crc press - proton exchange membrane fuel cells - materials properties and performance - 2010

Proton ExchangeMembrane Fuel Cells

Materials Propertiesand Performance

GREEN CHEMISTRY AND CHEMISTRY ENGINEERINGSeries Editor: Sunggyu Lee

Missouri University of Science and Technology, Rolla, USA

Proton Exchange Membrane FuelCells: Materials Properties and Performance

David P. Wilkinson, Jiujun Zhang, Rob Hui, Jeffrey Fergus, and Xianguo Li

Solid Oxide Fuel Cells: Materials Properties and Performance Jeffrey Fergus, Rob Hui, Xianguo Li, David P. Wilkinson, and Jiujun Zhang

CRC Press is an imprint of theTaylor & Francis Group, an informa business

Boca Raton London New York

Proton ExchangeMembrane Fuel Cells

Materials Propertiesand Performance

Edited by

David P. WilkinsonJiujun Zhang

Rob HuiJeffrey Fergus

Xianguo Li

CRC PressTaylor & Francis Group6000 Broken Sound Parkway NW, Suite 300Boca Raton, FL 33487-2742

© 2010 by Taylor and Francis Group, LLCCRC Press is an imprint of Taylor & Francis Group, an Informa business

No claim to original U.S. Government works

Printed in the United States of America on acid-free paper10 9 8 7 6 5 4 3 2 1

International Standard Book Number: 978-1-4398-0664-7 (Hardback)

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Library of Congress Cataloging-in-Publication Data

Proton exchange membrane fuel cells : materials properties and performance / editors, David P. Wilkinson ... [et al.].

p. cm. -- (Green Chemistry and Chemistry Engineering)Includes bibliographical references and index.ISBN 978-1-4398-0664-7 (hardcover : alk. paper)1. Proton exchange membrane fuel cells. I. Wilkinson, David P. II. Title. III. Series.

TK2931.P785 2010621.31’2429--dc22 2009036642

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v

Contents

Series Preface ........................................................................................................ viiPreface ......................................................................................................................ixEditors ......................................................................................................................xiContributors ......................................................................................................... xiii

1. Recent Developments in Electrocatalyst Activity and Stability for Proton Exchange Membrane Fuel Cells ...............................................1David Thompsett

2. Catalyst Layers and Fabrication ................................................................. 61Zhong Xie, Chaojie Song, David P. Wilkinson, and Jiujun Zhang

3. Proton Exchange Membranes .................................................................. 107Timothy J. Peckham, Yunsong Yang, and Steven Holdcroft

4. Diffusion Layers ......................................................................................... 191Mauricio Blanco and David P. Wilkinson

5. Bipolar Plates and Plate Materials ..........................................................305Tim Cheng

6. Physical Modeling of Materials for PEFCs: A Balancing Act of Water and Complex Morphologies .....................................................343Michael H. Eikerling and Kourosh Malek

Index .....................................................................................................................435

vii

Series Preface

The subjects and disciplines of chemistry and chemical engineering have encountered a new landmark in the way of thinking about, developing, and designing chemical products and processes. This revolutionary phi-losophy, termed green chemistry and chemical engineering, focuses on the designs of products and processes conducive to reducing or eliminating the use and/or generation of hazardous substances. In dealing with hazard-ous or potentially hazardous substances, there may be some overlaps and interrelationships between environmental chemistry and green chemistry. Environmental chemistry is the chemistry of the natural environment and the pollutant chemicals in nature; however, green chemistry proactively aims to reduce and prevent pollution at its very source.

In essence, the philosophies of green chemistry and chemical engineer-ing tend to focus more on industrial application and practice rather than academic principles and phenomenological science. However, as both a chemistry and chemical engineering philosophy, green chemistry and chemical engineering derive from and build upon organic chemistry, inorganic chemistry, polymer chemistry, fuel chemistry, biochemistry, analytical chemistry, physical chemistry, environmental chemistry, thermo-dynamics, chemical reaction engineering, transport phenomena, chemical process design, separation technology, automatic process control, and more. In short, green chemistry and chemical engineering are the rigorous use of chemistry and chemical engineering for pollution prevention and environ-mental protection.

The Pollution Prevention Act of 1990 in the United States established a national policy to prevent or reduce pollution at its source whenever feasible. Adhering to the spirit of this policy, the Environmental Protection Agency (EPA) launched its Green Chemistry Program in order to promote innova-tive chemical technologies that reduce or eliminate the use or generation of hazardous substances in the design, manufacture, and use of chemical products. The global efforts in green chemistry and chemical engineering have recently gained a substantial amount of support from the international community of science, engineering, academia, industry, and governments in all phases and aspects.

viii

Some successful examples and key technological developments include the use of supercritical carbon dioxide as green solvent in separation technolo-gies, application of supercritical water oxidation for destruction of harm-ful substances, process integration with carbon dioxide sequestration steps, solvent-free synthesis of chemicals and polymeric materials, exploitation of biologically degradable materials, use of aqueous hydrogen peroxide for effi-cient oxidation, development of hydrogen proton exchange membrane (PEM) fuel cells for a variety of power generation needs, advanced biofuel produc-tion, devulcanization of spent tire rubber, avoidance of the use of chemicals and processes causing generation of volatile organic compounds (VOCs), replacement of traditional petrochemical processes by microorganism-based bioengineering processes, replacement of chlorofluorocarbons (CFCs) with nonhazardous alternatives, advances in design of energy-efficient processes, use of clean, alternative and renewable energy sources in manufacturing, and much more. This list is only a partial compilation, but, undoubtedly, is grow-ing exponentially.

This book series on green chemistry and chemical engineering by CRC Press/Taylor & Francis is designed to meet the new challenges of the twenty-first cen-tury in the chemistry and chemical engineering disciplines by publishing books and monographs based upon cutting-edge research and development to the effect of reducing adverse impacts upon the environment of chemical enterprise. In achieving this, the series will detail the development of alternative sustain-able technologies, which will minimize the hazard and maximize the efficiency of any chemical choice. The series aims to deliver an authoritative information source in the field of green chemistry and chemical engineering to readers in academia and industry. The publisher and its series editor are fully aware of the rapidly evolving nature of the subject and its long-lasting impact upon the quality of human life in both the present and the future. As such, the team is committed to making this series the most comprehensive and accurate literary source in the field of green chemistry and chemical engineering.

Sunggyu Lee

ix

Preface

A fuel cell is an electrochemical device that provides efficient conversion of the chemical energy of fuels directly into electricity for power generation. Fuel cells are expected to play a significant role in the strategy to affect posi-tive global change, increase fuel efficiency, and decrease dependency on tra-ditional fossil fuels. Fuel cells and direct electrochemical fuels (particularly hydrogen) provide the promise of being one of the long-term solutions to the improvement of energy efficiency, energy sustainability, and energy secu-rity and the reduction of greenhouse gases and urban pollution. Significant environmental benefits are expected for fuel cells, particularly for energy conversion for transportation and electric power generation.

The polymer electrolyte fuel cell (PEFC) or proton exchange membrane fuel cell—also known as the polymer electrolyte membrane fuel cell (PEMFC)—is a lower temperature fuel cell (typically less than 100°C) with a special poly-mer electrolyte membrane. This lower temperature fuel cell is well suited for transportation, portable, and micro fuel cell applications because of the importance of fast start-up and dynamic operation. The PEMFC has applica-bility in most market and application areas.

Technical progress as well as investments in PEMFCs for transportation, stationary, portable, and micro fuel cell applications has been dramatic in recent years. The present view is optimistic for fuel cell power generation; the status is presently at the field trial level, or early commercialization stage, moving into volume commercialization. Although commercially viable, niche PEMFC applications exist today, the first commercial mass markets for fuel cells are expected to be for handheld electronic devices, PCs, and other portable devices.

However, the PEMFC will need to be competitive on an economic and con-sumer basis with the established and highly developed internal combustion engine and other forms of power generation. Even though much progress has been made with the PEMFC, significant technical challenges still remain today in a number of areas, including reliability, durability, cost, operational flexibility, technology simplification and integration, fundamental under-standing, and life cycle impact. Fundamental understanding, new advanced materials, and associated engineering design and modeling will be required to close these technical gaps.

With the continued extensive progress in PEMFC technology and science, there is a need for updated information—particularly in the area of material properties and performance. Given the highly interdisciplinary nature of the fuel cell field, a wide spectrum of relevant scientific, engineering, and tech-nical aspects needs to be covered. This book will provide updated, detailed background material on key developments in the PEMFC area. In particular,

x Preface

the book reviews the progress and current aspects of materials and perfor-mance for PEMFCs.

The book’s chapters are divided between the major components of the unit fuel cell with a strong materials focus; however, they also include design and modeling aspects. The first two chapters focus on catalysts and catalyst layers; the next three chapters on the major components of membranes, dif-fusion layers, and bipolar plates; and the last chapter on materials modeling for the PEMFC. In all cases, it is clear that production of a commercially viable PEMFC will require a compromise of materials with adequate prop-erties, design interaction, and manufacturability. This book will provide a perspective on the status of PEMFC fuel cell technology today, research and development directions, and the scientific and engineering challenges that the fuel cell community faces.

David P. Wilkinson

xi

Editors

David P. Wilkinson received his BASc degree in chemical engineering from the University of British Columbia in 1978 and his PhD degree in chemistry from the University of Ottawa in 1987. In 2004, after 20 years of industrial experience, Dr. Wilkinson was awarded a Tier 1 Canada research chair in clean energy and fuel cells in the Department of Chemical and Biological Engineering at the University of British Columbia. He pres-ently maintains a joint appointment with the university and the Canadian National Research Council Institute for Fuel Cell Innovation.

Prior to this appointment, Dr. Wilkinson was the director, and then vice president of research and development at Ballard Power Systems and involved with the research, development, and application of fuel cell tech-nology for transportation, stationary power, and portable applications. Until 2003, Dr. Wilkinson was the leading all-time fuel cell inventor by number of issued U.S. patents. Dr. Wilkinson’s main research interest is in electrochem-ical power sources and processes to create clean and sustainable energy. He is an active member of the Electrochemical Society, the International Society of Electrochemistry, the Chemical Institute of Canada, and the American Chemical Society.

Jiujun Zhang is a senior research officer and PEM catalysis core competency leader at the National Research Council of Canada Institute for Fuel Cell Innovation (NRC-IFCI). Dr. Zhang received his BS and MSc degrees in elec-trochemistry from Peking University in 1982 and 1985, respectively, and his PhD in electrochemistry from Wuhan University in 1988. Starting in 1990, he carried out three terms of postdoctoral research at the California Institute of Technology, York University, and the University of British Columbia. Dr. Zhang has over 27 years of R&D experience in theoretical and applied elec-trochemistry, including over 13 years of fuel cell R&D (among these, 6 years at Ballard Power Systems and 5 years at NRC-IFCI) and 3 years of electro-chemical sensor experience.

Dr. Zhang holds several adjunct professorships, including one at the University of Waterloo and one at the University of British Columbia. His research is mainly based on fuel cell catalysis development. He has coauthored more than 200 publications, including 150 refereed journal papers and three edited books. He also holds over 10 U.S. patents and patent publications. Dr. Zhang is an active member of the Electrochemical Society, the International Society of Electrochemistry, and the American Chemical Society.

Shiqiang (Rob) Hui is a senior research officer in high-temperature fuel cells at the National Research Council of Canada Institute for Fuel Cell Innovation.

xii Editors

He is an adjunct professor at the University of British Columbia, Canada, and three other major universities in China. Dr. Hui received his PhD in materials science and engineering from McMaster University in 2000. He has conducted research and development in materials, processing, and charac-terization for more than 20 years.

Dr. Hui has worked on various projects, including chemical sensors, solid oxide fuel cells, magnetic materials, gas separation membranes, nanostruc-tured materials, thin film fabrication, and protective coatings for metals. He has more than 80 research publications, one worldwide patent, and one U.S. patent (pending). He is currently leading and involved in several projects for the development of metal-supported solid oxide fuel cells (SOFCs), ceramic nanomaterials as catalyst supports for high-temperature PEM fuel cells, pro-tective ceramic coatings on metallic substrates, ceramic electrode materials for batteries, and ceramic proton conductors. Dr. Hui is also an active mem-ber of the Electrochemical Society and the American Ceramic Society.

Xianguo Li received his BS degree in internal combustion engineering from Tianjin University, Tianjin, China, in 1982, and his MSc in 1986 and PhD in 1989, respectively, in mechanical engineering from Northwestern University, Evanston, Illinois. Dr. Li’s academic career formally began in 1992 when he was appointed an assistant professor in the Department of Mechanical Engineering, University of Victoria. In 1997, he joined the University of Waterloo, and was promoted to the rank of full professor in 2000.

Dr. Li’s current research involves both experimental and theoretical analyses in the areas of fuel cells, green energy systems, liquid atomization, and sprays. He has published extensively, including journal and conference articles, confiden-tial contract reports, and invited seminars and presentations. Dr. Li is also active in the professional community, serving as editor-in-chief for the International Journal of Green Energy and on the editorial board for a number of international journals and an encyclopedia on energy engineering and technology. He has also served as guest editor for a number of international journals.

Jeffrey W. Fergus received his BS degree in metallurgical engineering from the University of Illinois in 1985 and his PhD degree in materials science and engineering from the University of Pennsylvania in 1990. He was a postdoc-toral research associate at the Center for Sensor Materials at the University of Notre Dame and, in 1992, joined the materials engineering program at Auburn University, where he is currently a professor.

His research interests are generally in high-temperature and solid-state chemistry of materials, including electrochemical devices (e.g., chemical sen-sors and fuel cells) and the chemical stability of materials (e.g., high-tempera-ture oxidation). Dr. Fergus is an active member of the Electrochemical Society, the Metals, Minerals and Materials Society, the American Ceramics Society, the Materials Research Society, and the American Society for Engineering Education.

xiii

Contributors

Mauricio BlancoChemical and Biological

Engineering DepartmentandClean Energy Research Center

(CERC)University of British ColumbiaVancouver, British Columbia,

CanadaInstitute for Fuel Cell Innovation

National Research Council of Canada

Vancouver, British Columbia, Canada

Tim ChengAutomotive Fuel Cell Cooperation

(AFCC)Vancouver, British Columbia,

Canada

Michael H. EikerlingDepartment of ChemistrySimon Fraser UniversityBurnaby, British Columbia, CanadaInstitute for Fuel Cell InnovationNational Research Council of

CanadaVancouver, British Columbia,

Canada

Steven HoldcroftDepartment of ChemistrySimon Fraser UniversityVancouver, British Columbia,

Canada

Kourosh MalekInstitute for Fuel Cell InnovationNational Research Council of


Canada

Timothy J. PeckhamDepartment of ChemistrySimon Fraser UniversityVancouver, British Columbia,

Canada

Chaojie SongInstitute for Fuel Cell InnovationNational Research Council of


Canada

David ThompsettJohnson Matthey Technology CenterReading, England

David P. WilkinsonChemical and Biological

Engineering DepartmentandClean Energy Research Center

(CERC)University of British ColumbiaVancouver, British Columbia,

CanadaInstitute for Fuel Cell InnovationNational Research Council of


Canada

xiv Contributors

Zhong XieInstitute for Fuel Cell InnovationNational Research Council of


Canada

Yunsong YangBallard Power SystemsVancouver, British Columbia,

Canada

Jiujun ZhangInstitute for Fuel Cell InnovationNational Research Council of


Canada

1

1Recent Developments in Electrocatalyst Activity and Stability for Proton Exchange Membrane Fuel Cells

David Thompsett

CONTENTS

1.1 Introduction ....................................................................................................21.2 The State of the Art ........................................................................................21.3 Application Targets ........................................................................................41.4 Electrocatalyst Discovery .............................................................................5

1.4.1 High-Throughput Screening ............................................................51.4.2 Computation Studies .........................................................................7

1.5 Electrocatalyst Preparation ...........................................................................91.5.1 Conventional Routes .........................................................................91.5.2 Colloidal Routes ............................................................................... 101.5.3 Molecular Precursor Routes ........................................................... 111.5.4 Vapor Phase Routes ......................................................................... 121.5.5 Surface Modification Routes .......................................................... 12

1.6 Electrocatalyst Testing ................................................................................ 131.7 Enhanced Activity Cathode Catalysts ...................................................... 14

1.7.1 Pt Alloy Catalysts ............................................................................. 141.7.1.1 Model Surface Studies ...................................................... 151.7.1.2 Pt Alloy Nanoparticles and Particle Size Effects .......... 16

1.7.2 Pt Core-Shell Catalysts .................................................................... 201.7.2.1 Use of Alternative Promoters to Pt ................................. 24

1.7.3 Non-Pt Catalysts .............................................................................. 241.7.4 Pd-Based Catalysts ...........................................................................25

1.7.4.1 Fe- and Co-Based Materials .............................................251.7.4.2 MeOH-Tolerant Oxygen Reduction Catalysts ............... 27

1.8 Cathode Catalyst Stability .......................................................................... 291.8.1 Pt Electrochemical Area Loss ........................................................ 291.8.2 Stabilization of Pt Catalysts toward Potential Cycling ............... 311.8.3 Effect of High-Cathode Voltages on Catalyst Stability ............... 321.8.4 Stabilization of Pt Catalysts toward High-Voltage

Excursions .........................................................................................34

2 Proton Exchange Membrane Fuel Cells

1.1 Introduction

The aim of this chapter is to review recent developments in electrocatalyst tech-nology for proton exchange membrane fuel cells (PEMFCs) fueled by hydrogen, impure hydrogen (reformate), and methanol. Efforts will be made to summarize catalyst activity and stability targets for the emerging commercial applications, review progress against these targets, and identify remaining challenges. The aim here is not to provide a comprehensive review of recent work, but rather to provide selected examples to illustrate the main developments.

1.2 The State of the Art

The current practical catalyst technology of choice for H2 PEMFCs is Pt nano-particles (typically 2–3 nm) supported on a high surface area carbon black, with Pt loadings typically 40–60 wt% (see Figure 1.1). This type of catalyst is

1.8.5 Alternative Supports .......................................................................351.8.5.1 Oxides .................................................................................351.8.5.2 Carbides and Nitrides ......................................................361.8.5.3 Nonconductive Whiskers .................................................36

1.9 Carbon Support Materials .......................................................................... 371.9.1 Conventional Carbon Blacks .......................................................... 37

1.9.1.1 Modification of Carbon Blacks ........................................381.9.2 Synthetic Carbon Materials ............................................................38

1.9.2.1 Carbon Nanotubes and Nanofibers ...............................381.9.2.2 Synthetic Mesoporous Carbons ...................................... 39

1.10 Reformate-Tolerant Anode Catalysts ........................................................ 411.10.1 Mechanistic Studies .........................................................................421.10.2 Improved Reformate-Tolerant Catalysts .......................................43

1.10.2.1 PtRu Variants .....................................................................431.10.2.2 PtMo Catalysts ...................................................................44

1.11 Reformate-Tolerant Catalyst Stability .......................................................451.11.1 Ru and Mo Stability .........................................................................451.11.2 Cell Reversal Tolerance ...................................................................46

1.12 MeOH Oxidation Catalysts ........................................................................ 471.12.1 Mechanistic Advancements ........................................................... 471.12.2 PtRu Variants ....................................................................................481.12.3 Alternative Pt Catalysts...................................................................50

1.13 MeOH Oxidation Catalyst Stability .......................................................... 52References ...............................................................................................................53

Recent Developments in Electrocatalyst Activity and Stability 3

111

A

B

C

4

3

2

1

[110]

Step Vacancy

A Step

B Step

Kink

11–1

0.10Pt(s)-5(111) × (100)

2 nm

Step

4

3

2

1

B

A

111/100Edge

Pt(s)-4(100) × (111)

111/111Edge

Terrace

IslandB Step

A Step

FIGURE 1.1Restored high-resolution transmission electron micrograph of a 6 nm Pt particle supported on carbon. Also shown: best-fitted simulation and three-dimensional atomic model used to calculate B. (L. Cervera Gotard et al., Angewandte Chemie International Edition (2007), 46, 3683. Copyright Wiley–VCH Verlag GmBH & Co. KGaA. Reproduced with permission.)


used in anode and cathode applications. Catalyst electrode loading studies have shown that Pt loadings on the anode can be reduced to at least 0.05 mg (Pt) cm–2 without significant performance losses.1 Anode loadings as low as 0.017 mg (Pt) cm–2 have been demonstrated using pulsed laser deposition.2 Cathode loadings are typically at ~0.40 mg (Pt) cm–2 and reductions in load-ing result in activity losses consistent with kinetic losses due to the oxygen reduction reaction.

