for the methanol oxidation.
All the samples showed a methanol oxidation current hysteresis, which has recently been
explained by a change in the rate determining step, from water dissociation to the methanol
dehydrogenation step. In the forward scan, oxygenated species are adsorbed on the electrode,
changing the surface nature of the electrode and thus generating a peak in the backward scan at
different potentials attributed to the oxidation of fresh methanol adsorbed on the Pt surface [31,32].
Among the studied catalysts, only Pt/gCMK-R8-BM demonstrated a high oxophilicity as the peak
current in the forward scan, If , was higher than that of the backward scan, Ib [32], while the catalyst
supported on gCMK-3-R2 and Pt/gCMK-R8-EG and Pt/gCMK-R8-FAM developed a higher peak
current during the backward scan, indicating lower oxophilicity of these materials that improved
the adsorption of methanol on the oxidized Pt surface, allowing its oxidation on oxidized Pt active
sites [31].
285

Surfaces 2019, 2



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catalysts. Experiments were performed at 0.020 V s−1 and 20 ◦ C.

Similar results were observed for the Pt–Ru catalysts. The catalysts reduced with formic acid
showed higher current densities than the commercial Pt–Ru/C and the catalyst reduced with ethylene
glycol (see Figure 8). The best performance was obtained with Pt–Ru/gCMK-3-R8-FAM, which
can again be related to the high content of Ru oxides in these catalysts, which promote the CO
oxidation [21].
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Figure 8. Cyclic voltammograms in 2 M CH3 OH + 0.5 M H2 SO4 aqueous solution on the
Pt–Ru/gCMK-3 catalysts. Experiments were performed at 0.020 V s−1 and 20 ◦ C.

Figure 9a depicts the j–t curves for the methanol oxidation on the Pt and Pt–Ru catalysts supported
on graphitized ordered mesoporous carbons recorded at 0.6 V vs. RHE, a typical DMFC work potential.
In the case of the Pt catalysts, the commercial one presented the highest current densities, a fact
explained by an improved reoxidation of the intermediate species formed during the methanol
oxidation, which could be favored by the presence of the micropores of the carbon black, which
delays their diffusion and allows their reoxidation [33]. Nevertheless, the difference in the stationary
286

Surfaces 2019, 2

current densities among the commercial catalyst and the gCMK-3-supported catalyst with the highest
performance (Pt/gCMK-3-R8-EG) was around 11–12 μA cm−2 , low values that allowed us to conclude
that the use of Pt/gCMK-3 catalysts could still be an alternative as anodes in DMFCs. Among the
synthesized catalysts, no significant differences in the magnitude of the stationary current densities
were found, except in the case of the catalyst Pt/gCMK-3-R2-BM, which displayed the lowest current
density (0.0279 mA cm−2 ).
The behavior of the stationary current densities for the Pt–Ru catalysts was different to that
observed for the Pt materials, as can be seen in Figure 9b. The synthesized catalysts developed higher
current densities than those observed for the commercial Pt–Ru/C catalyst from E-TEK and the studied
Pt catalysts, making evident the effect of Ru and their surface oxides determined by XPS analysis on
the enhancement of the methanol oxidation stationary current densities obtained on these catalysts,
with Pt–Ru/gCMK-3-R8-FAM being the catalyst displaying the highest stationary current density, as
previously observed in the cyclic voltammetry experiments.
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Figure 9. j–t curves in 2 M CH3 OH + 0.5 M H2 SO4 aqueous solution at 0.60 V vs. RHE on the
Pt/gCMK-3 (a) and Pt–Ru/gCMK-3 (b) catalysts.

Considering the results presented here, it is possible to establish a comparison between the
performance of the catalysts presented in this work and those reported for materials supported on
ordered mesoporous carbons. In the case of Pt catalysts supported on ordered mesoporous carbons,
Su et al. prepared materials using sodium borohydride as reducing agent, obtaining catalysts that
achieved current densities close to 1.6 mA cm−2 , surpassing the values generated for a commercial
Pt/C catalyst from E-TEK [34]. The authors explained this high performance by the improved transport
of methanol toward the catalytic nanoparticles through the pores of the support. As mentioned above,
Salgado et al. [30] synthesized Pt catalysts supported on ordered mesoporous carbons, which displayed
CO oxidation peak potentials and methanol oxidation current densities comparable to those presented
in this work. The improved diffusion of the electroactive species through the carbon support was
stated as the main cause for the remarkable performance of these catalysts. It is important to highlight
that the results presented here are quite similar to those described in Reference [6], although in this
case, the Pt nanoparticles were supported on graphitized ordered mesoporous carbons. This fact
corroborates that the materials studied here possess a catalytic activity that follows the trend of the
catalysts already reported in the literature.
The behavior Pt–Ru catalysts supported on ordered mesoporous carbons toward the methanol
oxidation have also been studied. Salgado et al. [30] reported the activity of Pt–Ru catalysts supported
on non-graphitized ordered mesoporous carbons, which displayed lower methanol oxidation current
densities than those reported by Calvillo et al. [6] as well as that obtained in the present study. The
increased current densities obtained on graphitized CMK-3 can be explained by the high conductivity
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Surfaces 2019, 2

of these materials achieved after the heat treatment applied on them. Other OMCs have been
demonstrated to play a key role in the increase in the performance of the as-supported nanoparticles.
Maiyalagan et al. compared the activity of highly stable Pt–Ru nanoparticles supported on ordered
mesoporous carbons and carbon blacks, finding higher current densities for the methanol oxidation
on the OMC-supported materials [35]. Nitrogen doping of ordered mesoporous carbons has also
been suggested as an alternative to improve the performance of Pt–Ru catalysts, generating methanol
oxidation current densities higher than 6 mA cm−2 in acidic and alkaline media, which were attributed
to both the presence of nitrogen as dopant and the increase in the porosity of the carbon supports as
a consequence of the treatment employed to dope the ordered mesoporous carbons [36]. From the
comparisons presented here, it is possible to suggest that the Pt and Pt–Ru catalysts supported on
graphitized CMK-3 are suitable materials to be employed as anode catalysts in DMFCs. However,
further studies are necessary to fully understand and characterize the mechanism of methanol oxidation
on these catalysts and therefore finding the effect of carbon support and composition of the catalysts
on this reaction. Studies at different temperatures can be useful to determine the rate determining step
on the different catalysts, along with differential electrochemical mass spectrometry to identify the
onset potentials for both CO and methanol oxidations.
4. Conclusions
Pt and Pt–Ru catalysts supported on graphitized ordered mesoporous carbons (gCMK-3) were
synthesized using sodium borohydride, formic acid, and ethylene glycol as reducing agents. Metal
contents close to 20 wt% and Pt:Ru atomic ratios close to 1:1 were obtained with well-dispersed
metal nanoparticles on the carbon materials and a narrow particle size distribution around 5 nm.
The Pt/gCMK-3 catalysts displayed similarities in terms of the surface relative abundance for Pt (0),
Pt (II), and Pt (IV) oxidation states, with values near 60%, 20%, and 20%, respectively. On the other
hand, the Pt–Ru/gCMK-3 materials exhibited a higher abundance of surface Pt and Ru oxides. As
expected, the Pt–Ru materials were the most tolerant to CO as they were able to oxidize this adsorbate
at the lowest potentials, an effect attributed to the presence of Ru oxides present on the surface of
the catalyst. In agreement with the CO oxidation results, the Pt–Ru catalysts, in particular the one
supported on gCMK-3-R8, were the most active ones toward the methanol oxidation. Pt/gCMK-3
catalysts displayed oxidation current densities with some hysteresis, indicating the adsorption of
methanol on oxidized Pt to react on this surface. Moreover, the stationary current densities at 0.6 V
developed by the Pt–Ru/gCMK-3 catalysts were higher than those of the Pt materials, suggesting a
profitable effect of this metal and their oxides on these catalysts. The results found in this research
indicate that the studied catalysts can be good candidates to be used as anodes in direct methanol
fuel cells.
Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/2/1/1/s1,
Figure S1: Pt 4f XPS spectra for the synthesized Pt/gCMK-3 catalysts. Red line: Pt (0), blue line: Pt (II), green line:
Pt (IV), black dotted line: fitted curve, black full line: spectrum. Figure S2: Pt 4f XPS spectra for the synthesized
Pt–Ru/gCMK-3 catalysts. Red line: Pt (0), blue line: Pt (II), green line: Pt (IV), black dotted line: fitted curve, black
full line: spectrum. Figure S3: Ru 3p XPS spectra for the synthesized Pt–Ru/gCMK-3 catalysts. Red line: Ru (0),
blue line: Ru (IV), green line: Ru (IV) hydrated, black dotted line: fitted curve, black full line: spectrum.
Author Contributions: J.C.C.G. performed the experiments and wrote the paper. V.C. and L.C. synthesized the
carbon supports, prepared the catalysts, and performed the physical characterization (EDX, XRD, TEM). D.S.
performed the XPS experiments and analysis. R.M. and M.J.L.E. contributed reagents/materials/analysis tools.
All the authors conceived and designed the experiments and revised the manuscript.
Funding: This work was carried out with the support of the Conversion Group of Gobierno de Aragon T06
and FEDER.
Acknowledgments: V.C. kindly thanks the UK Catalysis Hub for resources and support provided via the
membership of the UK Catalysis Hub Consortium and funded by EPSRC (EPSRC grants EP/K014706/1 and
EP/K014714/1).
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Ir-Ni Bimetallic OER Catalysts Prepared by
Controlled Ni Electrodeposition on Irpoly and Ir(111)
Ebru Özer 1 , Ilya Sinev 2 , Andrea M. Mingers 3 , Jorge Araujo 1 , Thomas Kropp 4 ,
Manos Mavrikakis 4 , Karl J. J. Mayrhofer 3 , Beatriz Roldan Cuenya 2 and Peter Strasser 1,5, *
1

2
3
4
5

*

The Electrochemical Energy, Catalysis and Materials Science Laboratory, Department of Chemistry,
Technical University Berlin, Straße des 17. Juni 124, Berlin 10623, Germany;
ebru.oezer@campus.tu-berlin.de (E.Ö.); ferreiradearaujo@campus.tu-berlin.de (J.A.)
Fritz Haber Institute of the Max Planck Society, Department of Interface Science, Faradayweg 4-6,
Berlin 14195, Germany; ilya.sinev@rub.de (I.S.); roldan@fhi-berlin.mpg.de (B.R.C.)
Max-Planck-Institut für Eisenforschung GmbH, Max-Planck-Strasse 1, Düsseldorf 40237, Germany;
mingers@mpie.de (A.M.M.); k.mayrhofer@fz-juelich.de (K.J.J.M.)
University of Wisconsin-Madison, 1415 Engineering Drive, Madison WI 53706, USA;
tekropp@wisc.edu (T.K.); emavrikakis@wisc.edu (M.M.)
Ertl Center for Electrochemistry and Catalysis, Gwangju Institute of Science and Technology,
Gwangju 500-712, Korea
Correspondence: pstrasser@tu-berlin.de; Tel.: +49-030-314-29542

Received: 26 November 2018; Accepted: 13 December 2018; Published: 14 December 2018

Abstract: The alteration of electrocatalytic surfaces with adatoms lead to structural and electronic
modifications promoting adsorption, desorption, and reactive processes. This study explores the
potentiostatic electrodeposition process of Ni onto polycrystalline Ir (Irpoly ) and assesses the
electrocatalytic properties of the resulting bimetallic surfaces. The electrodeposition resulted in
bimetallic Ni overlayer (OL) structures and in combination with controlled thermal post-deposition
annealing in bimetallic near-surface alloys (NSA). The catalytic oxygen evolution reaction (OER)
activity of these two different Ni-modified catalysts is assessed and compared to a pristine,
unmodified Irpoly . An overlayer of Ni on Irpoly showed superior performance in both acidic and
alkaline milieu. The reductive annealing of the OL produced a NSA of Ni, which demonstrated
enhanced stability in an acidic environment. The remarkable activity and stability improvement of Ir
by Ni modification makes both systems efficient electrocatalysts for water oxidation. The roughness
factor of Irpoly is also reported. With the amount of deposited Ni determined by inductively coupled
plasma mass spectrometry (ICP-MS) and a degree of coverage (monolayer) in the dependence of
deposition potential is established. The density functional theory (DFT) assisted evaluation of H
adsorption on Irpoly enables determination of the preferred Ni deposition sites on the three low-index
surfaces (111), (110), and (100).
Keywords: oxygen evolution reaction; iridium; nickel; electrodeposition; model catalyst; water
oxidation; CO oxidation; DFT; hydrogen adsorption

1. Introduction
Iridium is one of the most important electrode materials for basic surface science and industrial
applications especially those in the field of electrochemistry due to its enhanced catalytic properties
in water splitting electrolyzers [1,2]. The efficiency of electrolyzers and, thus, the catalytic properties
are highly dependent on the interface between the electrode surface and the electrolyte. However,
the main impediments to an effective application of water electrolysis as a sustainable method for
hydrogen production are the high potential losses at the anode side and the high Ir material costs.
Surfaces 2018, 1, 165–186; doi:10.3390/surfaces1010013

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Surfaces 2018, 1

In this context, studies on low index surfaces of single crystals offer a meaningful opportunity for
a fundamental understanding and, based on it, the rational design of enhanced electrocatalytic
materials. A wealth of information on structural and electronic effects can emerge from these studies
particularly in the context of electrochemical deposition of metal adsorbates at a sub-monolayer
coverage including underpotential deposition (UPD). Potential-controlled deposition results in the
growth of a new two-dimensional phase, which is believed to have electronic properties that differ
significantly from those of the respective bulk materials. In addition, the electronic structure of the
substrate itself and, thus, the interfacial reactivity is altered by the metal adatoms [3,4]. The over-layer
and near-surface alloy systems, therefore, offer a great variability in optimizing the binding energies to
reaction intermediates and adsorption properties [5–9]. The deposition of metal adatoms on foreign
metal substrates provides attractive systems with highly controllable surface coverages by using the
electrode potential and deposition time. In particular, the resulting structure and surface properties
of the new electrochemical materials are dependent on the variable surface coverage and, thereby,
allow for a precise and reproducible control of the coverage-dependent properties. Due to its practical
importance in surface chemistry and electrocatalysis, a number of ex situ and in situ techniques such
as cyclic voltammetry (CV), scanning tunneling microscopy (STM), X-ray photoelectron spectroscopy
(XPS), and low-energy electron diffraction (LEED) have been applied to augment the process of
deposition at the molecular and atomic level [10].
In this study, the deposition of Ni on polycrystalline Ir and well-characterized Ir(111) electrodes
in both sulfate-containing electrolytes and phosphate-buffered solutions will be presented employing
experimental techniques. In situ cyclic current-potential curves are used to detect the redox potentials
of the electroactive species or rather the potential range in which deposition takes place. Ex situ XPS
data provides information on the chemical composition of the surface. In addition, a more quantitative
trace metal analysis was performed with inductively coupled plasma-mass spectrometry (ICP-MS)
to determine the adsorbed metal concentration based on high sensitivity and detection capability.
Surface sensitive CO stripping measurements provide valuable insights into the morphological and
structural composition of the deposit. Theoretically calculated hydrogen binding energies were taken
into consideration. The Faraday charge efficiencies (FE) towards oxygen evolution of the modified
electrodes were measured with differential electrochemical mass spectrometry (DEMS) and it was
confirmed that the entire charge is transferred to facilitate the oxygen evolution reaction (OER).
In this case, a majority of the work has been carried out on polycrystalline Ir electrodes to study the
multi-faceted Ni deposition process and to establish profound procedures and protocols for the precise
control of surface coverage and structural composition. The intent is to present a broad overview and
to provide detailed observations of the deposition process of Ni on Ir using not only electrochemical
methods but also knowledge obtained from various surface sensitive techniques. The following studies
and conclusions, therefore, present the first step in the electrodeposition of Ni layers on Ir electrodes
and their electrochemical properties in the oxygen evolution reaction (OER). The results of this research
expand the fundamental knowledge about the deposition process on polycrystalline Ir and show initial
studies on Ir(111) single crystal electrodes under well-defined conditions of cleanliness thanks to a
custom-made inductive and electrochemical crystal chamber [11].
2. Materials and Methods
2.1. Substrates
Ir(111) Single Crystal and Polycrystalline Ir (Irpoly )
Ir(111) single crystal electrodes with a diameter of 10 mm and a precision of orientation of
≤0.1◦ (MaTecK GmbH, Jülich, Germany) were applied as working electrodes for the study of surface
modification. To mount the cylinder into the inductive and electrochemical setup in which the surface
was refined thermally, two lateral holes with a diameter of 0.5 mm were machined into the shell surface
through which loops made of Ir wire (∅ = 0.25 mm, 99.9%; Alfa Aesar, Germany) run. The detailed
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Surfaces 2018, 1

description of mounting can be found elsewhere [11]. A heating protocol consisting of several heating
cycles per heating phase is applied until the determined electrochemical surface characterization is
consistent with a perfectly clean and well-ordered (111) single-crystalline Ir surface [12].
2.2. Electrochemical Characterization
2.2.1. Ni Sweep Voltammetry
The electrodeposition of Ni on polycrystalline Ir was conducted in a standard rotating disk
electrode (RDE) setup. Measurements on Ir(111) single crystals were performed in a custom-made
setup [11].
Initially, the Ni electrodeposition was studied in 0.05 M H2 SO4 (diluted from 95% H2 SO4 ,
Suprapure, Carl Roth, Germany) and in a phosphate buffer solution of pH 6.85 (0.1 M KH2 PO4 and
0.1 M K2 HPO4 , 1:1 mixture by volume, Merck, Germany) using the cyclic voltammetry (CV) technique.
The electrodes were transferred into the degassed solution containing 2 mM NiSO4 . A voltammetric
scan was then performed from the immersion potential of 1.0 VRHE for Irpoly and 0.8 VRHE for Ir(111) to
the negative potential limit of 0.05 VRHE at a scan rate of 50 mV s−1 . Under the same conditions, CVs of
the electrodes were recorded in the absence of Ni2+ ions to locate the Ni adsorption-desorption region.
These solutions contained 2 mM Na2 SO4 to have the same anion concentration. To trace conductivity
changes of the electrode, impedance spectroscopy was performed to determine the ohmic drop of
the electrode. In both electrochemical measurements, the voltammetric scan ended at the positive
potential limit to ensure the oxidation of all deposited Ni. All electrolytes were purged with either Ar
(Ir(111) in the inductive setup) or N2 (Irpoly in the RDE) for at least 15 min to remove dissolved oxygen
prior to use.
2.2.2. Controlled-Potential Technique
The potential range of Ni deposition on Irpoly in phosphate buffer was examined more thoroughly.
A dependence of the deposited Ni amount on Irpoly as a function of deposition potential was established.
The total metal content of the deposition process was determined at six potentials between 0.05 and
0.5 VRHE using ICP-MS. The given potential window was determined in the Ni sweep voltammetry
experiments. The electrodes were immersed under a potential control at the deposition potential
and the respective potential was held for a time of 10 min. Subsequently, the coverage of Ni was
determined by ICP-MS after dissolving all Ni by immersing the electrode in 5 mL 0.1 M HNO3 for one
hour. Beforehand, the electrodes were rinsed thoroughly with milli-q water.
2.2.3. Preparation of Ni Overlayers (OLs) and Ni (near-) Surface Alloys (NSAs)
To produce an overlayer of Ni (Ni/Ir OL), the electrode was immersed at the desired deposition
potential (Edep ) into a phosphate buffer solution of pH 6.85 (0.1 M KH2 PO4 , 0.1 M K2 HPO4 ) containing
2 mM NiSO4 and held for 10 min. The Ni/Irpoly OL electrode catalyst for further electrochemical
testing was produced at a deposition potential of 0.05 VRHE .
For the preparation of a Ni/Irpoly near surface alloy (NSA), the Ni/Irpoly OL electrode
(Edep = 0.05 VRHE ) was subsequently annealed for 1 min in 4% H2 /Ar at 400 ◦ C in a custom-made
inductive setup [11]. The Ni-modified catalysts were then tested for OER activity in 0.1 M KOH and
0.05 M H2 SO4 .
2.2.4. Electrochemical OER Protocol at RDE
Two types of OER protocols were used to test the performance of unmodified and Ni-modified
Irpoly electrodes at the RDE. First, the OER activities of Irpoly , Ni/Irpoly OL, and Ni/Irpoly NSA were
tested in both 0.05 M H2 SO4 and 01 M KOH. In a separate protocol, the stability of the Ni/Irpoly OL
catalyst was subject to greater stress. The catalyst was first treated in 0.05 M H2 SO4 before continuing

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Surfaces 2018, 1

the OER measurement in 0.1 M KOH. In between measurements, the electrode was rinsed thoroughly
with milli-q water.
The electrolytes were prepared by dilution of 95% H2 SO4 (Suprapure, Carl Roth, Germany) with
ultrapure water. Purified Fe-free 0.1 M KOH was prepared according to the method reported by
Trotochaud et al. [13].
For both procedures, the immersion of the electrodes at 1.00 VRHE was followed by 10 polarization
curves into the potential region of OER at a scan rate of 5 mV s−1 from the immersion potential up
until a current of 0.785 mA (current density of 1 mA cm−2 ) was reached. The measured currents
are normalized to the geometric area of the electrodes to provide comparable current densities (jgeo ).
The geometric surface area of the electrodes is 0.785 cm2 corresponding to the diameter of 10 mm.
After 10 consecutive cycles, electrochemical impedance spectroscopy (PEIS) was measured to obtain
the ohmic resistance. The respective potentials are referred to as EiR .
2.2.5. CO Stripping Protocol
The CO stripping measurements were carried out in a rotating disk (RDE) electrode setup with
a reversible hydrogen electrode (RHE) as a reference electrode and a Pt-mesh as a counter electrode.
The measurements were recorded with a Bio-Logic SP-200 potentiostat (Bio-Logic, France). The CO
stripping measurements were performed on an unmodified Irpoly and two Ni-modified Irpoly electrodes
in purified Fe-free 0.1 M KOH (semiconductor grade, 99.99% trace metals basis, Merck, Germany).
Ni/Irpoly OL and Ni/Irpoly NSA electrodes were prepared as described in Section 2.2.2.
2.3. Theoretical Characterization
Density Functional Theory (DFT) Calculated H-Ir Binding Energies
Theoretical H-Ir binding energies were calculated using the projector augmented wave
(PAW) method [14,15] as implemented in the Vienna ab initio simulation package (VASP) [16,17].
Exchange-correlation energies were obtained by using the functional by Perdew, Burke, and Ernzerhof
(PBE) [18]. A plane wave cutoff of 600 eV was used. Structure optimizations were performed until total
energies were converged to 10−6 eV and forces acting on the relaxed ions were below 0.02 eV Å−1 .
Surface models were obtained by cutting bulk iridium (lattice constant a = 3.877 Å) in (100), (110),
and (111) orientation. The resulting slab models consist of five atomic layers separated by a 12 Å
vacuum layer. The lowest two atomic layers are frozen to simulate the bulk. For Ir(100) and Ir(111),
the Brillouin zone was sampled at 6 × 6 × 1 k points while 4 × 4 × 1 k points were used for the larger
Ir(110) model.
2.4. Physicochemical Characterization
2.4.1. Inductively Coupled Plasma Mass Spectrometry (ICP-MS)
ICP-MS of the HNO3 aliquots taken to dissolve the electrodeposited nickel was measured with a
NexION 300 X (PerkinElmer® , Shelton, CT, USA) equipped with a concentric PTFE Nebulizer-PFA-ST
(PerkinElmer® , Shelton, CT, USA). and a spray chamber.
2.4.2. Atomic Force Microscopy (AFM)
The AFM images were done in a tapping mode using a Bruker MultiMode 8 microscope (Bruker,
Berlin, Germany) on the samples after transport in Ar.
2.4.3. X-ray Photoelectron Spectroscopy (XPS)
The pristine and Ni-modified Ir(111) surfaces were analyzed by X-ray photoelectron spectroscopy
in an ultra-high vacuum (UHV) setup equipped with Phoibos 150-MCD 9 hemispherical analyzer
(SPECS GmbH, Berlin, Germany). A monochromatic Al Kα X-ray source (hν = 1486.5 eV) at 12 kV

294

Surfaces 2018, 1

of anodic voltage and 300 W power was used to perform the characterization. For calibration
of the binding energy, the C1s signal of graphitic-like carbon (285 eV) was used. For peak
deconvolution, Gaussian-Lorentzian product functions and a Shirley background were applied in the
Casa XPS software.
2.5. Mass Spectrometry
Differential Electrochemical Mass Spectrometry (DEMS)
Differential electrochemical mass spectrometry (DEMS) was acquired in a custom-made dual
thin-layer electrochemical flow cell based on the concept reported elsewhere [19,20]. A Prisma
quadrupole mass spectrometer (QMS 200, Pfeiffer-Vacuum, Germany) equipped with two
turbomolecular pumps (HiPace 80, Pfeiffer, Germany) operating under 10−6 mbar were used to
detect any volatile products. The instrument allows a continuous flow of the electrolyte of 10 μL s−1
through the hydrophobic membrane to the mass spectrum ion source. We used the hydrophobic PTFE
membrane with a pore size of 30 nm and a thickness of 150 μm (Cobetter, Cat. No. PF-003HS) as the
separation interface for the electrolyte and vacuum. A high mass transport rate is guaranteed due to
the low volume of electrolyte flowing over the electrode surface as well a stable ionic current for the
electrochemical process.
The electrochemistry was controlled using a BioLogic potentiostat, a leak-free Ag/AgCl as
reference electrode (Warner Instruments, USA), and a Pt mesh as a counter electrode. Cyclic
voltammograms in 0.1 M KOH and 0.05 M H2 SO4 were recorded in a potential range of 0.6 and 1.6 VRHE
(prior to iR-compensation) at a scan rate of 50 mV s−1 . The flow cell setup was used for evaluating the
faradaic charge efficiencies toward oxygen evolution for unmodified and Ni-modified polycrystalline Ir
electrodes. The preparation of the Ni-modified electrodes is described above. In the first step, the ionic
current baseline was recorded in a steady-state measurement to obtain a catalyst-specific calibration
factor K* . A constant purge of N2 gas over the electrodes’ surface was done during integration of the
electrode with the flow cell system. In addition, during DEMS experiments, the purge N2 was kept
over the non-electrolyte contact areas of the electrode to ensure that the electrochemistry is running in
an oxygen-free environment. DEMS measurements were monitorized with control product formation
species with m/z values of 2, 16, 18, 28, 30, 36, 32, 34, 44 during the CV scan. The previously mentioned
catalyst-specific calibration factor Kj * was applied for quantification of the faradaic efficiency (FE)
of O2 (m/z 32). Chronopotentiometric steps were applied in the linear region of the ionic current
baseline during the steady-state measurement protocol. In Equation (1), the relation between the mass
spectrometric ion current (iMS,j ) and the faradaic current (iF ) to the volatile product j is given as:
Kj∗ =

iMS,j ·zj
.
iF

(1)

In this case, zj is the number of transferred electrons per molecule of volatile product j. The division
of the integrated mass spectrometric ion current (QMS ,j ) by the integrated capacitive current-corrected,
faradaic current (QF ) allows for the determination of the faradaic efficiency (FE) (charge selectivity)
according to:
QDEMS
QMS,j ·zj
F,j
FE(%) =
·100 =
·100.
(2)
∗
Q F · Kj
QF
In the analysis, only anodic faradaic currents are taken into account since solely the anodic
processes of molecular oxygen evolution is of interest. The limits of integration were given by the
detection signal of oxygen evolution (m/z 32) and the initial baseline level. For the different samples,
the limits varied due to distinct onset potentials for oxygen evolution.

295

Surfaces 2018, 1

3. Results and Discussion
3.1. Roughness Factor (RF) of Irpoly
For the subsequent evaluation of the accurate amount of deposited Ni on Irpoly (which will be
transferred into fractions of a single monolayer), the following calculations aim at determining the
surface roughness of a polycrystalline Ir electrode. Equation (3) establishes a relationship between the
double-layer charging capacitive current density jc and the double-layer capacitance Cdl (per geometric
area) [21].
jc = ν·Cdl .
(3)
For Irpoly , the potential window between 0.4 and 0.5 VRHE lacks any faradaic process (see Figure S1)
and was, therefore, chosen to be the electrochemical double-layer capacitance. The roughness factor
(RF) of Irpoly is then calculated from the double-layer capacitance according to the Equation (4) below.
RF =

Cdl
.
CS

(4)

CS in Equation (4) is the electrochemical double-layer capacitance of a smooth and planar electrode
surface of the same material measured under the same conditions. These measurements have been
conducted for a single crystal Ir(111) electrode. The double-layer capacitance measurements for Irpoly
and Ir(111) are depicted in Figure 1. Given the slopes of both electrode surfaces, a roughness factor of
4.3 of the polycrystalline Ir electrode can be defined according to the following calculation.
RF =

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Figure 1. Double-layer capacitance measurements of (a) Ir(111) and (b) Irpoly displaying the anodic
charging current densities measured at 0.42 VRHE .