For reformate applications, cathodes are similar to H2 PEMFCs; how-ever, anode catalysts are typically high-loading PtRu on high surface blacks used at higher electrode loadings than H2 PEMFC (typically, 0.4–0.6 mg (Pt) cm–2). A Pt:Ru ratio > 1 is favored.3 MeOH-fueled PEMFCs (i.e., direct methanol fuel cells—DMFCs) also use PtRu catalysts for the anode. A Pt:Ru ratio of 1:1 is favored for these catalysts. Until recently, unsup-ported PtRu catalysts (“blacks”) were favored as high electrode loadings were used to maximize activity. However, more recently, high loading PtRu on carbon blacks has been used to maximize active surface areas in an effort to reduce anode Pt loadings. The DMFC cathode typically used unsupported Pt blacks, but is being replaced with high-loading Pt/C catalysts. All these catalysts are commercially available from a number of suppliers (e.g., Johnson Matthey, Tanaka, Umicore) up to at least kilogram quantities.

1.3 Application Targets

As application markets become closer, more emphasis has been placed on defining performance and stability targets for membrane electrode assem-bly (MEA) technology and individual components including catalysts. In particular, the U.S. Department of Energy (DoE) has been prescriptive in defining MEA targets for automotive, stationary, and portable applications as the basis of its long-term hydrogen and fuel cell research and develop-ment (R&D) program.5 Similar targets have been established by both NEDO (New Energy and Industrial Technology Development Organization) and the European Union for their respective R&D programs.6,7 In addition, commercial fuel cell system developers have also established MEA perfor-mance and stability targets, although these are usually not available in the public domain.

The DoE has been the most detailed in translating MEA targets into com-ponent targets for automotive applications. These will be summarized in the following Table 1.1 for electrocatalysts.

Catalyst performance targets for stationary and portable applications have not been as consolidated and are usually embedded into MEA performance


targets. Similarly, durability targets for these applications have not been detailed for catalysts, but MEAs are expected to survive for >40,000 and 3,000 hours, respectively, with low degradation rates.

1.4 Electrocatalyst Discovery

Advanced discovery of new electrocatalyst formulations is increasingly dominated by two techniques: high-throughput screening of both model and practical catalyst materials and computational approaches to identify new active surfaces through theory.

1.4.1 High-Throughput Screening

As well as theoretical approaches to identify more active surfaces for elec-trocatalysis, there has been a recent strong interest in using high-throughput and combinatorial approaches for electrocatalyst discovery. In general, these approaches have used thin-film deposition and high surface area material preparation techniques to explore composition phase space for binary and more complex alloys. In particular, the thin film deposition approach has been used to create low-surface and high-surface materials; the 3M Company’s nanostructured thin-film substrate family of materials is the best known (see Figure 1.2). This has the advantage of creating combinatorial libraries of films

TABLE 1.1

Stack Targets

Property Units 2005 Status 2010 Target 2015 Target

MEA PGMa content Grams/kilowatt 1.1 0.3 0.2MEA PGM content Milligrams PGM/cm2

geometric area0.8 0.3 0.2

MEA catalyst cost Dollars/kilowatt 9 5 3Durability with cycling (≤80 C)

Hours >2,000 5,000 5,000

Cathode ECA loss Percent 90 <40 <40Cathode support loss

Millivolts after 100 h at 1.2 V

>>30 <30 <30

Cathode mass activity

Ampere/mg Pt at 900 mViR-free

0.11 0.44 0.44

Cathode specific activity

Microampere/cm2 Pt at 900 mViR-free

180 720 720

Cathode non-Pt volume activity

Ampere/cm3 at 800 mViR-free

8 >130 300

a PGM platinum group metal.


FIGURE 1.2Scanning electron micrographs, at different magnifications, of a 3M Company Pt-coated thin film catalyst. (Reprinted from M. K. Debe et al., Journal of Power Sources, 161, 1002. Copyright 2006, with permission from Elsevier.)


with sufficient surface area that can be tested as practical catalyst layers in PEM fuel cells. This overcomes one objection of the combinatorial approach that argues that only model surfaces are created and screened, rather than practical materials.

Early combinatorial screening focused on anode catalyst discovery, with Symyx using sequential Rf magnetron sputtering to create binary and ter-nary Pt alloy libraries that were screened for CO tolerance and MeOH oxi-dation. Gurau et al. used arc-melted alloys and chemical reduction of salts to produce libraries that were screened for MeOH oxidation.9 A PtRuOsIr quaternary composition was identified to have significantly higher activity than PtRu.

More recent work has been applied to the oxygen reduction reaction by a number of groups. In collaboration with Dahn at Dalhousie University, the 3M group has used co-sputtering to investigate the activity and acid contact stability of a range of binary and ternary Pt alloys.10 Other groups have used other techniques to deposit multielement films. Honda has used physical vapor deposition to deposit libraries of binary Pt alloys of varying compo-sition and screened them using rotating disk electrode (RDE) techniques.11 Workers at Southampton University have used physical vapor deposition to co-deposit ternary multicompositional metal films of, for example, PtAuPd, and screened for oxygen reduction activity in static electrochemical cells. This approach has also been extended to non-Pt compositions12 (see Figure 1.3).

Although these approaches appear successful in identifying alternative alloy compositions with significant oxygen reduction activity, no claims have been made that significantly higher oxygen reduction than binary Pt alloy surfaces has been achieved. In addition, it is known that the majority of base metals present in Pt alloys are sensitive to acid dissolution, so recent emphasis has been on using the thin-film approaches to screen for acid-re-sistant formulation. In particular, the work of 3M/Dalhousie has shown that many Pt binary and ternary thin films show acid dissolution of base metals (e.g., Co, Ni, Mn, Fe), which may limit their usefulness in PEMFCs.13 The thin-film approach has also been useful in screening for alloy formulations that show acid resistance. Dalhousie has shown that films of PtTa alloys are stable toward acid contact; however, the addition of even low amounts of Ta appears to reduce the oxygen reduction activity of the alloy.14

1.4.2 Computation Studies

Computational chemistry techniques have increasingly been used to inves-tigate electrocatalytic reactivity. Although strong emphasis has been placed on probing mechanistic pathways for oxygen reduction, methanol oxidation, and hydrogen oxidation in the presence of CO, the role of different surfaces has also been studied. A key activity has been to determine relevant descrip-tors that allow the screening of different surfaces to predict reactivity. For example, for oxygen reduction, the calculated stabilities of likely surface


intermediates such as O and OH have been used to correlate against experi-mental activity. This has allowed the screening of pure metals, alloys, and metal overlayer surfaces. Efforts have also been made to introduce the effect of electric field to explain observed overpotential effects; this has been com-bined with intermediate stability to predict more active surfaces. Specific examples of work will be surveyed within each electrocatalytic section.

It would seem opportune to combine both these computational approaches with the thin-film combinatorial approaches described earlier to identify

Working Electrode(s)

1 cm

2 cmTop View

Cell Lid

Cell Top

Reference Electrode

Counter Electrode

O-ring

F

Expanded Cutaway View

of a Single Electrode

A

C

D

E

Cutaway Side View 1 cm

FIGURE 1.3Schematic of a 64-electrode cell for electrochemical screening of electrocatalysts. (Reprinted with permission from Journal of Combinatorial Chemistry, 6, 149 (2004). Copyright 2004 American Chemical Society.)


more active and stable formulations. However, to date, the literature contains relatively few examples of this, although the study of Strasser et al. on the modeling of CO oxidation activity to predict MeOH oxidation verified using thin-film deposition.15

1.5 Electrocatalyst Preparation

The preparation of commercial Pt containing electrocatalysts typically uses deposition-precipitation methods in which metals are deposited onto a car-bon support from soluble species by pH swing or chemical reduction. These routes have the advantage of generally being aqueously based and batch pro-cessible to multikilogram scale. One disadvantage of these routes, however, is that metal particle dispersions are somewhat dictated by the surface char-acteristics of the support. For example, the catalyzation of highly graphitic carbon supports is limited at high Pt loadings by the low surface areas and surface site densities of these supports. The preparation of bimetallic and more complex formulations often requires a further annealing step to ensure good alloying. This can lead to excessive sintering of the Pt alloy particles, leading to low active surface areas and limiting activity.

Many alternative preparative routes have been investigated in recent years in attempts to overcome the limitations of the conventional routes and to investigate the effect of these alternate routes on catalyst structure and activ-ity. However, providing a comprehensive review of alternative routes is not the intent here. Rather, the discussion seeks to highlight the range of routes reported and identify potential advantages and disadvantages, and it will be limited to Pt-containing catalysts. The alternative deposition methods can be arranged into five general types: conventional, colloidal, molecular precur-sor, vapor phase, and surface modification.

1.5.1 Conventional Routes

As noted before, many conventional methods rely on hydrolysis/precipi-tation chemistry to deposit Pt and other metals onto carbon, followed by chemical or gas-phase reduction. One common method is the use of metal sulfito chemistry. This method involves the preparation of metal sulfito com-plexes (e.g., Na3[Pt(SO3)2(OH)2]) in water, the addition of carbon, and precipi-tation of metal by oxidation to deposit metal oxide particles.16 The route has advantages in that alkali metals and halides are excluded from the prepara-tion. The method has been extended to bimetallics such as PtRu.

The preparation of Pt alloys with base metals usually relies on a further annealing step to ensure good mixing of metals. Thompsett17 and Antolini18 have recently reviewed this area.


In recent years, many academic workers have used impregnation methods as a simple means to deposit metal onto carbon supports. Typically, a solu-tion of a metal salt (or salts) is mixed with the carbon support, the solvent removed by evaporation, and the resulting solid reduced chemically or by gas-phase treatment. Variations on this method include the use of a chemical reducing agent (e.g., NaBH4) in the slurry phase and the use of sonication19 or microwaves20 to encourage deposition and reduction.

1.5.2 Colloidal Routes

These routes involve the formation of (usually) prereduced metal particles that are then adsorbed or deposited onto the support. They have the advan-tage that the particle size of the particles is predetermined by the chemistry of the colloids and that resulting catalysts have narrow particle size distri-butions. However, the colloidal particles often are surface stabilized by sur-factant molecules, which can be difficult to remove once the particles are adsorbed onto the support. One further disadvantage is that the colloidal particles are prepared at high dilution (typically millimolar concentrations—for example, 0.2 g Pt l–1), which is a disadvantage in terms of scale-up.

Early work in this area was performed by the Bönnemann group at the Max Planck Institute, Mülheim, Germany.21 Using both tetraalkylammonium and trialkylaluminium stabilizers in nonaqueous solvents, researchers prepared colloidal PtRu and Pt3Sn particles with small average particle sizes (~2 nm) and narrow particle size distributions. The colloids could be adsorbed onto carbon to at least 20 wt% loading and the adsorbed surfactant molecules could be removed by careful thermal treatments under sequential inert, oxidizing, and reducing conditions. Other groups have shown similar results using a range of alternative surfactant stabilizers—for example, dodecyldimethyl (3-sulfo-propyl) ammonium hydroxide (SB12) in aqueous/alcoholic mixtures22,23 (see Figure 1.4). A further related approach is the use of surfactants to produce microemulsions in water in oil mixtures, in which controlled water microdrop-lets are formed in nonpolar solvents by use of non-ionic and ionic surfactants.24 The droplets allow reaction of metal salts with reducing agents and carbon to produce supported nanoparticles (e.g., PtRu/C).

In addition to surfactant-stabilized colloids, there has been work on form-ing metal colloids without stabilizers. Lee et al. have shown that direct reduction of metal chloride salts in tetrahydrofuran with LiBH4 gives small nanoparticles that can be impregnated onto carbon.26 No further treatment is required after the removal of the solvent. This route was applied to the preparation of PtRu, PtNi, PtMo, and PtW particles.

Liquid crystals as structure-directing agents have been used to produce mes-oporous Pt and PtRu materials.27,28 These unsupported materials (akin to metal blacks) have regular arrays of mesopores of ~2.5 nm separated by ~2.5 nm walls (see Figure 1.5). For Pt, high surface areas (60 m2 g–1) have been produced; these are significantly higher than those shown by commercial Pt blacks.


1.5.3 Molecular Precursor Routes

These routes rely on the direct transformation of soluble molecular spe-cies into supported metal (or mixed metal) particles. One method that has recently become popular is the “polyol” method. This takes a solution of metal salts, the carbon support, and a polyalcohol such as ethylene glycol. On heating, the polyol acts as both stabilizer and reductant, forming reduced metal particles on the carbon. It has been used successfully to prepare Pt and PtRu catalysts.30

Another route that has been used recently to prepare Pt bimetal-lics (PtRu,31 PtNi,32 and PtCr33) is via metal carbonyl chemistry. In this method, metal carbonyl complexes (either preformed or formed in situ)

3.7 nm3.7 nm

(a) (b)

20 nm

FIGURE 1.5(a) Schematic diagram of the geometry of the mesoporous H1-Pt catalyst; (b) TEM image of mesoporous H1-Pt catalyst. (Reprinted from A. Kucernak and J. Jiang, Chemical Engineering Journal, 93(1), 81. Copyright 2003, with permission from Elsevier.)

30 nm 10 nm

FIGURE 1.4HRTEM images of surfactant-stabilized PtRuOsIr oxide colloids at different magnification. (Reprinted with permission from M. T. Reetz et al., Journal of Physical Chemistry B, 107, 7414 (2003). Copyright 2007, American Chemical Society.)


in organic solvents (e.g., xylene, MeOH) are heated in the presence of car-bon, resulting in deposition and decomposition to form metal particles. The route has been shown to produce good particle mixing and, in the case of PtCr, remarkably good alloying after only a mild heat treatment (500°C). A similar route was investigated by Ding et al. for the prepara-tion of PtTi alloy catalysts.34 Carbon was added to Pt2(bis-dibenzylidene)3 and TiCl4(THF)2 in THF and the solvent removed by evaporation. The resulting powder was then heated in H2/N2 at temperatures up to 950°C to induce alloying.

True bimetallic molecular precursors have been used to prepare PtRu cata-lysts. Steigerwalt, Deluga, and Lukehart impregnated the complex (h-C2H4)(Cl)Pt(h-Cl)2Ru(Cl)(h3:h2-2,7-dimethyloctadienediyl) on graphitic nanofibers via solvent evaporation.35 The deposited complex was subsequently decom-posed by reductive annealing up to 650°C.

1.5.4 Vapor Phase Routes

These routes rely on vapor phase preparation of catalysts by deposition of metal precursors onto carbon or by direct formation of the catalyst in the vapor phase. Direct vapor deposition of volatile molecular precursors such as acetylacetonate complexes onto carbon has been demonstrated by Sivakumar and Tricoli for PtRu and PtRuIr.36

Spray pyrolysis routes have been extensively investigated to prepare Pt-based catalysts. Typically, a liquid feed of metal precursor and carbon is atomized into an aerosol and fed into a continuous furnace to evaporate and heat-treat to form a collectable powder. The method has good control over final aggregate particle size and metal particle size distributions, as well as producing powder without further isolation or separation. Hampton-Smith et al. have reviewed efforts of Superior MicroPowder (now Cabot Fuel Cells) in this area.37

Direct metal deposition from metallic sources has been extensively used for model catalyst deposition for high-throughput and combinatorial stud-ies. However, these methods are also increasingly used to deposit practical electrocatalyst materials. The best known approach is the one developed by 3M; researchers have used physical vapor deposition to deposit Pt and Pt alloys onto nanostructured (NS) films composed of perylene red whiskers. The approach has been recently been reviewed by Debe.38

1.5.5 Surface Modification Routes

An alternative approach to the preparation of bulk Pt bimetallic particles has been the preparation of surface modification of Pt particles with second metals. This approach has been used to prepare partially coated and fully coated Pt par-ticles to give “core-shell” structures. Core-shell structures of Pt or PtM bimetal-lics on alternative metal cores have also been prepared. This approach argues


that the surface composition of a Pt particle is key to reactivity and that the bulk of the particle has little influence. The use of a different core with a Pt shell has been shown to impart geometric influences due to lattice mismatches.

Early work concentrated on depositing partial shells of a second metal (e.g., Ru, Sn, Mo) on Pt and evaluating for oxidation reactions (CO oxidation, MeOH oxidation). Lee and Bergens used a Pt-catalyzed hydrogenation reac-tion of a Ru solution precursor ((1,5-cyclooctadiene)Ru(h3-C3H5)2) to selec-tively deposit Ru on Pt black electrodes.39 Ru coverage ranged from 0.05 to 3.5 monolayer (ML) equivalents with a coverage of 0.05 showing the highest activity for MeOH oxidation at 25°C. Crabb, Marshall, and Thompsett used a similar methodology to deposit Sn onto Pt/C using Bu4Sn as the Sn pre-cursor.40 Sn coverage up to a monolayer were deposited, but it was shown not to cover the entire Pt surface. This type of hydrogen-mediated, surface-controlled modification onto Pt/C has been shown for Ge,41 Mo,42 and Ru.43 A number of workers have shown that Ru absorbs spontaneously onto Pt surfaces from both simple RuCl3 solutions and complex precursors (summa-rized in Cao and Bergens44).

Pt partial and full monolayers have been deposited onto alternative metal particle cores. Brankovic, Wang, and Adzic found Pt spontaneously depos-ited onto hydrogen-reduced Ru particles at a surface coverage of 0.11 and 0.50.45 An alternative method of depositing Pt onto Pd, Pd alloy, and Au par-ticles has been developed by Sasaki et al.46 and Zhang et al.47 This involves the deposition of a Cu overlayer by underpotential deposition followed by galvanic displacement by Pt. Complete monolayers of Pt have been claimed and the technique has been extended to preparation of mixed monolayers of Pt and other precious metals.48

1.6 Electrocatalyst Testing

The evaluation of catalysts for PEMFCs in recent years has been well defined experimentally and numerically. In particular, studies have correlated that observed in liquid electrolytes to that observed in MEAs (e.g., for cathode activity, see Gasteiger et al.49).

The evaluation of catalysts typically uses two techniques. The first is evalua-tion as a thin layer on a bulk electrode (e.g., glassy carbon) in dilute liquid elec-trolyte (e.g., H2SO4) either as a static electrode or an RDE. In the study of oxygen reduction, there has been much discussion as to the most appropriate electro-lyte to use. In general, dilute perchloric acid (HClO4) is preferred; because of its noncoordinating nature, it is thus closest to the environment found within a PEM catalyst layer with perfluorosulfonic acid ionomer. A possible alterna-tive is trifluoromethylsulfonic acid (CF3SO3H), which mimics perfluorosulfo-nic acids closely, but there are relatively few studies with this acid.50 Rotating


disk electrode evaluations are preferred over static experiments because they provide a means to quantify and partially remove effects due to diffusion of reactants to active sites. However, deposition of practical catalyst materials (i.e., mesoporous particular samples) onto flat disk electrodes may still induce dif-fusion limitations of reactants due to the microstructure of the sample, which cannot be eliminated by high rotation. It has been suggested that the use of catalysts supported on microelectrodes can overcome these limitations.

The second approach is to test catalysts as layers in full MEA structures. This has the advantage of testing catalysts under realistic conditions and in realistic environments. However, this approach depends on creating a near-optimal catalyst layer structure that shows high utilization of the catalyst, together with a structure that allows adequate hydration and reactant/product transport.

With both approaches, it is key to establish the current regions where samples are under kinetic control to allow the correct comparison. Many reported comparisons of catalysts in MEA structures point to differences in performance, which are attributed to intrinsic catalyst differences when it is clear that differences are due to mass transport effects because of catalyst layer structure. To help overcome these difficulties, it is recommended that, for catalyst evaluation, pure reactants be used (e.g., O2 instead of air) and at relatively high stoichiometries. Use of current-voltage curves should be corrected for electrolyte or membrane resistances and Tafel analysis used to identify the kinetically controlled current regions.