For calculating the mass of Ni needed to produce a monolayer on a polycrystalline Ir cylinder
with a diameter of 10 mm, it is legitimate to assume that, on an Ir polycrystal, only a (111) face exists
as the terraces of a polycrystal, which are presumed to consist of the three low-index surfaces (111),
(110), and (100) nearly to the same extent.
In Table 1, the individual steps of our estimation are summarized. In the first step, the atoms
per unit area are determined by dividing the number of Ir atoms per (111) face in the unit cell by the
surface area of the (111) face. By using the Avogadro constant NA , the amount of substance per surface
area is introduced. Given a geometric projected surface area of 0.785 cm2 , the molar or atomic amount
of Ni needed to form a complete monolayer is obtained. Analogously, the mass of Ni per monolayer
can be obtained by introducing the molar mass of Ni (M = 58.68 g mol−1 ). With a given RF of 4.33,
the estimated Ni amount to produce a monolayer on Irpoly is 0.52 μg.
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Surfaces 2018, 1

Table 1. Calculations for determining the mass needed to produce a monolayer of Ni on a polycrystalline
Ir cylinder with a diameter of 0.785 cm2 . (Lattice parameter Ir: a = 3.8394 Å, according to PDF:
00-006-0598).
Steps
surface area calculation
surface area/10−15 cm2
atoms per (111) face in unit cell
atoms per unit area/1015 cm−2
amount of substance per surface area/nmol cm−2
amount of substance for monolayer/nmol ML−1
mass of Ni per monolayer/μg ML−1
mass of Ni per monolayer for Irpoly /μg ML−1

Calculation
√
3/4· a2
1.277
2
1.566
2.6
2.04
0.12
0.52

3.2. Surface Chemical Characterization
Ni deposition was conducted and analyzed on polycrystalline Ir (Irpoly ) and single-crystalline
Ir(111) electrodes. The experiments with Irpoly have served to gain a better understanding of the
deposition process and for determining the influence of adsorbed Ni on the surface and the resulting
catalytic properties.
3.2.1. Ni Deposition on Irpoly and Ir(111) by Sweep Voltammetry
The Ni deposition process on a polycrystalline Ir electrode was first studied by using the cyclic
voltammetry technique in order to determine the potential area in which Ni deposition occurs.
Figure 2a depicts the cyclic voltammogram of a polished and electrochemically cleaned polycrystalline
Ir cylinder in a phosphate buffer solution with and without 2 mM NiSO4 . The solution without Ni2+
contains 2 mM Na2 SO4 to guarantee the same concentration of anions in the electrolyte and to make
the measurements comparable.

a



b

Irpoly







Ir(111)








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Figure 2. Cyclic voltammograms of (a) a polycrystalline Ir electrode and (b) a pristine Ir(111) in a
phosphate-buffered solution (pH = 6.85) containing either 2 mM Na2 SO4 (solid lines) or 2 mM NiSO4
(dashed lines) at a scan rate of ν = 0.05 V s−1 . The electrodes were immersed at a potential of 1.0 VRHE
and 0.8 VRHE for Irpoly and Ir(111), respectively.

In the presence of Ni, the hydrogen adsorption-desorption region is reduced significantly, which
suggests a blocking of Ir sites by Ni ad-atoms. By contrast, in the negative going sweep, there is only
a slight distinction in the total electric charge between the cyclic voltammograms in the potential
area of 0.5 to 0.0 VRHE , which implies that Ni electrodeposition on Ir is not recognizable as strongly
defined deposition peaks. It cannot be determined with absolute certainty if oxygen desorption in
the Ni-free solution (at ~+0.7 VRHE ) is shifted toward lower potentials (at ~+0.65 VRHE ) due to the
so-called ligand effect [22], which is triggered by adsorbed Ni adatoms or if deposition of Ni occurs on
297

Surfaces 2018, 1

distinct adsorption sites, which are merged into one broad redox wave. In the positive going sweep,
desorption of Ni ad-atoms begins beyond the hydrogen underpotential deposition (HUPD) region and
attains its maximum value at around +0.65 VRHE .
Figure 2b displays the cyclic voltammogram of a pristine Ir(111) single crystal electrode in
phosphate buffer of pH 6.85 containing 2 mM NiSO4 . A cyclic voltammogram of the same electrode in
the absence of Ni is also shown for comparison. Both voltammograms were recorded in the potential
window of 0.05–0.8 VRHE , which is a potential range adequate for a sufficient characterization of the
Ir(111) surface [12]. The CV of Ir(111) in phosphate buffer without Ni is mostly featureless and only
displays the HUPD redox peak at ~+0.1 VRHE . The significant difference between both voltammograms
is located in the HUPD region. The characteristic HUPD peak of the (111) face is mostly masked in the
presence of Ni. The voltammogram is presumably dominated by a broad deposition and stripping
region of deposited Ni between 0.50–0.05 VRHE and 0.3–0.8 VRHE , respectively. The peak in the positive
going sweep at ~+0.1 VRHE can either be ascribed to stripping off Ni or to a hydrogen desorption
charge due to remaining hydrogen adsorption on Ir(111) sites or both. The latter case presupposes
that not all Ir sites are being blocked by Ni adatoms during deposition. This might be a kinetically
controlled problem if the deposition process by cyclic voltammetry requires some time to form a
complete Ni layer. Free Ir sites may be explained by the deposition of Ni on existing Ni adatoms.
In the negative going sweep, the charge is found to be higher (≈375.9 μC cm−2 ) than the sum of
the charges in the positive going sweep. From this observation, it seems that Ni deposition at a single
crystalline Ir(111) is not entirely reversible, which means that Ni is not completely stripped off the
surface in the anodic scan up to 0.8 VRHE .
3.2.2. Ni Deposition on Irpoly by the Controlled-Potential Technique
The formation of a Ni layer is likely a process that requires some time since the cyclic
voltammogram in Figure 2a does not display defined Ni redox peaks. Hence, the cycling technique
may not be the most suitable approach. On this account, the potential range Edep in Figure 2a was used
to establish a relationship between Ni coverage and the applied deposition potential (Edep ). Therefore,
the electrode was immersed at the desired potential under the potential control and the potential was
held for 10 min. The deposited amount of Ni was determined chemically by ICP-MS. The process
was conducted for six distinct deposition potentials. In Figure 3, the effect of the deposition potential
Edep on the amount of adsorbed Ni is shown. The obtained values from ICP-MS and the resultant
extent of Ni monolayer formation on polycrystalline Ir are presented in Table S1 As described earlier,
the analyzed amount of Ni with ICP-MS was translated into a degree of coverage (ML) by assuming
that (111) facets are prevalent on a polycrystalline Ir surface.


























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electrode and deposition potential Edep . The electrode was held for 10 min at each potential in
a phosphate-buffered solution (pH = 6.85) containing 2 mM NiSO4 . The amount of deposited
Ni was analyzed by ICP-MS. The right axis (green) gives the determined value of the monolayer
(ML) formation.
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Surfaces 2018, 1

At first, a slight increase in Ni coverage and a decreasing deposition potential is observed.
The linear rise becomes larger until the Ni amount approaches a limiting value at a potential of ≤+0.1
VRHE . Since the surface of a polycrystalline Ir electrode is not well-defined, the calculated values of Ni
ML coverage provide approximate guidance for the relationship between absolute molar Ni amount
and deposition potential. The assumption of a prevalent (111) facet on Irpoly is thereby the most severe
assumption as a polycrystalline surface generally consists of a number of different types of facets.
Some of these orientations may be less suitable or have less potential for Ni deposition. However,
with an estimated average roughness factor of 4.33 for the polycrystalline Ir electrode, a maximum
coverage of ~0.36 ML was obtained.
3.2.3. Surface Chemical State of Ni/Ir OL
The X-ray photoemission spectra of the Ni-modified single and polycrystalline Ir electrodes are
presented in Figure 4.

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Figure 4. XP spectra of (a) Ni-modified Ir(111) single crystal (Ni/Ir(111) OL) and (b) a Ni-modified Ir
polycrystal (Ni/Irpoly OLt=20min ). (c) XP spectra from the Ir 4f core level of Ir(111), Ni/Ir(111) OL, and
Ni/Irpoly OLt=20min . The deposition potential was 0.05 VRHE . The deposition time for the formation of
the over layer on Irpoly was twice as long as for Ir(111).
299

Surfaces 2018, 1

The deposition was performed at a potential of 0.05 VRHE to form an over layer (OL). The Ni
over layer on polycrystalline Ir (Ni/Irpoly OLt=20min ) was formed by deposition for 20 min, which is
twice as long as for the formation of an OL on Ir(111). For comparison, the XP core level spectra of
the well-defined and unchanged Ir(111) single crystal are given in Figure S2. In this case, the survey
spectrum over the entire energy range (Figure S2a, top view) is characterized by intense photoemission
peaks distinctive for Ir and weak O1s and C1s lines. The high-resolution spectra in Figure S2b (4 panes)
excludes the presence of Ni in the sample. During transfer of the single crystal into the UHV chamber
and until vacuum conditions were reached, the sample was shortly exposed to a laboratory atmosphere.
The detected carbon and oxygen signals can, therefore, be assigned to chemisorbed species during the
loading process. After extended exposure time, a Ni-free near the surface region could be affirmed.
The binding energy and line shape of Ir(111) and Irpoly in the Ni-modified samples in Figure 4c
are in accordance with the pure Ir spectrum. Thus, the deposition process does not affect the Ir state.
This observation is, moreover, consistent with images of the topography of the Ni-modified Ir(111)
surface analyzed by AFM. The AFM images and their respective height profiles are presented in Figure
S3 with increasing magnifications from Figure S3a–c. The Ni/Ir(111) OL surface was found to exhibit
well-resolved flat terraces vertically separated by ~1.53 Å tall steps, which is coherent with single or
double atomic steps. The Ni 2p region of all prepared samples in Figure 4a,b indicate the prevalence of
Ni2+ species with their main peaks at ~857 and 874 eV and their typical shake-up satellite peaks at
~863 and 880 eV. The intensity of the Ni lines increase for longer deposition times while the underlying
Ir lines decrease in intensity (see Figure 4c), which implies higher Ni coverages. In Ni/Irpoly OLt=20min ,
the Ni content determined by XPS was three times greater than on the Ni-modified single crystalline Ir
sample (Ni/Ir(111) OL). The bright spots in the AFM images correspond with very high probability to
contamination by particles of dirt since these measurements have been performed in air.
3.3. Structure-Activity Correlations
3.3.1. Oxygen Evolution Reaction (OER) and Faraday Efficiencies (FE)
The electrochemical OER activity of Ni-modified Irpoly electrodes have been tested using the
rotating disk electrode (RDE) technique in both acidic and alkaline environments. The respective OER
polarization curves are depicted for the first cycle in Figure 5a,c and are compared to the pristine and
unmodified Irpoly catalyst. The 10th OER cycle of the electrode materials is depicted (Figure S4a,b) to
monitor the stability of the catalysts. The Ni/Irpoly NSA sample was formed by reductive annealing of
Ni/Irpoly OL for 1 min and shows a very weak pronounced Ni(OH)2 /NiOOH redox peak in 0.1 M
KOH in the 10th cycle (see Figure S4b). The reductive treatment of the over layer for 2 min was
sufficient to produce a near-surface alloy of Ni in the subsurface region [11]. Thus, a reduced time was
chosen to design a surface in which Ni is both at the surface and in the subsurface of the catalyst.

300

Surfaces 2018, 1


Figure 5. Electrocatalytic OER activities of Ni-modified Ir electrodes compared to pristine
polycrystalline Ir at scan rates of 5 mV s−1 . (a) 1st OER cycle in 0.05 M H2 SO4 and (b) corresponding
Tafel plots (with ohmic-drop correction) of Irpoly , Ni/Irpoly OL, and Ni/Irpoly NSA. (c) 1st OER cycle
in 0.1 M KOH and (d) corresponding Tafel plots of all catalysts. (e) Current densities of electrodes
measured at 1.53 VRHE normalized to the electrode’s geometric surface area of 0.785 cm−2 for the 1st
and 10th cycle. The 10th cycle is depicted in Figure S4 in the Supplementary Information.

As can be seen from Figure 5a,c, the geometric OER activity of the Irpoly reference catalyst
was improved by the addition of Ni in the form of an over layer (OL) in both electrolyte systems.
The superior catalytic performance of Ni/Irpoly OL in an acidic solution is not only displayed via
the higher current density of the cyclic voltammogram (CV) curves but is also reflected by its lower
Tafel slope at higher potentials of ~52 mV dec−1 (Figure 5b), which is significantly smaller than the
slope of the unmodified (~70 mV dec−1 ) and NSA (~65 mV dec−1 ) electrocatalysts. Ni atoms in the
subsurface layer (NSA) formed by the annealing of the OL do not seem to have an enhanced effect
301

Surfaces 2018, 1

on the OER activity in acidic and alkaline media. In acidic solution, the OER performance of the
NSA is almost identical to the unmodified Irpoly . These observations indicate that the performance
of the modified material is strongly influenced by the relative position of the second component in
the electrode material. In order to correlate the observed reactivity with the structure of the catalyst,
the surface composition of a polycrystalline surface has to be considered in detail. Given the high
roughness factor of the polycrystalline Ir surface (due to the grain boundaries), the Ni/Irpoly OL might
have a similar surface structure to an Ir-Ni surface alloy in which Ni atoms are adjacent to Ir surface
atoms. The result is an optimization of the binding energy towards critical reaction intermediates of
the proposed oxygen evolution mechanism such as O, OH, and OOH.
In an alkaline environment, the mechanism of the OER has been proposed to occur via the
following steps [23–26]: (M represents the catalytic active site).
M + OH− → M − OH + e−

(6)

M − OH + OH− → M − O + H2 O + e−
−

M − O + OH → M − OOH + e
−

M − OOH + OH → M − OO + H2 O + e
M − OO → O2 .

(7)

−

(8)
−

(9)
(10)

Annealing the OL system in a reductive atmosphere significantly reduces the amount of Ni
adatoms at the surface. Since the characteristic Ni(OH)2 /NiOOH redox peak can be recognized
very weakly in the cyclic voltammogram of the 10th curve in 0.1 M KOH (see Figure 5c), Ni is still
present at the surface albeit to a considerably smaller extent. The electronic effect of underlying Ni
atoms on surface Ir atoms seems to have a slightly negative impact on the binding energy of reaction
intermediates to Ir. However, the activity behavior of the three examined catalysts is somehow similar
in KOH (see Figure 5c) whereby the pristine Irpoly is slightly more active than the NSA. The Tafel
slopes of the electrocatalysts differ significantly in both examined electrolyte solutions. In acidic
solution, the OL system exhibits the lowest slope followed by the NSA and the unmodified catalyst.
Under alkaline conditions, the order of the Tafel slopes is as follows: NSA > Irpoly > OL. The analysis
of the experimentally observed Tafel slopes is a powerful tool for determining the reaction steps
that are rate-limiting at the reaction interface. It must be kept in mind that these parameters are
greatly influenced by the concentration of the surface species formed during the reaction. In an acidic
medium, the Tafel slopes of the catalysts differ far more than in alkaline milieu, which suggests a
greater influence of Ni on the reaction mechanism. In 0.1 M KOH, the Tafel slopes lie in the range of
~60 mV dec−1 . This predicts a slow chemical step to be rate-limiting [27]. The above presented reaction
mechanism (6)–(10) proposes electron-transfer steps that determine the overall kinetics. However,
the reaction step in Equation (7) can also be divided into an electron transfer process subsequent to a
proton transfer step as:
(11)
M − OH + OH− → M − O− + H2 O
M − O− → M − O + e− .

(12)

Thus, if the rate-determining step involves the purely chemical step in Equation (11), a Tafel slope
of 60 mV dec−1 is predicted.
The electrode processes were also measured in a dual-thin layer flow cell setup in combination
with differential electrochemical mass spectroscopy (DEMS). The respective Faraday efficiencies (FE)
and faradaic current densities are given in Figure S5. In this case, the same order of activity performance
was measured in the alkaline solution (OL > Irpoly > NSA). The superior OER performance of the
Ni/Irpoly OL catalyst can be deduced to surface enrichment of Ir atoms promoting positive synergies
between Ir and Ni atoms. The increase in activity with Ir-Ni binary systems has been observed in
various studies in alkaline and acidic electrolytes [23,28–30]. The higher activity of Ni/Irpoly OL in
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Surfaces 2018, 1

acidic electrolyte may be assigned to an increase of surface OH groups initiated by leaching off Ni
adatoms from the surface [23]. At these high potentials, partial oxidation of the catalyst surface cannot
be precluded. The oxygen atoms, which were previously bound to Ni, will be protonated to become
electroneutral. As OH is considered to act as a crucial intermediate in the oxygen evolution mechanism,
a higher fraction of OH at the surface will have an advantageous effect on the reactivity.
As can be seen from Figure 5e, all Ni-modified catalyst systems display a higher OER reactivity
in 0.1 M KOH compared to the acidic milieu, which might be partly attributed to the already higher
activity of polycrystalline Ir in alkaline media when compared to the acidic substance. During repetitive
cycling, the OER performance is declining for all catalysts in both systems (compare Figure S4a,b).
In H2 SO4, the OL system shows the highest drop in activity, which is most likely due to dissolution of
Ni surface atoms. After 10 cycles into the potential regions of OER, the activities of the OL and NSA in
an acidic solution approach each other, or in other words, the OL catalyst approaches the more stable
NSA system. The relative position of Ni is critical for the stability of the catalyst material. The addition
of Ni in the form of an over layer significantly enhances the electrocatalytic activity in comparison
with an unmodified reference system. However, the stability of the Ni OL system is not given in acidic
media. By contrast, Ni in the subsurface layer of Irpoly seems to positively affect the stability of the
respective catalyst material since it shows the lowest drop in current density during the 10 cycles.
The OER activity of polycrystalline Ir has been enhanced in both an acidic and an alkaline solution
with the addition of Ni at the topmost layer (OL) using the controlled-potential technique. Based
on the significantly higher OER performance of the OL by contrast with the NSA and unmodified Ir
electrode, the Ni/Irpoly OL system was investigated in a second OER protocol. In Figure 6, the OER
polarization curves of two Ni OL catalyst systems in 0.1 M KOH are depicted. The electrochemical
performance of each sample was tested by 10 consecutive cycles. However, both samples differ in their
electrochemical pretreatment. The Ni-modified Irpoly electrode OLdi was tested in KOH directly after
the process of Ni deposition. The electrochemical activity of the catalyst OLpr was initially recorded in
0.05 M H2 SO4 before it was transferred into the alkaline setup.



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Figure 6. (a) OER polarization curves of Ni/Irpoly OL in N2 -saturated 0.1 M KOH without (lilac) and
with (green) electrochemical OER pre-treatment in 0.05 M H2 SO4 . Scan rate ν = 5 mV s−1 . 1st (solid
lines) and 10th cycle (dotted lines) included. The over layer (OL) was formed by electrodeposition of
Ni at 0.05 VRHE for 10 min. (b) Cathodic shift of the Ni(OH)2 /NiOOH redox peak of the OL implies
destabilization of Ni2+ .

For the OLdi catalyst, the characteristic Ni(OH)2 /NiOOH redox transition [31,32] at ~+1.35 VRHE
(see Figure 6b) is clearly visible and grows with an increasing cycle number. A slight cathodic shift
of the precatalytic Ni2+/3+ redox wave is observed with an increasing cycle number, which indicates
a destabilization of Ni2+ under OER catalytic potential. The redox charge obtained by integration
of the area under the cathodic wave from the last scan amounts to 137.15 μC (with consideration
of geometric surface area), which is equivalent to 1.42 nmol redox active Ni atoms. Taking into
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Surfaces 2018, 1

consideration that the Ni amount determined by ICP-MS for deposition at 0.05 VRHE was 3.178 nmol
(see Table S1), this result would imply that not all Ni sites in the Ni/Irpoly OL system contribute equally
to the OER current. The pre-treated OLpr catalyst shows no sign of a Ni(OH)2 /NiOOH redox peak,
which clearly confirms the instability of a Ni OL in acidic media. The characteristic redox peak of
the OLdi , however, demonstrates the successful implementation of the electrochemical deposition of
Ni on the polycrystalline Ir electrode and forms the basis of further atomic structure modifications.
The prior OER treatment of the OL catalyst in sulfuric acid seems to activate the catalyst by ~31%.
The enhanced electrochemical property of the pre-treated catalyst might be attributed to an increased
presence of electrophilic oxygen species [33–35]. Pfeifer et al. identified the reactive OI− species by
characteristic near-edge X-ray absorption fine structure (NEXAFS) features in Ir(III/IV) oxohydroxides.
The dissolution of Ni adatoms along with the introduction of Ir vacancies in highly acidic solution and
the accompanying formation of electronic defects in the cationic and anionic framework might lead to
the formation of highly electrophilic oxygen atoms. Since these oxygen atoms are highly susceptible
to a nucleophilic attack, they are suggested to facilitate the O-O bond formation, which is a key step
during electrocatalytic OER. Another conclusion that can be drawn from the higher activity of the
OLpr catalyst is roughening of the surface by repetitive cycling.
The selective leaching of a sacrificial component is a promising approach to design alloy
systems with enhanced activity and stability towards the OER. In the work of Seitz et al., a highly
active IrOx /SrIrO3 electrocatalyst was formed by leaching strontium in the near-surface during
electrochemical testing from the surface of SrIrO3 thin-films [36]. In comparison to known IrOx
catalyst materials, these IrOx surface layers induced by Sr deficiencies showed significantly higher
OER activities, which supports the theory of increasing active sites by selective leaching of sacrificial
components from surface alloys.
The observed order of activity of the Irpoly catalysts measured in KOH in the RDE setup could be
confirmed with an electrochemical flow cell setup, which directly interfaced to a mass spectrometer
(MS). In Figure S5a, the OL system displays the highest measured current density at 1.56 VRHE followed
by Irpoly and the NSA system, which are quite close to one another. The subsequent OER measurements
in acidic solution slightly differ in the sequence of activity. These were measured directly after the
catalysts were treated in alkaline solution. The unmodified polycrystal experiences a drastic loss in
current density. The NSA catalyst system, however, offers a greater stability and activity compared
to Irpoly . The enhanced activity behavior of the NSA system was also observed in the 10th cycle
directly measured in acidic solution (see Figure S4a). In Figure 5e, almost no activity loss is observed
in acidic solution. Within the context of the measuring accuracy, the faraday efficiency (FE) of O2
evolution in Figure S5b reaches ~100% for Irpoly and for the OL catalyst in both electrolytes and in
KOH, respectively. The FE over 100% for the OL catalyst in acidic solution is significant and may
be deduced to Ni dissolution in the highly corrosive environment, which simultaneously produces
additional oxygen. For both Ni-modified NSA catalysts, the FE amounts to ~104% within the precision
of measurement. It should be noted that the faradaic efficiencies estimated using the calibration factor
K* are assuming 100% efficiency in the steady state protocol and, therefore, the OER efficiencies reflect
processes during non-equilibrium conditions.
3.3.2. CO Stripping Voltammetry
The surface-sensitive electrochemical oxidation of adsorbed CO on pristine Irpoly and Ni-modified
Irpoly electrodes has been studied in alkaline solution (0.1 M KOH) using stripping voltammetry.
In this scenario, the Ni/Irpoly NSA catalyst was prepared by annealing the OL for only 1 min in the
reductive atmosphere. Due to the reduced exposure to the reductive treatment, the characteristic
Ni(OH)2 /NiOOH feature is still visible in the cyclic voltammogram. The CO oxidation mechanism
is examined in the alkaline electrolyte in order to integrate the Ni electrochemistry and to identify if
CO adsorption is favored in the presence of Ni since the (sub)surface layers of Ni on Ir have not been
stable in an acidic environment. Furthermore, the concept of OER in alkaline media may be of interest

304

Surfaces 2018, 1

from the catalytic point-of-view since the pretreatment of the Ni/Irpoly OL catalyst in H2 SO4 seems to
have a beneficial effect in the subsequent activity measurement in KOH.
Figure 7 compares the anodic stripping voltammetry observed for a CO adlayer on the three
distinct electrode surfaces in 0.1 M KOH solution at 50 mV s−1 , which is followed by the second
sweep cycle between the potentials of 0.05 and 1.40 VRHE . The presence of a CO surface adlayer can be
confirmed since, for each electrocatalyst, a CO stripping peak was recorded. The voltammetric profiles
further validate the complete electrooxidation of the adsorbed CO in the first anodic sweep since the
second cycle produces a voltammogram displaying HUPD peaks characteristic for Ir.

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Ni/Irpoly OL, and Ni/Irpoly NSA measured in 0.1 M KOH. (a) Anodic curves over the entire potential
range and (b) anodic and cathodic curves in a potential range of 1.25–1.44 VRHE .

It can be noticed that there is a clear distinction between the Ni-modified electrocatalysts and the
unmodified, pristine polycrystalline Ir electrocatalyst with regard to the stripping peak shape and
position in Figure 7a. The Ni/Irpoly OL and Ni/Irpoly NSA electrodes produced a stripping peak at
much lower potentials than the pure Irpoly , which suggests that the oxidation of surface CO to CO2 is
promoted by the presence of Ni. In accordance with the enhanced CO oxidation of the Ni-modified
catalysts, the onset of OH adsorption on the OL and NSA starts at more negative potentials than on
pristine Irpoly [37]. The oxidation of Ir sites is favored, or in other words, the adsorption of HO at
free metallic sites is highly favored in the presence of Ni. An additional voltammetric feature of the
CO-loaded Ir polycrystal is a broad peak at ~+0.3 VRHE (Figure 7a), which is only weakly defined for
the OL and no longer present for the NSA. The possible origin of this peak will be discussed below.
To gain a quantitative understanding of the effect of Ni modification on the CO oxidation activity
in alkaline electrolyte, the CO stripping profiles were deconvoluted into multiple peaks. In order to
place hidden peaks, a linear background-correction was performed to the CO stripping peaks and,
subsequently, the individual peaks were fitted. The respective CO stripping curves are presented in
Figure 8 for Irpoly , Ni/Irpoly OL, and Ni/Irpoly NSA catalysts. At least two peaks can sufficiently fit
all profiles.

305

Surfaces 2018, 1



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Figure 8. CO stripping curves for (a) the unmodified polycrystalline Ir electrode and the Ni-modified
(b) Irpoly OL and (c) Irpoly NSA catalysts obtained at 50 mV s−1 in N2 -saturated 0.1 M KOH (purified
Fe free) at a rotation of 1000 rpm. A linear-background correction was performed through baseline
subtraction for the anodic CO stripping voltammetry in Figure 7a between 0.4 and 1.2 VRHE . The peaks
were fit by Gaussian functions and the individual contributions are shown in colors. The cumulative
curve is given in turquoise.

The major component is centered at ~+0.85 VRHE for Irpoly and is shifted toward more cathodic
potentials for Ni/Irpoly OL and Ni/Irpoly NSA. The effect of a more facile CO oxidation with respect
to the unmodified Irpoly is quite likely due to a Ni richer surface, which has been observed in earlier
studies [38–40]. The hypothesis is further reinforced by the fact that no Ni was removed by the
CO stripping procedure since the characteristic Ni(OH)2 /NiOOH redox couple was observed for
both Ni-modified catalysts in Figure 8b after CO was stripped off the surface. The early onset
of CO oxidation for the OL and NSA is initiated by highly hydrophilic Ni-rich areas forming
Ni(oxo)hydroxides, which further react with adsorbed CO on neighboring Ir sites as early as
~+0.72 VRHE for Ni/Irpoly OL. The most likely reaction scheme to the CO oxidation process in alkaline
solution can be assigned to a Langmuir-Hinshelwood mechanism involving the competitive adsorption
of the reactants CO and OH [41].
COads + OHads → COOHads

(13)

COOHads → CO2 + H+ + e− .