1.7 Enhanced Activity Cathode Catalysts

As discussed previously, the DoE has set a target of catalyst activity of four times that of pure Pt/carbon. This can be expressed as activity per mass of Pt or activity per cost equivalent. Three general approaches have been investi-gated to achieve this target: Pt alloys, Pt core-shells, and non-Pt catalysts

1.7.1 Pt Alloy Catalysts

Since the late 1970s, it has been known that alloying Pt with other metals enhances oxygen reduction activity in acid electrolytes. Strong commercial interest in the phosphoric acid fuel cell led to an extensive development of carbon-supported Pt alloy catalysts in the 1980s. These catalysts typically showed mass and specific activity benefits of 1.5–2 and 5–6, respectively, over corresponding Pt catalysts. Together with superior durability aspects, Pt alloy catalysts were utilized as the cathode catalyst of choice for the first generation of PAFC power plants, such as United Technology Corporation’s PC25C.51 With the emerging interest in PEMFCs in the 1990s, similar Pt alloy


catalysts were evaluated for oxygen reduction performance; it was claimed that they had benefits similar to if not greater than pure Pt.

1.7.1.1 Model Surface Studies

In recent years, there has been a strong effort to use model catalyst surfaces combined with atomistic modeling approaches to understand why Pt alloys show activity enhancements and to predict new surfaces that show even greater activity for oxygen reduction.

The study of model Pt alloy surfaces has been pioneered by the work of Ross and Markovic in the United States, first at the Lawrence Berkeley National Laboratory and more recently at Argonne National Laboratory. A series of bulk Pt3M (M V, Ti, Fe, Co, Ni) alloy polycrystalline materials were prepared and their surfaces characterized by surface science techniques.52 The composition of the surface was found to be dependent on posttreatment. Surfaces sputtered with Ar ions showed compositions similar to the bulk, while surfaces that had been annealed in vacuum showed strong segrega-tion of Pt to give Pt skins.

This segregation has been predicted by density functional theory (DFT) calculations of Pt alloys by Norskov’s group.53 Evaluating these surfaces for oxygen reduction activity in diluted acid showed a range of enhance-ments over Pt ranging from 1.4 to 3.4, with surfaces with Pt skins over PtFe and PtCo bulks being the most active. The activity was correlated with a shift of the d-band center as measured using Ultra-Violet Photoelectron Spectroscopy (UPS). This was shown to lead to a weaker Pt–OHads interac-tion and a reduction in OH coverage at a given potential. This reduction in OH coverage allows a greater number of sites to be available for O2 adsorp-tion, dissociation, and reduction. Analysis of the Tafel slopes of the Pt alloy surfaces showed similar values to pure Pt, indicating no change in the O2 reduction mechanism but rather an increase in the number of active sites at a given potential leading to a higher oxygen reduction rate.

The work has been extended very recently to the examination of single-crystal surfaces of Pt3Ni.54 Study of the three low-index planes of Pt (<100>, <110>, <111>) has previously shown that the oxygen reduction reaction is somewhat structure sensitive in dilute noncoordinating acid (HClO4), with an activity ranking of (100) < (111) ~ (110). In contrast, the Pt3Ni surfaces show an activity ranking of (100) < (110) << (111); the Pt3Ni(111) is approximately 10 times as active as polycrystalline Pt and is claimed to be the most active Pt surface yet observed (see Figure 1.6). The activity enhancement was attrib-uted to a large shift in the d-band center (0.34 eV) in comparison with Pt(111). This was manifested in a 100 mV shift in the onset of Pt-OH formation in cyclic voltammetry experiments and had a direct effect on the sites available for O2 adsorption (1– OHads).

The investigation of model surfaces for OR has been closely followed by theoretical calculations with the aim to develop robust thermodynamic


models based on possible reaction mechanisms and energies of surface intermediates. Norskov’s group at DTU has been in the forefront of these approaches. In this work, the strength of adsorption of O and OH interme-diates on the metal surfaces is considered the key descriptor in predicting oxygen reduction activity, resulting in a volcano plot (see Figure 1.7).55

In this plot, pure Pt is shown to have the highest activity with an interme-diate M–O binding energy. Metals with higher O binding energies are less active than Pt and activity is limited by slow removal of adsorbed O and OH species. Metals with lower O binding energies (Ag and Au) also show lower activity due to low O2 dissociation rates and protonation of absorbed O2. However, the plot predicts that Pt is not at the peak of the plot and sug-gests more active surfaces could exist. The approach has been extended to Pt3M alloy surfaces, showing that although oxygen reduction activities are increased, the alloys sit beyond the peak of the volcano with O binding energies that are too weak56 (see Figure 1.8). This approach appears to pre-dict a maximum activity of at least double that shown by the most active Pt alloys.

1.7.1.2 Pt Alloy Nanoparticles and Particle Size Effects

Although model catalyst studies show the maximum possible activity obtain-able, practical catalyst systems use nanoparticles of Pt or Pt alloys usually

SurfaceMorphology

d-bandcenter (eV)

Sp

ecifi

c A

ctiv

ity

: ik (

μA

/cm

2)

0.1

M H

ClO

4 a

t 0

.9 V

vs.

RH

E(111)

15

10

5

0

–3.1

Pt3Ni(111)

|Δd(111)| = 0.3 eV

|Δd(100)| = 0.3 eV |Δd(110)| = 0.2 eV

Pt3Ni(100) Pt3Ni(110)Pt(111) Pt(100) Pt(110)

–2.8 –3.2 –2.9 –2.7 –2.5

(100) (110)

FIGURE 1.6Oxygen reduction specific activities of the low index planes of Pt and Pt3Ni in 0.1 M HClO4 at 0.9 V versus RHE at 60 C. (From Science, 315, 493 (2007). Reproduced with permission from AAAS.)


1.6

Pt

–0.30

–0.28

A/e

V

–0.26

–0.22

–0.24

–0.20

1.7 1.8 1.9

ΔEO/eV

Pt3Ni

Pt3Ni

Pt3Ti

Pt3Ti

Pt3CoPt3Fe Pt3Fe

Pt3Co

FIGURE 1.8Trends in calculated and measured oxygen reduction activity of bulk Pt and Pt3M alloys as a func-tion of oxygen binding energy. (V. Stamenkovic et al., Angewandte Chemie International Edition (2006) 45, 2897. Copyright Wiley–VCH Verlag GmbH & Co. KGaA. Reproduced with permission.)

FIGURE 1.7Trends in calculated oxygen reduction activity plotted as a function of the oxygen binding energy. (Reprinted with permission from Journal of Physical Chemistry B, 108, 17886 (2004). Copyright 2004 American Chemical Society.)


supported on carbon black supports to maximize the surface area per mass of Pt. It is therefore important to determine the surface intrinsic (i.e., specific) activity of these materials to determine the relationship to model “ideal” sur-faces. It is well known that Pt nanoparticles exhibit a particle size effect for oxygen reduction in a range of electrolytes, with the specific activity decreas-ing with decreasing particle size.

Kinoshita originally explained this by a change in ratio of single-crystal fac-ets as a function of particle size, together with an increased proportion of edge and corner sites.57 These sites have lower coordination than terrace sites and therefore bind O more strongly. It has been shown that as the particle size is decreased from 30 to 1 nm, the onset of oxidation is shifted to cathodic poten-tials, together with an increase in the irreversibility of the oxide formation and reduction. This increase in the oxide coverage with decreasing particle size at a given potential reduces the sites available for O2 adsorption and reduc-tion activity. Comparing the activity in aqueous acid of supported small Pt particles and single crystals shows a 13× improvement in area specific activity between 1 nm Pt particles and Pt(110) surfaces.58,59 Comparing Pt particles of increasing size only when particles reach ~30 nm (3M NS catalyst films) was the specific activity comparable to bulk polycrystalline Pt (see Figure 1.9).

Carbon-supported Pt alloys have also been studied in both dilute acid and MEA environments for oxygen reduction activity. In dilute HClO4 spe-cific activity enhancements of 1.5–3 were found over high surface area Pt, while in MEAs specific activity enhancements of up to 10 times that of Pt have been observed.61,62 As with pure Pt particles, Pt alloy particles also show an effect of particle size on oxygen reduction activity. In particular, large carbon-supported Pt3Co alloy particles (~10 nm) can show specific activities in excess of those for single-crystal Pt surfaces. However, when specific activities are compared to Pt particles of the same size, an activ-ity enhancement of 2.5 times is found, similar to the 3 times enhancement observed with polycrystalline Pt3Co surfaces over Pt. However, the specific activity of the best carbon-supported Pt alloys is still 10 times less than that of the Pt3Ni(111) surface, suggesting that if this surface could be cre-ated with nanoparticles, then cathode performances could be improved by 60–70 mV in the kinetic region.

A strong caveat to this prediction is the role of active surface area to the actual activity observed in fuel cells. This is often termed the mass activity—that is, the activity per mass of active metal (usually Pt), which has a direct correlation to cost. Mass activity is a combination of specific activity and surface area:

Mass activity (A/gPt) specific activity

(A/m2actual active area) × surface area (m2/gPt)

Mass and specific activities are often quoted at 900 mV cell potential with pure O2 to indicate the kinetic activity in a region as free of mass transport


0.30

0.28

0.26

pzk

(V

RH

E)

0.24

0 50 100

ECA |m2/gPt|

150 200

8.0

T = 333K

T = 293K

RDE

1600 rpm

0.1M HCIO4

4.0

0.0(a)

(c)

j h |

mA

/cm

2 real

| @

0.8

50

VR

HE

2.0

1.0

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Pt-

pol

y

NST

F

5 n

m

2 n

m

1 n

m

Pt-

pol

y

NST

F

5 n

m

2 n

m

1 n

m

(b)

j h |

mA

/cm

2 real

| @

0.9

00

VR

HE

FIGURE 1.9Comparison of oxygen reduction specific activity from polycrystalline Pt, nanostructured thin film Pt, and carbon-supported Pt from RDE measurements in 0.1 M HClO4. (Reprinted from K. J. J. Mayrhofer et al., Electrochimica Acta, 53(7), 3181. Copyright 2008 with permission from Elsevier.)


effects as possible. The DoE targets further define the temperature, pressure, and relative humidity at which the value should be measured (at least in MEAs). Surface area is usually measured electrochemically in situ by hydride adsorption or adsorbed CO oxidation charge.

Surface area has a strong effect on mass activity, especially as the surface area becomes lower due to increasing Pt particle size. Figure 1.10 shows how both the mass and specific activity change as a function of surface area for a series of carbon-supported Pt catalysts as cathodes in MEAs. Although the specific activity of the Pt steadily increases as the surface area decreases (increasing Pt particle size), the mass activity remains constant until below 20 m2 g–1 (~8 nm), where it decreases toward zero (see Figure 1.10). This indicates that although larger particles have better intrinsic activity, this increase in activity is not enough to compensate for reducing surface area.62

Therefore, the challenge is not only to create surfaces more similar to Pt3M(111) surfaces for maximum intrinsic activity, but also to create these on particles smaller than ~10 nm.

1.7.2 Pt Core-Shell Catalysts

As well as improving intrinsic activity of Pt nanoparticles, there has also been a strong interest in achieving similar activity improvements by intelligent

0

0.05

0.1

0.15

0.2

0.25

0 10 20 30 40 50 60 70 80

Electrochemical Area/m2 g–1 Pt

Mas

s A

ctiv

ity/

mA

mg

–1 P

t

0

200

400

600

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1000

1200

1400

Sp

ecifi

c A

ctiv

ity/

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cm

–2 P

t

Mass activity

Specific activity

FIGURE 1.10Comparison of oxygen reduction mass and specific activity of a series of carbon supported Pt catalysts with varying metal area as cathodes in MEAs. Activity at 900 mV, 150 kPa O2, 100% RH, 0.4 mg (Pt) cm–2.


thrifting of Pt to reduce cost. The most interesting of these approaches has been the synthesis of core-shell materials that has been pioneered by the group of Adzic at Brookhaven National Laboratory in the United States. This approach argues that because only surfaces of nanoparticles actually par-ticipate in the catalysis, Pt only needs to be at the surface. In addition, by supporting Pt (and Pt alloys) on other metal cores, modification of the elec-tronic properties of the Pt shell results, thus allowing further improvements in oxygen reduction activity.

Initial work showed that preparing Pt monolayers on both Pd(111) substrates and Pd/C nanoparticles gave improvements in oxygen reduction activity over pure Pd and Pt.63 In terms of mass activity of the Pt present, enhancements of 10 were found for the Pt/Pd/C systems for 9 nm particles and greater for 5 nm Pd particles when compared to Pt particles of 3.5 nm. Even account-ing for differences in specific activities due to the different particle sizes, these are remarkable improvements in activity. Characterization of the Pt/Pd(111) system showed interconnected Pt islands that were monoatomically thick but with holes allowing some access to the underlying substrate. It was thought that the Pt/Pd particles would show similar structures.

The improvement in oxygen reduction activity was proposed to be due to the formation of Pt-OH species that, through repulsion effects, retarded the formation of Pt-OH until higher potentials. This was manifested in higher Tafel slopes (–95 mV decade–1), indicating no poisoning of the oxygen reduction mechanism by adsorbed OH species. This approach was extended to investi-gate other substrates for Pt monolayers (i.e., Ru, Ir, Rh, Au); using DFT calcula-tions, a rationale was developed to explain the results.64 With Pt monolayers on single-crystal substrates, a volcano relationship was found, with Pt(111) and PtML/Pd(111) showing the highest activities. It was felt that maximum activity occurred when the energies of O2 dissociation and O hydrogenation were bal-anced. When the energy of dissociation is too high (as with Pt/Ir), the activity will drop because O-O bond breaking becomes too difficult. Conversely, when the energy of O hydrogenation becomes too high (as with Pt/Au), the activity will also drop because strong OH formation will block sites. Platinum and Pt/Pd show moderate energies for both processes.

The realization that holes in the Pt monolayer allow the core metal to oxidize and retard the formation of Pt-OH, thus leading to enhanced activity, led to the investigation of mixed metal monolayers.65 A range of PtM (M Re, Ru, Os, Rh, Ir, Pd, Au) on Pd(111) was prepared and it was found that second metals with easier formation of M-OH (-O) species gave significant activity benefits over Pt or PtPd (see Figure 1.11). Of these, PtRe and PtOs monolayers on Pd(111) gave enhancements of five times over Pt. Investigation of ratio indicated that Pt0.8M0.2 showed the best activity, which appeared to be a balance between enough M to give sufficient surface M-OH groups to retard Pt-OH formation and enough Pt to provide sites for O2 adsorption and turnover. It was shown that these enhancements could be translated to Pd/C cores, with the Pt0.8Re0.2/Pd/C showing 20 times greater activity than Pt/C based on Pt content alone.


This general approach of designed core-shell materials shows that large activity benefits can be achieved by efficient presentation of Pt to the reaction medium. However, challenges remain in terms of large-scale manufacture of these materials and whether the Pt-containing shells are durable for long-term fuel cell use. One long-term durability test has been reported using a Pt/Pd/C core-shell cathode.66 The cathode survived for 2,900 h; however, significant loss in performance was observed (>150 mV), even though this was a steady-state test without significant cycling. However, the approach of core-shell materials remains an attractive approach and warrants further development.

Koh and Strasser67 and Srivastava et al.68 have recently reported an alterna-tive approach based on core-shell structures. In this work, a series of base metal rich Pt alloys supported on carbon were prepared and the base metal deliberately leached out by electrochemically cycling in acid. The leached cat-alysts were found to have oxygen reduction activities significantly enhanced over Pt. In particular, leached Pt25Cu75 catalysts showed mass activities up to five times the rate of Pt. Interestingly, the increase in specific activities was only four times, some of which appeared to be due to the larger particles shown by the alloy particles. Characterization of the leached catalysts shows that much of the Cu has been removed during the leaching, leaving a pseudo core-shell structure with a PtCu alloy core surrounded by a skeletal Pt shell.

Similar results were found with dealloyed Pt ternary alloys (e.g., Pt20Cu20Cu60) and showed translation to MEA testing, with some samples showing activities in excess of the DoE performance target of 0.44 Amg–1 Pt

50

40

30

20

-jk a

t 0

.8V

mA

/cm

2

10

AuPd

Pt

–0.2 0 0.2 0.4

OH-OH or OH-O Repulsion eV

0.6 0.8 1

M0.2 Pt0.8/Pd(111)

Rh

Ru

Ir

Re

Os

FIGURE 1.11Oxygen reduction specific activities at 0.8 V of a range of (Pt0.8M0.2)ML/Pd(111) surfaces as a function of the calculated interaction energy between two OH or OH and O. (Reprinted with permission from Journal of the American Chemical Society, 127, 12480 (2005). Copyright 2005 American Chemical Society.)


(see Figure 1.12). As with the more conventional core-shell materials, this approach shows that significant improvements in Pt activity for oxygen reduction can occur and gives confidence that practical active and stable sys-tems can be developed.

0.94

0.92

0.88

0.86

E/V

0.84

Im/A mgPt–1

(a)

(b)

10.1

0.90

Pt20Cu60Co20

Pt20Cu20Co60

Pt20Cu40Co40

Pt25Co75

Pt 45 wt%

Pt 30 wt%

0.94

0.92

0.88

0.86

E/V

0.84

Is/μA cmPt–2

1000100

0.90

Pt20Cu60Co20

Pt20Cu20Co60

Pt20Cu40Co40

Pt25Co75

Pt 45 wt%

Pt 30 wt%

FIGURE 1.12Oxygen reduction mass activities of dealloyed ternary Pt alloy catalysts as cathodes in MEAs at 80 C, 150 kPa O2. (R. Srivastava et al., Angewandte Chemie International Edition (2007), 46, 8988. Copyright Wiley–VCH Verlag GmbH & Co. KGaA. Reproduced with permission.)


1.7.2.1 Use of Alternative Promoters to Pt

As well as using alloying and core-shell approaches to enhance the reactivity of Pt for oxygen reduction, the use of alternative promoters has been consid-ered. Although use of promoters in controlling activity and selectivity is com-mon in gas- and liquid-phase heterogeneous catalysis, relatively few examples have been reported for electrocatalysis. One approach reported by Bouwman et al. has shown that dispersing Pt in a hydrous iron phosphate matrix results in increased mass activity of oxygen reduction activity.69 Characterization of the as-prepared and electrochemically cycled catalyst showed no evidence of reduced metallic Pt. Sasaki et al. recently reported that Pt deposited onto NbO2 and mixed with carbon gave a catalyst that showed a mass activity three times that of Pt/C.70 A strong electronic interaction between the Pt and NbO2 was suggested as the possible cause of the enhancement.

1.7.3 Non-Pt Catalysts

Given the perceived expense and scarcity of Pt (~$1,300/troy ounce $40/gram as of August 2009), there has been a long and extensive effort to iden-tify much lower cost materials with significant oxygen reduction activity. This effort has been given a recent emphasis with the definition of catalyst requirements for automotive PEMFCs. A key feature of the use of alternative catalysts is that they do not lead to a significant increase in the volume of a fuel cell stack. Using a catalyst with poorer activity would lead to an increase in the number of MEAs required to achieve a certain power rating, as well as increasing the quantities of membrane, gas-diffusion media, and bipolar plates—thus alleviating any benefit of using a cheaper catalyst.

General Motors has assessed the required activity of a catalyst that costs less compared to the current state-of-the-art Pt activity based on these con-straints.71 Assuming that the catalyst layer thickness could be increased to ~100 μm from the currently used 10 μm, GM has calculated that the mini-mum volume activity (i.e., Acm–3) for a cathode catalyst that costs less should be at least 10% of the current Pt activity. In reality, this seems rather gener-ous, given the recent trend to reduce catalyst layer thicknesses to optimize high-current performances. The DoE has developed a series of volume activ-ity targets for nonprecious metal catalysts, with the 10% of Pt activity target (300 Acm–3 at 0.8 V, H2/O2) necessary by 2015.

There has long been interest in investigating Fe- and Co-based catalysts for oxygen reduction because of their role as highly effective enzymes for oxy-gen transport and conversion in biological systems. More recently, additional interest has been centered on alternative precious metals, metal oxides, and metal carbides and nitrides as possible oxygen reduction catalysts.

Good progress has been made in improving the activity of non-Pt cata-lysts. The most promising systems will now be reviewed. However, very little work has been reported on the stability of these systems and virtually


nothing on applying the accelerated ageing protocols, which have shown that even Pt catalysts are susceptible to dramatic activity loss. This should be a key feature of future work with non-Pt catalysts once adequate activity targets have been met.

1.7.4 Pd-Based Catalysts

The use of Pd as an alternative to Pt is attractive because of the cost differ-ences between the two metals. As of August 2009, Pd is 20% of the cost of Pt per weight and 11% per mole. However, it is well known that Pd oxides are readily soluble in acid and that even well-divided Pd metal is soluble in oxidizing acids. Despite these caveats, the oxygen reduction activity of Pd is closest to Pt in terms of activity per unit area (specific activity). The activity of Pd(111) is 0.15 that of Pt(111). As with Pt, Pd alloys as well as Pd overlayers have been investigated as a means to improve Pd activity. Adzic has investigated Pd overlayers on a range of single crystal sub-strates and found that only Pd on Pt(111) shows enhanced activity over Pd(111), although it is still inferior to Pt(111).72 Alloying Pd with based metals such as Fe, Co, and Ni has been shown to enhance activity over pure Pd; the best materials have shown similar activity to Pt (Pd2Co, Pd3Fe, PdCoAu, PdCoMo).73–76

It has been predicted that a Pd skin on a Pd3Fe core would sit close to the top of the activity/O binding energy “volcano” curve and thus have signifi-cantly higher activity than Pt and Pt/Pd core-shell materials.77 If these mate-rials could be shown to be stable to long-term PEM conditions, then these could represent viable replacements for Pt. As with other alternative non-Pt catalysts, very few stability studies have been reported.

1.7.4.1 Fe- and Co-Based Materials

Fe- and Co-based catalysts have long been explored for their oxygen reduc-tion activity in acidic media. In general, these materials were based on molecular macrocyclic compounds (e.g., porphyrins, phthalocyanines, tet-raazaannulenes) and supported on carbon blacks. Activity and stability were often enhanced after annealing/pyrolysis at high temperatures. The nature of the active sites for oxygen reduction stimulated great debate and has not been fully resolved. In general, it is agreed that the transition metal is still bound to N functionalities, even after pyrolysis, and other species such as metal and oxide particles are spectators to the reaction. This macrocyclic work was reviewed by Zhang et al.78 More recently, there has been a move away from preformed macrocyclic compounds to routes that create surface M-N species in situ. A study by Bouwkamp-Wijnoltz et al. showed that mix-ing a Co salt, carbon, and a N donor (e.g., 2,5-dimethylpyrrole), followed by pyrolysis, can catalyze with similar oxygen reduction activity to heat-treated Co tetraphenylporphyrin on carbon.79


Gasteiger et al. reviewed the best performing Fe-based catalysts in the lit-erature up to 200471. Even the best of these catalysts (Fe on pyrolyzed peryle-netetracarboxylic dianhydride) showed a corrected turnover frequency of 7% and a volume activity density of 0.2% of Pt. More recent work has focused on optimizing the metal, nitrogen, and carbon composition of the materials.

Work by the Los Alamos National Laboratory (LANL) has focused on pre-treating carbon black with molecular organic species such as pyrrole and then adding Co, followed by reduction and (optionally) pyrolysis.80 These gave catalysts with similar activity to the best reported pyrolyzed macro-cycle catalysts, but with crucially much improved stability. When tested as an MEA at 10%, the Co catalyst (at an electrode loading of 0.2 mg (Co) cm–2) showed a stable activity of 0.12 Acm–2 at 0.4 V (80°C, H2/air) for 110 h (see Figure 1.13). Somewhat similar work at the University of South Carolina showed that high surface area carbons could be modified by N and X (a non-disclosed nonmetallic element) and that significant oxygen reduction took place without any transition metals.81 This gave a similar stable performance to that of the LANL Co compound (0.12 Acm–2 at 0.4 V, 75°C, H2/O2, 200 h). Using these modified carbon materials as supports for transition metals (not specified) increased performance considerably (e.g., 0.22 Acm–2 at 0.5 V, 75°C, H2/O2, 80 h).

A further promising approach has been investigated by the 3M/Dalhousie group. Using vacuum deposition routes to prepare model Fe-C-N and Co-C-N films and investigating the effect of annealing, researchers found the most active surfaces during the transition from amorphous films to those contain-ing graphite and well-defined Fe3C and b-Co species.82 It was concluded that

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although these species were not active or stable toward oxygen reduction, they were necessary to maximize the number of active sites during initial carbon graphitization.

Subsequently, catalysts were fabricated using nanotechnology routes to yield more practical materials. Both carbon and TiC supports were modified with nitroaniline precursors and transition metals (not disclosed). The best of these materials showed remarkably high activity. At 0.7 V, a current density of 0.1 Acm–2 was achieved on H2/O2 (80°C) in MEAs. At very low current densities, a similar Tafel slope to Pt was observed and achieved a volume current density of 19 Acm–3 at 0.8 V—twice as high as previously reported. Additionally, a TiC-supported version has shown 1,000 h durabil-ity (0.15 Acm–2, 0.6 V, H2/air, 75°C).

This recent work shows that the activity of nonprecious metals (particularly Fe and Co) can be significantly improved by careful design and optimiza-tion of catalytic sites. In particular, the durability of these materials has been improved to show significant steady-state stability. However, despite these improvements, volume activity needs to be further improved by over a mag-nitude to allow consideration of replacement of Pt in practical applications.

1.7.4.2 MeOH-Tolerant Oxygen Reduction Catalysts

For DMFC systems, Pt cathodes are also used as the catalyst of choice; how-ever, given Pt’s ability to reduce oxygen and oxidize methanol, this lack of selectivity makes them sensitive to methanol crossover from anode to cath-ode via the membrane. This methanol crossover can have a depolarizing effect on cathode performance, reducing overall cathode activity. To combat this, an extensive effort has been made to identify and develop selective oxygen/reduction catalysts unaffected by MeOH crossover.

Initial work by Alonso-Vante and Tributsch83 and Alonso-Vante, Bogdanoff, and Tributsch84 showed that mixed metal chalcogenides such as Mo4Ru2Se8 and Ru1–xMoxSeOz displayed good oxygen reduction activity that was unaf-fected by the presence of MeOH. This has led to a large body of work inves-tigating the MeOH-tolerant properties of precious metal chalcogenides (mainly sulfides and selenides). In terms of a direct comparison with Pt, Schmidt et al. showed that a Ru1.92Mo0.08SeO4/C gave a mass activity 50 times poorer than that for Pt85 (see Figure 1.14). However, when low levels of MeOH (30 mM) were added, Pt activity fell by over 200 mV, while the oxygen reduc-tion activity of the RuMo catalyst was unchanged even with 0.5 M MeOH. Schmidt et al. also showed that a Ru/C catalyst gave similar activity to the RuMo catalyst with similar MeOH tolerance.

More recent research has focused on the binary Ru sulfides and sele-nides. Schulenberg et al. showed that modifying a Ru/C with Se (via H2SeO3) improved activity by a factor of three.86 It was concluded that the Se inhibited surface oxide formation that limits active sites with Ru/C. Both catalysts showed some H2O2 formation at lower potentials (e.g., 3% at


0.7 V). Cao et al. investigated the effect of Se- and S-modified Ru and Rh catalysts.87 Se-modified Ru showed the highest activity; however, S barely promoted the activity of Ru. The addition of Se and S to Rh led to a signifi-cant decrease in oxygen reduction activity. It was also found that cycling the modified Ru catalysts to 1.2 V removed Se and S; however, Se was found to be stable if the potential was limited to 0.85 V. Sulfur has been used to promote the selectivity of Pt for oxygen reduction over methanol oxidation. Gochi-Ponce et al. have reported that carbon-supported platinum sulfide (PtxSy/C ) showed significant MeOH tolerance toward oxygen reduction.88 However, overall oxygen reduction activity was reduced by an order of magnitude from that of Pt/C.

Ruthenium/carbon catalysts have also been promoted by the addition of Fe. Bron et al. reported the addition of Fe to a preformed Ru/C catalyst via adsorption of Fe complexes, followed by heat treatment.89 They found an increase in oxygen reduction activity of three to five times over unmodified Ru/C. It was suggested that the surfaces of Ru particles were covered with FeNxCy sites. As discussed previously, the Pd alloys have shown significant MeOH tolerance toward oxygen reduction and appear to have activities clos-est to that of Pt.

The development of MeOH-tolerant cathode catalysts for DMFC is a key component for certain system arrangements. In particular, it is critical for

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FIGURE 1.14Oxygen/reduction polarization curves from RDE measurements for Ru1.92Mo0.08SeO4, Ru/C, and Pt/C in 0.6 M H2SO4, 60 C. (T. J. Schmidt et al., Journal of the Electrochemical Society, 47, 2620 (2000). Reproduced by permission of The Electrochemical Society.)


mixed reactant flow cells where MeOH and air are fed as a two-phase stream. However, the majority of DMFC cell configurations are separate feed designs where MeOH and air are fed separately. In these systems, although MeOH crossover can occur and is detrimental to cell performance, MEA design configurations minimize this effect. In particular, the use of hydrocarbon membranes, which have intrinsically low MeOH permeation properties, can severely limit the amount of MeOH that does reach the cathode. In these situations, it is apparent that a MeOH-tolerant catalyst would not provide an advantage, particularly if its intrinsic activity for oxygen reduction is lower than that of Pt.

1.8 Cathode Catalyst Stability

The stability of electrocatalysts for PEMFCs is increasingly a key topic as commercial applications become nearer. The DoE has set challenging near-term durability targets for fuel cell technology (automotive: 5,000 h by 2010; stationary: 40,000 h by 2011) and has detailed the contribution of the (cath-ode) catalyst to these. In particular, for automotive systems as well as steady-state stability, activity after simulated drive cycles and start–stop transients has been considered. In practice, both these treatments have been found to lead to severe degradation of the standard state-of-the-art Pt/C catalyst, as detailed next.

To simulate a typical drive cycle over the lifetime of an automotive fuel cell stack, a voltage cycle of 0.7–0.9 V (80°C, 150 kPa, 100% RH, 30 s per voltage step, square wave) has been defined. The test is run with H2/N2 or H2/air for at least 30,000 cycles. It has been considered that, over the stack lifetime (~5,500 h), there will be 300,000 large voltage cycles. However, it is expected that activity degradation rates will have been established by 30,000 cycles and can be extrapolated to 300,000 cycles. A catalyst-derived degradation rate of <3 μV h–1 is considered the acceptable target. It has also been esti-mated that the fuel cell stack will be subject to 30,000 start–stop transients during its lifetime. It has been found that the cathode potential rises to at least 1.2 V during a start–stop transient due to a local cell reversal event as H2 replaces air or air replaces H2 in the anode. Given that each start–stop event lasts 10 s, it has been estimated that a steady-state condition of 1.2 V for 100 h will give the cumulative effect over the lifetime of the stack.

1.8.1 Pt Electrochemical Area Loss

It has been found that the stability of the supported Pt nanoparticles (typi-cally, 2–3 nm) and the carbon black support are affected by the accelerated stability treatments described before. Even under steady-state operating


conditions, degradation of the Pt catalyst is observed. Ferreira et al. showed significant loss in Pt electrochemical surface area (ECA) after operating short stacks containing Pt/C cathodes at 0.2 A cm–2 or open circuit voltage (OCV) for 2,000 h.90 Under electrochemical load, the cathode catalyst lost about 45% of its ECA; at OCV, it lost about 80% of its ECA. The decay in cell voltage with time was correlated with ECA loss with time. This ECA loss was translated into a decay rate of 10–20 μV h–1.

Use of voltage cycling has been shown to exacerbate Pt ECA loss. Ferreira et al. showed that cycling an MEA containing a Pt/C cathode between 0.6 and 1.0 V for 10,000 cycles (H2/N2, 80°C) resulted in a loss of 65% ECA. This was translated into a voltage decay rate of 300 μV h–1 in 100 h (~10,000 cycles). Other workers have seen similar results in MEAs and in liquid elec-trolytes. The cause of this Pt ECA loss has been extensively investigated. Ferreira et al.90 and Borup et al.91 have reviewed this area in detail, so only the general conclusions will be given here.

Three general mechanisms have been proposed to explain the Pt ECA loss observed on steady-state and voltage cycling conditions:

Pt is dissolved under oxidizing potentials and is reprecipitated on neighboring Pt particles leading to increased particle growth (Ostwald ripening). This mechanism is supported by Pt dissolu-tion studies that show that Pt is increasingly dissolved at oxidizing potentials (>0.8 V). Evidence of gross soluble Pt migration is found by appearance of Pt particles within the membrane, formed by reduc-tion of soluble Pt species by H2 gas crossover from the anode.Pt particle coalescence is due to migration. This mechanism is sup-ported by observations that, upon cycling, Pt particle size distribu-tions are shifted toward larger sizes, indicating that smaller particles are more mobile. It is noted that this observation could also result from the effects of Ostwald ripening.Pt particle agglomeration is due to carbon support corrosion. Electrochemical carbon corrosion is known to occur above 0.9 V. It has been suggested that loss of carbon causes Pt particle agglomera-tion and electrical isolation, leading to loss in activity.

It has been well established that Pt can dissolve under oxidizing condi-tions, although the exact manner of how the species formed is a matter of debate at present. The formation of Pt crystallites in the membrane (or at the anode if no H2 is present) would indicate that micrometer transport of solu-ble Pt occurs. However, careful analysis of the Pt particle size distributions in the cathode after testing suggested that purely Ostwald ripening could not explain the observed distributions. Therefore, at present, it is concluded that a mixture of Pt dissolution/reprecipitation and Pt particle coalescence is responsible for Pt ECA loss.


1.8.2 Stabilization of Pt Catalysts toward Potential Cycling

Given the satisfactory initial activity of Pt/C for oxygen reduction (although still too low for final automotive cost targets), a key objective is to preserve this activity to the end of life of the MEA/stack. Given the poor stability of standard high area Pt/C catalysts to steady-state and (especially) dynamic cycling effects, much recent effort has been devoted to stabilizing Pt particles.

Supported Pt alloys have long been used as cathodes for phosphoric acid fuel cells to improve catalyst durability and stability. As discussed else-where in this review, Pt alloys have also been used as PEMFC cathodes and have shown activities typically double those of pure Pt catalysts. Platinum alloys (primarily PtCo) have also shown greater cycle durability than Pt catalysts. Yu, Pemberton, and Plasse showed a relatively low degradation rate (4 μV h–1) during a voltage cycle of 0.87–1.20 V versus RHE for 2,400 cycles (H2/air) with only a 35% loss in ECA.92 Gasteiger et al.71 also showed that a PtCo alloy showed modest losses in Pt ECA; however, much larger losses in mass and specific activity were found than would be expected from the correlation of activity and ECA loss developed for Pt catalysts. This result was confirmed by Ball et al., who showed that although ECA was stable over long-term cycling (0.7–0.9 V, square wave, H2/air, 80°C, 83,000 cycles), significant loss in activity was found over time, showing Pt-like activity at end of life (see Figure 1.15).62 This was correlated with loss of surface Co from the alloy particles and it was suggested that a surface alloy was necessary for enhanced activity.

Although PtCo alloy catalysts appear not to show activity stability over time, they do show ECA stability. This appears to be at least partially related to the larger particle sizes shown by these materials. These larger sizes come from their methods of synthesis, where unalloyed precursors are annealed at relatively high temperatures to ensure mixing of components. This also has the effect of growing particle sizes to typically 5–10 nm, depending on the loading of metal and the carbon support used.

This observation has led to the investigation of the stability of larger Pt par-ticles. Makharia et al. found that heat-treating a Pt/C catalyst to give particles of 4–5 nm gave similar cycle resistance to a PtCo/C catalyst with a similar particle size.93 Interestingly, the direct preparation of a Pt/C catalyst with simi-lar particle sizes (i.e., without subsequent heat treatment) gave only modest improvements over a standard 2–3 nm Pt/C catalyst. Due to the particle size effect, increasing the Pt particle size to 5 nm and higher does not significantly affect oxygen reduction activity; Ball et al. have shown that heat-treating Pt/C to give 8 nm particles does not significantly affect oxygen reduction activity.62

The use of larger Pt particles does appear to produce catalysts stable toward voltage cycling without any significant loss in oxygen reduction activity. The challenge is to produce Pt alloys with similar voltage cycling stability and enhanced activity.


1.8.3 Effect of High Cathode Voltages on Catalyst Stability

It has been found that during normal operating conditions (0.6–0.85 V), the currently used cathode carbon supports (e.g., XC72R, Ketjen) appear suffi-ciently stable for long-term use. However, it has been recently shown that during start-up and shutdown, short-term potential excursions of 1.2–1.5 V are possible; this leads to significant corrosion of high surface area carbon blacks.94,95 Stack control strategies have been shown to limit this potential excursion to ~1.2 V; however, it has been estimated that the cathode will experience 100 h at 1.2 V over the lifetime of the stack.

In addition, stacks are also expected to sit at idle (~0.9 V) for much of the time. Mathias et al. studied the effect of these voltages on cathode carbon stability.96 Holding a standard 50% Pt/C catalyst at 1.2 V caused 15% loss of its carbon in 20 h and it was predicted not to survive the required 100 h. At 0.9 V, the catalyst was expected to lose 5% over a few thousand hours, which may be acceptable for long-term use (see Figure 1.16). The effect on MEA per-formance was also studied. After 20 h at 1.2 V, a 30 mV loss in performance was observed and it became progressively worse at longer times. The loss in

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FIGURE 1.15Change in electrochemical area and oxygen reduction specific activity for Pt and PtCo cath-odes during a 0.9–0.7 V voltage cycling as a function of log cycle number; 900 mV, H2/O2 2/10 stoichiometry, 150 kPaabs, 80 C. (S. C. Ball et al., Electrochemical Society Transactions, 11, 1267 (2007). Reproduced by permission of The Electrochemical Society.)


performance was ascribed to increasing mass transport losses induced by carbon corrosion.

The mechanism of carbon corrosion has been investigated in MEAs and in liquid electrolytes. Carbon itself is thermodynamically unstable toward oxidation at higher potentials, but this oxidation is kinetically limited:

C H2O CO2 4H 4e–, E0 0.207 VRHE

It is generally thought that carbon corrosion proceeds in three steps97:

1. oxidation of the carbon lattice, Cs Cs e–; 2. hydrolysis, Cs ½H2O CsO H ; and 3. gasification to carbon dioxide, 2CsO H2O CsO CO2 2H

2e–.

Ball et al. investigated the effect of carbon surface area on carbon corrosion at 1.2 V for 24 h and found that, for commercial carbon blacks, cumulative carbon corrosion correlated with carbon BET (Brunauer Emmett Teller) area, although when analyzed as specific carbon corrosion (weight of carbon cor-roded per unit of carbon area), some variation was observed.98 The effect of Pt on carbon corrosion has also been studied and conflicting results have been reported. Roen, Paik, and Jarvi found that Pt did increase carbon corrosion

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FIGURE 1.16Carbon weight loss as a function of time for two different catalysts in MEAs at both 0.9 and 1.2 V, 80 C, 100% RH, 120 kPaabs. (M. F. Mathias et al., Electrochemical Society Interface, 14, 24 (2005). Reproduced by permission of The Electrochemical Society.)


during voltage excursions,99 but Ball et al. found relatively little effect of Pt on carbon corrosion during 1.2 V steady-state holds.98

The stability of Pt particles during the 1.2 V hold has also been investi-gated. At 1.2 V and 80°C in 1 M H2SO4, up to 35% of the ECA was lost after 24 h. Transmission electron microscopy analysis of the tested catalysts found a growth in the Pt particle size distribution, suggesting that small Pt particles (~2 nm) are particularly susceptible to dissolution/agglomeration under steady-state voltage holds at 1.2 V.

1.8.4 Stabilization of Pt Catalysts toward High-Voltage Excursions

The instability of high surface area carbon blacks toward high-voltage excur-sions has prompted investigations of more stable forms of carbon black. Earlier developments with PAFC catalysts had identified graphitized carbon blacks as suitable supports for Pt cathode catalysts. These are typically stan-dard carbon blacks (e.g., XC72R, Ketjen, BP2000) that are subjected to high-temperature graphitization conditions (1,800–2,800°C). This has the effect of reducing surface area by removal of micropores (<300 m2 g–1), inducing greater graphitic domains with the carbon particle structure, and removing much of the oxidized surface chemistry. Given the lower surface areas and surface group density, graphitized blacks are more difficult to catalyze at higher Pt loadings to give small particle sizes and are often used with lower Pt loadings (<40 wt%).