(14)

The oxidative removal of adsorbed CO on Ir in the presence of Ni is presumed to proceed
according to the following equations.
Niads + H2 O → Niads · · · OHads + H+ + e−

(15)

Niads · · · OHads + COads → Niads + CO2 + H+ + e−

(16)

In the CO stripping voltammetry of the pure Ir polycrystal in Figure 7a, an anodic redox wave is
visible at ~+0.3 VRHE . The origin of this peak cannot be adequately verified. However, a rather small
‘pre-peak’ observed by Spendelow et al. [41] and García et al. [42] was attributed to oxidation of CO
at step sites. In the study of the mechanism and kinetics of CO oxidation on stepped surfaces by in
situ infrared spectroscopy, Lebedeva et al. found that the step trough is the most active site for CO

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oxidation [43]. After deposition of Ni, this peak is weakly identifiable, which suggests that Ni adatoms
adsorb to some extent on Ir step sites. After annealing, however, the peak disappeared probably due
to healing of a majority of defect sites on the polycrystalline Ir surface.
The CO stripping peaks in Figure 8 all consist of a small peak or shoulder at higher potentials
(dark green fit) arising very likely from the adsorption of HO. Ir oxidation is prevented until the
oxidative removal of adsorbed CO. On the Ni-modified catalysts, OH adsorption is more favored and,
thereby, shifted to lower potentials as discussed earlier. The main peak positioned at ~+0.85 VRHE
for Irpoly , ~+0.72 VRHE for Ni/Irpoly OL, and ~+0.74 VRHE for Ni/Irpoly NSA can be attributed to CO
oxidation at terrace sites of Ir [44]. Based on the fitted data, the cathodic shift of the CO stripping peak
is slightly more pronounced for Ni/Irpoly OL (~+0.72 VRHE ) than for Ni/Irpoly NSA (~+0.74 VRHE ).
This finding can be interpreted under consideration of the CO oxidation charge of both Ni-modified
catalysts. The CO stripping charge is frequently used to determine the electrochemical active surface
area (ECSA), which is required to assess the stability and activity properties of electro-catalytic
materials [38,45–47]. The baseline-corrected deconvoluted peaks (light green) were integrated for each
individual catalyst. As expected, the net CO stripping charge of Irpoly was found to be the highest
with 803 μC cm−2 . The electrodeposition of Ni on the surface of Ir reduces the surface area available
to CO, as Ir sites are blocked by Ni adatoms. A net charge of 569 μC cm−2 can be assigned to the OL
catalyst. The CO charge for the NSA is slightly reduced to 529 μC cm−2 . However, in addition to purely
geometric blocking, electronic effects due to the Ni atoms need to be taken into consideration as well
when discussing the CO coverages. Since the outer layer of the NSA in this study consists of both Ir and
Ni atoms, the lower CO coverage on the NSA must be deduced to the modified electronic properties of
the Ir surface atoms due to Ni atoms in the subsurface layer. A lowering of the chemisorption energy
leads to a weaker adsorbate binding. A similar behavior has been reported by Bandarenka et al. in
the study of the CO oxidation on alloys of Cu/Pt (111) [48]. It is, therefore, feasible that the electronic
effect of Ni in the subsurface region on Ir either has a negative effect for the oxidation of CO or does
not compensate for the reduced amount of surface Ni atoms. Consequently, the peak position of the
main CO stripping peak is not only a function of the pure composition but depends strongly on the
specific position of the non-noble metal. Apparent from Figure 8b, the CO stripping curve of the OL
seems to consist of a second feature at ~+0.54 VRHE , which might be deduced to an OH adsorption
on Ni adatoms prior to the facilitated electrochemical oxidation of CO. Even after intense fitting, this
rather small peak was not found on the NSA most likely due to the comparatively small amount of Ni
surface atoms. However, the absence of this peak on the NSA system strengthens the observed higher
activity of the OL towards CO oxidation.
In the foregoing discussion, a distinction between distinct active Ir sites has not been addressed
yet. The behavior of hydrogen on the polycrystalline Ir surface, however, may provide important
information on the deposition process of Ni. In Figure 7a, there are three distinct redox features in the
anodic branch of the HUPD region of the pristine Irpoly in the second scan after the complete oxidation
of adsorbed CO. These redox peaks indicate distinct Ir sites at which oxidation of UPD hydrogen
occurs. The well-defined HUPD peaks can be assigned to the three low-index crystallographic facets
(111), (110), and (100) of Ir. A thorough theoretical study of the adsorption/deposition of hydrogen
on various close-packed facets of noble and non-noble metals including the face-centered cubic (fcc)
facets of Ir(111), Ir(110), and Ir(100) was performed. The energetically most favored adsorption sites for
Ir(111) and Ir(100) were adopted by Ferrin et al. [49]. The binding energies (BE) for the three distinct
low-index surfaces along with the coordination numbers (CN) of surface Ir atoms and calculated
geometries for surface hydrogen (Site) are given in Table 2.

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Surfaces 2018, 1

Table 2. DFT-calculated binding energies of hydrogen and H-Ir bond distances on the three low-index
surfaces of Ir.
Face

CN

Site

H BE/eV

H-Ir Bond Distance/pm

(100)
(111)
(110)

6
9
7

bridging
atop
bridging

−2.91
−2.73
−2.74

2 × 181
159
2 × 181

The binding energies were calculated relative to the hydrogen atom in the gas phase.
The computational details reveal that the Ir(100) surface (−2.91 eV) binds hydrogen the strongest and
is followed by Ir(111) (−2.73 eV) and Ir(110) (−2.74 eV). Based on these calculations, Ir(100) binds
hydrogen ~0.20 eV more strongly than the Ir(111) and (110) surfaces. The difference in binding energy
between those facets can be deduced to the hydrogen-metal bond distances and binding geometries.
It has been reported that, for Ir(111), adsorption on the top site [50,51] is favored while, on Ir(100),
the bridge site [52] is more favorable in terms of energy. Due to the distinct geometries of the adsorption
configuration for surface hydrogen, the H-Ir bond distance on (100) is larger (1.81 Å) than on (111)
(1.59 Å). The very similar hydrogen binding energies of the (111) and (110) sites hamper the assignment
of the HUPD peaks. For this, an additional DFT study on Ni binding energies was performed revealing
that Ni binds the strongest on the Ir(110) face (−4.55 eV), less strongly on the Ir(100) face (−4.47 eV),
and the weakest on Ir(111) (−3.76 eV). Given the stronger binding of Ni to the (110) sites, we assume
the first peak to originate from HUPD on (110) sites. The combined results from the DFT studies are
presented in Figure 9.

Figure 9. The second cycle of the CO stripping CV (from Figure 7) displaying the HUPD peaks of the
three distinct catalyst systems. The allocation of the HUPD peaks are based on DFT calculations of the
hydrogen binding energy on the three low-index surfaces (111), (110), and (100) of Ir.

As the HUPD peak attributed to the (110) sites is significantly reduced by forming an over layer
of Ni, it appears that Ni adsorbs rather on (110) sites and that (100) sites are less favored. For the NSA
catalyst system, the (110) peak is slightly visible whereas the (100) is more pronounced due to Ni in the
subsurface layers. The HUPD charge is the lowest for the NSA catalyst (0.10 mC cm−2 ). For pristine
Irpoly and Ni/Irpoly OL, the HUPD charge was determined to be 0.20 mC cm−2 and 0.13 mC cm−2 ,
respectively. The given order was also observed for the CO charge and might provide important
indications for the activity of the Ni-modified catalysts. The smaller HUPD charge of the NSA catalyst

308

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in comparison to the system in which Ni is present at the surface might be set into relation with its
lower activity performance toward catalyzing the oxygen evolution.
4. Conclusions
The analysis of the three distinct catalyst systems with respect to their electrocatalytic performance
towards oxygen evolution revealed that the Ni-modified polycrystalline Ir electrodes display distinct
electrochemical properties. Incorporating Ni in the form of an over layer significantly reduces the
overpotential to evolve oxygen in both the acidic and the alkaline solution. The promoting effect of Ni
in IrNi bimetallic catalysts has been observed in various studies [23,29,35]. In acidic milieu, the higher
activity might be assigned to the sacrificial leaching of Ni leading to a higher content of hydroxyl groups
(OH). The Ni-depleted active catalyst structure is expected to be composed of oxygen atoms protonated
from the electrolyte to form surface hydroxyl groups, which presumably act as crucial intermediates in
the OER mechanism. The beneficial effect adding Ni to the OER activity, however, is no longer given to
the same extent if Ni is present in the subsurface layer of the electrode. Thus, we can conclude that Ni
in the NSA configuration is not able to perform as a sacrificial catalyst component as effectively as Ni
in the OL configuration. Furthermore, the electronic effect of Ni on Ir might exercise a negative effect
on the binding energy of Ir towards reaction intermediates in the OER. However, it cannot be neglected
that the reductive treatment, which formed the NSA, led to partial healing of the Ir surface with a
smaller density of defect and step sites. Since the environment of Ir was significantly altered, so was
its activity. In alkaline media, the highly hydrophilic Ni adatoms are sufficient to achieve the desired
effect of increasing the OER performance of the catalyst material. The performance of the OL system
increases in KOH after it was pretreated in acid electrolyte beforehand. This observation confirms the
hypothesis of Ni dissolution in acid. It appears as if the preceding OER treatment in 0.05 M H2 SO4
has an advantageous impact on the electrocatalytic properties in KOH and can be deduced to the
previously mentioned leaching of Ni in an acidic solution, which raises the concentration of surface
OH atoms. The performance during OER of the catalysts including the shift of CO oxidation peaks is
not only a function of the pure composition of the material but does strongly depend on the relative
position of the metals to each other. However, it must be kept in mind that the interpretation of the
peaks with regard to the CO stripping measurements especially of the shoulders is quite complex due
to the presence of surface defects.
Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/1/1/13/s1.
Figure S1. Double-layer capacitance measurements displaying cyclic voltammograms of Irpoly measured in a
potential window in which no faradic currents occur. Following scan rates were adjusted: 500 mV s−1 , 200 mV s−1 ,
100 mV s−1 , 50 mV s−1 , 20 mV s−1 and 10 mV s−1 ; Figure S2. X-ray photoelectron spectroscopy measurements
of (a) Ir(111) over the entire energy range. (b) High-resolution spectra of Ir(111) displaying the Ir4f, C1s, O1s
and Ni 2p spectra. (c) XP spectra of Ni/Irpoly OLt=20min over the entire energy range; Figure S3. (a) Atomic force
microscopy (AFM) images of Ni/Ir(111) OL with respective height profiles. (b) Magnification of a terrace assembly
of (a) and (c) higher magnification. Scale bar is 400 nm. Arrows give the direction of the calculated height profile;
Figure S4. 10th cycle of the electrocatalytic oxygen evolution reaction (OER) activities of Ni-modified Ir electrodes
compared to pristine polycrystalline Ir at scan rates of 5 mV s−1 in (a) 0.05 M H2 SO4 and (b) in 0.1 M KOH of Irpoly ,
Ni/Irpoly OL and Ni/Irpoly NSA. Current densities of the electrodes are normalized to the electrodes’ geometric
surface area of 0.785 cm−2 ; Figure S5. Electrocatalytic OER activities of Ni-modified electrocatalysts and pristine
Irpoly measured by differential electrochemical mass spectrometry (DEMS) in a dual-thin layer flow cell setup at a
scan rate of 10 mV s−1 . (a) Faradaic current densities measured at an OER overpotential of 300 mV (1.53 VRHE )
normalized to the electrodes’ geometric surface area of 0.785 cm−2 . The electrocatalytic OER activity in H2 SO4
was measured directly after the treatment in KOH. (b) Faraday efficiency of oxygen evolution (FEO2 ) determined
for the catalysts; Table S1. Calculated amount of adsorbed Ni on a polycrystalline Ir electrode in dependence of
deposition potential Edep. The coverage of Ni (ML) was determined from ICP MS based on the assumption of a
predominant (111) face on Irpoly.
Author Contributions: Conceptualization, E.Ö and P.S. Methodology, E.Ö. Validation, E.Ö. Formal analysis, E.Ö.,
J.A., A.M.M., I.S., and T.K. Investigation, E.Ö., J.A., and I.S. Resources, P.S., K.M., M.M., and B.R.C. Data analysis,
E.Ö., J.A., I.S., and T.K. Writing—original draft preparation, E.Ö. Writing—review and editing, E.Ö., P.S., B.R.C.,
K.J.J.M., T.K., J.A., and A.M.M. Visualization, E.Ö. and P.S. Supervision, P.S. Project administration, E.Ö.

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Surfaces 2018, 1

Funding: This research was funded by the Deutsche Forschungsgemeinschaft under the priority program SPP1613,
grant number STR 596/3-1. The U.S. Department of Energy (DOE) through the Office of Basic Energy Sciences
(BES) by grant DE-FG02-05ER15731 supported work at UW–Madison. Computational work was performed using
supercomputing resources at the Center for Nanoscale Materials (CNM) at the Argonne National Laboratory
under contract number DE-AC02-06CH11357 and at the National Energy Research Scientific Computing Center
(NERSC) under contract number DE-AC02-05CH11231. Thomas Kropp is grateful for partial financial support by
the Alexander von Humboldt Foundation.
Acknowledgments: We thank Benjamin Paul and Zarina Pawolek for helpful discussions and help with the
set up.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Metallic Iridium Thin-Films as Model Catalysts
for the Electrochemical Oxygen Evolution Reaction
(OER)—Morphology and Activity
Ebru Özer 1,† , Zarina Pawolek 1,† , Stefanie Kühl 1 , Hong Nhan Nong 2 , Benjamin Paul 1 ,
Sören Selve 3 , Camillo Spöri 1 , Cornelius Bernitzky 1 and Peter Strasser 1,4, *
1

2
3
4

*
†

The Electrochemical Energy, Catalysis and Materials Science Laboratory, Department of Chemistry,
Technical University Berlin, Straße des 17. Juni 124, D-10623 Berlin, Germany;
ebru.oezer@campus.tu-berlin.de (E.Ö.); zarina@mailbox.tu-berlin.de (Z.P.);
stefanie.kuehl@tu-berlin.de (S.K.); benjamin.paul@tu-berlin.de (B.P.); camillo.spoeri@tu-berlin.de (C.S.);
cornelius.bernitzky@campus.tu-berlin.de (C.B.)
Max Planck Institute for Chemical Energy Conversion, Department of Heterogeneous Reactions, Stiftstr. 34,
D-45470 Mülheim, Germany; hongnhan.nong@gmail.com
Zentraleinrichtung Elektronenmikroskopie, Technische Universität Berlin, D-10623 Berlin, Germany;
soeren.selve@tu-berlin.de
Ertl Center for Electrochemistry and Catalysis, Gwangju Institute of Science and Technology,
Gwangju 500-712, Korea
Correspondence: p.strasser@tu-berlin.de; Tel.: +49-(0)30-314-29542
These author contributed equally to this work.

Received: 3 November 2018; Accepted: 5 December 2018; Published: 6 December 2018

Abstract: Iridium (Ir) oxide is known to be one of the best electrocatalysts for the oxygen evolution
reaction (OER) in acidic media. Ir oxide-based materials are thus of great scientific interest in current
research on electrochemical energy conversion. In the present study, we applied Ir metal films as
model systems for electrochemical water splitting, obtained by inductive heating in a custom-made
setup using two different synthesis approaches. X-ray photoelectron spectroscopy (XPS) and selected
area electron diffraction (SAED) confirmed that all films were consistently metallic. The effects of
reductive heating time of calcined and uncalcined Ir acetate films on OER activity were investigated
using a rotating disk electrode (RDE) setup. The morphology of all films was determined by
scanning electron microscopy (SEM). The films directly reduced from the acetate precursor exhibited
a strong variability of their morphology and electrochemical properties depending on heating time.
The additional oxidation step prior to reductive heating accelerates the final structure formation.
Keywords: oxygen evolution reaction; water splitting; iridium; thin-films; spin-coating; model
systems; electrocatalysts

1. Introduction
Water electrolysis as a non-polluting way of producing hydrogen is considered as one of the key
technologies for an affordable and sustainable energy cycle [1]. Proton exchange membrane (PEM)
electrolyzers are the most promising candidates for the supply of high purity hydrogen, because
of their operation at high current densities and low gas crossover rates [2]. Among electrocatalysts
applied in acidic media of water electrolysis, Ir oxide represents the state-of-the-art oxygen evolution
catalyst with excellent activity and stability [3–5]. As a noble metal, it is less sensitive to corrosion in
acidic media and further combines a high activity and stability.
The conditions under which the oxide electrodes are formed have a significant influence on the
oxygen evolution reaction (OER) performance. In relation to thermally prepared oxides, this comprises
Surfaces 2018, 1, 151–163; doi:10.3390/surfaces1010012

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the applied temperature [4], gas atmosphere [6], and particularly the precursor [6], which in turn
has a significant effect on the morphology. These oxides were reported to have a higher stability, but
lower activity in contrast to electrochemically oxidized Ir [7,8]. Accordingly, further investigation
of electrochemically formed oxides on metallic Ir surfaces is of great scientific interest. In order to
understand the role of morphology and structure of these materials, an adequate model system
of well-defined Ir metal thin films is required. The fundamental research with model systems
can elucidate the structure–activity–stability challenges of OER materials and, therefore, enable the
systematic development and design of enhanced catalysts for industrial applications [9,10]. Given the
high sensitivity of the OER kinetics towards the catalyst’s preparation, the variability in either film
thickness or film composition in a model system should be kept to a minimum, which facilitates future
performance comparisons and performance enhancements.
Here, we present a study of novel Ir metal films on Ti substrate obtained by a precursor-based
synthesis route with subsequent reductive heating in a custom-made inductive chamber. The highly
accurate temperature and time regulation in the inductive setup guarantees the formation of
homogeneous metallic Ir thin films. Design details and capabilities of the setup can be found
elsewhere [11]. Herein, the metallic Ir film was obtained by reduction of two distinct Ir species,
namely Ir acetate and Ir oxide. The latter was formed from the acetate by an additional calcination
step prior to reduction in order to study the templating effect of Ir oxide on the resulting catalyst film.
Three different heating times were applied during synthesis to monitor possible sintering processes
during heat treatment. In the setup, the morphology of the metallic films can be varied under controlled
heating conditions. The catalytically active Ir oxide species was obtained by subsequent electrochemical
oxidation. It is well accepted that the synthetic oxidative route from metallic Ir to Ir oxide, that is,
either electrochemical or thermal oxidation, has a great influence on the OER performance. To our
knowledge, the influence of the synthesis route of the initial metallic Ir, however, has never been
taken into account. The present work addresses this issue. It presents a comparative investigation
of electrochemically oxidized Ir oxide species, which originate from distinct Ir metal thin films with
distinct morphology. In this way, we were able to influence the stability and activity of the catalytic
systems to a certain degree by the different arrangement and environment of the active sites.
A wide range of physical and electrochemical characterization, including cyclic voltammetry,
stepped potential voltammetry (SPV), X-ray photoelectron spectroscopy, transmission electron
microscopy (TEM)/selected area electron diffraction (SAED), and scanning electron microscopy
(SEM)/energy-dispersive X-ray spectroscopy measurements (EDX), was performed on the as-prepared
and the electrochemically oxidized films to uncover links between structural and catalytic properties.
2. Materials and Methods
As-prepared (AP) iridium acetate films (Chempur, Karlsruhe, Germany) on titanium were
prepared via spin coating as described by the authors of [10]. From here, two different synthesis
routes were applied. In the direct reduction route, the films were directly reduced in a custom-made
inductive cell directly from the acetate in 4% H2 /Ar (purity 5N, 99.999%, Linde, Munich, Germany)
at 400 ◦ C, while in the oxide route, the acetate films were annealed at 450 ◦ C for 15 min prior to the
reduction in the inductive cell. Samples will be referred to as Act (samples reduced directly from the
acetate) and Oxt (samples reduced from the oxide), where t is replaced by the time (in min) that the
sample was exposed to reduction in the inductive cell. For each synthesis route and each heating time,
several samples have been produced to ensure the reproducibility and to provide sufficient samples for
the various physical analysis methods, so not every characterization was made with the same sample.
Electrochemical activity and stability measurements were performed at room temperature in
0.05 M H2 SO4 (diluted from 95% H2 SO4 , Suprapure, Carl Roth, Germany) using a potentiostat (SP-240,
BioLogic, Seyssinet-Pariset, France) and a common three-electrode rotating disk electrode (RDE)
setup. A platinum mesh was used as counter electrode and a mercury/mercury sulfate electrode
(Ametek, Berwyn, PA, USA, sat., calibrated against reversible hydrogen electrode, RHE) was used

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as reference electrode. All potentials were converted into RHE scale (referred to as VRHE ). Nitrogen
was used for degassing. The electrodes were immersed in the electrolyte at a potential of 0.4 VRHE .
First, for surface characterization, cyclic voltammetry was performed in different ranges below the OER
onset. The lower turning potential was 0.05 VRHE , the upper turning potential was varied between
0.8 and 1.2 VRHE in 0.2 V steps (five cycles each), followed by 50 cycles between 0.05 and 1.4 VRHE for
electrochemical oxidation (samples referred to as _CV). After every 10th cycle of the latter, three cycles
in a range of 0.4 to 1.4 VRHE were performed to obtain the anodic charge. A scan rate of 50 mV s−1 was
applied. The anodic charge of the last scan was used for charge normalization of the current density
curves obtained from the following activity measurements. A rotation speed of 1600 rpm was applied
to the RDE working electrode and chronoamperometries (CA) were performed for 15 min at different
potentials between 1.534 and 1.784 VRHE , resulting in a stepped potential voltammetry (SPV) curve.
In this potential range, adequate currents were obtained in pre-measurements. At each potential,
electrochemical impedance spectroscopy was measured to obtain the ohmic resistance. After the last
potential, the electrolyte was degassed; the rotation of the RDE was turned off; and again, three cycles
at 50 m V s−1 in the potential range of 0.4 to 1.4 VRHE were conducted. The protocol was performed
twice, and the respective samples are referred to as _OER.
For electrochemical stability testing, the potential steps in the OER protocol were replaced by one
long chronoamperometry at circa 1.734 VRHE (without correction of ohmic resistance) for 20 h.
Scanning electron microscopy (SEM) images were obtained using a JEOL 7401F instrument
(Tokyo, Japan) with an accelerating voltage of 10 kV. Corresponding EDX analysis was established by
a Bruker XFLASH 4010 (Bruker, Billerica, MA, USA) at an accelerating voltage of 6 keV. The EDX was
normalized by the sample current, which was measured in a Faraday cup if not stated otherwise.
XRD patterns were obtained by grazing-incident X-ray diffraction (GIXRD) measurement in
a Bruker D8 Advance diffractometer (Bruker AXS). The analysis was performed at a grazing incidence
of 1◦ with a Goebel mirror, a Cu Kα source, a 0.23◦ secondary Soller, and a scintillation counter as
detector. The identification of the individual phases in the measured diffractograms was enabled
with reference patterns. Rietveld refinement with the software TOPAS (Bruker AXS) allowed a clear
definition of the individual phases present in the sample. A quantitative analysis, however, is not
given with Rietveld refinement, as it is primarily aimed at examining powder samples. The structure
and composition of film samples thus do not allow an exact determination of concentrations as
the individual phases might occupy distant layers, which consequently affects the intensity of the
incident beam.
For transmission electron microscopy (TEM), Cu grids with a 200 mesh and a holey carbon layer by
Quantifoil (Großlöbichau, Germany) were coated with the dispersed catalyst. The measurements were
performed at a FEI Tecnai G2 20 S-TWIN transmission electron microscope (Hillsboro, OR, USA) with
a LaB6 cathode operated at an accelerating voltage of 200 kV. The SAED evaluation was performed by
PASAD [12,13] plugin in Digital Micrograph™, and the obtained reciprocal distances were transferred
to Cu Kα scale to obtain a calculated diffraction pattern.
For the investigation of a cross section of Ir film on Ti, a TEM lamella was prepared by focused
ion beam (FIB). A FEI Helios NanoLab 600 workstation equipped with (1) an Omniprobe lift-out
system, (2) a Pt & W gas injector system, and (3) a scanning electron microscope (SEM) detector.
High resolution TEM (HRTEM) was performed using an FEI Titan 80-300 TEM electron microscope
with a high brightness X-FEG electron source and a Cs corrector for the objective lens. The microscope
was operated at 200 kV.
Laboratory-based X-ray photoelectron spectroscopy (XPS) measurements were carried out at
room temperature in an ultrahigh vacuum (UHV) setup using a monochromatic Al Kα source
at 1486.6 eV (SPECS, Berlin, Germany). The XP spectra were analyzed using CasaXPS software.
Fitting of O 1s spectra was performed using three singlets corresponding to lattice O (at ~ 530.0 eV),
OH (at ~ 531.4 eV), adsorbed water, and organic O (at ~533.0 eV), respectively. The Ir 4f spectra were
fitted with one component (for Ac10_AP and Ox10_AP samples, corresponding to Ir0 species) or three

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components (for the electrochemically treated samples Ac10_CV, Ac10_OER, Ox10_CV, and Ox10_OER,
corresponding to Ir0 , IrIV , and IrIII species). The Ir 4f 7/2 positions of the Ir0 , IrIV , and IrIII species were
located at ~60.9 eV, ~61.7 eV, and ~62.3 eV, respectively. The peak models consist of either a doublet
(Ir0 component), a doublet and three satellite peaks (IrIV component), or a doublet and two satellite
peaks (IrIII component) [14]. The peak models (including asymmetry parameter) were derived from the
fits of the Ir 4f XPS of the polycrystalline metallic Ir standard (for the Ir0 component) or of a rutile-type
IrO2 standard sample (for the IrIV component) [15].
3. Results and Discussion
In the study of the morphology, structure, and catalytic activity of supported metallic Ir films,
two different synthesis routes were established. The synthesis of the ‘Ox’ films was based on the
reductive treatment of Ir oxide films, which were prepared according to a previously reported
precursor-based spin coating method [10]. The ‘Ac’ films, however, skipped the calcination step
and were directly reduced from the spin coated acetate film.
3.1. Morphology of Ir Metal Thin-Film Catalysts
3.1.1. Scanning Electron Microscopy (SEM)
In order to investigate the morphology of the Ox and Ac films, micrographs of the samples before
(_AP) and after OER (_OER) measurements were obtained using a scanning electron microscope
(SEM). The respective results are displayed in Figure 1. The Ox_AP films (Figure 1, second row)
display a homogeneously cracked surface with clear porosity. Their morphology (Ox10, Ox30,
and Ox60) does not change significantly with heating time, in contrast to the morphology of the
Ac_AP film series. The surface of the Ac films (Figure 1, third row) displays larger cracks compared
with that of the Ox films, and in addition, they exhibit a larger crack size distribution. Thus, a clear
tendency towards a more cracked surface with a longer annealing time was observed for the Ac_AP
films. The morphological restructuring (and probable sintering process) produces a highly cracked,
‘sponge-like’, and more fibrous morphology. It can be assumed that in contrast to the Ac samples,
the calcination step in the Ox route creates a defined and crystalline rutile structure [10] from the
beginning, which consequently leads to a faster formation of the final metallic Ir structure through
reductive heating. After OER measurements, the Ac_OER samples (Figure 1, fourth row) exhibit
a swollen and filled surface with rounded edges on the cracks. A similar effect is only weakly visible
for the Ox films (Figure 1, first row) as a result of their less pronounced cracks.
3.1.2. Energy-Dispersive X-ray Spectroscopy Measurements (EDX)
EDX measurements performed on the _AP samples during SEM investigation in Figure S1
(in the Supplementary Information) show the atomic percentage (at%) ratio of O to Ir (not normalized
by sample current) depending on the reductive heating time. The results reveal a significantly higher
amount of O in the Ox samples than in the Ac samples, regardless of the heating time. This indicates
the presence of an O species in the Ox films that cannot be removed by reductive heating under applied
conditions, or that occurs while handling the samples under air, respectively. This species might either
originate from a small fraction of (re)oxidized Ir in the film; physically adsorbed or trapped O within
the pores of the film; or, most likely, a Ti oxide layer formed during the calcination step [10] and not
being reduced at applied reduction conditions.
The position of O within the film, however, cannot be obtained by a single EDX measurement. For this
purpose, several EDX measurements have been conducted along the radius of an Ox10_AP sample and
IrOx film on Ti (before reductive heating). The results are given in Figure 2. Here, the intensity (standardized
to the sample current and time) of each element (left y-axis) along with the calculated O/Ir count ratio (right
y-axis) is plotted against the radius of the film.

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Figure 1. Scanning electron microscopy (SEM) images of Ox (top) and Ac (bottom) film series before
(_AP, second and third row) and after oxygen evolution reaction (OER) (_OER, first and fourth row).
The cracks in the Ac films enlarge with increasing heating time. The weakly cracked morphology of
the Ox films does not show any change depending on the reductive heating time.

Figure 2. Counts per charge of the three main elements detected in energy-dispersive X-ray
spectroscopy measurements (EDX) of (a) IrOx and (b) Ox10_AP. The different intensities are caused by
different film thicknesses, mainly due to border effects during spin coating.

Within a radius of 3.5 mm from the center, both samples show no variation in film thickness.
Closer to the edge, the film thickness is altered as a result of the surface tension/boundary effects of the
precursor solution during the spin coating process. The uneven film thickness at the edge of the sample
and the ability of Ir to partially shield radiation emitted from deeper layers may at least give a hint on
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the O distribution, considering two extreme cases. If O was (exclusively) homogeneously distributed
within the Ir(oxide) film (regardless of the chemical nature of O), the O/Ir ratio would be independent
of the film thickness. In other words, at spots with lower Ir counts (lower film thickness), the O counts
should also be lower. If oxygen only existed within a Ti oxide layer underneath the film, the O/Ir
ratio would increase with decreasing film thickness as a result of lower shielding by Ir. With respect
to the IrOx film on Ti (Figure 2a), both cases seem to apply here; a lower amount of Ir accompanied
by a lower amount of O (lower film thickness, 3.6 and 4.3 mm from center) points to an oxidized Ir
film. At the same time, a higher O/Ir ratio speaks for a Ti oxide or Ir/Ti mixed oxide layer underneath
the Ir oxide film. The homogenous part of the Ox10_AP film (Figure 2b) exhibits a drastically lower
O/Ir ratio compared with the oxide film (Figure 2a), so a considerable amount of O has been removed
from the sample by reductive heating. Even though inductive heating has successfully reduced the
film, the location of remaining O has not changed during reduction; areas with a lower Ir amount
in Figure 2b (lower film thickness, 4.4 mm) exhibit a larger O/Ir count ratio, indicating a significant
amount of O beneath the film incorporated within a TiOx or Ir–Ti mixed oxide layer. At areas with
a higher Ir amount (3.6 mm), the O total counts increases with Ir, pointing to a fraction of O existing
inside the film. However, the metallic state of the as-prepared film, which was confirmed by XPS
(as discussed later in the manuscript), contradicts that the O is chemically bound to Ir.
3.2. Structural Characterization of Ir Metal Thin-Film Catalysts
3.2.1. X-Ray Diffraction (XRD)
To examine if the crystallinity and atomic structure of the Ac and Ox films changes depending
on the reductive heating time, a grazing-incident X-ray diffraction (GIXRD) analysis was conducted
prior to OER measurements (Figure S2). For both preparation methods, two crystalline phases are
detected in the as-prepared state, where the dominating reflections (35◦ and 40◦ , purple sphere) can
be assigned to the hexagonal phase of the Ti substrate. The weak reflections (47◦ and 68◦ , red star),
however, are assigned to metallic Ir phases, evidencing crystallinity of the films. The low intensity of
the reflections is a result of a low amount of Ir, or rather, a low thickness of the film in contrast to the
bulk Ti cylinder.
The Rietveld refinement technique was used to analyze the XRD patterns of the as-prepared
samples (Figure S3). The analysis evidences the absence of a crystalline IrO2 phase, but suggests
the presence of a TiO2 rutile phase to a very small extent. Still, the existence of an amorphous or
very nanocrystalline Ir oxide phase cannot be ruled out. To achieve a reasonable Rietveld refinement,
an additional hexagonal phase had to be incorporated into the pattern fitting. The fittings further
incorporate a metallic Ti phase and a nanocrystalline phase of metallic Ir. This supplementary reflection
will be referred to as TiOx and can be assigned to a Ti phase with larger lattice spacing rather than
a rutile-type TiO2 . The presence of this phase might be deduced as a result of inclusion of oxygen into
the lattice of Ti prior to the annealing process.
3.2.2. Selected Area Electron Diffraction (SAED)
To investigate the structure of the films at the nanoscale, selected area electron diffraction (SAED)
was performed on the Ac_AP and Ox_AP catalyst films, as well as on the electrochemically treated
Ac10_OER and Ox10_OER samples. The SAED patterns are given in Figure S4, and the corresponding
calculated diffraction patterns (Cu Kα scale) are shown in Figure 3a–d and were made from dispersion
of scraped off films. The reciprocal diffraction patterns were converted back into a diffraction pattern
on the Cu scale.