The corrosion stability of a series of graphitized carbon blacks was investi-gated by Ball et al. at 1.2 V in 1 M H2SO4 for 24 h and compared to a series of con-ventional carbons.98 The graphitized blacks show up to 25 times less corrosion than large-area carbons and also showed less sensitivity to surface area (i.e., decreasing specific corrosion with increasing surface area) (see Figure 1.17).

Several groups have investigated the stability of Pt supported on graphi-tized carbon under PEMFC-type conditions (typically, 1.2 V holds). Stevens et al. showed that Pt on graphitized BP2000 showed much greater ECA sta-bility and performance at 800 mV at 1.2 V over 50 h than either Pt/XC72R or Pt/Ketjen catalysts.100 This has been confirmed by measurements of carbon corrosion of graphitized carbon catalysts at 1.2 V in MEAs by Yu et al.101 These show that graphitized carbons can be 35 times more stable than con-ventional carbon blacks (as assessed by the time taken to lose 5% carbon by weight at 1.2 V). It was also shown that although the graphitized carbon cata-lyst showed a constant corrosion rate with time, the standard carbon catalyst showed a decreasing corrosion rate. This was correlated with the corrosion of the amorphous core of the conventional carbon, which is removed more easily than the graphene shell of the carbon particles.

The use of graphitized carbons can impart significant stability to high-volt-age excursions and appears suitable for automotive fuel cell use. However, a number of workers have reported that the performance of Pt on these sup-ports is poorer than that found for conventional Pt catalysts. For example,


Mathias et al. showed that a 30% Pt alloy on a corrosion-resistant support gave an initial performance 50 mV below that of a standard 50% Pt/C cata-lyst.96 Although the Pt alloy catalyst showed much greater stability at 1.2 V, this large discrepancy in activity is unacceptable. Therefore, the challenge is to develop Pt (and Pt alloy) catalysts with high stability and similar activities to that shown by state-of-the-art Pt/C catalysts.

1.8.5 Alternative Supports

Prompted by the thermodynamic instability of carbon, efforts have been made to identify and develop alternative support materials that show good stability up to voltages of 1.4–1.6 V. Attempts to replace carbon as a support for Pt-based catalysts for fuel cells is not a new activity. General Electric investigated mate-rials such as boron carbide as early as 1968.102 During the development of PAFC catalysts, alternative supports were considered; for example, titanium carbide was found to have better corrosion resistance to graphitized carbon blacks. More recently, a wide range of ceramic materials, such as oxides, carbides, and nitrides, has been investigated as alternatives to carbon.

1.8.5.1 Oxides

A number of oxides have significant electronic conductivity at low tempera-ture and have been studied as possible fuel cell catalyst supports. Reduced forms of TiO2 (e.g., Ebonex, Ti4O7) have been catalyzed with Pt; although

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FIGURE 1.17Comparison of BET surface area and cumulative carbon weight loss following corrosion at 1.2 V versus RHE, 1 M H2SO4, 80 C, 24 h for a range of carbon blacks. (Reprinted from S. C. Ball et al., Journal of Power Sources, 171, 18. Copyright 2007 with permission of Elsevier.)


it has been shown that they can have improved corrosion resistance com-pared to carbon above 0.9 V, they also are susceptible to irreversible oxida-tion at high potentials, losing conductivity.103 Better results were found for Nb-doped TiO2, which was found to be more stable under oxidative condi-tions. One alternative approach to overcome the general insulating nature of TiO2 is to mix a Pt/TiO2 catalyst with a carbon black.104 The absence of direct contact of Pt with carbon would alleviate any Pt-promoted carbon corrosion at higher temperatures.

1.8.5.2 Carbides and Nitrides

Due to its wide commercial availability and its lowest electrical resistivity of any interstitial carbide (s 105 S cm–1 at 20 C), tungsten carbide has been the most studied carbide. Chhina, Campbell, and Kesler recently reported the thermal and electrochemical stability of WC with and without deposited Pt and assessed its suitability as a corrosion-resistant support for PEMFCs.105 Both WC and Pt/WC were found to be stable on cycling to 1.8 V in 0.5 M H2SO4, despite some indications that the surface of the WC was oxidizing to semiconducting WOx. Even though the initial oxygen reduction activity of the Pt/WC was poorer than that of Pt/C, it was concluded that WC had potential to replace carbon for applications where high-voltage transients would be observed.

One interesting approach that has been suggested is to improve the cor-rosion resistance of carbon blacks by coating them with WC.106 Depositing Pt onto WC-treated carbon has been shown to improve cycle degradation resistance to 1.8 V.

Titanium carbide has also been widely studied. Vinod and Frost prepared relatively high surface area forms of TiC (25–125 m2 g–1) and showed lower corrosion currents at 1.0 V in 100% H3PO4 at 200°C than graphitized XC72.107 The oxygen reduction specific activity of Pt/TiC was superior to that of Pt/C, although this may have been influenced by the larger Pt particles deposited onto the TiC.

Pt on TiC and TiN has also been reported to show good cycle stability (to 1.2 V), especially when mixed with uncatalyzed carbon.108 Steady-state corrosion tests of Pt/TiC at 1.2 and 1.4 V showed gradual oxidation to TiO2, although this was not dependent on potential. TiC has also been used by 3M as a support for a series of non-PM catalysts and has showed very stable per-formance over 1,000 h at 0.6 V. However, the overall activity was four times lower than that of an equivalent catalyst on carbon.

1.8.5.3 Nonconductive Whiskers

A unique approach has been taken by 3M, which has developed a family of nanostructured whisker-like nonconductive materials as supports for Pt or Pt alloys. Electronic conductivity is achieved by sputter-coating a continuous


Pt nanoparticular film. The whiskers are formed by the thermal sublimation and subsequent annealing of an organic pigment, N,N-di(3,5-xylyl)perylene-3,4,9,10-bis(carboximide) (“perylene red”). The resulting crystalline whiskers are 500–1,000 nm in length and approximately 50 nm in diameter. Given the nonconductivity of the whiskers and that Pt essentially covers their surface, the films have shown excellent corrosion resistance at high potentials (1.5 V) for 3 h.109

1.9 Carbon Support Materials

Carbon black is the traditional support material for fuel cell catalysts for the following reasons:

high electrical and thermal conductivity;available with a wide range of surface areas and porosities;high stability in acid environments;stable in reducing and reasonably oxidizing environments;available with low impurity levels; andcheap (less than or equal to a few dollars per kilogram) and available in

at least multikilogram quantities.

However, this has not stopped the investigation of alternative carbon forms as supports for fuel cell catalysts.

1.9.1 Conventional Carbon Blacks

A wide range of carbon blacks is available from a number of suppliers (e.g., Cabot, Columbian Chemicals, Azko Nobel, Denka, Timcal, Degussa, Mitsubishi). Blacks for fuel cell use fall into three general categories:

furnace blacks (e.g., XC72, BP2000) made from pyrolysis of heavy aro-matic fuel oils;

extra conductive blacks (e.g., Ketjen EC300J, Printex XE-2) made using the Shell gasification process, a by-product of heavy oil cracking; and

acetylene blacks (e.g., Shawinigan) made from pyrolysis of acetylene (ethyne).

These types of carbon blacks are characterized by particularly high surface area (50–1,500 m2 g–1)-to-volume ratios and have much lower impurity levels. However, despite the many black products available, only a small number have been reported for fuel cell use (XC72, BP2000, Ketjen EC300J). Currently,


high area supports such as Ketjen appear favored for PEMFC use, especially for cathode applications.

The main issue with these materials is oxidative stability at potentials above 1.0 V. The development of graphitized blacks to address this is covered elsewhere in this chapter. The thermal stability of high area blacks is also an issue. Stevens et al. showed that high area blacks such as BP2000 suffer excessive gas-phase oxidation at 150°C in dry air when high loadings of Pt are deposited onto them.110 Oxidation rates increase in the presence of H2O/air mixtures, indicating that Pt/C can self-catalyze the water-gas shift reac-tion at temperatures as low as 150°C.

Relatively few recent studies have investigated systematically alternative carbon blacks against the commonly used ones. Takasu et al. investigated a range of blacks from Mitsubishi Chemicals Corporation as supports for PtRu for MeOH oxidation.111 The carbons had surface areas ranging from 29 to 194 m2 g–1. They were characterized as mesoporous with pores in the 40–70 nm region. Rao et al. investigated a range of Sibunit carbons prepared by pyrolysis of methane as supports for PtRu for MeOH oxidation.112 Surface areas ranged from 6 to 415 m2 g–1 and pore sizes and distributions were varied systemati-cally, although all of the carbons had low microporous volume. The catalysts were tested as MEAs, and catalysts on the lower surface area carbons gave the highest mass activities for MeOH oxidation. This was ascribed to the high utilization achieved with these catalysts.

1.9.1.1 Modification of Carbon Blacks

There have been a few reports of modification of carbon blacks or supported black catalysts to enhance activity or performance. Simple oxidative treat-ments on preformed catalysts with the aim of increasing the acidic surface group density have been reported by Wilkinson et al.113 Once fabricated into the cathode, improved MEA performance was found. Direct chemical modi-fication of catalysts has also been reported. Xu, Qi, and Kaufman reported that both sulfonation (with 2-aminoethanesulfonic acid) and phosphona-tion (with 2-aminoethylphosphonic acid) of preformed Pt/XC72 catalysts result in improved MEA performance when the treated catalysts are used as cathodes.114,115 It was argued that the surface modifications aided proton transport and as such allowed reduced levels of ionomer to be used. Cabot reported similar work with directly sulfonated carbons.116

1.9.2 Synthetic Carbon Materials

1.9.2.1 Carbon Nanotubes and Nanofibers

A wide range of nanostructured carbons has been discovered since the origi-nal discovery of carbon nanotubes (CNTs) by Iijima in 1991.117 Carbon nano-tubes and nanofibers are nanoscale cylinders of rolled up graphene sheets,


with an extensive range of variations now available; for example, single-walled nanotubes (SWNTs), double-walled nanotubes (DWNTs), multiwalled nano-tubes (MWNTs), graphitic carbon nanofibers, carbon nanocoils, and carbon nanohorns. The potential benefits for fuel cell use that have been suggested include higher utilization of active metal due to the lack of smaller porosity and higher corrosion resistance due to the (theoretically) inert surfaces.

A wide range of these materials has been investigated for fuel cell use, usually as supports for PtRu particles for DMFC testing (presumably due to the ease of experimentation). The theoretically inert surfaces of CNTs pose some difficulties for metal cluster deposition because no sites exist for depo-sition or stabilization. Therefore, clusters tend to be deposited onto defect and amorphous portions of samples (see Figure 1.18).

To overcome this limitation, pretreatments such as oxidation have been used. Liu et al. showed that treated Pt/CNT gave O2 reduction activity supe-rior to that of untreated Pt/CNTs.118 Steigerwalt et al. examined a range of PtRu/CNTs and found that a PtRu/narrow herringbone graphitic carbon nanofiber gave 65% higher MeOH oxidation activity than a PtRu black in MEA testing.119 The enhancement was ascribed to the stabilization of a smaller particle size and presenting a more optimized particle composition to the reaction. Hyeon et al. showed that PtRu supported on a carbon nano-coil also showed better MeOH oxidation activity to PtRu/XC72.120 This was ascribed to the nanocoil having a higher surface area and thus being able to disperse the PtRu particles more effectively.

Although the majority of authors who have investigated CNTs as supports for Pt and PtRu particles claim higher activity or performance compared to conventional catalysts, it is not clear why these enhancement arise. It seems unlikely that the CNTs provide any electronic enhancement to Pt(Ru) reactiv-ity, so it is likely that CNTs provide benefits for catalyst layer structure. Part of this may be related to surface area because CNTs can have relatively high surface areas and are often compared to XC72 supported catalysts that have only a moderate surface area (~250 m2 g–1). Given the current high expense of these materials (>>$10 kg–1), further benefits of their use need to be identified before they can be practically considered as candidates for fuel cell catalyst supports.

1.9.2.2 Synthetic Mesoporous Carbons

An alternative type of synthetic carbon that has started to be widely investi-gated for fuel cell use is ordered mesoporous carbons (OMCs) with tunable pore sizes from 2 to 50 nm.121 It has been suggested that ordered mesopores offer better mass transport properties than the range of mesopores shown by conventional carbon blacks. Ordered mesoporous carbons are synthesized by a templating procedure starting with a highly ordered silica support such as MCM-41 or SBA-15. Negative carbon structures are made by filling the SiO2 pores with a carbon source such as sucrose or phenol/formaldehyde


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mixtures, followed by carbonization at elevated temperatures. The SiO2 is then removed by acid or basic etching.

As with CNTs, OMCs are often evaluated as supports for PtRu particles for MeOH oxidation. The range of materials tested to 2003 was reviewed by Chan et al., who found a number of examples that showed superior activity to conventional Pt and PtRu catalysts.122 More recent work—for example, use of PtRu catalysts derived from mesoporous SiO2 spheres by Chai et al.—also showed enhancements over PtRu/XC72 catalysts for MeOH oxidation.123

Although the use of synthetic OMCs appears a promising approach to optimize porosity for more effective mass transport to active sites, the cur-rent synthetic routes appear complex and wasteful in terms of the removal of the high-value mesoporous SiO2 template. Interestingly, nontemplated mesoporous carbons synthesized from the self-assembly of starch molecules have been recently reported by Budarin et al.124 After carbonization, the so-called “Starbons” contain negligible microporosity, with average pore diam-eters ranging from 6 to 17 nm and surface areas from 150 to 500 m2 g–1. These materials may provide interesting alternatives to carbon blacks as fuel cell catalyst supports.

1.10 Reformate-Tolerant Anode Catalysts

Although pure H2 is the preferred fuel for PEMFCs, H2 is usually produced by the reforming of natural gas to give reformate (a mixture of H2, CO2, and CO). Although CO levels can be progressively reduced through the use of water-gas shift and selective oxidation (with added air), it is difficult to reduce CO to levels below 10 ppm or higher. CO is a strong adsorbate to Pt at low temperatures and even 10 ppm of CO present in a H2 feed is enough to poison a Pt anode toward H2 electro-oxidation.

Originally, it was considered that most H2-based fuel cell applications would use reformate as the fuel due to the difficulty of purifying and trans-porting H2; thus, they would use integrated fuel processors to produce the reformate. However, after a number of years of development, the DoE announced in 2003 that it was cancelling its fuel processor research pro-gram for automotive applications. It is now accepted that automotive fuel cell systems will run on pure H2 as a fuel and, as a consequence, emphasis on developing effective H2 storage materials is much higher. Because the use of reformate as fuel is still favored for small, stationary fuel cell systems, the need to achieve reformate tolerance is still critical.

The definition of reformate tolerance is that, compared to running on pure H2, a fuel cell stack can run on reformate and show no change in perfor-mance, apart from that expected for dilution effects (of H2 due to CO2, N2, H2O). This requires the development of reformate-tolerant anode catalysts capable of tolerating the remaining levels of CO and CO2 in the fuel feed.


Wilkinson and Thompsett reviewed the area of reformate tolerance in 1995 and showed that a combination of PtRu catalysts and the use of an air bleed gave practical CO tolerance for CO levels below 100 ppm.125

The use of a separate gas-phase selective oxidation layer positioned under-neath the electrocatalyst layer was also shown to offer significant durability benefits when operating with an air bleed. They also highlighted the need to consider the effect of CO2 in the fuel because it has been shown by a number of workers that CO2 can reduce on Pt at low potentials to give CO, leading to poisoning. More recently, Tada, Inoue, and Yamamoto reported on the devel-opment of PtRu catalysts for reformate tolerance and claimed that Ru-rich formulations appeared to give the best tolerance.126

1.10.1 Mechanistic Studies

How Pt alloy catalysts achieve CO tolerance has been much debated. Two mechanisms have been proposed:

Ligand effect: CO adsorption is lowered by alloying, thus decreas-ing CO coverage and increasing sites available for H2 adsorption/dissociation and oxidation.

Bifunctional effect: CO oxidized by the alloying element is effective at dissociating H2O and providing OH to react with CO adsorbed on Pt and thus decreasing CO coverage.

It is likely that both mechanisms are active and dependent on potential. At low potentials (<200 mV) on PtRu, the bifunctional mechanism is not active because Ru is unable to dissociate adsorbed H2O to produce OH. However, above 250 mV, this does occur and CO oxidation by adsorbed OH becomes the dominant reaction in achieving CO tolerance. This is strongly related to the use of PtRu as a MeOH oxidation catalyst because CO oxidation is also the rate-determining step for this reaction.

A recent theoretical study by Liu, Logadottir, and Nørskov showed that alloy-ing Ru and Sn with Pt has two effects: providing additional sites for H2O dis-sociation at lower potential than Pt and decreasing the CO adsorption energy on Pt, thus reducing CO coverage and increasing sites for H2 oxidation.127 PtSn was predicted to be more CO tolerant than PtRu; this has been confirmed experimentally.128 Hayden, Rendall, and South showed that promotional effect of Sn was due to an OH adsorption redox process at low potentials.129

The effect of CO2 poisoning has been well understood. Gut et al. showed that CO is produced by a Pt anode when fed with CO2/H2 mixtures, indicat-ing that the reverse water-gas shift reaction is active at cell operating tem-peratures of 70°C.130 As previously observed, this has low performance over and above that expected for dilution. Gut et al., Ball et al.,131 and de Bruijn et al. showed that PtRu anodes were less poisoned by CO2.132 Using in situ IR and mass spectroscopy measurements, Smolinka et al. showed that CO2


is reduced on Pt at H upd (under potential deposition) potentials to give adsorbed CO.133 The maximum coverage reached was only 0.45 of an equiva-lent CO monolayer and this dropped as the hydride coverage dropped. The low CO coverage resulted in only modest poisoning of anode performance and it was concluded that CO2 reduction is kinetically limited through reac-tion with adsorbed hydride species rather than H2.

1.10.2 Improved Reformate-Tolerant Catalysts

In recent times, efforts have been made to optimize PtRu tolerance through the addition of third and fourth metals, as well as to identify alterative Pt-based catalysts with much greater reformate tolerance, particularly at much higher CO levels. Many of the reported studies are concerned with CO rather than reformate tolerance, and few long-term stability measurements have been reported.

1.10.2.1 PtRu Variants

As with the development of MeOH oxidation catalysts (of which PtRu is also the practical catalyst of choice), numerous attempts have been made to improve PtRu catalysts with third or fourth metals. Common additives are Sn, Mo, and W. However, few studies appear to show significant improve-ments in CO tolerance. These conclusions are supported by depositing Ru, Sn, and Mo onto Pt(111) and (533) single-crystal surfaces, as reported by Massong et al.134 Whereas Ru had an overall effect promoting CO oxidation over the whole surface (although Sn and Mo did promote CO oxidation to very low potentials), the effect was localized, with the majority of adsorbed CO oxidized as potentials similar to pure Pt. Co-depositing Sn and Ru onto Pt gave only a superposition of the binary surface responses indicating no synergy.

One of the most interesting studies reported has been a combined com-putational and high-throughput screening of PtRuM ternaries by Strasser et al.15 A series of Pt-containing ternaries was defined (with Ru, Co, Ni, and W), prepared as thin films by rf (radio frequency) sputtering, and evaluated for MeOH oxidation (to assess CO oxidation activity). Theoretical calcula-tions were performed on similar model structures and a series of PtRuM (M Fe, Co, Ir, Rh, Sn) ternaries was predicted to have better CO tolerance than PtRu. The experimental screening found that the PtRuCo (20:20:60) gave the best activity for MeOH oxidation and stood out from the other PtRuM ternaries studied.

An alternative approach in structuring a PtRu catalyst for CO tolerance has been reported by Brankovic, Wang, and Adzic,135 who prepared core-shell Pt-Ru catalysts with submonolayer coverage of Pt (10–50%). In the pres-ence of 100 ppm CO/H2, the 10% ML Pt catalyst showed much higher CO tolerance than a conventional PtRu catalyst with three times less Pt loading.


The 10% Pt ML core-shell catalysts were tested in an MEA and gave satisfac-tory performance with 10 ppm CO in H2 with an electrode loading of 0.018 mg Pt cm–2.136 A similar surface modification of a Pt core with Ru has been investigated by Crabb et al.137 They reported that, on reduction, a surface PtRu alloy formed and showed similar CO tolerance to a conventional nano-particle PtRu alloy.

1.10.2.2 PtMo Catalysts

Of the alternative Pt formulations, the PtMo system has been the most stud-ied in recent years. Work on bulk Pt3Mo alloys by Grgur, Markovic, and Ross showed similar CO tolerance to PtRu in the presence of H2.138 This tolerance was correlated with the ability of PtMo to oxidize CO at potentials as low as 0.05 V. However, unlike Ru but similar to Sn, the Mo appeared to oxidize CO just at neighboring Pt sites, with the majority of CO oxidized at potentials typical of pure Pt. The surface Mo atoms were found to be oxidized even at 0.0 V. Therefore, it was postulated that H2O dissociation to form OH was mediated by a Mo(IV)/(VI) couple. Carbon-supported PtMo catalysts were reported to have better CO tolerance than PtRu in MEA testing up to CO concentrations of 100 ppm.139

However, Ball et al.131 reported that the presence of CO2 in the H2 feed resulted in greater performance loss than 100 ppm CO. Testing on full refor-mate mixtures showed that PtMo anodes offered no advantages over PtRu anodes. Later work showed that PtMo anodes produced CO with H2/CO2 mixtures at low overpotentials, indicating that PtMo was active for the reverse water-gas shift reaction.140 Ball and Thompsett also investigated the CO/CO2 tolerance of PtMo at higher CO concentrations (up to 1,000 ppm) and found that MEA performance was similar to PtRu with 100 ppm CO/H2.141

The opposite CO/CO2 tolerance and intolerance of PtRu and PtMo led to the development of an electrochemical bilayer with separate PtMo and PtRu layers. The PtMo electrochemically oxidized the CO of the reformate to a level where it could be tolerated by the PtRu layer. It was shown that the bilayer could operate using a reformate containing 5,000 ppm CO and give similar performance to a PtRu anode operating on a 40 ppm CO reformate142 (see Figure 1.19). Commercial application has been hampered by the medi-um-term instability of the PtMo catalyst. Lebdeva and Jannsen found that cycling a PtMo catalyst in acid and MEAs showed loss of Mo and migration of Mo to the cathode, leading to loss in CO tolerance.143

Although great efforts have made to improved reformate-tolerant cata-lysts, no intrinsic reformate-tolerant catalysts have yet been discovered. The PtMo system appears to offer the greatest possibilities, especially at higher CO concentrations and at higher temperatures. However, the corrosion sen-sitivity of Mo over time needs to be addressed before these catalysts become practical systems.


1.11 Reformate-Tolerant Catalyst Stability

1.11.1 Ru and Mo Stability

Although durability studies of PEFC stacks and MEAs have increased, relatively few reports have been offered on the stability of anode catalysts operating on reformate. An early study by Ballard Power Systems showed that MEAs with PtRu anodes showed poor stability when operating with 40 ppm CO containing reformate and a 3% air bleed over 1,000 h. Stability was improved by adding an extra gas-phase catalyst layer to the anode. More recently, Cleghorn and colleagues at W. L. Gore showed a 26,300 h durability test of an MEA with a PtRu anode; the first 3,000 h were operated on 50 ppm CO reformate and 4% air bleed. Postmortem analysis showed no evidence of Ru migration from the anode to the cathode.144

A similar study by workers at Osaka Gas Company showed that MEAs with PtRu anodes operating on reformate up to 47,000 h showed significant degradation of CO-tolerant properties, although no MEA/catalyst postmor-tems were reported.145 The stability of PtMo catalysts under limited dura-bility conditions has been reported by Ball.142 Operating on 5,000 ppm CO containing reformate, a PtMo/PtRu bilayer anode showed significant loss in

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FIGURE 1.19Polarization curves of a PtMo/PtRu bilayer anode (0.45 mg (Pt) cm–2) operating on 63% H2, 23% CO2, 14% N2, 5,000 ppm CO reformate, and a PtRu anode (0.25 mg (Pt) cm–2) operating on 75% H2, 25% CO2, 40 ppm CO reformate. 80 C, 200 kPa, reformate/air stoichiometry 1.5/2. (Crown copyright 2004.)


anode performance at constant current over 140 h, which has been attributed to loss of Mo from the anode.

Clearly, further work is necessary to elucidate the effect of long-term opera-tion on PtRu and PtMo durability. However, given the length of experiments necessary to show such effects, the need to establish accelerated durability protocols for reformate anodes is clear.

1.11.2 Cell Reversal Tolerance

As well as stability under load conditions under normal cell operation, anode catalysts need to show stability under a range of transient operat-ing events. One of the most demanding is stability under conditions of low fuel flow or starvation. Under this condition, if insufficient hydrogen is present to sustain the current load, the anode potential will rise to that required to oxidize water (see Figure 1.20). Knights et al. showed that anode potentials can reach values above 1.2 V and potentials where car-bon is oxidized (>1.4 V).146 Taniguchi et al. showed that, under cell reversal conditions, PtRu/C anodes lost Ru through dissolution and showed particle agglomeration with increasing cell reversal times.147 Various approaches have been suggested to minimize the effects of cell reversal; one successful solu-tion has been the incorporation of water oxidation catalysts into the anode catalyst layer. Ralph, Hudson, and Wilkinson showed that the most effective

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material in their testing was Ir-doped RuO2, which was stable under cell reversal conditions for over 26 h.148

1.12 MeOH Oxidation Catalysts

The PtRu bimetallic system has been the catalyst of choice for MeOH oxida-tion in acid electrolytes since its discovery by workers at Shell in the early 1960s.149 In practice, PtRu lowers the overpotential for MeOH oxidation by >200 mV compared to pure Pt. The MeOH oxidation reaction on Pt and PtRu is probably the most studied reaction in fuel cell electrocatalysis due to its ease of study in liquid electrolytes and the many possible mechanistic path-ways. In recent years, the deposition of PtRu particles onto novel carbon sup-ports and the novel PtRu particle preparation routes have proved popular as a means to demonstrate superiority over conventional PtRu catalysts.

A number of recent reviews of DMFC technology are available. See those by McNicol, Rand, and Williams for earlier developments of catalysts for DMFC,149 Thomas et al. for cathode catalyst development at LANL,150 and Liu et al. for a summary of anode catalyst preparation and support development.151

1.12.1 Mechanistic Advancements

Study of the mechanism of MeOH oxidation over Pt and PtRu surfaces has recently been given new insights using a combination of experimental and theoretical approaches. The use of electrochemically linked mass spectros-copy techniques (e.g., differential electrochemical mass spectroscopy—DEMS) has allowed the quantification of the MeOH oxidation reaction in terms of comparing CO2 yields with electrons passed. In addition, detection and quantification of reaction intermediates has also been demonstrated. In addition, use of theoretical techniques such as DFT has allowed calculation of adsorbate energies, probing reaction pathways, and activation of H2O to provide active OH species.

The general mechanism of MeOH on Pt and PtRu is well established. First, MeOH is adsorbed and subjected to multiple dehydrogenation steps to give adsorbed CO. This dehydration step is known to occur at low potentials. The adsorbed CO is then oxidized by active OH species produced by the dissocia-tion of H2O. This is the potential-driven rate-determining step because OH formation does not occur on Pt until higher potentials. The addition of Ru pro-motes the reaction because it is able to produce OH species at lower potentials. This promotional effect is known as the “bifunctional” mechanism:

CH3OH CH3OHads

CH3OHads COads 4H 4e–


H2O H2Oads

H2Oads OHads H e–

COads OHads CO2 2H 2e–

This is known as the indirect mechanism because it requires the removal of CO. The direct mechanism goes through formaldehyde (HCHO) and formic acid (HCO2H) intermediates and allows for desorption at each step.

Modeling of PtRu surfaces has focused on the CO oxidation reaction. Koper, Shubina, and van Santen used DFT methods to model CO and OH adsorption energies on Pt, Ru and Pt2Ru, and PtRu2 surfaces.152 It was found that alloying Ru with Pt weakened and strengthened CO and OH adsorption on Pt and Ru, respectively. Desai and Neurock investigated the role of H2O on OH forma-tion and resulting CO oxidation.153 They found that the presence of hydration favored the formation of OH on Ru and induced the adsorption and activation of H2O on neighboring Pt sites. The surface OH continues to diffuse across the surface via proton transfer, allowing CO oxidation remote from the Ru sites.

The use of DEMS has allowed quantification of the MeOH oxidation reac-tion on different catalysts. Using Pt/C, Jusys and Behm showed that nine electrons per CO2 molecule are produced compared to the theoretically expected six electrons, together with detectable quantities of methyl formate (HCOOCH3).154 The HCOOCH3 is formed by reaction of MeOH and HCO2H. Therefore, it was concluded that up to one-third of the current was being used to form side products such as HCHO and HCO2H. Similar work was performed on unsupported PtRu catalysts and the conversion efficiency of MeOH to CO2 was found to be very close to six electrons per CO2 molecule. Only small amounts of HCOOCH3 were found with these catalysts.155

The kinetics of MeOH oxidation of a 1:1 PtRu in an MEA has been well established by Vidakovic, Christov, and Sundmacher.156 At low overpoten-tials, the MeOH oxidation reaction was found to be zero order in MeOH con-centration, indicating that CO oxidation is the rate-determining step. A Tafel slope of 50–60 mV dec–1 was found at 60°C. At higher overpotentials, positive reaction orders were found, suggesting that MeOH adsorption becomes rate determining. An activation energy of ~55 kJ mol–1 was found; this agrees well with the values found for similar bulk PtRu electrodes.

1.12.2 PtRu Variants

There have been many reports of variants of PtRu based on novel preparation chemistry or novel support materials showing superior activity to commer-cially available PtRu catalysts. These have been recently reviewed by Liu et al.151 One interesting feature of this work is that the Pt:Ru atomic ratio used has been fixed at 1:1 (e.g., Chu and Gilman157 and Takasu et al.158). However, this ratio dis-agrees with the optimal ratios determined from bulk PtRu alloys.


Gasteiger et al. determined that at 25°C the optimal surface ratio of Pt:Ru for MeOH oxidation in 0.5 M H2SO4 is approximately 90:10, with this surface showing a 30-fold improvement over bulk Pt.159 Increasing the temperature to 60 C changed the optimal surface composition to approximately 70:30.160 Somewhat similar results were found for deposition of Ru onto single and polycrystalline Pt surfaces.161 Characterization of the surfaces of practical PtRu particles has not been widely reported. Bock, MacDougall, and Le Page prepared a series of unsupported PtRu powders (9–46% Ru) and character-ized their surface composition by XPS.162 Similar values to the bulk composi-tions were found.

The development of techniques able to deposit controlled amounts of Ru onto the surface of Pt particles has allowed probing of these ratios with practical catalysts. Lee and Bergens deposited Ru onto Pt Black gauzes up to 3.5 ML equivalents using the hydrogenation of a Ru organometallic com-plex.163 MeOH oxidation activity was the highest at a coverage of 0.05ML Ru at room temperature. This work was extended to depositing onto Pt black and Pt/C particles and tested as DMFC MEAs at higher temperatures.164 At 60°C, the optimal Ru coverage was approximately 33%, while at 90°C, a broad range from 30 to 60% Ru gave similar and optimal performance. Similar results were obtained by Fachini et al. with the deposition of Ru onto Pt/C via a carbonyl precursor.165 Waszczuk et al. deposited Ru onto Pt black particles via spontaneous deposition and showed that a surface cover-age of 0.4–0.5 gave the maximum MeOH oxidation at room temperature.166

Although bulk- and surface-decorated samples agree broadly in terms of optimal Pt:Ru surface ratios for MeOH oxidation, there is less agreement with practical PtRu catalysts, although the data are sparse. This would sug-gest that PtRu particles show Pt-segregated surfaces as predicted by theoreti-cal calculations.

Increasingly, the investigation of PtRu containing ternary and quaternary catalysts has been reported with the aim to improve the MeOH oxidation activity of PtRu. Ley et al. showed that adding Os to PtRu at concentra-tions close to its solubility limit (10 at%) showed modest improvements in activity167 (see Figure 1.21). Using a rapid combinatorial screening method, Reddington et al.168 and Gurau et al.169 showed that adding small amounts of Ir to PtRu improved activity still further, with PtRuOsIr (47:29:20:4) showing an overpotential 150 mV lower than PtRu (50:50).

More recent work has focused on the identification of alternative ternary PtRuM formulations with enhanced activity. Strasser et al. used a thin-film approach to screen 64 PtRuM formulations and found that PtRuNi and PtRuCo showed eight times the activity of PtRu (60:40).15 In particular, PtRuCo (20:20:60) was found to be significantly higher in activity than any other alloy. Also, a series of PtWNi and PtWCo formulations were found to show similarly high activity. Other workers have claimed similar activ-ity enhancements for PtRuNi,170 PtRuFe,171 and PtRuIr.172 In addition, sev-eral reports have indicated that the addition of phosphorus to PtRu has


enhanced MeOH oxidation activity. Nitani et al.173 and Xue et al.174 showed that the addition of P to PtRu causes a reduction in PtRu particle size and an increase in activity.

Although many claims of improved activity from PtRu-containing ternary and quaternary formulations have been offered, none have yet been devel-oped commercially. One key feature that has yet to be studied is the stability of the additional elements added. Considering the reported poor stability of Ru at higher potentials, it is essential that these extra elements do not lead to extra instability.

1.12.3 Alternative Pt Catalysts

In the search for catalyst formulations superior to PtRu, many alternative Pt binary alloys have been investigated. In recent years, strong interest in PtOs alloys and bimetallics has been shown.175 Although Os does promote MeOH oxidation on Pt, it is somewhat less active than Ru, apart from high overpotentials (>0.50 V), where Os is less susceptible to overoxidation com-pared to Ru.176,177

Other combinations of Pt with precious metals have been studied. Platinum/rhodium was found to be only slightly more active than pure Pt, although adding Rh to PtRu (5:4:1) was found to have better activity than PtRu at higher currents when tested as an MEA.178 Similarly, PtIr has been investigated and

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shown to have much higher activity than PtRu, although with no shift in the onset of MeOH oxidation.179 PtSn formulations have long been studied for MeOH oxidation activity and, although activities similar to those of PtRu have been reported, this has not translated into practical catalysts.

Honma and Toda studied the temperature dependency of MeOH oxida-tion of a PtSn alloy and found that the onset of oxidation occurs at 200 mV lower voltage than with Pt.180 The stability of Pt3Sn catalysts as a function of cycling voltage has been studied. Liu et al. found that the voltammetric pro-file of Pt3Sn/C was relatively stable on cycling to ~0.75 V; however, on cycling to ~1.20 V, the profile became similar to Pt after 100 cycles.181 Interestingly, other alloys of Pt with p-block elements have also been recently investigated. Although PtBi intermetallics show only modest activity for MeOH oxida-tion, PtPb alloys have been shown to have superior activity to PtRu in liquid electrolytes182,183 (see Figure 1.22). The addition of W oxide species to Pt has also been studied for the effect on MeOH oxidation.

The addition of WO3 via electrodeposition was studied by Jayaraman et al., who found that it showed modest improvements in activity over Pt.184 Adzic and Marinkovic claimed that adding NiWO4 or CoWO4 to Pt gave better activ-ity than Pt, although inferior to PtRu.185 One intriguing report concerned the use of perovskite-based catalysts as MeOH oxidation catalysts in acid media. Lan and Mukasyan evaluated a series of ABO3 oxides (A Ba, Ca, Sr, La; B Fe, Ru) and found that the Ru-containing oxides gave modest activities in a parallel channel MEA screening array.186 However, adding LaRuO3 to Pt gave higher steady-state MeOH oxidation currents at 0.4V than PtRu.

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–100

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PtPb/C

PtRu/C

0.4 0.6

E(V) vs Ag/AgCl

0.8 1.0

FIGURE 1.22Cyclic voltammograms of PtPb and PtRu catalysts in 0.1 M H2SO4, 0.5 M MeOH at room tem-perature. (Reprinted with permission from S. Maksimuk et al., Journal of the American Chemical Society, 129, 8684 (2007). Copyright 2007 American Chemical Society.)


Although no alternative system has yet shown consistently higher activities than PtRu, recent research does point to some interesting possibilities (e.g., PtPb intermetallics). However, as with any practical electrocatalyst formulation, the stability under realistic operating conditions must be evaluated before formula-tions can be considered as alternatives to the well-established PtRu system.

1.13 MeOH Oxidation Catalyst Stability

With the development of portable DMFCs toward commercial applications, interest in the durability of DMFC MEAs has increased.187 Of the various degradation mechanisms identified, the loss of Ru from PtRu anode catalysts has been found to have a significant effect on MEA stability. Piela et al. first showed that Ru could be removed from PtRu black anode catalysts under a range of DMFC operating conditions, including open circuit and cell reversal modes188 (see Figure 1.23). They showed that Ru was transported across the membrane and deposited on the cathode, inhibiting oxygen reduction activ-ity. More detailed studies by Jeon et al. and Chen et al. showed that Ru loss was restricted at lower current densities (i.e., lower anode overpotentials) and was accelerated at higher current densities or at short circuit.189,190

One further degradation mode related to catalysis is a consequence of oper-ating at low current densities typical of portable power application. Under these conditions, overoxidation of the Pt cathode catalyst occurs, reducing cathode and overall MEA performance. Zelenay has shown that starving the cathode of air flow lowers the cathode potential to low values, causing reduc-tion of Pt oxides and restoring cathode activity.187

It is clear that to guard against DMFC performance degradation, improve-ments in PtRu stability are required.

0.8 40 mVCathode Loss

0.7

0.6

0.5

0.4

0.3

iR-C

orr

ect

ed C

ell

Vo

ltag

e (V

)

0.2

0.00 0.05 0.10

Current Density (A cm–2)

0.15

DMFC DMFC (Extreme Ru Contamination)

0.20

Ru-contaminated cathodePt-Ru cathode

Ru-free cathode

0.8>200 mV

Cathode Loss0.7

0.6

0.5

0.4

0.3

iR-C

orr

ect

ed C

ell

Vo

ltag

e (V

)

0.2

0.00 0.02 0.04 0.06


0.08 0.10

Extremely Ru-contaminated cathodeRu-free cathode

FIGURE 1.23DMFC performance losses caused by average (left) and extreme (right) contamination of Pt cath-odes by crossover Ru. (P. Zelenay, Electrochemical Society Transactions, 1, 483 (2006). Reproduced by permission of The Electrochemical Society.)


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61

2Catalyst Layers and Fabrication

Zhong Xie, Chaojie Song, David P. Wilkinson, and Jiujun Zhang

CONTENTS

2.1 Introduction .................................................................................................. 622.2 Catalyst Layer Components and Their Corresponding Functions ......64

2.2.1 Overview of Catalyst Layer Components and Functions ..........642.2.2 Properties of the Catalyst Layer.....................................................66

2.2.2.1 Catalyst Loading and Catalyst Utilization ....................662.2.2.2 Nafion Loading..................................................................682.2.2.3 Hydrophobicity and Hydrophilicity ..............................682.2.2.4 Porosity ............................................................................... 692.2.2.5 Ionic (Proton) Conductivity and Electronic

Conductivity ...................................................................... 702.3 Types of Catalyst Layers ............................................................................. 70

2.3.1 CCGDL .............................................................................................. 702.3.1.1 Uniform CCGDL ............................................................... 702.3.1.2 Gradient CCGDL ............................................................... 712.3.1.3 Dual-Bound Composite Catalyst Layer .........................75

2.3.2 CCM ................................................................................................... 762.3.2.1 Conventional CCM ........................................................... 762.3.2.2 Nanostructured Thin-Film Electrode ............................77

2.3.3 Novel Structural Catalyst Layer ....................................................772.3.3.1 CNT-Based Catalyst Layer ...............................................772.3.3.2 Columnar Oxide Supported Catalyst Layer ..................772.3.3.3 Nanowire-Based Three-Dimensional Hierarchical

Core/Shell Catalyst Layer ................................................ 792.3.3.4 Self-Supported Catalyst Layer .........................................802.3.3.5 Catalyst Layer with Additives .........................................802.3.3.6 Catalyst Layer with Novel Ionomers .............................. 81

2.4 Catalyst Layer Fabrication .......................................................................... 812.4.1 First-Generation Catalyst Layer Fabrication ................................ 81

2.4.1.1 Pt Black Catalyst Layer Fabrication ................................ 812.4.1.2 PTFE-Bound Catalyst Layer Fabrication Using

Supported Catalyst ........................................................... 82


2.1 Introduction

The concept of the catalyst layer (CL) can be traced back to the 1840s, when Grove found that the three-phase boundary was important in improving fuel cell reaction rate [1]. The first practical gas diffusion electrode was developed by Schmid in 1923 [1,2], significantly increasing the electrode active surface area and thus representing a revolutionary improvement in fuel cell electrode technology. Since then, great progress has been made in fuel cell CL design and performance optimization in terms of both technological advancement and commercialization.