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Figure 3. Cont.

Figure 3. (a–d) Diffraction patterns evaluated from selected area electron diffraction (SAED) patterns
(Figure S4) of Ox10 (a and c) and Ac10 films (b and d) before (_AP; a and b) and after electrochemical
OER protocol (_OER; c and d). The patterns only show metallic Ir. (e,f) TEM cross-section of Ac10_AP
film shows crystalline Ir particles on an amorphous interlayer.

For the calcined IrOx film (prior to any reductive treatment), the clear diffraction pattern reveals
a rutile-type phase corresponding to IrO2 (Figure S5). According to SAED, after 10 min heating in
the reducing atmosphere at 400 ◦ C, the rutile-type reflections are replaced by a metallic Ir phase
(Figure 3a), which is still present upon cycling into OER potential regions (Figure 3c). The same applies
for the Ac10 catalyst, where the evaluation of the internal order of the catalyst films confirms no
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IrOx phase at several positions of Ac10, as well as Ac10_OER after polarization into OER potential
regions. Accordingly, SAED evaluation conclusively confirms that all investigated films are consistently
metallic and that the oxidation observed by XPS analysis (see Section 3.3.3) is limited to near-surface
Ir. Furthermore, the morphology of the thin-films was selectively studied on a FIB-prepared TEM
cross-section of an Ac10 sample. The resulting images are displayed in Figure 3e and 3f. After 10 min
of heating in the reducing atmosphere, the catalyst film including the crystalline particles and
the amorphous layer is about 10 nm thick. The high-resolution TEM (HRTEM) images reveal the
presence of an amorphous interlayer between the crystalline Ti substrate and crystalline Ir particles.
The particles, which are believed to be caused by cracking of the film, cover the amorphous layer
unevenly. These particles have also been found in the TEM images of the scraped film (Figure S6).
3.3. Surface Electrochemical Characterization
3.3.1. OER Activity
The electrochemical activity curves of the stepped potential voltammetry are depicted in Figure 4a.
Ac and Ox samples show different electrochemical behaviors. A longer heating period of the
Ac samples is accompanied with a slight decrease in the electrocatalytic activity. The distinct activities
of the directly reduced films are likely the result of sintering or reorganization of the surface structure
occurring during heating. The Ox samples, however, do not show remarkable differences between
different heating times, which agrees with a stable morphology of the films observed by SEM (Figure 1,
top row). This indicates the formation of a thermic stable structure, which is possibly benefited by the
predefined structure of the former Ir oxide.
A comparison between the current density jgeo (Figure 4a) and the charge normalized current
density jspec (Figure 4b) also supports the prediction of a sintering process on the Ac films; the charge
normalized curves of the Ac films approach each other in comparison with the unnormalized current
density, indicating a decrease of active surface area by sintering. Accordingly, the decrease in
electrocatalytic activity with longer heating time mainly results from the smaller active surface area as
a result of sintering. On the other hand, the charge normalization barely changes the relative position
of the current density curves of the Ox films, which is further proof of the stable morphology of these
films. Compared with polycrystalline bulk iridium (measured on a polished Ir cylinder), all films
possess higher charge normalized current densities. The slightly non-linear/bend behavior of the
Tafel plot (Figure 4c) and the loss of current density during each potential step of the SPV (Figure S7)
is an interplay between the approach of the steady-state activity after a sudden potential step and the
dissolution of the electrochemically obtained iridium oxides, as was evidenced by Cherevko et al. [16].
3.3.2. OER Stability
The stability performance was investigated by CA at approximately +1.734 VRHE in 0.05 M H2 SO4 ;
the results are shown in Figure S8. First, an abrupt drop of the current can be observed, assumed to be
an approach to steady-state after quick potential change. The current then decreases slowly and almost
linearly, presumably associated with the dissolution of oxidic Ir. After about 3 h, the current becomes
almost constant. There are two reasonable explanations for this behavior. Either an equilibrium of
formation and dissolution of the active iridium oxide species (oxidation and dissolution rates are the
same) is maintained, or a more stable active species is present on the catalyst surface, as was reported
by Li et al. [17].
To examine the stability of the catalysts, EDX was measured on all samples, which have undergone
electrochemical treatment including OER, stability testing, and Faraday efficiency. Here, a loss of
roughly 50 % Ir is observed, with an accuracy deviation of 10 % at most. This behavior agrees with the
already mentioned lack of stability of electrochemically oxidized iridium.

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Figure 4. Electrochemical OER activity curves of Ac and Ox films obtained by stepped potential
voltammetry. (a) Geometrical current densities in 0.05 M H2 SO4 . The shorter heated Ac films exhibit
a slightly higher activity compared with the Ox films and longer heated Ac films. (b) Specific current
density in 0.05 M H2 SO4 based on the anodic charge obtained from the last scan (between 0.4 and
1.4 VRHE ) before SPV. Because of charge normalization, the activities of the Ac films no longer differ
significantly from each other. The relation of the Ox films’ activities to each other is not affected. (c) Tafel
plot of the _OER films extrapolated from the stepped potential voltammetry (SPV) measurements.

3.3.3. X-Ray Photoelectron Spectroscopy (XPS)
The changes in the surface chemical state of Ir after activation and oxidation were traced with
X-ray photoelectron spectroscopy (Figure 5).
The Ir 4f XP spectra in Figure 5a of the as-prepared Ac10 and Ox10 samples confirmed the
presence of only metallic Ir in the XPS accessible area, indicating that the applied annealing step in
H2 is sufficient to reduce Ir on the surface to its metallic state. After the electrochemical treatments,
the Ir 4f core-level spectra shifted to higher binding energy as a result of the formation of surface
Ir oxide. The line shapes of the Ir 4f signals differ significantly depending on the potential range.
Fits of the Ir 4f spectra revealed three different Ir species presented in these samples: metallic Ir0 ,
IrIV , and IrIII species. The Ir species with a binding energy located at ~0.6 eV higher than that of the
IrIV species was assigned as IrIII in this work, although the true origin of this component (IrIII or IrV )
is still a controversy [14,18]. However, it is unlikely that IrV can exist in an Ir simple oxide under
ex situ conditions.

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Figure 5. X-ray photoelectron spectroscopy of the Ac10 and Ox10 films at as-prepared states (_AP),
after cyclic voltammetry between 0.05–1.4 VRHE (_CV) and OER activity testing (_OER) with fitted
components of (a) Ir 4f and (b) O 1s species. The as-prepared samples only show metallic Ir (Ir0 ). Through
potential cycling, the major part of metallic Ir was oxidized to IrIII and IrIV . After stepped potential
voltammetry with potentials up to 1.60 VRHE , almost no Ir0 is left and the ratio of IrIII to IrIV is increased.

After cycling in the potential range of 0.05–1.4 VRHE , both Ac10 and Ox10 films were oxidized to
a major extent; however, there is still discernible contribution of metallic Ir to the line shapes. On closer
inspection, the Ir4f XP spectra of both films differ in the IrIII /IrIV species ratio, highlighting different
oxidation behaviors. The Ac10_CV sample shows a higher IrIII /IrIV ratio than the Ox10_CV sample,
which might be the origin of its higher OER activity. In previous studies by Pfeifer et al. [14,19],
electronic defects in the form of O 2p holes (OI- species) were identified in amorphous IrOx .
These highly electrophilic OI- species are created by iridium vacancies in the IrOx framework and
possess a great activity in catalyzing the OER. The formation of these species is in turn accompanied
by the reduction of neighboring IrIV to IrIII species, which appear at an Ir 4f7/2 binding energy of
~0.6 eV higher than that of the IrIV species, according to their density functional theory calculations.
Drawing upon these findings, the enhanced activity of the Ac10 catalyst may be deduced to
coordination defects in the amorphous morphology facilitating the O-O formation in the OER [19].
Upon exposing the catalysts to higher OER potentials (Ac10_OER and Ox10_OER), the metallic
components diminished significantly to below 3 %. Two major oxidic components contribute to the Ir
4f spectrum. The IrIII component in the Ac10_OER sample is now much higher compared with that
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of Ox10_OER, indicating that the dominant species is IrIII . The higher specific OER activity based on
charge normalization gives reason to believe that the OI- species are present to a higher concentration
on the Ac films. The difference in oxidation behavior of the Ac10 and Ox10 samples is plausibly
explained by the difference in the morphology of the as-prepared metallic samples, as discussed
previously. The Ac10 film with a highly cracked, sponge-like morphology is prone to defect formation
during electrochemical treatment, in contrast to the Ox10 film with a smoother surface and more
defined structure, as the starting point of the Ox film series is a rutile phase with Ir in the oxidation
state +4.
The O 1s spectra in Figure 5b of the Ac catalysts prior to and after OER testing revealed
three different oxygen species, which can be assigned to adsorbed water (~533.1 eV), lattice oxygen
(~530.1 eV), and hydroxyl groups (~531.5 eV). All samples are distinguished by a high surface hydroxyl
group and very small lattice oxygen contents, typical for electrochemically grown iridium oxide [20].
The O 1s XP spectra indicate that the oxide layers in the electrochemically treated samples are largely
amorphous, which explains the absence of any crystalline IrOx phase in the SAED of these samples,
as discussed earlier.
4. Conclusions
In the study of the morphology, structure, and catalytic activity of supported metallic Ir films,
we presented the successful synthesis of metallic Ir thin-films on Ti substrate using two different
synthesis approaches. It was found that the catalyst structure was significantly influenced by the
synthesis routes. We emphasize that the choice of the synthesis route to the initial metallic Ir strongly
affects the performance of the subsequent electrochemically obtained Ir oxide. Further, these catalytic
films exhibited an enhanced stability and activity in the OER compared with the surface of a polished,
massive polycrystalline Ir cylinder.
Regardless of the heating time, the Ox route consistently resulted in the identical structure,
creating a system insensitive against fluctuations in synthesis parameters. The Ox route might thus
also be conducted in a common tube furnace, with a less sharp heating window (heating-up and
cool down). In contrast, SEM and TEM revealed that the morphology of the Ir films synthesized
over the Ac route exhibited a significant sensitivity towards the heating treatment. The application
of a sharp inductive heating accompanied by a defined heating time thus enables one to control the
progress of morphological restructuring, like sintering processes. These distinctive morphological
properties were found to have a strong effect on the catalytic properties. A slight decrease of the OER
activity with increasing heating period was observed. The difference between the catalytic properties
of both catalyst types is reduced, as the Ac films approach the Ox samples with a longer heating
time. Although this effect is not observable in the given scale of the morphology analysis, it cannot be
excluded that structural similarities occur on a smaller scale. As expected, a complete oxidation of
surface near Ir occurs during OER, while the majority of the film stays in its metallic state.
In the search for cost-effective materials catalyzing the OER, modifications of Ir with metallic
additives are under broad investigation. On the basis of our results, we propose the Ox route as
a promising synthesis approach for such bimetallic catalysts, as the final surface structure of the
resulting products is quickly formed and is tolerant to different heating periods.
Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/1/1/12/s1.
Author Contributions: Conceptualization, E.Ö. and Z.P.; methodology, E.Ö and Z.P.; validation, E.Ö, Z.P.,
and P.S.; formal analysis, E.Ö., Z.P., S.K., B.P., C.S., and S.S.; investigation, E.Ö., Z.P., S.K., H.N.N, B.P., S.S., and
C.S.; resources, P.S., S.S., and C.B.; data curation, E.Ö, B.P, and S.K.; writing—original draft preparation, E.Ö. and
Z.P.; writing—review and editing, E.Ö., Z.P., S.K., H.N.N, and C.S.; visualization, E.Ö. and Z.P.; supervision, P.S.;
project administration, E.Ö., Z.P., B.P., and P.S.
Funding: This research was funded by Deutsche Forschungsgemeinschaft, grant number SPP1613.
Acknowledgments: We thank the Fritz-Haber-Institut of the Max-Planck-Gesellschaft for access to XPS and
ZELMI for access to TEM/SAED and the focused ion beam preparation of a cross-section lamella.

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Surfaces 2018, 1

Conflicts of Interest: The authors declare no conflict of interest. The funders had no role in the design of the
study; in the collection, analyses, or interpretation of data; in the writing of the manuscript; or in the decision to
publish the results.

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Article

Diazonium Salts: Versatile Molecular Glues for
Sticking Conductive Polymers to Flexible Electrodes
Momath Lo 1,2,3 , Rémi Pires 3 , Karim Diaw 1 , Diariatou Gningue-Sall 1, *, Mehmet A. Oturan 2 ,
Jean-Jacques Aaron 2, * and Mohamed M. Chehimi 3, *
1
2
3

*

Faculté des Sciences, Université Cheikh Anta Diop, BP 5005 Dakar-Fann, Sénégal;
momath.lo@ucad.edu.sn (M.L.); karimdiaw@hotmail.com (K.D.)
Laboratoire Géomatériaux et Environnement, Université Paris-Est, 5 Bd. Descartes,
77454 Marne-la-Vallée CEDEX 2, France; mehmet.oturan@u-pem.fr
Université Paris Est, CNRS, ICMPE (UMR 7182), 2-8 rue Henri Dunant, 94320 Thiais, France;
pires@icmpe.cnrs.fr
Correspondence: diariatougningue.sall@ucad.edu.sn (D.G.-S.); jeanjacquesaaron@yahoo.fr (J.-J.A.);
chehimi@icmpe.cnrs.fr (M.M.C.); Tel.: +33-149-781-160 (M.M.C.)

Received: 20 June 2018; Accepted: 22 July 2018; Published: 8 August 2018

Abstract: Adhesion of polymers to surfaces is of the upmost importance in timely applications
such as protective coatings, biomaterials, sensors, new power sources and soft electronics. In this
context, this work examines the role of molecular interactions in the adhesion of polypyrrole
thin films to flexible Indium Tin Oxide (ITO) electrodes grafted with aryl layers from various
diazonium salts, namely 4-carboxybenzenediazonium (ITO-CO2 H), 4-sulfonicbenzenediazonium
(ITO-SO3 H), 4-N,N-dimethylbenzenediazonium (ITO-N(CH3 )2 ), 4-aminobenzenediazonium (ITO-NH2 ),
4-cyanobenzenediazonium (ITO-CN) and 4-N-phenylbenzenediazonium (ITO-NHPh). It was
demonstrated that PPy thin layers were adherent to all aryl-modified surfaces, whereas adhesive
failure was noted for bare ITO following simple solvent washing or sonication. Adhesion of
polypyrrole was investigated in terms of hydrophilic/hydrophobic character of the underlying aryl
layer as probed by contact angle measurements. It was found that sulfonic acid-doped polypyrrole
(PPy-BSA) thin films were preferably deposited on the most hydrophobic surfaces. More importantly,
the redox properties and electrochemical impedance of PPy were closely related to the hydrophobic
character of the aryl layers. This work demonstrates that diazonium compounds are unique
molecular glues for conductive polymers and permit to tune their interfacial properties. With robust,
diazonium-based architectured interfaces, one can design high performance materials for e.g., sensors,
printed soft electronics and flexible thermoelectrics.
Keywords: polypyrrole; diazonium salts; flexible ITO; adhesion; redox properties

1. Introduction
In the recent decades, conductive organic polymers have attracted a great deal of interest due to
their salient features, namely electric, thermoelectric, electronic, optical, ion exchange, and reinforcing
properties, to name but a few [1–3]. Particularly, polypyrrole (PPy) constitutes an excellent material
for building electronic devices [4], because of its ease of preparation by a variety of chemical [5],
electrochemical [6], biocatalytic [7], and radiation-induced [8,9] methods. In addition, it is possible to
tune its electrical conductivity upon doping, copolymerization, blending and nanostructuring [10,11].
Several studies have been focused on the effects solvent nature [12], electrolyte [13], potential
or current density [14,15], concentration of monomer [16], composition of working electrode
(metals, glassy carbon, ITO) [17,18], on the electrochemical deposition and mechanical properties of PPy
films. However, the performances as well as the long-term stability of polymer thin films is often limited
Surfaces 2018, 1, 43–58; doi:10.3390/surfaces1010005

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Surfaces 2018, 1

by various factors, including poor adhesion to the substrate and migration of metal contamination
from the substrate to the polymer [19]. To overcome these adhesion problems, coupling agents such as
organosilanes and aryl layers from diazonium salts proved to efficiently anchor PPy [20–23] and other
insulating polymers [24–26] to a large variety of materials surfaces. These coupling agents not only
impart strong adhesive bonding to the conductive polymer top layer, but also limit or suppress the
diffusion in the film of elements leached from the substrates [21].
While organosilanes are the main coupling agents used in adhesion science and technology [27–29],
studies of diazonium salts as adhesion promoters remain sparse. This is due to the fact that the
diazonium interface chemistry is still incipient. The interest of diazonium salts lies in their fast reaction
with practically all surfaces, triggered by chemicals, reactive surfaces or radiations [30]. In addition,
diazonium compounds are versatile as they impart functionalities to the materials surfaces depending
on the nature of the substituent borne by the benzene ring in either para-, meta- or ortho- position of the
diazonium group. In this way, one can design new surfaces possessing a plethora of properties such
as hydrophilic/hydrophobic character [31], electrocatalytic activity [32], initiation of surface radical
or oxidative polymerization [33–35], resistance to biofouling [36], sensing [37,38], reactivity towards
polymer matrices [39,40].
However, to the best of our knowledge, no systematic study comparing the effect of several
functional groups on the adhesion of diazonium-modified surfaces has been undertaken. Therefore,
the objective of this work is to examine the propensity of a series of diazonium compounds to provide
ultrathin adhesive layers to conductive polymer top coatings and to check to which extent the said aryl
adhesive layers govern the physicochemical properties of the polymer top coatings. To illustrate the
actual study, polypyrrole was selected as a reactive and functional polymer coating, and flexible ITO
as a substrate which is generally difficult to coat with conductive polymers for obtaining long-term
robust adhesion.
Herein, we investigated the electrografting of a series of diazonium compounds of the general
formulae N2 -C6 H4 -R, with R in para position =CO2 H, SO3 H, N(CH3 )2 , NH2 , CN and NH-Ph.
The hydrophobic/hydrophilic character of aryl-modified ITO electrodes was probed by means of
contact angle measurements of water drops. The modified flexible ITO sheets further served as
platforms for the electrodeposition of benzene sulfonic acid-doped polypyrrole (PPy-BSA), and the
electrochemical properties were interrogated by cyclic voltammetry and electrochemical impedance
spectroscopy. The electrochemical properties of PPy-BSA coatings will be discussed in terms of surface
composition and wettability of the underlying flexible aryl-modified ITO sheets.
2. Materials and Methods
2.1. Reagents
4-Aminobenzoic acid, 4-aminobenzene sulfonic acid, N,N-dimethyl-1,4-phenylenediamine,
1,4-phenylenediamine, 4-aminobenzonitrile, N-phenyl-4-phenylenediamine, were purchased from
Sigma Aldrich. Sodium nitrite (≥99.0%, Sigma-Aldrich, St. Louis, MO, USA), chloric acid (37%),
ferrocene (Sigma-Aldrich) and potassium chloride were used without further purification. Pyrrole
(purity ≥ 98%, Aldrich) and benzene sulfonic acid were refrigerated prior to synthesis. All aqueous
solutions were prepared using Milli-Q ultrapure water (MQ 18.2 MΩ·cm). ITO-coated polyethylene
naphtalate (PEN) sheets (Peccel, Yokohama, Japan) were cut to 5 × 10 mm with scissors and served as
working electrodes.
2.2. Diazonium Modification of Flexible ITO
The diazonium cations were generated in situ from the parent aromatic amines. The aromatic
amines and the corresponding modified electrodes ITO-R (R = functional group) were the following:
4-aminobenzoic acid (ITO-CO2 H), 4-aminobenzene sulfonic acid (ITO-SO3 H), N,N-dimethyl-1,4phenylenediamine (ITO-N(CH3 )2 ), 1,4-phenylenediamine(ITO-NH2 ), 4-aminobenzonitrile (ITO-CN),

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N-phenyl-p-phenylenediamine (ITO-NH-Ph). The ITO surface was cleaned by sonication in a mixture
of water/ethanol 80:20 (v/v) for 2 min. After sonication, the electrode was washed with DI water.
The electrografting of aryl layers from the diazonium salts generated in situ was performed using
a three-electrode system, including an ITO-coated polyethylene naphthalate (PEN) working electrode
(area ~ 5 × 5 mm), a stainless steel grid as counter electrode and a Ag/AgCl (saturated KCl)
reference electrode.
The diazonium salt was produced by mixing 0.1 mmole of aromatic amine (equivalent to
13.7, 17.3, 13.6, 10.8, 11.8, or 18.4 mg of 4-aminobenzoic acid, 4-aminobenzene sulfonic acid,
N,N-dimethyl-1,4-phenylenediamine, 1,4-phenylenediamine, 4-aminobenzonitrile, or N-phenyl-4phenylenediamine, respectively) with 90 mL of 0.5 M HCl in a glass bath for 30 min, before
adding 10 mL of 10 mM NaNO2 aqueous solution (this 10 mL solution contains 0.1 mmol NaNO2 ).
The resulting mixture was stirred for 1 h. The electrografting was then carried out in potentiostatic
conditions (−0.8 V/SCE for 45 s), using a Biologic SP 150 potentiostat (Seyssinet-Pariset, France),
computer-controlled with an EC-lab software. The blocking effect of the bare and modified ITO
electrodes was investigated in a 1 mM Fe(CN) 3−/4− redox solution by cyclic voltammetry (CV) in
the −0.4 and 0.6 V/SCE range at a scan rate of 50 mV·s−1 . The relative electroactivity of the modified
electrode, which is inversely related to its blocking effect properties, was quantified by CV in the
presence of the redox probe species. The Parameter Irel is defined as:
Irel =

Ipa with a f ilm
× 100
Ipa for a bare ITO

(1)

where Ipa is the intensity of the anodic peak current. Electrochemical impedance spectroscopy
measurements were performed at open circuit potential in 1 mM K3 Fe(CN)6 and 0.1 M KCl.
2.3. Preparation and Electrochemical Characterization of Polypyrrole Films on Diazonium-Modified ITO
Prior to the electropolymerization, diazonium-modified ITO electrodes were sonicated in
a water/ethanol 80:20 (v/v) mixture for 2 min. The PPy on the ITO modified aryl layer was prepared by
electropolymerization in a 0.01 M BSA + 0.1 M pyrrole aqueous solution (20 mL) by CV in a potential
range of −1.0 to 1.2 V. The fresh PPy films were characterized by CV under the experimental conditions
indicated in Section 3.
The impedance spectra were recorded in the 0.1 Hz to 100 kHz frequency range.
2.4. Surface Analysis
The X-ray photoelectron spectra (XPS) were recorded using a K Alpha apparatus (Thermo
Fisher Scientific, East Grinsted, UK) fitted with a monochromatic Al X-ray source (hν = 1486.6 eV;
spot size = 400 μm). A flood gun was used for static charge compensation.
An Easy Drop Krüss instrument (Hamburg, Germany) was used to determine the contact angles
of water drops deposited on diazonium-modified ITO and on PPy samples.
3. Results and Discussion
3.1. Brief Description of the Strategy and the Objectives of the Work
Six diazonium compounds of the general formulae N2 -C6 H4 -R, with R = CO2 H, SO3 H, N(CH3 )2 ,
NH2 , CN and NH-Ph, were selected for the preparation of aryl layers on flexible ITO surfaces.
Electrochemistry was selected as a means of attaching aryl monolayers, because it is a simple and
straightforward method which can be conducted in aqueous medium. Diazonium compounds were
generated in situ from the commercially-available parent aromatic amines; this adds up to the versatility
of chemical surface engineering with diazonium compounds. The choice of the functional group in para
position of the diazonium depends on two criteria: the electron-donor or electron-acceptor character of

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the functional group, on the one hand, and its propensity to govern the wetting of ITO surfaces, on the
other hand.
Figure 1 depicts the chemical structures of the in situ generated diazonium cations (Upper panel)
and the simple strategy adopted to attach PPy in two steps (Lower panel). Attachment of an aryl
adhesive layer was followed by electrodeposition of PPy. It is to note that no physicochemical study
will be provided on the adhesion of PPy on bare ITO, which is known to be very poor, as previously
reported [22]. Instead, we will concentrate the present investigation on the diazonium susbstituent
effect on the PPy adhesion and its redox behaviour.
Hereafter, the study is split into three main sections: (i) attachment and properties of aryl layers,
(ii) electropolymerization and behaviour of the PPy topcoat, and (iii) adhesion aspects of PPy top coat
to modified ITO sheets.

Figure 1. Upper panel: Chemical structures of the diazonium compounds under test. Lower panel:
electrodeposition of benzene sulfonic acid-doped polypyrrole on diazonium-modified flexible
ITO electrodes.

3.2. Electrochemical Modification of Flexible ITO Sheets with Diazonium Salts
Aryl layers were electrografted by chronoamperometry by setting the potential to −0.8 V/SCE
for 45 s. This electrografting period of time was sufficient for good polypyrrole adhesion without
inducing any major blocking effect that hampers the deposition of the polymer [22,41]. The blocking
effect of the aryl layers was assessed using Fe(CN)6 3−/4− , as depicted in Figure 2. The quasi-reversible
redox wave obtained with bare ITO (Figure 2A(a), peaks at 0.25 and 0.1 V) suffered a decrease in the
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Surfaces 2018, 1

intensity for ITO-N(CH3 )2 (Figure 2A(b), peaks at 0.38 and 0.1 V), but disappeared quasi-completely
in the case of ITO-SO3 H (Figure 2A(c)), indicating a blocking effect of the aryl layer as reported by
several authors [22,42–47]. The relative Fe(CN)6 3−/4− redox peak intensity (Irel ) decreased in the order
ITO > ITO-N(CH3 )2 > ITO-NH2 > ITO-NH-Ph > ITO-CN > ITO-CO2 H > ITO-SO3 H (see Table 1).
The influence of the aryl grafting to ITO electrodes is thus straightforward since the response of
Fe(CN)6 3−/4− is flattened for modified ITO plates compared to bare ITO strips. From measurements
similar to those of Figure 2A, Irel values were computed and reported in Table 1 for all aryl layers.
The ITO-N(CH3 )2 , ITO-NH-Ph and ITO-NH2 surfaces exhibit a blocking effect for Fe(CN)6 3−/4− redox
system reactions with Irel values of 49, 46.7, 23.6%, respectively. The highest blocking effects correspond
to ITO-CN, ITO-CO2 H and ITO-SO3 H, with Irel values of 15.7, 11.2 and 3.9%, respectively. Under the
same grafting conditions, the blocking effect appear to be more important for the ITO-SO3 H modified
electrode [48].

Figure 2. Electrochemical characterization of diazonium-modified ITO electrodes determined in
aqueous solutions of 1 mM K3 Fe(CN)6 and 0.1 M KCl. (A) Typical cyclic voltammograms (a) bare ITO,
(b) ITO-N(CH3 )2 , and (c) ITO-SO3 H. (B) Electrochemical impedance spectroscopy plots recorded with
(a) ITO-N(CH3 )2 , and (b) ITO-SO3 H.
Table 1. Irel , Rct and water drop contact angles (θ) obtained for aryl-modified ITO.
Surface

Irel (%)

Rct (kΩ/cm2 )

θ (◦ )

ITO-SO3 H
ITO-CO2 H
ITO-CN
ITO-NH-Ph
ITO-NH2
ITO-N(CH3 )2

3.90
11.2
15.7
23.6
46.6
48.6

68.8
47.3
44.0
9.32
8.97
7.15

48.7 ± 0.5
70.7 ± 0.4
81.2 ± 0.3
103.1 ± 0.3
116 ± 0.4
136.7 ± 0.1

The Nyquist plots for ITO-SO3 H and ITO-N(CH3 )2 electrodes (Figure 2B) present typical
semicircle shapes in the high-frequency domain, characteristic of an interfacial charge-transfer
mechanism. The surfaces exhibit straight lines in the low-frequency regime, characteristic of
a semi-infinite diffusion phenomenon. ITO-N(CH3 )2 displays the lowest charge transfer resistance
(Rct ), ca. 7.15 kΩ/cm2 compared to 68.8 kΩ/cm2 for ITO-SO3 H, a trend that is in agreement with
previously-reported studies on glassy carbon electrodes [48]. Table 1 reports Rct values for all
diazonium-modified ITO surfaces and clearly indicates the general following trend: low values
for electron-donor aryl groups (7.15–9.32 kΩ cm2 ) and ~5–10-fold higher Rct values (44.0–68.8 kΩ cm2 )
for electron-acceptor aryl groups.
Figure 3 shows images of water drops deposited on the various ITO-R surfaces. For R = N(CH3 )2 , NH2
and NH-Ph, the contact angles were found to be 136.7 ± 0.1, 116 ± 0.4 and 103.1 ± 0.3◦ , respectively.
Taking into account the accepted rule of thumb, the water contact angle (θ) values above 90◦ indicated
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Surfaces 2018, 1

hydrophobic surfaces. In contrast, more wetting was observed for the electron-acceptor groups
R = CN, COOH and SO3 H, the water contact angles of which are 81.2 ± 0.3, 70.7 ± 0.4 and 48.7 ± 0.5◦ ,
respectively. These surfaces can thus be ranked as relatively hydrophilic (θ < 90◦ ). As expected,
the layers with more ionizable groups COOH and SO3 H produced the lowest contact angles and thus
led to the highest hydrophilicity. It is thus possible to control the hydrophobic/hydrophilic character
of flexible ITO surfaces by simply changing the radical R in para position of the diazonium group.

Figure 3. Contact angles of water drops gently deposited on diazonium-modified ITO: (a) ITO-N(CH3 )2 ,
(b) ITO-NH2 , (c) ITO-NH-Ph, (d) ITO-CN, (e) ITO-CO2 H, (f) ITO-SO3 H.

Figure 4 displays plots of the electrochemical characteristics of the aryl layers versus their wetting
property, expressed by the water drop contact angle values. For the hydrophilic surfaces, the barrier
to the Fe(CN)6 3−/4− probe is most probably due to the favourable water-aryl interfacial interactions.
Under such wetting conditions, the probe molecules hardly reach the aryl-modified ITO surfaces
and undergo redox reactions. Higher Irel values indicate weaker barrier, and the probe interact more
favorably with the surfaces compared to water.

Figure 4. Correlations between the electrochemical and wetting properties of diazonium-modified
flexible ITO electrodes: (a) Irel -vs-water contact angle; and (b) Rct -vs-water contact angle.

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Surfaces 2018, 1

As far as Rct is concerned, this physical property decreases for hydrophobic aryl-modified ITO
surfaces (θ > 90◦ ). These results are due to the steric hindrance or the physical barrier role played by
the free pair of the organic molecules [48].
All in one, the results can be discussed in terms of hydrophobic/hydrophilic effects [49]. Indeed,
Fe(CN)6 3−/4− redox probe is relatively hydrophilic [49] and can rather interact with water instead of
the aryl-modified surface. In addition, the hydrophilic aryl groups are negatively charged which may
add up electrostatic repulsions at the ITO-aryl/probe interface.
3.3. Electrodeposition of Polypyrrole on ITO
3.3.1. Electrochemical Characterization of Poly(pyrrole-benzene sulfonic acid) Films on
Aryl-Modified ITO
The aryl layers served as a primer for the electrodeposition of PPy-BSA. We achieved the
electropolymerization on aryl-modified ITO strips by performing ten potential scanning cycles between
−1 and 1.2 V/ECS at 100 mV·s−1 in an aqueous solution of 0.1 M of pyrrole and 0.01 M BSA (Figure 5).
Table 2 reports the Epox and Epred values for the electropolymerization process together with the
oxidation peak density (Jpox ). Highly uniform and adherent PPy-BSA films were easily obtained on
all modified electrodes, regardless the nature of the diazonium compound. The films obtained on
the hydrophobic surfaces (ITO-N(CH3 )2 , ITO-NH2 , ITO-NH-Ph) were much more conductive than
those obtained on the hydrophilic surfaces (ITO-CO2 H, ITO-SO3 H, ITO-CN), as deduced from the
corresponding high current in the CV curves.

Figure 5. Cont.

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Surfaces 2018, 1

Figure 5. Electrosynthesis of PPy-BSA films by cyclic voltammetry on (A) ITO-SO3 H; (B) ITO-CO2 H;
(C) ITO-CN; (D) ITO-N(CH3 )2 ; (E) ITO-NH-Ph; and (F) ITO-NH2 .
Table 2. Electrochemical characteristics of electrodeposition of polypyrrole top layers: Eox , Ered and
Jpox values for the 6 ITO-aryl electrode materials.
Electrode Material

Eox (V)

Ered (V)

Jpox (μA/cm2 )

ITO-N(CH3 )2
ITO-NH2
ITO-NH-Ph
ITO-CN
ITO-CO2 H
ITO-SO3 H

0.25
0.27
0.2
0.25
0.41
0.24

−0.53
−0.52
−0.5
−0.58
−0.59
−0.43

972
960
936
208
908
228

As a matter of fact, the oxidation peak is centered at 0.25, 0.27, and 0.2 V/SCE for ITO-N(CH3 )2 ,
ITO-NH2 and ITO-NH-Ph electrodes, respectively. These peaks are positioned near those obtained
on the hydrophilic surfaces ITO-CN and ITO-SO3 H (0.25 and 0.24 V/SCE, respectively). In contrast,
on ITO-CO2 H the oxidation peak is centred at a higher potential value (0.41 V/SCE) and the reduction
peak at −0.59 V/SCE. This is the largest potential difference for the electropolymerization of pyrrole on
the actual series of diazonium-modified ITO electrodes. The density of Jpox of PPy-BSA on electrodes
materials takes the maximum values for hydrophobic surface (972 μA/cm2 on ITO-N(CH3 )2 ), hence the
highest conductivity achieved for electropolymerization of pyrrole on the actual electrode materials.
The results obtained so far show that thin films of PPy were preferentially deposited on
hydrophobic surfaces, in line with the literature [50,51].
Figure 6 shows the blocking effect and EIS results of PPy films on all aryl-modified ITO electrodes.
The blocking effect of PPy on all aryl layers was assessed with Fe(CN)6 3-/4- , as depicted in Figure 6A.
The quasi-reversible redox couple (located at 0.25 V and 0.1 V) obtained on the hydrophobic surfaces
was characterized by a higher current, which accounted for a more conductive surface. It is worth to
note that the PPy coating on ITO-N(CH3 )2 , ITO-NH2 , ITO-NH-Ph (Figure 6B) exhibited the lowest
charge transfer resistance (0.15, 0.36, 0.6 kΩ/cm2 , respectively). Moreover, the Rct increased for
PPy on ITO-SO3 H, ITO-CO2 H, and ITO-CN; the impedance curves were significantly broadened
and the semi-circle diameters were about 28.59, 21.44 and 10 kΩ/cm2 respectively. These results
indicated that ITO-SO3 H had the largest obstruction, which resulted in reducing electron-transfer
rate or increasing resistance to the flow of electrons. Furthermore, EIS showed that charge transfer
resistance of ITO-N(CH3 )2 -PPy was much lower compared to others.
The results reported so far stress the paramount importance of the hydrophobic character of
the substrate, here imparted by the aryl layers. As a matter of fact, PPy is relatively hydrophobic,

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Surfaces 2018, 1

and preferentially binds to hydrophobic surfaces [52]. Conversely, the strength of (macro) moleculepolypyrrole molecular interactions are maximized in water [53].

Figure 6. PPy-BSA film on (a) ITO-SO3 H, (b) ITO-CN, (c) ITO-CO2 H, (d) ITO-NH-Ph, (e) ITO-NH2 ,
(f) ITO-N(CH3 )2 . (A) Blocking effect of the PPy-ABS films on the ITO electrodes modified by the
different diazonium cations. (B) Impedance measurements recorded for PPy-ABS coatings on modified
ITO sheets.

3.3.2. XPS Characterization of PPy Top Coats and Reference Materials
Figure 7 displays survey regions of ITO, ITO-SO3 H, ITO-N(CH3 )2 , ITO-N(CH3 )2 -PPy samples in
the 100–600 eV range. S2p, C1s, N1s, and In3d/Sn3d doublets are centered at 168, 285, 400, 445–453 and
487–496 eV, respectively.

Figure 7. XPS survey regions of (a) ITO, (b) ITO-SO3 H, (c) ITO-N(CH3 )2 and (d) ITO-N(CH3 )2 -PPy.

Bare ITO (Figure 7a) exhibits a quasi-horizontal background in the 160–285 eV spectral region
and a sharp In3d doublet together with Sn3d peaks. Attachment of an aryl monolayer screened ITO
(Figure 7b,c) therefore resulting in an increase of the relative C1s peak intensity, and in the appearance

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Surfaces 2018, 1

of an S2p peak at ~168 eV for ITO-SO3 H (Figure 7b) and of a N1s peak at 400 eV for ITO-N(CH3 )2
(Figure 7c). After electropolymerization, the survey region resembled that of pure polypyrrole with
effective screening of the characteristic In3d and Sn3d features from ITO (Figure 7d).
All apparent compositions (in at. %) are reported in Table 2 for the main elements. The carbon
and nitrogen contents increased in the order ITO < ITO-aryl < ITO-aryl-PPy. Sulfur content is due to
the sulfonate group from ITO-SO3 H, and to all PPy-BSA topcoats. This sulfur content is invariably
higher for PPy-BSA than for ITO-SO3 H. In contrast, there is a clear decrease for the O at. %, as one
shifts from ITO to ITO-aryl, then PPy. Indium and tin atoms are detected for bare and aryl-modified
ITO surfaces only. After electropolymerization, In and Sn could not be detected since XPS is surface
specific technique (about 12–15 nm depth) while PPy topcoat thickness is in the micrometer scale range
(~3 μm, SEM image not shown). This is a much too large thickness for a topcoat to possibly detect the
buried layers using XPS (10–12 nm sampling depth for organic coatings).
Table 3 provides also the In4d/In3d intensity ratio for bare and aryl-modified ITO sheets.
Invariably, this ratio is higher for the modified ITO sheets. Indeed, an aryl layer screens ITO, but at
a different extent for the In4d and In3d core electrons. The sampling depth of In4d is about 12.6 nm
higher than that of In3d (10.8 nm) [47]. It follows that In3d peak attenuation is larger, hence In4d/In3d
intensity ratio increases upon aryl grafting. All aryl layers have nanometer scale thickness well below
10.8 nm (the sampling depth of In3d).
Table 3. XPS-determined apparent surface chemical composition of ITO, ITO-aryl and ITO-aryl-PPy samples.
Materials

C

N

S

O

In

Sn

In4d/In3d

ITO
ITO-SO3 H
ITO-SO3 H-PPy
ITO-COOH
ITO-COOH-PPy
ITO-CN
ITO-CN-PPy
ITO-NH-Ph
ITO-NH-Ph-PPy
ITO-NH2
ITO-NH2 -PPy
ITO-N(CH3 )2
ITO-N(CH3 )2 -PPy

32.5
45.2
73.9
58.6
78.7
71.4
75.2
49.2
75.5
42.9
77.3
52.7
76.5

0.77
9.29
1.68
10.5
4.91
9.23
1.97
9.83
1.31
12.2
1.13
6.94

1.00
1.38
1.13
2.12
1.71
2.22
1.18

39.1
30.0
15.4
25.1
9.71
16.8
13.5
29.6
12.9
33.1
8.22
28.2
15.4

25.4
20.5
11.3
5.88
16.2
16.2
15.5
-

3.06
2.43
3.34
1.04
3.05
6.55
2.48
-

0.155
0.167
0.174
0.179
0.171
0.159
0.181

As the blocking effect can be due to the thickness or grafting density of the top layer, we were
tempted to correlate Irel (in Table 1) to I4d /I3d XPS intensity ration (in Table 3). The correlation is poor
and the data corresponding to ITO-NH2 (I4d /I3d = 0.159; Irel = 46.6%) differs markedly from the trend
given by the 5 other surfaces. As a matter of fact, Baranton and Bélanger [48] clearly indicated that
blocking effect cannot be simply related to the film thickness or the surface coverage. Herein, we do
confirm such an observation previously reported [48] and confirm that the hydrophilic/hydrophobic
effects operate at the interface, which gives full credit to Downard and Prince [49]. It follows that barrier
properties cannot be interpreted simply in terms of coverage or thickness; herein we demonstrate that
barrier properties are very likely to be correlated to the hydrophobic/hydrophilic character of the
ITO-aryl surfaces as shown in Figure 4a,b.
3.4. Practical Adhesion Aspects of Polypyrrole Thin Layers
Practical adhesion of PPy thin layers has been qualitatively investigated by solvent washing,
ultrasonication and bend stress. Simple solvent washing using a laboratory wash bottle demonstrates
the efficiency of the diazonium adhesive layer in anchoring PPy to ITO. This is shown in a video
available in the Supplementary Material SM Video1. This video shows that without any aryl adhesive

335

Surfaces 2018, 1

layer, PPy is simply and easily removed upon washing with deionized water. The middle part of the
video (starting at 0.48 min) shows that another PPy film grown on bare ITO is very easily delaminated
using tweezers. In the last part of the video (starting at 1.07 min), prolonged washing of PPy topcoat
of ITO-NH2 electrode does not induce any adhesion failure. We have then compared adhesion of
PPy to bare ITO and ITO-NH2 using ultrasonication (video not shown): again, the aryl layer prevents
adhesion failure of the PPy, whereas on bare ITO the film is removed previously reported [22]. Actually,
the removal of PPy from ITO-NH2 was only possible through scratching. Such robust adhesion of
polypyrrole to aminophenyl-grafted glassy carbon electrode found by Patterson and Ignaszak [23]
confirms our previous findings on ITO-NH2 and supports the actual work which generalizes the strong
adhesion of polypyrrole to flexible ITO and other electrodes.
One of the important aspects of polymer adhesion is to define the type of adhesion bonding.
In a previous publication, Samanta et al. [47] have demonstrated that photochemically-prepared
PPy/silver nanocomposite film was removed using adhesive tape from bare ITO, whereas no failure
occurred in the case of pyrrole-functionalized ITO sheets. A failure within the polymer film is
defined as “cohesive failure”. Herein, ultrasonication of ITO-PPy resulted in the delamination of
the conductive polymer topcoat. XPS inspection of the ITO fracture surface, after ITO-PPy adhesion
failure, brought strong supporting evidence for a neat ITO surface without any attenuation of the
characteristic peaks from ITO (see Survey region in Figure 8a). It is worthwhile to note that the C1s
peak is weakly intense and is partly due to unavoidable adventitious contamination during the transfer
from atmosphere to the vacuum chamber of the XPS apparatus. The spectrum displayed in Figure 8a
is a clear sign of an “adhesive failure”, ca exactly at the ITO-PPy interface. Both materials are split
apart upon ultrasonication.

Figure 8. Adhesion testing of polypyrrole thin films on (a) bare ITO using XPS; and (b) ITO-NH2 using
SEM and XPS. Inset of 8b shows the bar graph of In/C atomic.

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Surfaces 2018, 1

We also examined the propensity of the ITO-NH2 -PPy hybrid electrodes to withstand bending
stress. The ITO-NH2 -PPy test electrode was prepared as described above and stuck on the outer wall of
a tube for 24 h bending stress; the bending radius was 3.2 mm as measured using an electronic digital
caliper (Fowler, model S225; Newton, MA, USA). This is a small radius and accounts for an important
stress [20]. Compared to the unbent sample, the PPy film withstood the very high bending stress and
remained stuck to the underlying ITO-NH2 substrate. However, we found a slight increase in the In/C
and Sn/C atomic ratios, which probably resulted from a possible diffusion of In and Sn in the top
coat. Indeed, SEM inspection of the ITO-NH2 -PPy surface after bending showed a crack in the ITO
layer (Figure 8b).
From these qualitative tests, clearly the conductive polymer films strongly adhere to the
ITO surface due to the molecular interactions that operate at interfaces between PPy and all
diazonium-modified ITO flexible electrodes. The aryl groups were shown to be bridged to the
oxide surface through metal-O-C bonds [54,55]. We have selected functional groups in para position,
that undergo either electrostatic interactions or electron-donor/acceptor bonds (Figure 9). Although not
covalent in nature, these interfacial interactions are sufficiently strong to keep PPy anchored to the
ITO via the aryl layer. The deposition of PPy on aryl-modified ITO is simple, fast, and provides strong
adhesion that withstands severe conditions, even under high mechanical stress. This contrasts with
the poor adhesion of polypyrrole to bare ITO, as previously reported in the literature [19].

Figure 9. Possible molecular interactions at aryl-polypyrrole interfaces.

337

Surfaces 2018, 1

4. Conclusions
In this work, flexible ITO sheets were grafted using in situ-generated diazonium salts bearing
electron acceptor (COOH, SO3 H and CN) or electron-donor (NH2 , N(CH3 )2 , NH-Phenyl) groups.
We have carefully grafted monolayers which were thin enough to prevent passivation but which
permitted to prepare adherent polypyrrole top coatings by electropolymerization. The aryl layers were
characterized by contact angle measurements and by XPS. Interestingly, the blocking effects and charge
transfer resistance of the aryl layers were correlated with their hydrophobic character. All aryl layers
served as remarkable coupling agents for the polypyrrole top coats which withstand sonication, solvent
washing and bend stress. On bare ITO, PPy could be removed by simple washing with water or using
tweezers, without any trace of polymer left on the surface as judged from XPS analysis of the fracture
surface. This accounts for adhesive failure of the coating. In contrast, on all aryl-modified ITO sheets
no delamination took place, and the polymer could be only removed by scratching the polypyrrole film.
The adherent polypyrrole exhibited a better redox activity in terms of high current output on the most
hydrophobic surfaces obtained with aminated, electron-donor groups, thus indicating the possibility
to tune both the adhesion strength and redox properties of the conductive polymer. Adhesion of
polypyrrole to ITO via aryl layers, is likely to be due to interfacial electrostatic interactions between
the electron acceptor groups COOH/COO− , SO3 H/ SO3 H and CN/C+ = N− and the polypyrrole
positively charged backbone. As far as the nitrogen-containing electron-donor groups, such as NH2 ,
N(CH3 )2 and NH-Phenyl, are concerned, the adhesion is attributed to n-σ* electron-donor-acceptor
interactions (the pyrrole N-H bond being the σ* acceptor).
This work sheds a new light on the use of diazonium salts to make adhesive primer layers
for electroactive polymers and other organic coatings. As silanes in the 20th century, so diazonium
salts could be regarded as magic molecular glues for tuning interfacial properties of next generation
advanced materials of relevance to sensors, printed soft electronics and flexible thermoelectrics to
name but these timely remarkable applications.
Supplementary Materials: The following are available online at http://www.mdpi.com/2571-9637/1/1/5/s1.
Author Contributions: Conceptualization of the research work by M.L., M.M.C., D.G.-S. and K.D.; Methodology
by M.L. and M.M.C.; Validation by M.M.C., K.D., D.G.-S., J.-J.A.; Formal Analysis by M.L., R.P., M.M.C.; Writing
of Original Draft was done by M.L. and M.M.C.; Writing: Review and Editing, M.L., M.M.C., J.-J.A. and M.A.O.;
Supervision, M.M.C. and D.G.-S.; Funding Acquisition, J.-J.A., M.A.O. and D.G.-S.
Funding: One of us (Momath Lo) gratefully thanks the Cooperation and Cultural Action Service of the French
Embassy in Senegal for a PhD grant.
Acknowledgments: All authors thank S. Chehimi (EFREI, Villejuif, France) for the video making presented as
Supplementary Material.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

On the Contrasting Effect Exerted by a Thin Layer of
CdS against the Passivation of Silver Electrodes
Coated with Thiols
Emanuele Salvietti 1, *, Walter Giurlani 1 , Maria Luisa Foresti 1 , Maurizio Passaponti 1 ,
Lorenzo Fabbri 1 , Patrick Marcantelli 1 , Stefano Caporali 2,3 , Stefano Martinuzzi 2,3 ,
Nicola Calisi 2,3 , Maddalena Pedio 4 and Massimo Innocenti 1,3, *
1

2
3
4

*

Dipartimento di Chimica, Università degli Studi di Firenze, Via della Lastruccia 3,
50019 Sesto Fiorentino (FI), Italy; walter.giurlani@unifi.it (W.G.); foresti@unifi.it (M.L.F.);
maurizio.passaponti@unifi.it (M.P.); lorenzo.fabbri2@stud.unifi.it (L.F.); patrick.marcantelli@unifi.it (P.M.)
Dipartimento di Ingegneria Industriale, Università degli Studi di Firenze, via S. Marta, 50139 Florence, Italy;
stefano.caporali@unifi.it (S.C.); stefano.martinuzzi@unifi.it (S.M.); nicola.calisi@unifi.it (N.C.)
Consorzio Interuniversitario Nazionale per la Scienza e Tecnologia dei Materiali, via Giusti 9,
50121 Florence, Italy
Istituto Officina dei Materiali, Consiglio Nazionale delle Ricerche, TASC Laboratory, 34012 Trieste, Italy;
pedio@iom.cnr.it
Correspondence: emanuele.salvietti@unifi.it (E.S.); m.innocenti@unifi.it (M.I.); Tel.: +39-055-457-3102 (M.I.)

Received: 14 June 2018; Accepted: 27 July 2018; Published: 31 July 2018

Abstract: The passivation of metal electrodes covered by self-assembled monolayers of long-chain
thiols is well known. The disappearance of the voltammetric peak of redox species in solution
is a classical test for the formation of full layers of thiols. Similar studies on semiconductors are
still very limited. We used silver surfaces covered by an ultrathin layer of CdS as substrate for
self-assembling of n-hexadecanethiol (C16 SH), and we compared the experimental results with those
obtained by using the bare silver surface as substrate. The strong insulating effect of C16 SH deposited
on Ag(III) is shown by the inhibition of the voltammetric peak of Ru(NH3 )6 3+/2+ . On the contrary,
the voltammogram obtained on CdS-covered Ag(III) is very similar to that obtained on the bare Ag(III)
electrode, thus suggesting that the presence of CdS exerts a contrasting effect on the passivation of
the silver electrode. A crucial point of our work is to demonstrate the effective formation of C16 SH
monolayers on Ag(III) covered by CdS. The formation of full layers of C16 SH was strongly suggested
by the inhibition of the stripping peak of Cd from the CdS deposit covered by C16 SH. The presence
of C16 SH was confirmed by electrochemical quartz crystal microbalance (EQCM) measurements as
well as by Auger electron spectroscopy (AES) analysis.
Keywords: ECALE; CdS; silver single crystals; alkanthiols; SAMs; EQCM; AES

1. Introduction
The assembly of alkanethiols on metals has been the focus of numerous studies in recent years [1,2].
Most of the work was carried out on gold substrates, although compact alkanethiols films form
spontaneously upon immersion of a variety of other metals like silver and copper. Similar studies on
semiconductors are still very limited. To date, semiconductive substrates used for alkanethiols SAMs
formation are InP [3–6], GaAs [7–9], and Ge [10].
Systematic investigations performed on n-alkanethiols SAMs on gold, with n ranging between 1
and 21, showed marked differences between long- and short-chain thiols monolayers [11]. In particular,
long-chain thiols were found to form a densely packed, crystalline-like assembly with fully extended
alkyl chains tilted from the surface normal by 20–30◦ . These long-chain thiols provide substantial
Surfaces 2018, 1, 29–42; doi:10.3390/surfaces1010004

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Surfaces 2018, 1

barriers to electron transfer and are strongly resistant to ion penetration. On the contrary, the barrier
becomes weaker while decreasing the chain length. At the same time, the structures become
less ordered.
Self-assembled monolayers of hydroxythiols have been used as insulating barriers between the
electrode surface and redox species in solution [12–15]. The closest approach of the redox species to
the electrode surface is limited by the thickness of the insulating SAM layer [16,17]. This barrier
decreases the electron transfer rate by increasing the separation between the electrode surface
and redox molecules. Electron transfer in these cases proceeds via electron tunneling through the
insulator, resulting in relatively slow kinetics. Electron tunneling through the full thickness of the
self-assembled layer is strongly suggested by the dependence of the electron transfer on the thickness
of the monolayer film [12]. From temperature-dependent current–voltage measurements carried out
on n-alkanemonothiol SAMs (with n = 8, 12 and 16), electron tunneling was also shown to be the
dominant transport mechanism [18]. More generally, the passivation of metal electrodes covered by
self-assembled monolayers of long-chain thiols is well known. The disappearance of the voltammetric
peak of redox species such as Ru(NH3 )6 3+/2+ or Fe(CN)6 3−/4− in solution is a standard test for the
formation of full layers of long-chain organic SAMs [11–14,19]. On the contrary, self-assembled
monolayers of alkanethiols on semiconducting surfaces seem to exhibit different properties. In this
respect, studies performed on electron tunneling at InP electrodes covered by alkanethiols showed that
the distance dependence is a factor of two softer than that of gold [4–6]. A softer distance dependence
for the electronic coupling for hole transfer from an alkanethiol covered n-InP electrodes to a redox
species in solution was also found. This observation can be explained by consideration of the change in
tunneling energy between the two types of electrodes. Earlier studies performed on organic SAMs on
gold explained the faradaic contributions by the electron transfer at defects sites and electron tunneling
at “collapsed” sites in the monolayer [20,21]. The hypothesis of possible contributions arising from
defects or pinholes of the alkanethiols films on InP was taken into account to explain the softer distance
dependence of the electron transfer. However, these contributions are likely to be very small [4].
The deposition of alkanethiols on semiconductors could open new perspectives in the vast field
of self-assembling phenomena. Our group has had great experience in the growth of II–VI [22–24]
and III–V compound semiconductors [25] on silver single crystals by the electrochemical atomic
layer epitaxy (ECALE) method [23,26–29]. This is a method based on surface-limited phenomena
such as underpotential depositions [30]. In the ECALE method, the UPD of the metallic element
is alternated with that of the non-metallic element to form a single monolayer of the compound
per cycle. The number of cycles determines how thick the deposit will be. The characteristics
of composition, morphology and structure of the compounds grown by ECALE, as well as the
bandgap values determined by photoelectrochemical measurements, indicate the high quality of these
compounds [23,24]. Hence, the ECALE technique holds the promise of being able to provide low-cost,
structurally well-ordered solids whose composition can be controlled at the nanoscopic level along the
direction perpendicular to the substrate. Therefore, the compound semiconductors grown by ECALE
seem to exhibit all of the needed characteristics for their use as a substrate for the self-assembling of
alkanethiols. In this paper we report on the self-assembly of 1-hexadecanethiol (C16 SH) on Ag(III)
covered by an ultrathin film of CdS deposited by ECALE, and we compare the experimental results
with those obtained by using the bare silver surface as substrate. Hexaaminerutenium ion was used as
a redox probe to test the electron transfer.
We chose CdS to exploit the high affinity of Cd towards the sulfur of thiol and due to the simplicity
of the deposition of this compound in a highly ordered form through ECALE. In fact, CdS grows on
Ag(III) along the Cd-terminated basal plane of the wurtzite [23], which is its most common crystalline
form in nature.
A long chain alkanethiol such as 1-hexadecanethiol (C16 SH) was chosen to ensure a strong
inhibition of the redox processes on the Ag(III) surface covered by the SAM.

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The self-assembling processes of thiols on a metal Me, such as Au, Ag and Cu, form stable
surface structures [1]. Complete electrodesorption of the n-alkanethiol SAMs occurs according to the
electroreduction reaction that yields a small peak preceding the hydrogen evolution reaction [31]:
Cn S− Me + e− → Cn S− + Me

(1)

2. Materials and Methods
2.1. Materials
Merck analytical reagent grade 3CdSO4 ·8H2 O and absolute ethanol, Aldrich analytical reagent
grade Na2 S, KCl, NaOH, Ru(NH3 )6 Cl3 and purity grade Fluka 95% hexadecanethiol were used without
further purification. Merck analytical reagent grade HClO4 and NH3 were used to prepare the pH 9.2
ammonia buffer. The working electrodes were silver single crystal discs grown in a graphite crucible,
oriented by X-rays and cut according to the Bridgman technique [32,33]. An automated deposition
apparatus consisting of Pyrex solution reservoirs, solenoid valves, a distribution valve and a flow-cell
was used under the control of a computer. The electrolytic cell was a Teflon cylinder with about a
5 mm inner diameter and a 30 mm outer diameter, whose inner volume, 0.5 mL, was delimited by the
working electrode on one side and the counter electrode on the other side. The inlet and the outlet for
the solutions were placed on the side walls of the cylinder. The counter electrode was gold foil, and the
reference electrode was an Ag/AgCl/sat. KCl placed on the outlet tubing. Both the distribution
valve and the cell were designed and realized in the workshop of the Department of Chemistry in
Florence [22].
2.2. ECALE Cycles for CdS Deposition
The procedure for CdS deposition is thoroughly described in reference [20]. Briefly, CdS growth
was obtained by depositing sulfur at −0.65 V from a Na2 S solution, washing the cell, injecting the
cadmium solution while keeping the electrode at the same potential, waiting 60 s to deposit Cd at
underpotential, washing the cell, and repeating this cycle as many times as desired. In this paper we
performed experiments on CdS deposited with 10 and 30 ECALE cycle. According to the thickness
found for CdS deposits obtained with 100 deposition cycles (20–25 nm) [24], the corresponding
thicknesses should range between 2 and 7.5 nm.
2.3. Self-Assembly of Alkanethiols on Ag(III) and on CdS-Covered Ag(III)
For the attainment of full layers formation, both substrates were immersed in a 0.3 mm solution of
alkanethiol in pure ethanol for at least 12 h. To avoid formation of thiols multilayers the substrates were
later kept for 1 h in pure ethanol. The effect of shorter treatments will be shown in the results section.
Alternatively, more dilute thiol solutions (0.03 mm) with longer times of modifications (more than
18 h) were used.
2.4. EQCM Measurement
EQCM measurements were carried out using the basic instrument supplied by Seiko EG&EG
(QCA917). The working electrode for the EQCM measurements was a 9 MHz AT-cut quartz crystal
with silver electrode furnished by Ditta Nuova Mistral (Latina, Italy). The diameter of the quartz
crystal was 14.0 mm, and the silver electrode diameter was 7.4 mm. The area of the working electrode
in contact with solution was limited to 0.43 cm2 by an O-ring. The silver electrode on the crystals
consisted of 300 nm Ag sputter deposited on an adhesion layer of 50 nm Ti. A suitable flow-cell
entirely made of Teflon was designed and realized in the workshop of the Department of Chemistry in
Florence. The counter electrode was gold foil, and the reference electrode was an Ag/AgCl sat KCl
placed on the outlet tubing of the cell.