Figure 2.1 shows a schematic structure of the fuel cell membrane electrode assembly (MEA), including both anode and cathode sides. Each side includes a catalyst layer and a gas diffusion layer. Between the two sides is a proton exchange membrane (PEM) conducting protons from the anode to the cathode.

The catalyst layer is located between the PEM and the gas diffusion layer (GDL). Protons transfer between the CL and the PEM, and electrons transfer between the catalyst layer and the GDL. Both require good interfacial contact.

In a PEM fuel cell, the CDLs are where the electrochemical reactions occur for electric power generation. For example, for H2/air (O2) PEM fuel cells, the reactions occurring at the anode and cathode catalyst layers are as follows:

Anode: H2 2H+ + 2e– (2.1)

Cathode: O2 + 4H+ + 4e– H2O (2.2)

For both reactions to occur, a three-phase boundary is required where the reactant gas, protons, and electrons react at the catalyst surface. The CLs should be able to facilitate transport of protons, electrons, and gases to the catalytic sites. Under normal PEM fuel cell operating conditions (≤80°C), the reactants are gaseous phase H2 and O2 (from air), and the product is water, primarily in the liquid phase. Water removal is a key factor affecting catalyst

2.4.2 Thin-Film Catalyst Layer Fabrication ...........................................832.4.2.1 Ink-Based Catalyst Layer Fabrication .............................832.4.2.2 In Situ Catalyst Layer Fabrication ...................................862.4.2.3 Other Methods .................................................................. 89

2.5 Catalyst Layer Optimization ...................................................................... 912.5.1 Catalyst Layer Composition Optimization .................................. 92

2.5.1.1 Modeling and Simulation to Optimize Catalyst Layers .................................................................................. 92

2.5.1.2 Experimental Studies on Optimization of CLs ............ 932.5.2 Catalyst Layer Microstructure Optimization .............................. 95

2.6 Prospects and Conclusion .......................................................................... 96References ............................................................................................................... 97

Catalyst Layers and Fabrication 63

layer performance. The presence of excess water in the catalyst layer can block gas transport, leading to reduced mass transfer and decreased fuel cell performance. On the other hand, a lack of water results in decreased proton conductivity of the membrane and the ionomer in the catalyst layers, leading to decreased fuel cell performance. The basic requirements for a CL include:

1. A large number of three-phase boundary sites; 2. Efficient transport of protons from the anode catalyst layer to the

cathode catalyst layer; 3. Facile transport of reactant gases to the catalyst surface; 4. Efficient water management in the catalyst layers; and 5. Good electronic current passage between the reaction sites and the

current collector.

The properties and composition of the CL in PEM fuel cells play a key role in determining the electrochemical reaction rate and power output of the system. Other factors, such as the preparation and treatment methods, can also affect catalyst layer performance. Therefore, optimization of the catalyst layer with respect to all these factors is a major goal in fuel cell development. For example, an optimal catalyst layer design is required to improve catalyst

ElectricallyConductive

Fibers

ElectricallyConductive

Fibers

e–

e–

H+

H2O

H2O

H2

O2

CarbonSupported

Catalyst

CarbonSupported

Catalyst

GOLGOL PEM Cathode

Catalyst

Layer

Anode

Catalyst

Layer

ProtonConducting

Media

FIGURE 2.1Schematic structure of a fuel cell membrane electrode assembly (MEA), including both anode and cathode catalyst layers. (Based on Lister. S. and McLean, G. Journal of Power Sources 2004; 130:61–76. With permission from Elsevier.)


(platinum [Pt]) utilization and thereby reduce catalyst loading and fuel cell cost.

In this chapter, we will focus on several important aspects of the PEM fuel cell catalyst layer, including the CL components and their corresponding functions, the types of catalyst layers, and catalyst layer fabrication and optimization.

2.2 Catalyst Layer Components and Their Corresponding Functions

2.2.1 Overview of Catalyst Layer Components and Functions

Two main types of catalyst layers are used in PEM fuel cells: polytetrafluo-roethylene (PTFE)-bound catalyst layers and thin-film catalyst layers [3]. The PTFE-bound CL is the earlier version, used mainly before 1990. It contains two components: hydrophobic PTFE and Pt black catalyst or carbon-sup-ported Pt catalyst. The PTFE acts as a binder holding the catalyst together to form a hydrophobic and structured porous matrix catalyst layer. This porous structure can simultaneously provide passages for reactant gas transport to the catalyst surface and for water removal from the catalyst layer. In the CL, the catalyst acts as both the reaction site and a medium for electron conduc-tion. In the case of carbon-supported Pt catalysts, both carbon support and catalyst can act as electron conductors, but only Pt acts as the reaction site.

In earlier research, no ionomer was used with this type of catalyst layer, and the Pt loading was very high, up to 4 mg/cm2. Later it was found that the addition of Nafion ionomer (by brush coating or spraying) to the PTFE-bound CL could lead to a 10-fold reduction in catalyst loading [4,5]. Figure 2.2 shows the effect that Nafion ionomer in the catalyst layer has on fuel cell per-formance. The Nafion ionomer provides proton conductive paths for proton migration to or from the catalyst and hence increases the number of active catalyst sites that meet the three-phase boundary requirement.

The PTFE-bound CL plays a dual role as both a gas diffusion layer and a catalyst layer, where gas transport, water removal, and electrochemical reac-tions occur in the same layer. For this type of structure without a proton-conducting ionomer, the Pt utilization is low. Although Nafion impregnation can reduce Pt loading, the reduction is limited, and a further decrease in catalyst loading is difficult without compromising cell performance. In addi-tion, the Nafion impregnated PTFE-bound CL has some disadvantages. For example, variation in Nafion impregnation depth could result in some areas not being fully impregnated and others being overimpregnated. In the latter, Nafion might penetrate to the substrate, leading to inefficient utilization of the impregnated ionomer and, at the same time, introducing an unnecessary transport barrier to gas diffusion. Evidently, achieving an adequate ionomer impregnation depth in the catalyst layer is difficult.


To overcome these disadvantages, a thin-film CL technique was invented, which remains the most commonly used method in PEM fuel cells. Thin-film catalyst layers were initially used in the early 1990s by Los Alamos National Laboratory [6], Ballard, and Johnson-Matthey [7,8]. A thin-film catalyst layer is prepared from catalyst ink, consisting of uniformly distributed ionomer and catalyst. In these thin-film catalyst layers, the binding material is not PTFE but rather hydrophilic Nafion ionomer, which also provides proton conductive paths for the electrochemical reactions. It has been found that the presence of hydrophobic PTFE in thin catalyst layers was not beneficial to fuel cell performance [9].

It is well known that Nafion ionomer contains both hydrophobic and hydrophilic domains. The former domain can facilitate gas transport through permeation, and the latter can facilitate proton transfer in the CL. In this new design, the catalyst loading can be further reduced to 0.04 mg/cm2 in an MEA [10,11]. However, an extra hydrophobic support layer is required. This thin, microporous GDL facilitates gas transport to the CL and prevents catalyst ink bleed into the GDL during applications. It contains both carbon and PTFE and functions as an electron conductor, a heat exchanger, a water removal wick, and a CL support.

In practice, the catalyst used in the thin-layer CLs for both anode and cathode is carbon-supported Pt catalyst (Pt/C) or Pt alloy, such as PtRu/C, although nonsupported catalysts can be used. In terms of the overall elec-trode structure, an electrode with a thin CL generally contains three layers:

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1

0 20 40 60 80 100 120 140 160 180 200

Current Density (mA/cm2)

Cel

l P

ote

nti

al (

V) A

B

C

FIGURE 2.2Polarization curves for H2/O2 fuel cells at 50 C, 1 atm pressure. Curve A: Nafion impregnated (brush coated) PTFE-bound electrode (0.35 mg/cm2 Pt loading); curve B: PTFE-bound catalyst layer (Pt loading: 4 mg/cm2); curve C: PTFE-bound electrode (Pt loading: 0.35 mg/cm2). (Based on Ticianelli, E. A. et al. Journal of the Electrochemical Society 1988; 135:2209–2214. By permission of The Electrochemical Society.)


carbon backing (paper), a thin carbon/PTFE microporous gas diffusion layer, and a thin-film ionomer/catalyst layer.

2.2.2 Properties of the Catalyst Layer

Because thin-film catalyst layers are the most commonly used in today’s PEM fuel cell technology, we will mainly focus in this section on the properties of the thin-film catalyst layer as well as its effect on fuel cell performance.

2.2.2.1 Catalyst Loading and Catalyst Utilization

In general, higher Pt loading leads to better performance, but it also results in higher cost, which is one of the key factors hindering PEM fuel cell com-mercialization. Therefore, one of the major goals in PEM fuel cell develop-ment is to reduce Pt loading without compromising fuel cell performance and durability. In terms of performance, great progress has been made in total Pt loading reduction, from several milligrams per square centimeter to 0.01–0.02 mg/cm2 in the laboratory [10,11]. Unfortunately, with such a low Pt loading, durability is an issue. At the present stage of technology, optimal Pt loading in terms of both practical fuel cell performance and durability is about 0.3 mg/cm2.

There is still significant room for Pt loading reduction because not all of the Pt catalyst in the CL is electrochemically utilized. Therefore, a parameter, called Pt utilization, is used to describe this CL property. The Pt utilization (Ptutilization) can be calculated according to the following equation:

Pt

ECAECAutilization

measured

calculated

100%

(2.3)

whereECAmeasured represents the electrochemical surface area of Pt measured

from the H2 adsorption/desorption peaks using cyclic voltammetry in either a half cell or a fuel cell

ECAcalculated represents the Pt surface area calculated from the Pt loading and the Pt catalyst particle size.

A typical cyclic voltammogram obtained from a fuel cell cathode is shown in Figure 2.3. From this figure, the ECAmeasured can be obtained using the hydrogen desorption charge (Qdesp) or the adsorption charge (Qadsp) mea-sured from the cyclic voltammogram and the well-established standard value of 0.21 mC/cm2 for planar polycrystalline Pt [12,13] (here, cm2 is the actual Pt surface), based on the following equation:

ECA

Qmeasured

desp

0 211000

. (2.4)


whereQdesp is the charge under the hydrogen desorption peak (μC/cm2) (here,

cm2 is the geometric CL surface)The unit of ECAmeasured is cm2/cm2, indicating the actual Pt surface area per

geometric area.

The following equation is used to obtain ECAcalculated:

ECA

Ldcalculated

Pt6 104

(2.5)

whereLPt is the Pt loading in the catalyst layer (mg/cm2)d is the Pt particle size (average particle diameter) in nanometersr is the density of Pt (g/cm3)the unit of ECAcalculated is also cm2/cm2.

–50

–40

–30

–20

–10

0

10

20

30

0 0.2 0.4 0.6 0.8 1

Potential (V) vs NHE

Cu

rren

t D

ensi

ty (

mA

/cm

2)

Qdesp

Qadsp

FIGURE 2.3Cyclic voltammogram recorded for the cathode of a membrane electrode assembly with a cath-ode Pt loading of 0.4 mg/cm2 in a fuel cell operated at 80 C and 100% RH. Cathode: N2; anode: H2; scan rate: 50 mV/s. (Unpublished data from the authors.)


For PTFE-bound type CLs, the Pt utilization is greater than 40%, reaching up to 80% [3,14,15]. Sasikumar, Ihm, and Ryu [14], for example, reported 52% Pt utilization and Li and Pickup reported 76% [15]. A variety of parameters affect Pt utilization and the CL preparation method, and will be discussed in the following sections.

2.2.2.2 Nafion Loading

Nafion content in the CL can significantly affect MEA performance by influ-encing gas permeability, ionic resistance, and catalyst utilization [16]. For example, Li found that the ionic conductivity of the CL was highest with a Nafion loading of 0.9 mg/cm2 (31 wt%) for an electrode with a Pt loading of 0.4 mg/cm2 [15]. Antolini suggested an empirical equation to calculate the optimum Nafion loading (LNafion in mg/cm2) in the CL [17]:

L

LPNafion

Pt

Pt

56

(2.6)

where LPt is the Pt loading (mg/cm2) and PPt is the weight percent of Pt sup-ported on carbon.

Using this equation, the optimum Nafion loading is around 36 wt% for all Pt loadings. Sasikumar et al. [14] summarized the literature results and found that the optimum Nafion content is in the range of 30–36 wt% for the CL.

2.2.2.3 Hydrophobicity and Hydrophilicity

The wetting property of the catalyst layer is an important parameter that may have a large impact on performance by affecting the transport of water and gas. This property is often characterized by the contact angle of water drops on it. The smaller the contact angle is, the more hydrophilic is the CL; the larger the contact angle is, the more hydrophobic is the CL. Contact angle is usually measured by an optical contact measurement system such as the sessile drop method. Others, such as the Wilhelmy method and the capillary rise method, have also been developed.

In the sessile drop method, a small drop of water is placed on the target surface and the contact angle is measured [18]. This method has some dis-advantages on rough and porous surfaces, where the water droplets may penetrate into the bulk and spread over a hydrophilic surface. Yu et al. [19] developed a method employing an environmental scanning electron micro-scope, whereby dynamic formation of water droplets on a CL and their contact angles can be measured. In the Wilhelmy method, a planar piece of material (catalyst layer, gas diffusion electrode, etc.) is vertically dipped into the water and the weight of the material is measured using a sensitive balance [19]. However, this method is not suitable for catalyst-coated membranes.


The capillary rise method, originally developed by Neumann and co-workers [20], is based on the Wilhelmy plate gravimetric technique, where the capillary rise for a vertical plate is measured as a function of the plate’s immersion or retraction rate into or from the liquid. Liam and Wang [21] modified the method using an optical technique to record and measure the capillary meniscus height directly. The contact angle between the liquid and the substrate specimen has the following relationship with respect to the meniscus height [20,21]:

sin 1

2

2gh

(2.7)

whereq is the contact angleΔr is the difference between the densities of liquid and vaporg is the gravitational constanth is the meniscus heights is the liquid-gas surface tension of water.

Therefore, the contact angle can be calculated from the experimentally mea-sured meniscus height.

Contact angle measurements on the CL may also be useful in the characteriza-tion of catalyst layer degradation in a fuel cell. Yu et al. [19] found that the contact angle of a degraded CL became smaller compared to that of an unused catalyst layer, indicating more hydrophilic behavior accompanying degradation.

2.2.2.4 Porosity

The micropores in the catalyst layer are necessary for gas transport to the catalytic sites. The porosity of a CL is usually measured using a mercury porosimeter, in which the mercury is forced into all the pores of the CL under pressure. This pressure is inversely proportional to the pore size. The volume of mercury penetrating into the pores is measured directly as a func-tion of applied pressure, which reflects the pore size and the CL porosity [22]. The latter can be increased when pore-forming reagents are added during catalyst layer fabrication [3,23,24]. Fischer, Jindra, and Wendt [23] reported that the original porosity of a CL after fabrication was 35%, but when the pore formers were added, the porosity could become as high as 65% depend-ing on the pore formers used.

Unfortunately, cell performance is not proportional to catalyst layer poros-ity. In order to achieve maximum fuel cell performance, the CL should have an optimal porosity [24]. With higher catalyst layer porosity, the mass transfer rate increases, while the electron and proton transport rates decrease. Gamburzev and Appleby [25] documented fuel cell performance with pore formers in the CL and found that optimum pore-former content was about 33%.


2.2.2.5 Ionic (Proton) Conductivity and Electronic Conductivity

The ionic or proton conductivity of the catalyst layer is an important factor affecting fuel cell performance. The value of the proton conductivity is usu-ally determined by the Nafion ionomer loading. Proton conductivity in the CL can be measured using AC impedance spectroscopy. However, this mea-surement usually gives an average value over the whole CL, whereas, in real-ity, the proton conductivity is not uniform in both directions (i.e., in plane and through plane). For example, Li and Pickup [15] reported that the ionic conductivity decreased at locations further away from the membrane. At the catalyst layer/membrane interface, with a Nafion loading of 0.9 mg/cm2, the ionic conductivity was 3 mS/cm, while at 5 μm away from the membrane, the ionic conductivity was 1 mS/cm [15]. It is therefore necessary to develop a tool for conductivity mapping of the CL.

The electronic conductivity can also be measured using AC impedance and is found to be higher than the ionic conductivity. Saab, Garzon, and Zawodzinski [26] reported that the electronic conductivity of the CL was ~0.025–0.1 S/cm.

2.3 Types of Catalyst Layers

As discussed in Section 2.2, there are two main types of catalyst layers: PTFE-bound CLs and thin-film catalyst layers. Because the latter are almost always used in current work, we will focus only on different types of thin-film CLs in the following sections.

There are two main types of thin-film catalyst layers: catalyst-coated gas diffusion electrode (CCGDL), in which the CL is directly coated on a gas dif-fusion layer or microporous layer, and catalyst-coated membrane, in which the CL is directly coated on the proton exchange membrane. In the follow-ing sections, these catalyst layers will be further classified according to their composition and structure.

2.3.1 CCGDL

2.3.1.1 Uniform CCGDL

Uniform CCGDLs have a uniform distribution of Nafion and catalyst through and over the catalyst layer and are prepared by methods such as spraying or screen printing catalyst ink (an ultrasonicated, uniform mixture of catalyst, Nafion solution, and solvent) on an electrode substrate. Catalyst loading and Nafion loading can be controlled by the amount or composition of the ink applied. Although this type of CCGDL demonstrates decent performance, it is not optimized for the reactant gas distribution and water management gradients in the CL that occur in practical fuel cells between the inlet and outlet of the active area.


2.3.1.2 Gradient CCGDL

A gradient CCGDL has a nonuniform distribution of catalyst (catalyst gra-dient CCGDL), Nafion (Nafion gradient CCGDL), or both (dual-gradient CCGDL) in the CL. Because the cathode side is the limiting factor in PEM fuel cells (slow O2 reduction reaction kinetics and significant water management issues), the majority of studies are focused on the cathode CL. Wilkinson and St-Pierre have shown that significant gradients exist in practical fuel cells between the reactant inlet and outlet, resulting in nonoptimized per-formance over the active area [27].

2.3.1.2.1 Catalyst Gradient CCGDL

The gradient CCGDL can be designed according to two major directions: the through-plane gradient (z,z-direction) across the catalyst layer—from the membrane/catalyst layer interface to the catalyst layer/gas diffusion layer interface—and the in-plane gradient (x,y-direction) along the CL correspond-ing to the path from the reactant gas inlet to the outlet (Figure 2.4).

Antoine et al. [28] investigated the gradient across the CL and found that the Pt utilization was dependent on the CL porosity. In a nonporous CL, cata-lyst utilization was increased through the preferential location of Pt close to the gas diffusion layer; in a porous CL, catalyst utilization efficiency was increased through the preferential location of Pt close to the polymer electro-lyte membrane. In PEM fuel cells, the CL has a porous structure, and better performance is expected if higher Pt loading is used at preferential locations close to the membrane/catalyst layer interface.

This concept was proved by Kim et al. [29] using a dual-gradient CL design. Wilkinson and St-Pierre [27] presented the first use of in-plane gradient CLs

FIGURE 2.4Schematic diagram of the catalyst loading gradient: through plane (left) and in plane (right).