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2.5. Auger Electron Spectroscopy (AES) Analysis
AES measurements were performed introducing the ex situ prepared samples, in an ultra-high
vacuum chamber equipped by a cylindrical mirror analyzer (CMA) with a coaxial electron gun.
The spectra were taken in the counting mode. The spectra in the range of 50–600 eV measure the
Auger features, namely C KLL, S MLL, Cd MNN and Ag MNN, and have been measured by 1000 and
3000 eV primary electron beams.
3. Results
3.1. SAM of C16 SH on Ag(III)
In Figure 1a a cyclic voltammogram (CV) of 0.1 mm C16 SH on Ag(III) electrodes using 0.1 M
NaOH in 95% ethanol and 5% water as recorded from −0.8 V to −1.6 V, after keeping the electrode in
the solution for 15 min a second voltammogram was recorded. The second CV overlaps completely
with the first, indicating that this time is insufficient to produce significant variation on the amount
of adsorbed C16 SH. However, even this reduced adsorption time is sufficient to strongly reduce the
capacitive current as shown by the flattening of the cyclic voltammogram in the supporting electrolyte
(inset of Figure 1a). Moreover, a significant inhibition of the Ru(NH3 )6 3+ reduction is also shown.
Figure 1b shows the cyclic voltammograms obtained from 1mm Ru(NH3 )6 3+ in 0.1 M KCl on the bare
Ag(III) (curve a), and on Ag(III) covered by C16 SH deposited from a 0.03 mm solution for 10 min
(curve b), 2 h (curve c) and 26 h (curve d). Similar dependence on the adsorption time was observed
for the cyclic voltammograms of Fe(CN)6 3−/4− on a gold electrode coated by thiols [34]. This behavior
is consistent with the kinetic model of thiol adsorption proposed by Nuzzo and coworkers [35].
This image cannot currently be display ed.

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(a)

(b)

Figure 1. (a) Cyclic voltammograms of 0.1 mm C16 SH on Ag(III) electrodes using 0.1 M NaOH in 95%
ethanol and 5% water as recorded from −0.8 V to −1.6 V/Ag/AgCl sat; inset: cyclic voltammogram
of a bare Ag(III) (solid curve) and on Ag(III) modified for 15 min in C16 SH solution (dashed curve);
(b) cyclic voltammograms obtained from 1 mm Ru(NH3 )6 3+ in 0.1 M KCl on the bare Ag(III) (curve a),
and on Ag(III) covered by C16 SH deposited from a 0.03 mm solution for 10 min (curve b), 2 h (curve c)
and 26 h (curve d). The scan rate was 50 mV/s.

It must be noted that the curves in Figure 1b were obtained from different experiments, and not
in succession. This means that the auto-assembling process was not disturbed at all. Analogous
experiments carried out in succession showed that the same degree of inhibition as that shown by
curve d is reached after three consecutive 10 min treatments of the electrode.

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Unfortunately, the reduction peak of C16 SH after SAM formation takes place at potentials similar
to those of Figure 1a. The partial overlap with hydrogen evolution hinders the accurate estimation of
the charge involved, and hence of the amount of C16 SH. Alternatively, the amount of the deposited
C16 SH can be obtained with electrochemical quartz crystal microbalance (EQCM) measurements by
measuring the mass decrease involved in thiol dissolution. Systematic EQCM measurements were
performed at increasingly longer modification times. The associated mass variation increases up to
reach a limiting value for times greater than 12 h. Repeated washings with ethanol are also necessary
to avoid formation of multilayers. Curve a in Figure 2 shows the frequency variation obtained
while recording a cyclic voltammogram from −0.1 to −1.8 V. Keeping in mind that a frequency
increase corresponds to a mass decrease, it is easy to verify that the SAM layer dissolution starts at
−1.25 V. However, in spite of the very low scan rate (2 mV/s), and the very negative final potential,
the dissolution process is not yet completed in the forward scan but continues during the backward
scan. Curve b in Figure 2 is obtained on a silver electrode not covered by SAM. Both curves were
recorded in a 0.1 M NaOH in 95% ethanol and 5% water. By comparison with curve a, the entire
frequency variation of the modified electrode seems to be ascribable to thiol dissolution. The mass
variation, Δm, is given by the Sauerbrey equation:
Δf = −

Cf Δm
A

(2)

where Δf is the frequency variation, A is the electrode area, and Cf is a coefficient that depends on the
quartz properties and on the fundamental resonance frequency. In our case, Cf = 0.183 Hz* cm2 /ng,
then, the frequency variation involved in curve a of Figure 2 (about 200 Hz for a mass coverage
of 3.9 Å2 /molecule) gives a mass variation higher than that deduced from the molecular coverage,
18.5 Å2 /molecule, found for alkanethiols on Ag(III) [1] and must be considered roughly approximate.
The EQCM measurements were performed using commercial polycrystalline silver electrodes (there is
no Ag monocrystalline (111) electrode for EQCM that are commercially available) with a very high
roughness factor, with the direct consequence of a larger real area respect to the geometrical one.
Therefore, the experimental mass variation involved in surface phenomena such as self-assembling is
scarcely indicative of what happens on a single crystal. Therefore, rather than to calculate the exact
amount of the deposited C16 SH, EQCM measurements were performed to establish the time necessary
for the complete formation of the SAM, and to constitute a reference for the analogous measurements
carried out in the presence of the CdS film, the measured amount of SAM in the two cases are consistent
meaning that the system is reproducible even if not approximal to an ideal flat surface.

Figure 2. Frequency change during a single potential scan from −0.1 to 1.8 V recorded on the silver
EQCM electrode modified with C16 SH for times greater than 12 h (curve a) and on a silver electrode
not covered by SAM (curve b). Both curves were recorded in a 0.1 M NaOH in 95% ethanol and 5%
water. The scan rate was 2 mV/s.

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The slowness of desorption process is confirmed by successive experiments: the silver electrode
was modified for 16 h. Then, it was repeatedly treated at E = −1.8 V up to a total time of 8 min. The effect
of each potential treatment was verified on a typical surface phenomenon such as the underpotential
deposition of sulfur. Figure 3a shows S UPD on a bare Ag(III) electrode, whereas curves a, b and
c of Figure 3b are obtained after each potential treatment. The S UPD peak does not increase with
further potential treatments. The attainment of a limiting behavior, together with the quasi-complete
disappearance of the thiol reduction peak in NaOH 0.1 M, suggests the quasi-complete desorption of
the thiol SAM layer. This experiment leads to two considerations. First, it confirms the incomplete
dissolution during the forward scan of Figure 2; in fact, at a scan rate of 2 mV/s, it takes about 5 min
to cover the potential range −1.25 V ÷ −1.8 V. Secondly, the partial inhibition of S UPD indicates that
traces of thiol are still present, even though the EQCM data suggest the complete dissolution.
It is interesting to note that the traces of thiol that are sufficient to inhibit surface phenomena
don’t significantly affect redox processes occurring in solution. To this purpose, we performed another
experiment: the electrode was modified for 2 h, and subsequently treated at E = −1.6 V for 3 min
and then for further 2 min (total time 5 min). Both the shorter time of modification and less negative
applied potential were chosen to better scale the experimental results. Then, the UPD of S (Figure 3c)
and the reduction curve of Ru(NH3 )6 3+ (Figure 3d) were alternatively recorded. The dashed curves
are recorded on the bare Ag(III) electrode; curves a’ and a” on the electrode modified for 2 h; curves b’
and b” after the first treatment at −1.6 V, and curves c’ and c” after the second treatment at −1.6 V.
From Figure 3d it is evident that curve c” is not far from the dashed curve, thus indicating that most of
the thiol has been removed. At the same time, curve c’ of Figure 3c is much lower than the dashed
one, thus again indicating that traces of thiols are sufficient to inhibit surface phenomena such as
underpotential depositions, providing to an over estimation of the thiol still present. This result is
in good agreement with the observation that n-alkanethiols with n > 6 could remain physisorbed on
the terraces [36]. However, our results seem to contradict the possibility of completely removing the
physisorbed micelles during electroreduction, at least for C16 SH.
The important findings of this experiment are that the criterion for SAM layer formation cannot
be based on surface limited phenomena (UPD), because of the interactions with other species on the
surface but must be based on the inhibition of the redox processes (Ru3+ /Ru2+ ) occurring in solution,
which seems be related only to the area actually available for the reaction. More precisely, curve d in
Figure 1b corresponds to the limiting shape of the inhibited process. Note that curve c in Figure 1b, albeit
not far from curve d, shows the typical inflection that is characteristic of a SAM layer not yet completed [37].
This image cannot currently be display ed.

This image cannot currently be display ed.

(a)

(b)

Figure 3. Cont.

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Surfaces 2018, 1

This image cannot currently be display ed.

(c)

(d)

Figure 3. (a) Oxidative UPD of S on Ag(III) from 1 mm Na2 S, as recorded from −1.1 to −0.68 V;
(b) oxidative UPD of S on Ag(III) modified with C16 SH for 16 h and successively treated at at E = −1.8
for 1 min (curve a), 3 min (curve b) and 4 min (curve c) up to a total time of 8 ; (c) oxidative UPD of S on
Ag(III) modified with C16 SH for 2 h (curve a’) and successively treated at E = −1.6 V for 3 min (curve
b’) and further 2 min (curve c’) up to a total time of 5 min; as comparison, the dashed curve is the curve
on the bare Ag(III); (d) cyclic voltammograms obtained from 1 mm Ru(NH3 )6 3+ in 0.1 M KCl on Ag(III)
modified with C16 SH for 2 h (curve a”) and successively treated at E = −1.6 V for 3 min (curve b”) and
further 2 min (curve c”) up to a total time of 5 min; as comparison, the dashed curve is the curve on the
bare Ag(III). All curves were recorded in pH 9.2 ammonia buffer solutions. The scan rate was 50 mV/s.

3.2. SAM of C16 SH on Ag(III) Covered by an Ultrathin Film of CdS
The first experiments were performed on Ag(III) covered by 10 ECALE cycles of CdS. As stated
in the experimental section, the thickness estimated for this deposit is about 2 ÷ 2.5 nm. Curve a
in Figure 4 shows the cyclic voltammogram of Ru(NH3 )6 3+ as obtained after having modified this
substrate with C16 SH for about 26 h. This curve is very close to that obtained on the bare Ag(III) (curve
b). The apparent lowering is actually due to the unavoidable decrease of the capacitive contribution
due to the SAM layer formation. For a comparison, curve c in the figure is the cyclic voltammogram
obtained on a Ag(III) covered by C16 SH, which is the same as curve d in Figure 1b.
3.3. C16 SH SAM Formation on CdS
Various evidences for SAM formation were obtained. EQCM measurements were performed
on a Ag(III) substrate covered by 10 ECALE cycles of CdS and modified with C16 SH in the same
experimental conditions as those of Figure 2. Figure 5 shows the comparison between the frequency
curves obtained in the presence of CdS (curve a) and in the absence (curve b). The almost similar
frequency variation suggests that the same amount of C16 SH is deposited on the two different
substrates. Incidentally, it is interesting to note that the only difference is in a more favored dissolution
process of C16 SH in the presence of CdS. The reduction process begins to occur at −0.8 V and it is over
at −1.7 V. Then, the dissolution process is completed during the forward scan. This observation is
in good agreement with the hypothesis of a contrasting behavior, exercised by the presence of CdS,
on surface reactivity limited by C16 SH. Other evidences for the SAM formation were given by AES
measurements. Table 1 reports the film thickness of CdS and C16 SH as estimated by the analysis of the
attenuated AES signal amplitudes [38,39] of substrates using the electron mean free path database of
NIST [40].

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This image cannot currently be display ed.

Figure 4. Cyclic voltammograms obtained from 1 mm Ru(NH3 )6 3+ in 0.1 M KCl on Ag(III) covered by
10 ECALE cycles and modified with C16 SH for 26 h (curve a), on bare Ag(III) (curve b), and on Ag(III)
modified with C16 SH in the absence of CdS (curve c).The scan rate was 50 mV/s.
Table 1. Film thicknesses of CdS and C16 SH as estimated by the analysis of the attenuated AES signal
amplitudes using the electron mean free path database of NIST.
Sample

Total Thickness (Å)

CdS Thickness (Å)

C16 SH Thickness (Å)

30 CdS/Ag(III)
C16 SH/Ag(III)
C16 SH + 10CdS/Ag(III)
C16 SH + 30CdS/Ag(III)

85 + 7
25 + 5
40 + 6
88 + 8

75 + 10
20 + 3
68 + 7

25 + 5
20 + 3
20 + 3

The thickness values estimated for the CdS films are in good agreement with the preceding
measurements [24]. It must be noted that the C16 SH SAM layer thickness on CdS is lower than that
on Ag(III), thus suggesting a higher tilt angle. The presence of traces of oxygen (1%) localized on
the surface and of a small amount of carbon (3%) probably arising from atmospheric contamination
were found. The latter value is difficult to be determined since the C signal is superimposed to the
silver peak.
The presence of the intact thiol molecule on Ag(III) and pre-covered CdS Ag(III) samples was
verified by X-ray absorption measurements on the C K-edge and S L2,3 edges and will be discussed in
a forthcoming paper.
Both EQCM and AES measurements confirm the presence of the C16 SH SAM layer. However,
none of them are able to detect details of the maximum coverage and the structure of SAM layer.
This latter measurement is provided by the inhibition of the stripping peak of Cd from the CdS deposit
covered by C16 SH. Curve a in Figure 6a is the stripping peak of Cd from CdS deposited on Ag(III) with
10 ECALE cycles. When such a substrate was modified with C16 SH for about 18 h, the dissolution of Cd
is completely hindered (curve b). For a comparison, curve c refers to only a partial formation of SAM,
as obtained when the substrate was modified for insufficiently long times. In this case, the stripping
peak is only partially inhibited. The stripping of Cd occurs at potentials progressively more positive as
the number of ECALE cycles is increased. Curve a in Figure 6b is the stripping curve of Cd from CdS
deposited with 30 ECALE cycles and curve b shows again the blocking effect of the thiol SAM Layer.

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It must be noted that the described behavior indicates that the thiol SAM physically blocks
the underlying CdS deposits. From this point of view, the possibility that curve a of Figure 4 was
simply due to a total or partial SAM removal and that, therefore, the effect of CdS was only apparent,
has to be disregarded. In fact, any partial Cd dissolution from the underlying CdS deposit should
distort the reduction curve of Ru(NH3 )6 3+ . As an example, Figure 7 shows the cyclic voltammogram
of Ru(NH3 )6 3+ on a silver electrode covered by 30 ECALE cycles of CdS not protected by the thiol.
The slightly positive current at the initial potential E = −0.08 V indicates that a small amount of Cd is
dissolved, what was reasonably expected on the basis of curve a of Figure 6b. This amount of dissolved
Cd is again reduced during the negative scan, giving rise to the small bump at about −0.33 V.

Figure 5. Frequency change during a single potential scan from −0.1 to 1.8 V recorded on the silver
EQCM electrode modified with C16 SH for times greater than 12 h in presence of CdS (curve a) and in
the absence (curve b). The scan rate was 2 mV/s.

This image cannot currently be display ed.

(a)

(b)

Figure 6. (a) Linear sweep voltammograms for the oxidative strippings of Cd from CdS deposited
with 10 ECALE cycles on Ag(III) not covered by thiol (curve a); modified for about 18 h (curve b) and
modified for short times (curve c); (b) linear sweep voltammograms for the oxidative strippings of Cd
from CdS deposited with 30 ECALE cycles on Ag(III) not covered by thiol (curve a); modified for about
18 h (curve b). The scan rate was 10 mV/s.

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Surfaces 2018, 1

This image cannot currently be display ed.

Figure 7. Cyclic voltammograms obtained from 1 mm Ru(NH3 )6 3+ in 0.1 M KCl on Ag(III) covered by
30 ECALE cycles in the absence of thiol. The scan rate was 50 mV/s.

3.4. Estimate of the SAM Layer Quality
An exhaustive discussion on the criteria used for defect level assessment is given by Miller et
al. in reference [12]. Here, electron tunneling through the full thickness of self-assembled organic
monolayers of ω–hydroxy thiols is demonstrated. Briefly, the possibility of either small pinholes
exposing the electrode surface or collapse sites [21] causing the monolayer to be significantly thinner
than the bulk monolayer SAM layer are taken into account. Because of the strong dependence of the
electron transfer kinetic on the thickness of the insulating SAM layer, even a small coverage of such
defects could be entirely responsible for the current measured. In order to assess the effect of such
defects, Miller et al. used the theoretical model of Amatore et al. for the redox kinetics at partially
blocked electrodes [37]. According to this model, the presence of pinholes or collapse sites separated by
distances greater than the characteristic diffusion length of the experiment (which, in our experiments,
is in the order of tens of microns) should give rise to sigmoidal voltammetric waves characteristic of an
array of microelectrodes. The behavior as an array of microelectrodes was also obtained after having
deliberately perforated monolayers-coated electrodes [41]. As a matter of fact, in our experiments,
sigmoidal waves were only observed with insufficient modification time (see for instance curve c
of Figure 1b). On the other hand, the presence of defects separated by distances smaller than the
characteristic diffusion length should give rise to voltammetric curves indistinguishable in shape from
those obtained at electrodes with continuous insulating layer.
Further evidences were obtained by modifying the Ag(III) covered by 30 ECALE cycles of CdS
for increasing time, in a way similar to the experiment of Figure 3. The thicker CdS film is necessary to
limit Cd dissolution during the recording of Ru(NH3 )6 3+ cyclic voltammogram. Figure 8a. shows the
curves of Ru(NH3 )6 3+ obtained after having modified the electrode for 20 min (curve a), 30 min (curve
b), 90 min (curve c) and 4 h (curve d). Surprisingly, the redox process shows an inhibition effect that
increases with the adsorption time. This result is in contrast with curve a of Figure 4 obtained after
having modified the substrate for 17 h. However, a further modification of 17 h yielded the dashed
curve in Figure 8a. This latter curve, although not coincident with curve a of Figure 4, indicates a
reversed trend. This behavior is consistent with a higher level of order ensured by the additional

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Surfaces 2018, 1

modification time. Then, it is reasonable to assume that the occurring of charge transfer is connected
to the order of the SAM rather than to the adsorption time.
This image cannot currently be display ed.

This image cannot currently be display ed.

(a)

(b)

Figure 8. (a) Cyclic voltammograms obtained from 1 mm Ru(NH3 )6 3+ in 0.1 M KCl on Ag(III) covered
by 30 ECALE cycles and modified with C16 SH for 20 min (curve a), 30 min (curve b), 90 min (curve
c) and 4 h (curve d); the dashed curve was obtained after having further modified the substrate for
17 h; (b) cyclic voltammograms obtained from 1 mm Ru(NH3 )6 3+ in 0.1 M KCl on Ag(III) covered by
30 ECALE cycles and modified with C16 SH for 3 h (curve a), and more than 12 h (curve b).The scan
rate was 50 mV/s.

As a matter of fact, curve a in Figure 8b shows the Ru(NH3 )6 3+ reduction peak obtained on
Ag(III) covered by 30 ECALE cycles and modified for a time (3 h), which, according to Nuzzo and
co-workers [35], is surely insufficient to produce an ordered monolayer. The Ru(NH3 )6 3+ reduction is
much less inhibited than in the experiments of Figure 8a. At the same time it is much more inhibited
that in the presence of a SAM formed on a similar substrate with an adsorption time longer than 12 h
(curve b). To avoid Cd dissolution, the initial potential was shifted to −0.05 V. Then, it is reasonable to
assume that the necessary condition for charge transfer is that the monolayer is ordered.
4. Conclusions
The behavior of self-assembled monolayers of 1-hexanedecanthiol (C16 SH) formed on Ag(III)
covered by ultrathin films of CdS was compared with that of analogous SAMs formed on the bare
Ag(III). The strong insulating effect of C16 SH deposited on Ag(III) is shown by the inhibition of
the voltammetric peak of Ru(NH3 )6 3+/2+ . On the contrary, the voltammogram obtained on Ag(III)
covered by an ultrathin film of CdS is very similar to that obtained on the bare Ag(III) electrode.
The crucial point of our work is to demonstrate the effective formation of C16 SH monolayers on
Ag(III) covered by CdS. A number of different experiments were performed to confirm the presence of
C16 SH. In particular, EQCM measurements revealed that an identical mass decrease is involved in
thiol dissolution both in the presence and in the absence of CdS. Then, AES measurements confirmed
the presence of C16 SH SAM layer both on Ag(III) and on CdS. The thickness of the SAM layer formed
on CdS, 20 Å, is lower than that of the SAM layer formed on Ag(III), 25 Å, thus indicating a higher
tilt angle.
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Surfaces 2018, 1

Neither EQCM nor AES are able to detect the formation of a complete SAM layer of thiol. However,
the evidence for a well-formed SAM layer is given by the physical blocking of the stripping peak
of Cd from the CdS deposit covered by C16 SH. For this reason, only partial inhibition of the Cd
stripping peak is observed with SAM layer of C16 SH obtained with not sufficiently long adsorption
times. Moreover, the presence of the intact thiol molecule on Ag(III) and pre-covered CdS Ag(III)
samples was verified by X-ray absorption measurements on the C K-edge and S L2,3 edges and will be
discussed in a forthcoming paper.
Our results suggest that self-assembly on semiconductors could become interesting whenever
a charge transfer is required. As an example, they could be used to anchor molecules that change
their properties following an electrochemical reduction or oxidation. The hypothesis that the observed
behavior was simply due to a total or partial SAM removal and that, therefore, the effect of CdS
was only apparent, has to be disregarded, since in this case the concomitant CdS dissolution should
be observed.
While the structure of thiols on Ag(III) is known to give well packed and ordered
structures [1,2,16,17,31], none of the measurements presented here allow us to determine the exact
structure or degree of crystallinity of the SAM layers obtained in the presence of CdS. Our results on
the occurrence of electron transfer are purely phenomenological, since they are limited to the system
studied, particularly Ag(III) covered by n-CdS using a reduction process such as that of Ru(NH3 )6 3+ .
Any interpretation would require the extension of the measurements to other substrates (for example
gold) and to other semiconductors (for example other II–VI or III–V compound semiconductors both
n-type and p-type). Finally, other electrochemical probes, such as Fe(CN)6 3− or Fe(CN)6 4− , should be
used to explore possible discrepancies towards reduction or oxidation processes.
Author Contributions: Funding acquisition, S.C.; Investigation, W.G., M.P. (Maurizio Passaponti), L.F. and P.M.;
Methodology, M.L.F., S.M. and N.C.; Supervision, M.P. (Maddalena Pedio) and M.I.; Writing-original draft, E.S.
Acknowledgments: The authors are grateful to Ferdinando Capolupo for the preparation of the silver single
crystal electrodes. The technical services of TASC laboratory are kindly acknowledged. The financial support
of the Italian CNR, of the Murst and of the Consorzio Interuniversitario per la Scienza e Tecnologia (INSTM) is
gratefully acknowledged. We are also grateful to Regione Toscana—Programma Operativo Regionale POR FESR
Toscana 2014–2020, BANDO N. 2: Progetti di ricerca e sviluppo delle MPMI, Project GADGET “Gioielli in Argento
da Galvanica, Ecologica e Tecnologica”.
Conflicts of Interest: The authors declare no conflicts of interest.

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(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

X-ray Absorption under Operating Conditions for
Solid-Oxide Fuel Cells Electrocatalysts: The Case
of LSCF/YSZ
Francesco Giannici 1, *, Giuliano Gregori 2 , Alessandro Longo 3,4 , Alessandro Chiara 1 ,
Joachim Maier 2 and Antonino Martorana 1
1
2
3
4

*

Dipartimento di Fisica e Chimica, Università di Palermo, Viale delle Scienze Ed. 17, 90128 Palermo, Italy;
alessandro.chiara31@gmail.com (A.C.); antonino.martorana@unipa.it (A.M.)
Max Planck Institute for Solid State Research, 70569 Stuttgart, Germany;
giulianogregori@hotmail.com (G.G.); s.weiglein@fkf.mpg.de (J.M.)
Netherlands Organization for Scientific Research (NWO), Dutch-Belgian Beamline, ESRF—The European
Synchrotron, CS40220, 38043, 71 Avenue des Martyrs, 38000 Grenoble, France; alessandro.longo@cnr.it
Istituto per lo Studio dei Materiali Nanostrutturati (ISMN)-CNR, UOS Palermo, Via Ugo La Malfa, 153,
90146 Palermo, Italy
Correspondence: francesco.giannici@unipa.it

Received: 5 December 2018; Accepted: 2 January 2019; Published: 8 January 2019

Abstract: We describe a novel electrochemical cell for X-ray absorption spectroscopy (XAS)
experiments during electrical polarization suitable for high-temperature materials such as those
used in solid oxide fuel cells. A half-cell LSCF/YSZ was then investigated under cathodic and anodic
conditions (850 ◦ C and applied electrical bias ranging from +1 V to −1 V in air). The in situ XAS
measurements allowed us to follow the LSCF degradation into simple oxides. The rapid deterioration
of LSCF is ascribed to the formation of excess of oxygen vacancies leading to the collapse of the mixed
perovskite structure.
Keywords: SOFC; cathode; XAFS; in situ

1. Introduction
Solid-oxide fuel cells (SOFC) represent a forefront theme in materials research, since they allow
the production of electrical energy from hydrogen or hydrocarbon fuels with very high thermodynamic
efficiency. SOFC electrodes act as catalysts (for oxidation of the fuel at the anode, and for reduction of
oxygen at the cathode), and as mixed electron/ion conductors. Although a great deal of performance
testing is routinely carried out on a wide variety of electrocatalysts to optimize their composition and
microstructure, advanced structural characterization studies have been relatively uncommon so far.
In this paper, we studied, under operating conditions (at high temperature and under an external
electrical potential), the atomic and electronic structure of La0.6 Sr0.4 Co0.8 Fe0.2 O3 (LSCF), which is a
technologically relevant electrode material for high-temperature and intermediate-temperature SOFC.
This material has been the subject of a large number of studies in the last several decades.
Kawada et al. [1] reported on the chemical stability of LSCF, ultimately concluding that it is
stable at high temperatures and under different atmospheres. However, cation segregation on one
surface was observed when an oxygen partial pressure gradient was applied to LSCF, which resulted
in Co-rich and Sr-rich precipitates on the surface exposed to higher P(O2 ).
In general, few examples of characterization in operating conditions have been reported in the
literature to date, regarding both in-house and large-scale facilities techniques. Among the few studies
in the literature on this topic, Baumann et al. [2] observed a strong activation for the oxygen reduction
reaction in an LSCF/YSZ model system, associated with a steep decrease of the resistivity of LSCF, due
Surfaces 2019, 2, 32–40; doi:10.3390/surfaces1010003

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Surfaces 2019, 2

to an applied electrical potential (either cathodic or anodic) of the order of several volts. This effect was
found to last for several hours after the polarization, which suggests a more profound transformation
rather than a simple change of oxidation state. X-ray photoelectron spectroscopy (XPS) performed on
films subject to different treatments were found to be enriched with Sr and Co on the surface.
An ever-increasing interest is developing at synchrotrons around the world to improve in
situ/operando experimental setups in order to control the chemical/physical environment and the
electrical loading applied to the sample under the beam. This allows for accurately monitoring
the chemical changes in technologically relevant materials under service conditions. This approach
was used, for instance, to study several SOFC cathode-electrolyte-air triple boundary surfaces using
photoelectron microscopy under polarization [3]. More recently, Nenning et al. [4] examined the
electrochemistry and XPS spectra of lanthanum strontium ferrite (LSF) and lanthanum strontium
cobaltite (LSC) thin film electrodes over YSZ, during polarization (around 0.5 V either cathodic or
anodic) and in different atmospheres (H2 or O2 ). Given the surface-sensitive conditions of such an
XPS investigation, a pronounced difference was found in the behavior of cobalt compared to all other
elements, which was attributed by the authors to the higher reducibility of the Co ions. In fact, in an
earlier soft X-ray absorption spectroscopy study of LSCF under different oxygen partial pressures at
800 ◦ C, it was already found that the Co absorption edge is modified by the oxygen non-stoichiometry,
while the Fe edge is not [5].
Among synchrotron techniques, X-ray absorption spectroscopy allows one to obtain information
about the atomic and electronic structure of the samples, so that it is particularly suited for SOFC
electrode materials, often constituted by variable valence elements that can modify their oxidation
state and the local structure in dependence of reaction environment and applied potential [6–9].
Regarding this concern, the development of reliable experimental protocols [10] is essential to obtain
significant information on SOFC materials [11–14]. Among the issues addressed by the above cited
authors, Lai et al. investigated catalyst poisoning, arguing that deactivation of an LSCF cathode
by CO2 is enhanced under a cathodic bias. Skinner et al. observed that cathodes constituted by
Ruddlesden-Popper phases did not undergo significative redox phenomena under polarization.
Irvine et al. studied the processes occurring at the electrode-electrolyte interface, and, in particular,
assessed the reversible exsolution phenomena occurring during the operation.
We have recently developed an electrochemical cell for the X-ray structural characterization
of SOFC electrocatalysts under operating conditions: high temperature (up to 800 ◦ C), controlled
atmosphere, and electrical loading. To this aim, we modified a thermochemical cell for in-situ X-ray
spectroscopy with an electrically insulating quartz sample holder, and current collectors in order to
impose a voltage bias on the sample.
In this paper, we present our first results with this setup on a LSCF/YSZ cathode/electrolyte
half-cell. By applying an external voltage bias, ranging from −1 V to +1 V, and employing X-ray
absorption spectroscopy at a high temperature, we were able to investigate the different role of iron and
cobalt concerning the oxidation state and structural rearrangement, as a function of time, temperature,
and electrical polarization in operating conditions.
2. Materials and Methods
A dense La0.6 Sr0.4 Co0.8 Fe0.2 O3 (LSCF) film of 200 nm thickness was deposited by pulsed laser
deposition on yttria-stabilized zirconia (100) oriented crystal (YSZ, 9.5 mol% Y2 O3 , 5 mm × 5 mm × 0.5 mm,
CrysTec GmbH, Berlin, Germany) from a commercial LSCF target (HITEC, Karlsruhe, Germany). The
deposition was carried out with a laser pulse energy of 1.6 J/cm2 and a pulse frequency of 5 Hz,
which keeps the substrate temperature at 770 ◦ C under an oxygen partial pressure of 0.4 mbar.
After deposition, the films were annealed at 640 ◦ C for 30 min in oxygen.
Platinum current collectors were deposited by sputtering on opposite sides of the bilayer, covering
half of the surface, and enabling the other half to be studied in the spectroscopy experiment (see
Figure 1). The sample was mounted on a fused quartz sample holder inside the “Microtomo”

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furnace developed at the European Synchrotron Radiation Facility (Grenoble, France) for in situ
X-ray absorption spectroscopy. The voltage across the electrode-electrolyte bilayer is applied through
an external I-V source, through two platinum wires ending with platelets, which are brought in contact
with the current collectors on the sample. For this experiment, a constant electrical bias was applied to
the half-cell, which ranged from −1 V (cathodic) to +1 V (anodic). In this convention, “cathodic” refers
to a negative polarization of the LSCF side electrode with respect to the counter-electrode on the YSZ
side (oxygen ions flowing from LSCF to YSZ).

Figure 1. (a) LSCF/YSZ half-cell. (b) Schematic of the X-ray absorption spectroscopy setup. (c) Inner
view of the cell.

A steady air flux flowed inside the cell during the measurement. In this cell, the sample holder
can be mounted at either 45 degrees, or parallel to the incoming beam, which allows X-ray absorption
spectroscopy in a transmission or a fluorescence mode. In particular, the X-ray absorption near edge
structure (XANES) data presented in this paper were acquired at the Fe K-edge (7112 eV) and at the
Co K-edge (7709 eV) in a fluorescence mode using a multi-element solid state Ge detector.
3. Results
During the heating of LSCF from RT to 800 ◦ C, the Co K-edge shows a gradual variation, which
is evident from Figures 2 and 3. In particular, from Figure 3, where the intensity ratio between the
features at 7729 eV and 7734 eV is plotted, we get an impression how continuous the evolution of
the XANES spectrum is with the temperature, without evident discontinuities due to the known
rhombohedral to a cubic structural transition below 500 ◦ C [15]. It is likely that the progressive changes
in the edge are linked to the structural rearrangement, but also to the expected increase of oxygen
vacancies concentration with the temperature.
The applied electrical bias as a function of time is plotted in Figure 4. The position of the
absorption K-edge of cobalt (i.e., its oxidation state) is plotted in the same way, and it follows almost
instantaneously the bias changes (Figure 5). On the other hand, it can be observed in the plot that,
within the first cycle of cathodic/anodic polarization, there is a continuous transformation of the
edge that takes about 30 min, which suggests a diffusive process and/or a chemical transformation
different from a simple change of oxidation state. One can appreciate that, after an initial polarization
at +1 V leading to a slight shift towards a larger Co K-edge value and an abrupt change to −1 V
yielding to a reduction of Co (lower values of the Co K-edge), the return to +1 V does not bring the

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Co K-edge to the initial energy values. Although the energy difference is rather small, this might be
indicative of irreversible processes occurring in the LSCF layer after only 60 min. This is supported
by the subsequent voltage step. In this case, during the application of −1 V, the edge energy does
not reach the initial values, which suggests that the LSCF film already underwent some permanent
structural modifications.

Figure 2. Co K-edge XANES spectra during heating, from room temperature (black line) to 800 ◦ C
(red line). Intermediate spectra are plotted in grey.

Figure 3. Intensity ratio between the Co K-edge XANES maximum at 7729 eV and the local minimum
at 7734 eV.

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Figure 4. DC bias applied during the in situ Co K-edge XANES experiment at 850 ◦ C.

Figure 5. Energy of the Co K-edge, measured as the maximum of the first derivative of the XANES
spectrum, as a function of time at 850 ◦ C.

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These edge transformations show a clear irreversibility, which is evident when comparing the
red spectrum in Figure 6 and corresponds to the initial anodic bias and the final (green) spectrum,
taken in anodic polarization, but after the cathodic step. The postmortem Co K-edge resembles that
of the spinel Co3 O4 [15]: it is completely different from the starting LSCF and this modification is
maintained irrespective of the applied potential. The room temperature spectra, taken before and after
polarization cycles, in Figure 7, confirm with higher resolution the final transformation of LSCF.

Figure 6. Representative Co K-edge XANES spectra at 850 ◦ C during two cycles of the electrical
polarization experiment. No applied bias (black), +1 V (red), −1 V (blue), +1 V (green).

Figure 7. Room temperature Co K-edge XANES spectra of LSCF before treatment (black line) and after
the in situ electrical polarization experiment (red line).

A subsequent polarization cycle (Figure 8) with simultaneous acquisition of the Fe K-edge spectra
highlights that: (1) the Fe edge is almost unchanged in response to the electric bias, (2) the spectrum is
different from LSCF, and resembles Fe2 O3 nanoparticles [16,17].

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Figure 8. Fe K-edge XANES spectra measured at 850 ◦ C with an applied electrical bias: +1 V (black),
−1 V (red), 0 V (blue), over the course of 40 minutes.

The evidence points toward the complete and irreversible transformation of the material. The post
mortem XANES spectrum on the iron K-edge shows a deep modification with respect to the fresh
sample (Figure 9), with the treated sample resembling Fe2 O3 [16,17]. It appears that the polarization
cycling has produced a total breakdown of the LSCF structure, with de-mixing of the starting material
in iron oxides, cobalt oxides, and other compounds. We propose that such a rapid structural breakdown
is induced both by the high cathodic voltage (−1 V) and by the abrupt cycling between cathodic and
anodic (+1 V) conditions. It has been noted in the past that the stability of the LSCF phase is tested and
optimized for SOFC operation (that is, cathodic polarization), and not for the inverse operation that is
characteristic of an electrolyzer. In fact, the current density curves show a pronounced asymmetry
between cathodic and anodic polarization for LSCF cathodes compared to other materials such as
lanthanum strontium manganites [18].

Figure 9. Room temperature Fe K-edge XANES spectra of LSCF before treatment (black line) and after
the in situ electrical polarization experiment (red line).

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4. Discussion
This paper presents one of the few structural studies of SOFC electrodes during operation.
In particular, we investigated a LSCF/YSZ half-cell under electrical polarization. On the basis of
the post mortem XANES spectra of iron and cobalt, we can state that, as a consequence of the high
temperature polarization cycling, the LSCF structure was disrupted, producing the de-mixing of the
spinel Co3 O4 , with a Co oxidation state intermediate between +2 and +3, and of hematite Fe2 O3 with
Fe oxidation state +3. This means that both cationic species are reduced compared to the LSCF mixed
oxide [19]. It has been reported before that an applied cathodic bias increases the concentration of
oxygen vacancies [20]. It appears then that the increased concentration of oxygen vacancies above a
certain level induces the collapse of the perovskite mixed oxide, which is also favored by the high
temperature. This eventually leads to a rearrangement of the mixed oxide toward the simple oxides
retaining an oxidation state lower than that of LSCF.
Author Contributions: Supervision, J.M. and A.M.; Writing—original draft, F.G., G.G., A.L. and A.C.
Funding: We acknowledge funding through MIUR projects Futuro in Ricerca “INnovative Ceramic and hybrid
materials for proton-conducting fuel cells at Intermediate Temperature: design, characterization, and device
assembly (INCYPIT)”, PON R&C 2007-1013 “Tecnologie ad alta Efficienza per la Sostenibilità Energetica ed
ambientale On-board (TESEO)”, PRIN2010 “Celle a combustibile ad ossido solido operanti a temperatura
intermedia alimentate con biocombustibili (BIOITSOFC)”.
Acknowledgments: We acknowledge the ESRF for provision of beamtime. We thank H. Vitoux (ESRF) for
assistance in the cell design and during the experiment, B. Stuhlhofer (MPI) for thin film deposition, R. Merkle
(MPI) for helpful discussion, and the Glass workshop of MPI for the sample holder fabrication.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
article distributed under the terms and conditions of the Creative Commons Attribution
(CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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Article

Switchable Interfaces: Redox Monolayers on Si(100)
by Electrochemical Trapping of Alcohol Nucleophiles
Long Zhang 1,2 , Ruth Belinda Domínguez Espíndola 1 , Benjamin B. Noble 3 ,
Vinicius R. Gonçales 4 , Gordon G. Wallace 2 , Nadim Darwish 1, *, Michelle L. Coote 3, *
and Simone Ciampi 1, *
1

2
3
4

*

School of Molecular and Life Sciences, Curtin Institute of Functional Molecules and Interfaces, Curtin
University, Bentley, Western Australia 6102, Australia; long.zhang@unsw.edu.au (L.Z.);
rdominguez@uaem.mx (R.B.D.E.)
ARC Centre of Excellence for Electromaterials Science, Intelligent Polymer Research Institute, University of
Wollongong, Wollongong, New South Wales 2500, Australia; gwallace@uow.edu.au
ARC Centre of Excellence for Electromaterials Science, Research School of Chemistry, Australian National
University, Canberra, Australian Capital Territory 2601, Australia; benjamin.noble@anu.edu.au
School of Chemistry, The University of New South Wales, Sydney, New South Wales 2052, Australia;
v.goncales@unsw.edu.au
Correspondence: nadim.darwish@curtin.edu.au (N.D.); michelle.coote@anu.edu.au (M.L.C.);
simone.ciampi@curtin.edu.au (S.C.); Tel.: +61-8-9266-9009 (S.C.)

Received: 29 June 2018; Accepted: 17 July 2018; Published: 20 July 2018

Abstract: Organic electrosynthesis is going through its renaissance but its scope in surface science
as a tool to introduce specific molecular signatures at an electrode/electrolyte interface is under
explored. Here, we have investigated an electrochemical approach to generate in situ surface-tethered
and highly-reactive carbocations. We have covalently attached an alkoxyamine derivative on an
Si(100) electrode and used an anodic bias stimulus to trigger its fragmentation into a diffusive
nitroxide (TEMPO) and a surface-confined carbocation. As a proof-of-principle we have used
this reactive intermediate to trap a nucleophile dissolved in the electrolyte. The nucleophile was
ferrocenemethanol and its presence and surface concentration after its reaction with the carbocation
were assessed by cyclic voltammetry. The work expands the repertoire of available electrosynthetic
methods and could in principle lay the foundation for a new form of electrochemical lithography.
Keywords: electrosynthesis; switchable surfaces; alkoxyamine surfaces; redox monolayers

1. Introduction
Synthetic organic electrochemistry traces its origin back to the work of Faraday and Kolbe and its
green credentials are currently prompting a renaissance [1–3]. Chemical reactions that are coupled to
the flow of electricity allow, for instance, chemists to generate unstable intermediates in situ, to control
very precisely and accurately the supply of reactants and to monitor reaction processes in real-time [4].
While the large majority of the work on synthetic electrochemistry has focused on bulk synthesis [5,6],
there is also a strong motivation to expand electrochemical synthetic methods toward the chemical
modification of interfaces [7–10]. Addressing the molecular details of a surface [11], especially those
of semiconductors, has been central to the development of fields such as molecular electronics [12],
sensing [13], energy conversion [14] and cell biology [15]. Semiconductor electrodes and in particular
silicon electrodes, have the advantage of being readily available in a crystalline form, they have unique
photo-electrochemical properties and can form covalently-bound monolayers [16]. Silicon like all
non-oxide semiconductors is thermodynamically unstable and especially under anodic polarization
in aqueous environments tends to grow an electrically-insulating silica layer [17,18]. A common

Surfaces 2018, 1, 3–11; doi:10.3390/surfaces1010002

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laboratory approach to attach covalently organic molecules on silicon substrates is a three-step wet
chemical process. Removal of the native silica layer with fluoride-containing solutions is followed by
chemical passivation [19] of the hydrogen-terminated silicon surface by means of hydrosilylation of
terminal alkenes or alkynes [20–22] and finally chemical derivatization of the aliphatic monolayer [16].
In this work we seek to expand the chemical repertoires of this last step and explore the synthetic
scope of reactive carbocations in the context of surface chemistry.
There are several methods to generate carbocations in solution [23] but only few examples are
available for electrochemical generation of carbocations at electrodes. We have recently reported
on the putative generation of carbocations at metallic electrodes after the fragmentation of anodic
intermediates of alkoxyamines. Alkoxyamines are heat-labile molecules, widely used as an in situ
source of nitroxides in polymer and materials sciences. We have shown that the one-electron oxidation
of an alkoxyamine leads to an anodic intermediate that rapidly fragments releasing a nitroxide species
at room temperature [24].
In the current work we have developed a surface model system to explore the feasibility of using
surface-tethered carbocations to trap solution nucleophiles. The carbocation is electro-generated in
situ from an alkoxyamine molecule that is exposed at the distal end of an organic monolayer grown
on a Si(100) electrode. By way of applying a positive bias to the Si(100) electrode, the alkoxyamine is
anodically cleaved to release a 2,2,6,6-tetramethylpiperidine-1-oxyl (TEMPO) molecule and generate
a surface-confined carbocation species. The latter is attacked by electron donors via nucleophilic
substitutions. As proof-of-principle we have prepared redox-active monolayers by generating the
surface carbocations in the presence of ferrocenemethanol (2) molecules.
2. Materials and Methods
2.1. Chemicals
All chemicals, unless specified otherwise, were of analytical grade and used as received.
Milli-Q™ water (>18 MΩ cm, Millipore Corporation, Burlington, MA, USA) was used to prepare
solutions and to clean all glassware. Anhydrous solvents used in chemical reactions were purified
under nitrogen by a solvent drying system from LC Technology Solutions Inc. Dichloromethane
(DCM), methanol (MeOH) and 2-propanol were redistilled prior to use. Hydrogen peroxide
(30 wt % in water), ammonium fluoride (PuranalTM , 40 wt % in water) and sulfuric acid
(PuranalTM , 95–97%) were used to clean the wafers and were of semiconductor grade. 1,8-nonadiyne
(Sigma-Aldrich, 98%) was redistilled from sodium borohydride (Sigma-Aldrich, 99+%) under
reduced pressure (80 ◦ C, 10−12 Torr) and stored under a high purity argon atmosphere prior
to use. Tetrabutylammonium hexafluorophosphate salt (Bu4 NPF6 , Sigma-Aldrich, ≥98%) was
used as a supporting electrolyte. 4-Vinylbenzyl chloride (90%), ammonium sulphite, ferrocene
(98%), ferrocenemethanol (2, 97%) and 2,2,6,6-tetramethylpiperidine-1-oxyl (TEMPO hereafter,
98%) were purchased from Sigma-Aldrich. The surface reactive (azide-tagged) alkoxyamine
1-(1-(4-azidomethyl)phenyl)ethoxy-2,2,6,6-tetramethylpiperidine (1) was synthesized according to
minor modification of previously reported procedure (see Scheme 1) [25]. Prime grade, single-side
polished silicon wafers of 100-orientation (<100> ± 0.5◦ ), p-type (boron-doped) of 100 mm diameter,
500–550 μm thickness and of a nominal resistivity of 0.001–0.003 Ω cm were obtained from Siltronix,
S.A.S. (Archamps, France).

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Scheme 1. Synthesis of the azide-tagged alkoxyamine 1. i. NaN3 , DMF, 82%. ii. TEMPO,
Mn(OAc)3 ·2H2 O, NaBH4 , Toluene/EtOH, 50%.

2.2. Synthetic Methods
Thin-layer chromatography (TLC) was performed on silica gel Merck aluminium sheets (60 F254 ).
Merck 60 Å silica gel (220–400 mesh particle size) was used for column chromatography. Nuclear
magnetic resonance (NMR) spectra were recorded on a Bruker Avance 400 spectrometer in deuterated
dimethyl sulfoxide (d-DMSO) using the residual solvent signal as internal reference. High-resolution
mass spectral data (HRMS, mass accuracy 2–4 ppm) of alkoxyamine 1 were obtained using a Waters
Xevo QTof MS via ESI experiments and infusing the sample at 8 μL/min.
Synthesis of 4-Vinylbenzyl azide (VBA). Sodium azide (1.30 g, 20 mmol) was added in one portion to a
stirred solution of 4-vinylbenzyl chloride (1.53 g, 10 mmol) in N,N-dimethylformamide (DMF, 20 mL).
The reaction mixture was stirred at room temperature for 12 h under argon then poured into a large
excess water (100 mL). The mixture was extracted with DCM (3 × 50 mL). The combined organic layers
were then washed with brine (2 × 100 mL), dried over MgSO4 , filtered and dried under vacuum to
afford the crude title compound as a brown oily residue. The crude material was purified by silica gel
column chromatography (hexane) to give VBA as a light yellow oil (1.31 g, 82%). 1 H NMR (400 MHz,
d-DMSO): δ 7.50 p.p.m. (d, Ar-H, J = 8.12 Hz, 2H), 7.35 (d, Ar-H, J = 8.04 Hz, 2H), 6.74 (dd, Ar-CH=CH2 ,
J = 17.66 Hz, 1H), 5.85 (dd, Ar-CH=CH2 , J = 17.68 Hz, 0.96 Hz, 1H), 5.28 (dd, Ar-CH=CH2 , J = 11.84 Hz,
0.92 Hz, 1H), 4.43 (s, Ar-CH2 -N3 , 2H); 13 C NMR (100 MHz, d-DMSO): δ 136.96, 136.11, 135.13, 128.69,
126.39, 114.64, 53.32.
Synthesis of 1-(1-(4-Azidomethyl) phenyl)ethoxy-2,2,6,6-tetramethylpiperidine (1). Alkoxyamine 1 was
synthesized from VBA via the following procedure. To an ice-cold solution of TEMPO (0.31 g, 2 mmol)
in toluene/ethanol (60 mL, 1:1, v/v) VBA (3.20 g, 20 mmol) and Mn(OAc)3 ·2H2 O (5.36 g, 20 mmol)
were added in one portion while stirring in air. Stirring was continued for one min and then a 15-fold
molar excess (with respect to TEMPO) of NaBH4 was added in portions over 15 min. After stirring
overnight under nitrogen atmosphere, the residue was isolated by filtration, the filtrate was suspended
in water and the aqueous solution was then extracted three times with DCM. The combined organic
layers were evaporated under vacuum and the crude material was purified by silica gel column
chromatography (ethyl acetate/hexane, 1:40, v/v) to yield alkoxyamine 1 as a colourless oil liquid
(0.32 g, 50%).
1 H NMR (400 MHz, d-DMSO): δ 7.36–7.29 p.p.m. (m, Ar-H, 4H), 4.75 (q, NO-CH-Ar, J = 13.50 Hz,
1H), 4.42 (s, N3 -CH2 -Ar, 2H), 1.54–1.38 (m, 6H), 1.37–1.19 (m, 6H), 1.12 (s, 3H), 0.97 (s, 3H), 0.57 (s, 3H);
13 C NMR (100 MHz, d-DMSO): δ 145.08, 134.08, 128.21, 126.76, 82.09, 59.22, 58.98, 53.40, 33.99, 33.76,
23.00, 20.04, 16.66; HRMS (1, m/z): [M + H]+ calcd for C18 H29 N4 O 317.2336, found 317.2335.

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2.3. Surface Modification
2.3.1. Light-Assisted Hydrosilylation of 1,8-Nonadiyne on Si(100) (S-1)
The assembly of the acetylenylated Si(100) surface by covalent attachment of 1,8-nonadiyne on
hydrogen-terminated silicon is based on a photochemical hydrosilylation method [26]. In brief, silicon
wafers were mechanically cut into pieces (approximately 10 × 10 mm in size), rinsed several times
with small portions of DCM, MeOH and Milli-Q™ water. The samples were then immersed in hot
Piranha solution (100 ◦ C, a 3:1 (v/v) mixture of concentrated sulfuric acid to 30% hydrogen peroxide,
Caution: piranha solution reacts violently with organic substances) for 20 min. The samples were then
rinsed with water and immediately etched with a deoxygenated 40% aqueous ammonium fluoride
solution for 5 min under a stream of argon. A small amount (ca. 5 mg) of ammonium sulphite was
added to the etching bath for building an anaerobic environment and hence avoiding the in situ
oxidation of the hydrogen-terminated silicon. The freshly etched samples were washed sequentially
with Milli-Q™ water and DCM and blown dry with argon before dropping a small deoxygenated
sample of 1,8-nonadiyne (approximate 50 μL) on the hydrogen-terminated wafer and covering it with
a quartz slide to minimize evaporation. The wafer was then rapidly transferred to an air-tight UV
reaction chamber and kept under positive argon pressure. A collimated LED source (λ = 365 nm,
nominal power output >190 mW, Thorlabs part M365L2 coupled to a SM1P25-A collimator adapter)
was fixed over the sample at a distance of about 10 cm. After illumination for a 2 h period, the resulting
acetylene-functionalized sample (S-1, Scheme 2) was removed from the reaction chamber, rinsed
several times with DCM and rested for 12 h in a sealed vial at +4 ◦ C under DCM before being further
reacted with alkoxyamine 1.

Scheme 2. Light-assisted (365 nm) hydrosilylation of 1,8-nonadiyne to passivate an hydrogenterminated Si(100) surface (S-1) and covalent attachment of alkoxyamine 1 via CuAAC “click” reactions
to yield an alkoxyamine monolayer (S-2). Anodization of S-2 in the presence of the alcohol nucleophile
2 leads to release of TEMPO in the electrolyte with formation of a redox-active monolayer (S-3) by
reaction of 2 with the putative surface-tethered carbocation intermediate.

2.3.2. Click Immobilization of Alkoxyamine 1 (S-2)
Surface S-1 was reacted with molecule 1 to yield the alkoxyamine monolayers (S-2) via a
copper(I)-catalysed “click” alkyne-azide cycloaddition (CuAAC) reaction. In brief, to a reaction
vial containing the alkyne-functionalized silicon surface (S-1) was added (i) the azide (alkoxyamine
1, 0.5 × 10−3 M, 2-propanol/water, 1:1, v/v), (ii) copper(II) sulphate pentahydrate (1.0 × 10−4 M)
and (iii) sodium ascorbate (5 mg/mL). The reaction was carried out without excluding air from the
reaction environment, at room temperature and under ambient light. The samples were removed from

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the reaction vessel after a reaction time of 2 h and were rinsed thoroughly with copious amounts of
2-propanol, water, 2-propanol and DCM and blown dry with argon before being analysed or further
reacted (Scheme 2).
2.4. Surface Characterization
2.4.1. X-ray Photoelectron Spectroscopy
X-ray photoelectron spectroscopy (XPS) characterization was performed on an ESCALab 250 Xi
(Thermo Scientific, Waltham, MA, USA) spectrometer with a monochromated Al Kα source to
characterize the formation of an alkoxyamine 1 monolayer on silicon. The pressure in the analysis
chamber during measurement was <10−8 mbar. The pass energy and step size for narrow scans were
20 eV and 0.1 eV respectively, with a take-off angle normal to the sample surface. Spectral analysis
was performed using Avantage 4.73 software and curve fitting was carried out with a mixture of
Gaussian–Lorentzian functions after background subtraction. Peaks were calibrated to C–C at 284.8 eV.
2.4.2. Electrochemical Measurements
Electrochemical experiments were performed in a single-compartment, three-electrode PTFE cell
with the modified silicon surface as the working electrode, a platinum mesh (ca. 200 mm2 ) as the
counter electrode and silver/silver chloride “leakless” electrode in 3 M sodium chloride as the reference
electrode. A rectilinear cross-section gasket defined the geometric area of the working electrode to
28 mm2 . Electrical contact between the silicon substrate and a copper plate was ensured by rapidly
rubbing a gallium indium eutectic onto the back of the silicon electrode. Cyclic voltammetry (CV)
measurements were performed using a CHI 910B electrochemical analyser. All potentials are reported
versus the reference electrode. Tetrabutylammonium hexafluorophosphate (Bu4 NPF6 , Sigma-Aldrich,
≥98%) was recrystallized twice from 2-propanol and used at the concentration of 1.0 × 10−1 M as the
supporting electrolyte in DCM solutions. The reaction between the anodically-cleaved S-2 samples and
ferrocenemethanol (2), as well as control experiments were 2 is replaced by ferrocene, were performed
under argon and at room temperature. S-3 samples (Scheme 2) were thoroughly washed with DCM
prior to analysis. Cyclic voltammetry of S-3 samples was performed under ambient conditions in
aqueous 1.0 M HClO4 electrolytes.
3. Results and Discussion
As a proof-of-principle, we have devised “trapping” experiments on an electrode surface modified
by alkoxyamine 1 (S-2, Scheme 2). The anodization of the S-2 monolayers is expected to generate
benzylic cations tethered on the electrode surface while the nitroxide fragment is released into the
electrolyte. The covalent attachment of alkoxyamine 1 to the silicon surface was first verified by the
XPS (Figure 1). Peaks at 99.3 and 100.0 eV for S-2 samples were respectively ascribed to the Si 2p3/2
and Si 2p1/2 emissions (Figure 1b). The presence of a band at 101.1-104.3 eV is due to SiOx formed
during the hydrosilylation and/or CuAAC processes. Two signals in the carbon envelope that are
centred at 284.6 and 286.3 eV indicates the formation of C-C and C-O/-N, respectively [27], as expected
for the putative structure of S-2 samples (Figure 1c). As shown in Figure 1d, refinement of the N
1s emission gives two fitted peaks with binding energies of 400.5 and 401.7 eV in an approximate
3:1 ratio. The position of the low binding energy line in the S-2 samples is in agreement with the
400.6 eV previously assigned to nitroxide nitrogen atoms in thin films [22]. Electrons from the triazole
heterocycle also contributed to the 400.5 and 401.7 emissions with an expected 2:1 ratio.

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Surfaces 2018, 1

b)
wide

Si 2p

Intensity / a.u.

Intensity / a.u.

a)

1200

900

600

300

108

0

104

100

96

92

Binding Energy / eV

Binding Energy / eV
d)

c)

N 1s

Intensity / a.u.

Intensity / a.u.

C 1s

292

288

284

280

276

408

Binding Energy / eV

404

400

396

392

Binding Energy / eV

Figure 1. XPS survey spectra (a) and high-resolution Si 2p (b), C 1s (c) and N 1s (d) scans of S-2 samples
assembled on high-doped p-type Si(100) surface by CuAAC reactions between alkoxyamine 1 and
alkyne-terminated monolayers (S-1).

As shown Figure 2, the anodization of the alkoxyamine monolayer (S-2) via a potentiostatic
experiment (electrode held at 1.0 V versus reference for 30 s) in 1.0 × 10−1 M DCM/Bu4 NPF6 in
the presence of an excess of 2 is effective in trapping the alcohol at the distal end of the monolayer.
The newly-formed redox-active surface (S-3) is washed with copious DCM and then analysed by
cyclic voltammetry in an electrolyte that does not contain a redox-active species. Figure 2a shows
the well-defined pair of redox waves (solid red lines) which are ascribed to the surface-confined
redox tag (S-3). The surface coverage of ferrocenyl units, derived from integration of the anodic
wave, ranges over almost on order of magnitude from 5.5 × 10−12 mol cm−2 , as the low end, up to
the value of 3.1 × 10−11 mol cm−2 for our highest current yield. Our best result is still only about
10% of close-packed packed layer of redox ferrocenes; defining the exact nature of any competitive
side-reactions is beyond the scope of this work. Two control experiments were carried out as follows.
We used ferrocene, which lacks the hydroxyl function, to rule out non-specific adsorption of the redox
probe as well as trying to remove the anodic step. That is, the working electrode modified bearing the
alkoxyamine 1 is rested at open circuit potential (OCP) conditions in the presence of the nucleophile
(2). In both controls (Figure 2b) we found no evidence of redox peaks when either the nucleophile is
removed from the system (+, Figure 2b) or when the S-2 surface is not anodized to trigger the switch
masked alkoxyamine/reactive carbocation (, Figure 2b).

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Surfaces 2018, 1

Figure 2. Cyclic voltammograms (CVs) of redox nucleophiles reacted with the anodized alkoxyamines
monolayer (conversion of S-2 into S-3). CVs were acquired at a voltage sweep rate of 100 mV s−1 .
(a) CVs analysis of the Si(100) electrode in 1.0 M HClO4 after the anodization of S-2 monolayers in the
presence of ferrocene methanol (2) with the surface coverage of ferrocene units ranging from 5.5 ×
10−12 mol cm−2 , as the low end, up to the value of 3.1 × 10−11 mol cm−2 . (b) Control experiments for
prolonged standing of S-2 electrodes at open circuit potential in the presence of 2 in the electrolyte ()
or CVs for the experiments as in (a) but with ferrocene replacing 2 (+).

4. Conclusions
We have demonstrated a method to achieve a switchable interface—from unreactive
to electrophilic—and used it to attach alcohol species as monolayers on Si(100) electrodes.
Unlike conventional esterification approaches on surfaces (e.g., DCC/DMAP), here we show an
electrochemical alternative that involves a surface carbocation. This reactive intermediate is generated
in situ under an external potential stimulus that switches a stable alkoxyamine into a reactive surface
trap for nucleophiles. This proof-of-principle expands to the realm of surfaces some of the recent
advances made in synthetic organic electrochemistry. This knowledge may aid the development of
cationic polymerization on surfaces and this surface chemistry platform can in principle be coupled
to photo-electrochemical reactions on semiconductors and photoconductors [26] in order to control
electron transfer, hence chemical reactivity, in two dimensions [28,29].
Author Contributions: Conceptualization, S.C.; Data curation, L.Z. and S.C.; Formal analysis, L.Z., B.B.N., M.L.C.
and S.C.; Funding acquisition, G.G.W., N.D., M.L.C. and S.C.; Investigation, L.Z., R.B.D.E., V.R.G. and S.C.;
Methodology, S.C.; Project administration, S.C.; Resources, G.G.W., M.L.C. and S.C.; Supervision, G.G.W. and S.C.;
Validation, R.B.D.E.; Visualization, L.Z.; Writing—original draft, L.Z.; Writing—review & editing, V.R.G., N.D.,
M.L.C. and S.C.
Funding: This work was supported by grants from the Australian Research Council (DE160100732 (S.C.),
DE160101101 (N.D.), CE 140100012 (L.Z., B.B.N., M.L.C. and G.G.W.)).
Acknowledgments: We acknowledge support from Curtin University’s Microscopy & Microanalysis Facility,
whose instrumentation has been partially funded by the University, State and Commonwealth. L.Z. would like to
thank AINSE Ltd. for providing financial assistance. M.L.C. gratefully acknowledges an ARC Georgina Sweet
ARC Laureate Fellowship.
Conflicts of Interest: The authors declare no conflict of interest.

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Review

Time-Resolved X-ray Absorption Spectroscopy in
(Photo)Electrochemistry
Martina Fracchia 1 , Paolo Ghigna 1 , Alberto Vertova 2 , Sandra Rondinini 2 and
Alessandro Minguzzi 2, *
1
2

*

Dipartimento di Chimica, Università degli Studi di Pavia, Viale Taramelli 13, 27100 Pavia, Italy;
martina.fracchia02@universitadipavia.it (M.F.); paolo.ghigna@unipv.it (P.G.)
Laboratory of Applied Electrochemistry, Dipartimento di Chimica, Università degli Studi di Milano,
Via Golgi 19, 20133 Milano, Italy; alberto.vertova@unimi.it (A.V.); sandra.rondinini@unimi.it (S.R.)
Correspondence: alessandro.minguzzi@unimi.it; Tel.: +39-025-0314-224

Received: 7 November 2018; Accepted: 1 December 2018; Published: 5 December 2018

Abstract: This minireview aims at providing a complete survey concerning the use of X-ray
absorption spectroscopy (XAS) for time-resolved studies of electrochemical and photoelectrochemical
phenomena. We will see that time resolution can range from the femto-picosecond to the second
(or more) scale and that this joins the valuable throughput typical of XAS, which allows for
determining the oxidation state of the investigated element, together with its local structure. We will
analyze four different techniques that use different approaches to exploit the in real time capabilities
of XAS. These are quick-XAS, energy dispersive XAS, pump & probe XAS and fixed-energy X-ray
absorption voltammetry. In the conclusions, we will analyze possible future perspectives for
these techniques.
Keywords: X-ray absorption spectroscopy; energy dispersive; quick-XAS; FEXRAV; free electron
laser; electrochemistry; photoelectrochemistry; photochemistry; pump & probe

1. Introduction
The research of efficient materials in chemistry and physics requires new criteria for their
rational design and thus the definition of new paradigms for the elucidation of structureperformance guidelines.
In catalysis, the investigation on reaction mechanisms and on the concomitant changes in the
catalyst nature becomes therefore crucial. Thus, the quest for analytical tools capable of directly
detecting any information on the reactant-intermediate-product sequence and on the catalyst role in
the course of the reaction is imperative.
In this respect, operando spectroscopies are obviously central, since they can assume the role of
an independent, uncoupled source of information with respect to the primary ones (such as products
yield, selectivity, or potential/current characteristics in the case of electrochemistry).
In the recent past, consequently, a flourishing use of operando spectroscopies in the study of
electrochemical and photoelectrochemical phenomena has emerged, particularly within the framework
of energy and environmental applications [1,2].
Among the different spectroscopic techniques, X-ray absorption (XAS) is particularly attractive
being able to provides a particularly complete picture of the system of interest. In fact, and despite
the need of large-scale synchrotron radiation facilities, it provides fundamental information on the
local electronic and atomic structure of a selected element. Indeed, the latter characteristic is one of
the most intriguing for XAS: tuning the energy of the incident X-ray beam, it is possible to select the
element under consideration. This is vital for operando experiments, since the spectroscopic response
is independent on any element in the beam different from that under investigation.
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Another enabling feature of XAS is the high energy of X-rays: especially in the so-called “hard”
X-ray regime (roughly, with energies higher than 8–10 keV), the absorption by all light elements
(e.g., C, H, O, N) is negligible. Thus, most of the electrolytes adopted in electrochemistry as well as
selected cell materials (especially if polymer-based) are transparent, and therefore operando studies
are quite undemanding.
The principles of X-ray absorption spectroscopy can be briefly described as follows. When an
X-ray photon with proper energy is shined over a sample, it can excite a core electron. Core electrons
are in localized states, the energy of which is almost independent of the chemical environment. This is
the reason for which XAS is element selective. At energy corresponding to the binding energy of the
core level, the absorption coefficient has an abrupt change, called absorption edge. The excited electron
can be promoted either to unoccupied bound levels or to continuum unbound levels, as pictorially
described in Figure 1.

Figure 1. Scheme of the excitation and de-excitation processes involved in X-ray absorption
spectroscopy (XAS). An X-ray photon (hνin ) can promote a core electron either to bound or continuum
empty states. Transitions to bound states result into spectral features in the X-ray Absorption Near
Edge Structure (XANES) region, while transitions to continuum states give rise to the Extended X-ray
Absorption Fine Structure (EXAFS) oscillations. Radiative de-excitation processes result into X-rays
fluorescence lines (hνout ). Core levels in the figure are named according to the Sommerfeld notation.

In this case, it can be back-diffused (scattered) by surrounding atoms. For a given photon
energy, the interaction between the emitted electron and back-scattered waves can be destructive or
constructive depending on the interatomic distance. This effect leads to an oscillation of the absorption
coefficient for X-ray energy, extending to several hundreds of eV above the absorption edge. This is
the so-called Extended X-ray Absorption Fine Structure (EXAFS), and allows for determining the local
structure (number, nature and distance of surrounding atoms) around the photoabsorber. At energies
close to the edge, the absorption coefficient has very strong modulations. This is the X-ray Absorption
Near Edge Structure (XANES), that is due to electronic transitions to bound unoccupied states, and/or
by multiple scattering events. XANES gives important information on the coordination environment
(type of coordination and its distortions). Finally, the exact energy at which the absorption edge falls
finely depends on the electron density on the photoabsorber, and thus reflects the oxidation state.
Therefore, XAS gives an independent determination of the oxidation state of atoms in compounds.
Time resolved spectroscopies are of fundamental importance in studying the kinetics and
dynamics of chemical reactions. With the introduction of third generation synchrotrons and, more
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recently, of free-electron lasers, time-revolved XAS became accessible thanks to the extremely high
photon flux that can be impinged on the sample.
Time-resolved XAS include different techniques that allow spanning between picoseconds/
femtoseconds and seconds/hours timescales, thus permitting the study of a wide range of phenomena,
from the formation of photogenerated charges to charge transfers, from reaction mechanisms to slowly
activity-losing systems.
In this review, we aim at summarizing the most important contributions achieved so far in the
use of time-resolved XAS in electrochemistry and photoelectrochemistry.
For a more comprehensive, extended and didactic source of information, the Reader is invited to
consider a recent book chapter [3] and earlier reviews [4–6].
Here we recall the fact that XAS is a bulk technique and averages on all atoms in the beam.
However, a proper sample choice allows being sensitive to surface atoms: the most adopted strategies
are (i) the study of small nanoparticles, where the surface/bulk atom ratio is particularly high and
(ii) the use of “bulk active” materials. A classic example of the last case are hydrous oxides where,
up to a certain load, all atoms (external surface and bulk) are addressable by charged and uncharged
species from/to the electrolyte, thus being in the condition to act as active sites [7].
For what concerns the different strategies that can be adopted for time resolved experiment, we
need to consider all variables that are into play: X-ray energy, X-ray absorption coefficient, applied
potential (or scan rate, if in coupling with electrochemical methods), presence/absence of illumination
and, obviously, time.
The most obvious approach is to record full XAS spectra at different times. This leads us to
discuss: quick-XAS, energy dispersive XAS and pump & probe XAS.
2. Quick-XAS
Usually, an XAS spectrum is recorded by stepping the monochromator and counting each point
for a relatively long time, typically in the order of seconds or several seconds. In the quick-XAS
approach, the monochromator is continuously moved over an energy interval and the absorption
coefficient, μ, is continuously measured. In this way, a typical XANES spectrum can be recorded in few
seconds or even in less than a second [8–11]. One of the very first examples of fast acquisition using a
quick-XAS approach is reported in Figure 2.

Figure 2. Cu-K edge spectra of a metal foil at room temperature measured in 72 s using a “conventional”
step-by-step scan (upper spectrum) and in 6 s using the quick-XAS approach (lower spectrum.
Reprinted from [9], with the permission of AIP Publishing.

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Quick XAS has been largely employed in electrochemistry for dynamic studies in energy storage
systems. For example, Yu et al. monitored the electrochemical delithiation of LiFePO4 , a well-known
cathode material in lithium batteries, by exploiting the capability of XANES to easily distinguish
between the LiFePO4 and the FePO4 phases [12]. They found that the delithiation proceeds with a
“two-phase reaction”, in agreement with other works in the literature. In addition, assuming that the
kinetic is limited by lithium diffusion, the authors could estimate an apparent Li+ diffusion coefficient
of 1.9 × 10−13 cm2 ·s−1 .
Ishiguro et al. monitored surface events occurring on a Pt3 Co/C cathode catalyst in proton
exchange membrane fuel cell (PEMFC), with a time resolution of 500 ms [13]. With quick XAS
(acquired both at the Pt LIII - and the Co K-edges) they could achieve a complete understanding of the
structural and electronic changes occurring in this material during the fuel-cell operating processes.
In addition, they compared the reaction mechanism of Pt3 Co/C with that of Pt/C catalyst. They found
that while the two catalysts have similar reactions mechanisms, large increases in the rate constants
are observed when Co is added to Pt.
Quick XAS has been employed in the study of electrocatalysts for various reactions, like Hydrogen
Evolution Reaction (HER), Oxygen Evolution Reaction (OER) or Oxygen Reduction Reaction (ORR).
The aim is to assess the effective chemical nature of an electrocatalyst under working conditions.
An example is the work by Gorlin Y. et al., where a bifunctional manganese oxide (MnOx ) was studied
as catalyst for both OER and ORR reactions [14]. At a potential value of 0.7 V vs. RHE, where the
ORR takes place, a disordered Mn(II,III)O4 spinel is found. When the potential is increased at 1.8 V,
relevant to the OER, 80% of the film is oxidized to Mn(III,IV) oxide, while the remaining part consists
of Mn(II,III)O4 , thus suggesting that a Mn(III,IV) oxide is the actual responsible for the OER. As quick
XAS cannot reach the ms timescale, the authors could not provide any insight into the dynamics of the
process. However, they were able to prove that the catalyst reached a stationary condition at all the
potentials under investigation.
3. Energy Dispersive XAS
While the quick-XAS approach opens the way to time-resolved XAS, the sub-second timescale is
not accessible, thus hindering the study of fast chemical reaction. In the energy dispersive approach,
an XAS spectrum can be obtained in a timescale of milliseconds, which is quite satisfactory for a large
part of fast chemical phenomena. In a dispersive XAS beamline, a bent monochromator is used to
focus a polychromatic (pink) beam on the sample. The beam then diverges after the sample and the
different wavelengths are spatially separated. A position sensitive detector placed after the sample is
therefore used to acquire the spectrum. Examples of the use of energy dispersive XAS (EDXAS) in
electrochemistry date back to late 1980s and early 1990s [15,16]. However, only after the introduction
of fast and linear detectors [17] it was possible to reach the sub-second timescale. Since then, the
technique has evolved into a powerful tool to investigate fast chemical phenomena [18–20]. From what
concerns electrochemistry, EDXAS was employed to monitor the dynamics of oxidation and reduction
on electrodes for different applications. In a seminal experiment in 1989, EDXAS was coupled to cyclic
voltammetry to investigate the XANES at the Ni K-edge in α-Ni(OH)2 and β-Ni(OH)2 in KOH [16].
A continuous shift to higher values of the edge energy position was observed in concomitance with
the oxidation from Ni(II) to Ni(III) and, at the same time, a change in the spectral profile related to a
distortion of the octahedral coordination of the six oxygen atoms surrounding Ni was detected.
Rose et al. investigated the kinetics of the electrochemical formation of palladium β-hydride
and β-deuteride in carbon-supported Pd catalyst nanoparticles [21]. The rate determining step of the
formation of β-hydride was found to be the diffusion of H through the bulk of the Pd particles; in case
of β-deuteride, the authors observed that the reaction was limited by the interfacial reduction of D+ .
In addition, they were able to determine in both cases the stoichiometry of the two phases.
It has to be noted that EDXAS can also be employed to study time-resolved processes occurring
in solution, particularly in the immediate proximity of the electrode. As an example, O’ Malley et al.

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considered the electrodesorption of copper on a platinum electrode in order to study the diffusion of
Cu2+ ions as a function of time and distance from the electrode surface [22].
Great effort has been devoted in the recent years to investigate the electrochemical behavior of
Pt-based electrodes. Allen et al. investigated both the oxidation and reduction of dispersed Pt catalyst,
studying in real time the modifications of the Pt oxidation state and of the number of O and Pt nearest
neighbors [15]. Moreover, they found evidences for a different mechanism of oxidation/reduction
in case of clusters with respect to bulk electrodes. In another paper, the same authors describe the
use of EDXAS to record full XAS spectra while the electrode potential follows a cyclic voltammetric
profile [23]. This approach allowed for determining structural parameters (such as the average
coordination number of Pt and O atoms around Pt) as a function of the applied potential. This, in turn,
allowed for distinguishing between oxygen adsorbing in the double-layer region from the Pt oxidation
that is typically observed in CV at higher potentials.
The electrochemical oxidation of the surfaces of Pt nanoparticles was further investigated by
Imai et al. with a sub-second time resolution [24]. As shown in Figure 3, after applying a potential of
1.4 V vs. RHE in acidic media, they exploited the high sensitivity of the EXAFS region towards the
local Pt-Pt and Pt-O distances, to monitor their change in time.

Figure 3. Example of outcome from an operando energy dispersive-XAS experiments on Pt
nanoparticles. Bond lengths and coordination numbers as a function of time for (a) Pt-Pt and (b) Pt-O
bonds during potential step at 1.4 V. In (a), it is possible to observe Pt-Pt bonds with 2.7 Å (metallic
platinum) plus two types of longer Pt-Pt bonds (Pt-Pt bonds in platinum oxides, 3.1 and 3.5 Å), the latter
appearing at about 40 and 100–120 s, respectively. In (b), longer Pt-O bonds (due adsorbed oxygen
species) turn into shorter Pt-O (Pt-O bonds in oxides). Reprinted with permission [24]. Copyright 2009
American Chemical Society.

The authors found that longer Pt-O bonds are initially formed, indicating the formation of Pt-OHH
and/or Pt-OH species; afterwards, these species turn to shorter Pt-O bonds, and the formation of
α-PtO2 was detected. Finally, after 100 s, a precursor of β-PtO2 is formed.
EDAXS was also used to study the oxidation of Pt nanoparticles in presence of competing halide
ions (Cl− and Br− ) [25]. The kinetic parameters obtained by the time resolved experiment confirmed
the two step model proposed by Zolfaghari, Conway and Jerkiewicz [26]. As evident from Figure 4,
the competing effect of the halides was found to be significant mainly on the first step. In addition,
the partial discrepancy, particularly at short times, between the integrated quantity of charge and the
occupancy of Pt 5d states obtained by XAS suggests a non-negligible role of anion redox activity.

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Figure 4. IR corrected results for operando ED-XAS in the case of Pt nanoparticles oxidation (and the
relevant following reduction) in aqueous in 0.1 M HClO4 (black), 0.1 M HClO4 + 10 mM KCl (red) and
0.1 M HClO4 + 10 mM KBr (orange). (a) Q/Qf (full lines) and Degree of reaction (DoR, by XAS) (dots)
and (b) (Q/Qf )−1 (full lines) and DoRXAS−1 (dots) as a function of log(t) for a 0.5 e 1.4 V (reversible
hydrogen electrode) step. Reprinted from [25]. Copyright 2018, with permission from Elsevier.

The oxidation and reduction of highly hydrated IrOx , one of the most active electrocatalysts for
OER, was recently investigated by EDXAS [27]. The complementary information between XAS and
chronoamperometry evidenced parasitic and side reactions. In addition, the role and time evolution of
hydration degree and persistence of iridium in different charge states was discussed.
Finally, a very recent paper discussed the stability of the copper(II) lactate complex in alkaline
solution, used as a precursor for the electrodeposition of Cu2 O, finding that the complex is stable in
a wide range of applied potentials [28]. The authors were also able to determine the structure of the
complex in solution by fitting the XANES data.
4. Pump & Probe XAS
This method can be applied in photochemistry or in photoelectrochemistry. It consists in exciting
the system with photons of desired energy (“pump”, typically in the visible or in the UV range) before
probing the system with X-ray photons. The X-ray beam from a synchrotron radiation source has a

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characteristic time pattern that reflects the pattern of the electron bunches in the storage ring. The
sub-μs timescale is then achievable with this approach. The probe is delayed with respect to the pump
in dependence on the timescale of the phenomena under investigation. This is evident in the study of
photochemistry, where pump & probe XAS was initially used instead of time-resolved X-ray diffraction
to have structural information on disordered or dissolved species [29]. The technique was recently
used to probe the charge carrier dynamics in semiconductors used as photocatalysts, such as TiO2 and
WO3 , as discussed in the following.
In the case of TiO2 , thanks to picosecond XAS, Rittmann and co-workers performed an experiment
at the Ti K-edge and the Ru LIII -edge to study the lifetime of photogenerated trapped electrons for
photoexcited bare and N719-dye-sensitized anatase and amorphous TiO2 nanoparticles [30]. The
outcomes point to the existence of different types of defects (either at the surface or in the bulk) in
dependence on the type of particles and on the presence/absence of the dye. In case of dye-sensitized
anatase or amorphous TiO2 they could observe the formation of Ti3+ defects in trapping sites that are
pentacoordinated, i.e. defective, meaning that the trapping center is localized at the outer surface of
the crystallite. The trap lifetime is in the order of nanoseconds. This is a typical result possible only
with XAS, which, being element-specific, allows for obtaining a direct information on the electronic
and local structure of Ti.
Pump & probe XAS was also used in photochemistry and photoelectrochemistry to investigate
different metal complexes, especially in the fields of energy storage and conversion. Many literature
works concern the usage of Cobalt-based complexes as water oxidation or water reduction catalysts
for artificial photosynthesis. In a work by Song et al. two cubane-like catalysts for water oxidation
were investigated through in situ XANES under photocatalytic conditions: it was observed that the
{Co(II)4 O4 } core undergoes a rapid oxidation from Co(II) to Co(III) or higher valent states, and then
Co(II) is slowly restored [31]. Moonshiram and co-workers investigated through transient XAS the
electronic and structural dynamics of [Ru(bpy)3 ]2+ /[LCo(III)Cl2 ]+ hybrid systems used as catalysts for
water reduction in the photosplitting of water [32]. When the chromophore is excited, the cobalt (III)
octahedral complex is reduced to Co(II). In addition, they repeated the experiment in the presence of
sodium ascorbate/ascorbic acid electron donor, finding the experimental proof of the formation of a
Co(I) square planar species, followed by the formation of the octahedral Co(III) starting complex with
two aquo-ligands.
Canton et al. employed ultrafast optical and X-ray techniques (coupled to DFT calculations) to
identify the light-induced electron transfer and the associated structural and spin changes occurring in
two photoexcited heterobimetallic ruthenium-cobalt complexes [33].
The quest for higher temporal resolution has been boosted in the very recent year by the
development of free-electron lasers (FEL) [34–36]. In FEL facilities, ultrashort pulses of X-rays
paved the way for spectroscopic experiment in the femtosecond timescale [37–39]. This allowed for
investigating the local electronic and structural modifications occurring in a transition metal compound
immediately after the photoexcitation. For example, Bressler et al. used femtosecond XANES to solve
a long-standing issue about the population mechanism of quintet states in iron(II)-based complexes,
which resulted to be a 1 MLCT → 3 MLCT → 5 T cascade from the initially excited state [29].
In the field of semiconducting oxides, Santomauro and co-workers extended the investigation of
Ritmann et al. by performing the XAS experiment at an XFEL facility to probe the dynamics of the
trapping of the photogenerated electrons in TiO2 [40]. They concluded that the electrons are localized
in correspondence of Ti atoms in a time lower than 300 fs, leading to the formation of Ti3+ centers.
In addition, they confirmed that electrons are localized at Ti penta-coordinate sites.
The dynamics of the photoexcitation was also investigated at an XFEL in the case of WO3
nanoparticles. Uemura et al. observed that immediately after the excitation the tungsten is initially
reduced from W(VI) to a mean oxidation state of 5.3 [38]. Subsequently, the tungsten oxide undergoes
a change in its local structure within the following 200 ps. Some of the results of this paper are reported
in Figure 5.

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Figure 5. Top, left: W LIII XANES spectra of WO3 in the ground state together with differential XANES
spectra of WO3 , each being the subtraction of the XANES spectrum of an excited state (at the indicated
time) and the spectrum of a ground state. Right: Absorption intensities of W LIII XANES for peaks A,
B, and C (as shown in the figure at the left) as a function of time. Bottom, left: A proposed scheme for
the photoexcitation process of WO3 . Reprinted from [38]. Copyright 2016, with permission from Wiley.

All the listed cases refer to ex-situ or in-situ (on solvent dispersed particles) experiments, where,
in most cases, the semiconductors or the complexes are photoexcited and the dynamics of the
photogenerated carriers is followed. As a matter of fact, pump & probe XAS, either at synchrotron
or at FEL facilities, is still in its very early days, to the point that studies under operando condition
are very rare, and are mainly obtained by adding a sacrificial donor or acceptor and not by applying
an external potential. In fact, to the authors’ best knowledge, only one contribution was published,
limited to synchrotron radiation pump & probe XAS [41].
In this work, the authors investigated the α-Fe2 O3 /IrOx architecture as a model photoelectrode,
and they could detect a charge transfer between hematite and IrOx occurring in the nanosecond
timescale, resulting in reduction of Ir or in an increased density of empty Ir 5d states depending on
the applied potential. Figure 6 represents schematically the adopted setup (left) together with the
results obtained when an anodic photocurrent is observed. In this case, difference spectra indicate a
hole transfer from the semiconductor to IrOx when no delay between pump & probe is applied. When
pump & probe are delayed for 600 ns, the Ir 5d state results to be depleted, indicating a more intense
hole transfer.

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Figure 6. Left: schematic setup for an operando pump & probe (P & P) experiment on a
photoelectrochemical system using synchrotron light. The UV-Vis pump is synchronized with the
probe pulses, the latter generated by single bunches from the synchrotron. Right: (A) Ir-LIII XANES
spectrum in the presence (red dotted line) and in the absence (black full line) of 400 nm light acquired
in presence of 600 ns delay between the pump and the probe; spectra are shifted along the y axis for the
sake of better clarity. (B) difference spectra (light on-light off, with error bars) in the presence (red line)
and in the absence (blue line) of delay; the two spectra are shifted for clarity, and the zero is defined by
black horizontal lines. Reprinted from [41]. Copyright 2016, with permission from Elsevier.

5. FEXRAV
Finally, an alternative approach is to fix the X-ray energy and record the corresponding absorption
coefficient as a function of time. In this case, in the framework of hyphenated techniques, it would
be convenient to collect joint information from electrochemical methods. This is the principle behind
the fixed energy X-ray absorption voltammetry (FEXRAV). This technique is an application of the
single energy X-ray absorption detection first proposed by Filipponi et al. [42]. It consists in recording
the X-ray absorption coefficient at a fixed energy while the applied potential in an electrochemical
cell is scanned according to a cyclic voltammetry triangular-wave program [43]. If the energy is
properly selected to give the maximum contrast between the spectral features of different oxidation
states, any shift from the original oxidation state determines a variation of the absorption coefficient,
and can then be detected. FEXRAV allows for quickly mapping the oxidation states of the element
under consideration within the selected potential window, and this can be preliminary to deeper X-ray
absorption spectroscopy (XAS) characterizations, like XANES or EXAFS. Moreover, the time-length of
the experiment is much shorter than a series of XAS spectra and opens the door to kinetic analysis.
Initially introduced for the study of IrOx materials, FEXRAV has been applied to operando studies of
the electrochemistry of a variety of materials, from Ag nanoparticles [44], to copper oxide materials [45],
to iron oxide and oxo-hydroxide materials [46], to palladium based electrocatalysts [47]. Figure 7
reports an example of a FEXRAV experiment in the case of Ag nanoparticles studied as electrocatalysts
for the electrodehalogenation of trichloromethane in aqueous media. Here the X-ray energy is set
at a value that guarantees the increase of μ when Ag is oxidized. The potential is swept from 0
to −1.1 V and there’s a visible decrease of μ due to the presence (adsorption) of trichloromethane
or of its reduction intermediates. The release of Cl− causes Ag to oxidize at about 0.2 V. This is
witnessed by XAS, by means of an increase of μ. The latter decreases when the potential is reversed, in
correspondence to the reduction peak at about 0.1 V.

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Figure 7. Example of a fixed energy X-ray absorption voltammetry experiment on Ag nanoparticles
in aqueous of 10 mM trichloromethane (chloroform): CV (black) and FEXRAV (red) at 1 mV·s−1 .
Reprinted from [44]. Copyright 2016, with permission from Elsevier.

6. Conclusions and Perspectives
This minireview aims at providing a comprehensive and updated overview of the state of the art
concerning use and capabilities of time resolved XAS in electrochemistry and photoelectrochemistry.
It is evident that the number of published papers on these topics is not very large, especially if
compared to other types of spectroscopies. This is due to two main factors: (i) the dependence on large
facilities, which limits the number of users that can approach and that have access to these techniques;
(ii) the general unawareness to electrochemists of the mere existence of XAS, notwithstanding its very
high potentialities. Of course, these two factors are strongly related.
Indeed, as we hope to have successfully demonstrated in this work, the potentialities of time
resolved XAS are very high. In fact, time resolved XAS can help to clarify reaction mechanisms,
metastable (transient) structures and to identify the nature of intermediate reacting states, and we
hope that the range of applications will extend progressively in the close future.
Fortunately, we are witnessing a progressive increase of available synchrotron radiation sources
(for a general outlook at the existing facilities, the reader is referred to the website https://lightsources.
org/lightsources-of-the-world/). In parallel, it is worth noting that innovative benchtop laboratory
X-ray spectrophotometers are commercially available. Even if the photon flux of these instruments is
still not comparable with that available at a synchrotron beamline, we believe that technical advances
will allow for rapid improvements which, together with an increasing knowledge on the correct design
of samples and cells [48], will likely make XAS available also in academics and companies worldwide.
Funding: A.M. is thankful to Università degli Studi di Milano through the “Piano di Sostegno alla Ricerca
2015/2017”.
Conflicts of Interest: The authors declare no conflict of interest.

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© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access
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