Inlet

Non-gradient0.3 mg-Pt/cm2 0.4 mg-Pt/cm2

0.34 mg-Pt/cm20.3 mg-Pt/cm2

0.25 mg-Pt/cm2

0.3

0.3

0.35

0.5

0.3

0.35

0.35

0.2

0.3

0.35

0.15

0.25

0.3

Outlet

5 cm

5 c

m

(a)

(b)

1.0

0.90.3 mg-Pt/cm2 (non-gradient)

Average Cathode Catalyst Loading

0.4 mg-Pt/cm2

0.34 mg-Pt/cm2

0.3 mg-Pt/cm2

0.25 mg-Pt/cm2

0.8

0.7

0.6Po

ten

tial

, V

0.5

0.4

0 200 400 600

Current Density, mA/cm2

800 1000 1200

FIGURE 2.5(a) Cathode catalyst loading distributions in a gradient electrode; (b) cell performance of non-gradient and gradient electrodes. (Reproduced from Prasanna, M. et al. Journal of Power Sources 2007; 166:53–58. With permission from Elsevier.)


in a variety of fuel cell hardware, demonstrating that these performed bet-ter than uniform catalyst layers. Prasanna et al. [30] also designed gradient catalyst layers for the oxygen reduction reaction in the gas inlet to outlet direction. At locations close to the gas inlet, O2 concentration was higher and low Pt loading was needed, while at the gas outlet, O2 concentration was lower and higher Pt loading was used. They observed that for average cata-lyst loadings of 0.25 mg/cm2 and 0.3 mg/cm2, the gradient CCGDLs showed better performance than their nongradient counterparts. The gradient struc-ture of the CL and the associated polarization curves of these electrodes are shown in Figure 2.5.

2.3.1.2.2 Nafion Gradient CCGDL

In the Nafion gradient CCGDL, unlike the catalyst gradient CCGDL, the gra-dient is usually in only one direction—that is, the through-plane direction of the catalyst layer. It is speculated that a gradient with higher Nafion con-tent at the membrane/catalyst layer interface and lower Nafion content at the CL/GDL interface should benefit proton migration and mass transport. Although contradictory results have been reported in the literature, theo-retical studies and some experimental results have shown that gradient CLs with higher Nafion loading at the membrane/catalyst layer give higher per-formance than those with uniform Nafion distribution [31,32].

Recently, Lee and Hwang [33] investigated the effect of Nafion loading and distribution on PEM fuel cell performance, and they found that a catalyst layer with Nafion ionomer on the surface (catalyst layer/membrane interface) exhib-ited better performance than a CL with Nafion inside. The best performance was obtained from a CL with a Nafion loading of 0.5 mg/cm2 inside the CL and 1.0 mg/cm2 on the surface (Figure 2.6). This can be explained by the dis-tribution of Nafion ionomer inside the CL playing a role in extending the elec-trochemical reaction zone more effectively when Nafion was loaded on the catalyst layer surface.

With respect to an in-plane Nafion gradient, Wu et al. designed a com-posite catalyst layer containing nonuniformly distributed Nafion and PTFE along the in-plane direction from reactant inlet to outlet [34]. In their design, for the area extending two-thirds of the way from the inlet, the Nafion load-ing and PTFE content were 0.29 and 0 mg/cm2, respectively; for the remain-ing third of the area close to the outlet, the Nafion loading and PTFE content were 0.6 and 0.37 mg/cm2, respectively. In this way, the latter area close to the outlet was more hydrophobic, facilitating water removal. Some enhanced performance in the mass transfer region was observed.

2.3.1.2.3 Dual-Gradient CCGDL (Catalyst and Nafion)

Frost et al. [35] anticipated the use of gradients in more than one component for the electrode (e.g., catalyst and Nafion) in the manufacture of electrodes by screen printing. Kim et al. [29] designed a dual CL as shown in Figure 2.7. Both anode and cathode contained two subcatalyst layers, each of which


Dual Catalyst

LayerAnode

Su

b-l

ayer

II

Su

b-l

ayer

II

Su

b-l

ayer

I

Su

b-l

ayer

I

Mem

bra

ne

Dual Catalyst

LayerAnode

FIGURE 2.7Schematic diagram of MEA using Nafion gradient catalyst coating method. (Reproduced from Kim, K. H. et al. International Journal of Hydrogen Energy 2008; 33:2783–2789. With permission from the International Association of Hydrogen Energy.)

Surface 1.0 mg/cm2 - inside none

Surface 1.0 mg/cm2 - inside 0.25 mg/cm2




10

0.9

0.8

0.7

0.6

0.5

0.4

Cel

l V

olt

age

(V)

0.3

0.2

0.1

0.00 200 400 600 800


1000 1200 1400 1600

FIGURE 2.6Cell voltage versus current density of the electrode with a Nafion content of 1.0 mg/cm2 on the surface and various Nafion contents inside the catalyst layer. (Reproduced from Lee, D. and Huang, S. International Journal of Hydrogen Energy 2008; 33:2790–2794. With permission from the International Association of Hydrogen Energy.)


contained a different catalyst and a different Nafion loading, as shown in Table 2.1. Figure 2.8 shows that the dual-gradient CL exhibited higher perfor-mance than the nongradient CL.

2.3.1.3 Dual-Bound Composite Catalyst Layer

Thin-film catalyst layers are usually hydrophilic, with no hydrophobic ingredi-ents added inside the CL. Although PTFE is generally unnecessary for thin-film catalyst layers, sometimes hydrophobicity may be required for better transport in the CL. Zhang et al. [11] designed a dual-bound composite CL that contained

TABLE 2.1

Gradient Pt and Nafion Loadings in MEAs

Pristine MEA

MEA 1 Sublayer

MEA 2 Sublayer

MEA 3 Sublayer

MEA 4 Sublayer

I II I II I II I II

Total Pt loading (mg) in catalyst layer (5 × 5 cm2)

10 7 3 7 3 7 3 7 3

Nafion loading in sublayer (wt%)

33 33 26.5 33 23 33 16.5 33 10

Sources: Kim, K. H. et al., International Journal of Hydrogen Energy 2008; 33:2783–2789; Lee, D. and Huang, S. International Journal of Hydrogen Energy 2008; 33:2790–2794. With per-mission from the International Association of Hydrogen Energy.

1800160014001200

Pristine MEA

MEA 1

MEA 2

MEA 3MEA 4

1000800


6004002000

0.2

0.3

0.4

0.5Vo

ltag

e (V

)

0.6

0.7

0.8

0.9

1.0

1.1

FIGURE 2.8Polarization curves for MEAs with dual gradients compared with a nongradient MEA. (Reproduced from Kim, K. H. et al. International Journal of Hydrogen Energy 2008; 33:2783–2789. With permission from the International Association of Hydrogen Energy.)


two layers: (1) a hydrophobic layer with PTFE as the binding material fabricated on the surface of the gas diffusion layer, and (2) a hydrophilic layer with Nafion as the binding material fabricated on top of the hydrophobic layer. This dual-bound composite CL is a combination of a PTFE-bound and a thin-film CL.

Zhang and Shi [36] found that the dual-bound composite catalyst layer exhibited higher performance than either a PTFE-bound CL or a thin-film CL, as shown in Figure 2.9. Optimization of the dual-bound CL showed that impregnation of Nafion between the two layers could lead to decreased cell performance [37]. Thus, the optimal structure for a dual-bound CL was a separate hydrophilic layer on top of a hydrophobic layer.

2.3.2 CCM

Catalyst layer ink can be deposited on gas diffusion layers to form a CCGDL, as discussed in the previous section. Alternatively, the catalyst ink can be applied directly onto the proton exchange membrane to form a catalyst-coated membrane (CCM). The most obvious advantage of the CCM is better contact between the CL and the membrane, which can improve the ionic connection and produce a nonporous substrate, resulting in less isolated catalysts. The CCM can be classified simply as a conventional CCM or as a nanostructured thin-film CCM.

2.3.2.1 Conventional CCM

The CCM was first developed in the 1960s [38]; it consisted of a Pt/PTFE mix-ture bonded on a membrane. This was similar to the PTFE-bound catalyst

1.61.41.2

PTFE-bound CL

Ionomer-bound CL

Dual-bound CL

1.00.8


0.60.40.20.0

0.4

Cel

l V

olt

age

(V)

0.6

0.8

1.0

FIGURE 2.9Polarization curves for a PEM fuel cell with different cathode catalyst layers. (Reproduced from Zhang, X. and Shi, P. Electrochemistry Communications 2006; 8:1229–1234. With permission from Elsevier.)


layer on a gas diffusion layer (e.g., carbon fiber paper). However, this type of CCM had high catalyst loading and low catalyst utilization. An early con-ventional CCM based on a Pt/Nafion mixture was developed by Wilson and Gottesfeld [6] at Los Alamos National Laboratory in the United States. They used a so-called decal method to prepare a thin-film CCM in which the cata-lyst ink was first applied to a Teflon blank and then transferred to the mem-brane by hot pressing.

Later it was found that the ink could be directly applied to the membrane [39]. However, for this technique, the membrane had to be converted to Na+ or K+ form to increase its robustness and thermoplasticity. With advancements in the technique, the total catalyst loading for the CCM could be reduced to 0.17 mg/cm2 without any compromise in cell performance [39]. Compared to CCGDL technology, the CCM approach seems the preferred method for CL fabrication.

2.3.2.2 Nanostructured Thin-Film Electrode

The nanostructured thin-film electrode was first developed at 3M Company by Debe et al. [40] and Debe [41], who prepared thin films of oriented crystal-line organic whiskers on which Pt had been deposited. The film was then transferred to the membrane surface using a decal method, and a nano-structured thin-film catalyst-coated membrane was formed as shown in Figure 2.10. Interestingly, both the nanostructured thin-film (NSTF) catalyst and the CL are nonconventional. The latter contains no carbon or additional ionomer and is 20–30 times thinner than the conventional dispersed Pt/carbon-based CL. In addition, the CL was more durable than conventional CCMs made from Pt/C and Nafion ionomer [40].

2.3.3 Novel Structural Catalyst Layer

2.3.3.1 CNT-Based Catalyst Layer

The carbon nanotube (CNT)-supported Pt-type catalyst layer is considered a promising electrode structure for PEM fuel cells [42–47]. Several reports have been published on CNT-based CLs in fuel cell applications. The attrac-tive features of CNT-based CLs include improved thermal and charge trans-fer, and maximum exposure of the catalyst sites to the gas reactant through uniform support geometry and parallel alignment. Figure 2.11 shows the structure of an aligned CNT-based MEA. A mass activity of 250 A/mg Pt at 0.6 μg Pt/cm2 was obtained by Tang et al. [43].

2.3.3.2 Columnar Oxide Supported Catalyst Layer

Similar to the 3M whisker support discussed earlier, electronic conduct-ing ceramic columnar supports have also been proposed. Bonakdarpour et al. [48] characterized columnar titanium structures on smooth glassy


carbon (GC) disks fabricated by using glancing angle deposition (GLAD) and physical vapor deposition techniques. This catalyst support consisted of posts 500 nm long with a nominal cross-sectional diameter of 100 nm. Platinum films, with an equivalent planar thickness of 10–90 nm, were

FIGURE 2.10Scanning electron micrographs of typical NSTF catalysts as fabricated on a microstructured catalyst transfer substrate, seen (top) in cross section with original magnification of ×10,000 and (bottom) in-plane view with original magnification of ×50,000. A dotted scale bar is shown in each micrograph. (Reproduced from Debe, M. K. et al. Journal of Power Sources 2006; 161:1002–1011. With permission from Elsevier.)


deposited onto these posts by magnetron sputtering. The electrochemical surface area of such catalysts was about 10–15 times higher than that of smooth Pt films. Although such structures were tested using a GC elec-trode, it should be possible to develop them into novel-structure catalyst layers for PEM fuel cells.

2.3.3.3 Nanowire-Based Three-Dimensional Hierarchical Core/Shell Catalyst Layer

Saha et al. [49,50] explored the feasibility of using nanowire supports for Pt-based catalysts in PEM fuel cells. A three-dimensional core/shell het-erostructure, consisting of a tin nanowire core and a carbon nanotube shell (SnC), was grown directly onto the carbon paper backing. Compared with the conventional Pt/C catalyst layer, the SnC nanowire-based CL showed a higher oxygen reduction performance and excellent stability in a fuel cell environment. The results demonstrated that the core/shell nanowire-based composites are very promising supports in making cost-effective electrocat-alysts for CLs in fuel cell applications.

CarbonBacking Layer

Anode

Microporous Layer

Catalyst Layer

CathodePEM

PTFE+C

e– e–

2e–

2e–

H+

O2

MWNT

Support

2H+

Nafion

Pt Nanoparticle

Carbon Black

Support

H2O

H2

H2

O212

FIGURE 2.11Schematic of the hydrogen fuel cell architecture using an ultra-low Pt loading thin-film Pt/MWNT catalyst layer (MWNT = multiwalled nanotube). (Reproduced with permission from Tang, J. M. et al. Journal of Physical Chemistry C 2007; 111:17901–17904. Copyright 2007 American Chemical Society.)


2.3.3.4 Self-Supported Catalyst Layer

Yamada et al. [51] prepared a Pt catalyst film with a dendritic structure by reducing a-PtO2. The PtO2 film was deposited on the microporous layer (MPL) surface of a GDL using a reactive sputtering process under 100% O2 at room tem-perature. The sputtering-deposited PtO2 film was then subjected to a hydrogen reduction process using 2% H2/He (0.1 MPa) at room temperature to obtain a dendritic Pt catalyst film, as shown in Figure 2.12. This self-supported dendritic Pt film exhibited a low density of 3.3 g/cm3. When it was applied as the cathode catalyst layer for fuel cells, higher performance, larger electrochemical surface area (ECA), and improved diffusion characteristics were observed compared to a conventional sputtered Pt film. The activity per unit ECA of the dendritic Pt was also higher than that of conventional sputtered Pt catalysts.

2.3.3.5 Catalyst Layer with Additives

Typically, Nafion ionomer is the predominant additive in the catalyst layer. However, other types of CLs with various hygroscopic or proton conductor additives have also been developed for fuel cells operated under low relative humidity (RH) and/or at elevated temperatures. Many studies have reported the use of hygroscopic g-Al2O3 [52] and silica [53,54] in the CL to improve the water retention capacity and make such CLs viable for operation at lower rela-tive humidity and/or elevated temperature. Alternatively, proton conducting materials such as ZrP [55] or heteropoly acid HPA [56] have also been added

(a) (b)

FIGURE 2.12Scanning electron microscopy photographs of the dendritic Pt film (a) before the reduction treatment and (b) after the reduction treatment. (Reproduced from Yamada, K. et al. Journal of Power Sources 2008; 180:181–184. With permission from Elsevier.)


into the CL to improve proton conductivity under low RH and/or elevated temperature. These types of CLs have the potential for applications in high-temperature PEM fuel cells at either cathode or anode.

2.3.3.6 Catalyst Layer with Novel Ionomers

In addition to Nafion-based catalyst layers, additional types have been devel-oped, including CLs with different ion exchange capacities (IECs) [57,58] or with other hydrocarbon-type ionomers such as sulfonated poly(ether ether ketone) [58–60], sulfonated polysulfone [61,62], sulfonated polyether iono-mers [63], and borosiloxane electrolytes [64], as well as sulfonated polyimide [65]. These nonfluorinated polymer materials have been targeted to reduce cost and/or increase operating temperature. Unfortunately, such CLs still encounter problems with low Pt utilization, flooding, and inferior perfor-mance compared with conventional Nafion-based CLs.

2.4 Catalyst Layer Fabrication

The fabrication of catalyst layers for PEM fuel cells involves maintaining a delicate balance between gas and water transport, and electron and proton conduction. The process of CL fabrication should be guided by both fuel cell performance and cost reduction.

2.4.1 First-Generation Catalyst Layer Fabrication

2.4.1.1 Pt Black Catalyst Layer Fabrication

Although thin-film catalyst layers are widely used in current fuel cell tech-nology, PTFE-bound CLs with Pt black as the catalyst, which were used in earlier generations of PEM fuel cells, have recently been revisited due to their excellent long-term durability [66,67]. A typical Pt black-based CL or electrode consisted of Pt black and hydrophobic PTFE as a binder. The CL formed by this Pt black catalyst has several disadvantages, including high platinum loadings (4 mg/cm2), large platinum agglomerates (~1 μm on average), lower electrochemical surface area (~25 m2/g), and poor access to the catalyst surface for gas, electrons, and protons. The unsupported Pt black can be directly deposited onto the GDL or membrane [68–71]. A procedure for the preparation of Pt black-based electrode includes the fol-lowing steps:

1. Mix Pt black and water (e.g., 3 mL H2O/g of Pt black) for 24 hours to deagglomerate the as-received catalyst particles.

2. Dilute the PTFE emulsion to 6%.


3. Add this diluted PTFE to the Pt/H2O slurry and mix them thor-oughly. The ratio of Pt to PTFE in the catalyst ink is controlled at ~85:15.

4. Apply the catalyst ink uniformly onto the GDL surface using a metering rod.

5. Sinter the electrode under nitrogen at 340°C for 90 minutes.

2.4.1.2 PTFE-Bound Catalyst Layer Fabrication Using Supported Catalyst

Although PTFE-bound Pt black catalyst layer (electrode) has demonstrated exceptional long-term performance, the catalyst costs are prohibitive for commercialization. A breakthrough was made by replacing pure Pt black with supported Pt catalysts [72], which can significantly reduce the Pt loading from 4 mg Pt/cm2 down to 0.4 mg/Pt cm2 [5,73]. In these CLs, the catalyst particles were still bound by hydrophobic PTFE. However, the impregnation of an ionomer (Nafion) into these CLs was found to be extremely effective in improving the three-dimensional reaction zone for fuel cell applications. The process employed for forming the PTFE-bound CL and MEA can be summarized as follows [3,74]:

1. Mix 20 wt% of Pt/C catalyst particles and the solvent for 30 minutes to form a slurry.

2. Add PTFE emulsion into this slurry until it constitutes 30% of the mixture.

3. Add a bridge-builder and a peptization agent into this mixture, fol-lowed by 30 minutes of stirring to form the catalyst ink.

4. Apply this ink onto the wet-proofed carbon paper, using a coating apparatus.

5. Dry this catalyst-coated carbon paper for 24 hours in ambient air and then bake it at 225°C for 30 minutes to form an electrode.

6. Sinter this electrode at 350°C for 30 minutes. 7. Brush 5 wt% Nafion solution onto the catalyst layer of this

electrode. 8. Dry this Nafion-impregnated electrode in an oven at 80°C for an

hour in ambient air.

Using a carbon-supported Pt catalyst to replace Pt black can reduce the platinum loading by a factor of 10—from 4 to 0.4 mg/cm2 [74]. However, the platinum utilization in this PTFE-bound catalyst layer still remains low: in the vicinity of 20% [75,76].


2.4.2 Thin-Film Catalyst Layer Fabrication

2.4.2.1 Ink-Based Catalyst Layer Fabrication

As mentioned earlier, the PTFE-bound catalyst layer electrode has some draw-backs. For example, due to the nonuniformity of the CL, large variations in impregnation with the recast ionomer occur; this can result in some areas of the CL being overimpregnated. These areas of overimpregnation can present a transport barrier to the diffusion of gas through the backing. Not surprisingly, it is difficult to achieve a uniform distribution in such a catalyst layer structure [6].

Approaches to improve the CL construction by significantly increasing the contact area between the polymer electrolyte and the platinum clusters can be achieved in two ways. First, the supported catalyst and ionomeric addi-tive are cast together with PTFE emulsion to form the CL. This assures that the thickness of the CL coincides with the depth of the ionomer. Second, the contact area between the ionomer additive and the catalyst is increased by completely removing the Teflon component and by improving the dispersion of the ionomer throughout the CL. The latter can be accomplished by blend-ing the solubilized ionomer and the platinized carbon into a homogeneous “ink” from which the thin-film CL of the electrode can be made.

Wilson, Valerio, and Gottesfeld [77] proposed an improved fabrication method for a low Pt loading (0.12 mg Pt/cm2) catalyst layer using a thermo-plastic ionomer. Low platinum loading CLs consist of a thin film of highly intermixed ionomer and a catalyst that is applied to the electrolyte membrane. The procedure for fabricating the thermoplastic catalyst layer is similar to the typical thin-film method, but with the addition of tetrabutylammonium (TBA) to the ink mixture. Subsequently, an additional ion-exchange process is needed to convert TBA+ ionomer into H+ form. The advantage of including large TBA cations is that this process can significantly improve the structural integrity of the CLs. Other benefits of thermoplastic catalysts include improv-ing the reproducibility and durability of the corresponding fuel cells.

Numerous efforts have been made to improve existing thin-film catalysts in order to prepare a CL with low Pt loading and high Pt utilization without sac-rificing electrode performance. In thin-film CL fabrication, the most common method is to prepare catalyst ink by mixing the Pt/C agglomerates with a solu-bilized polymer electrolyte such as Nafion ionomer and then to apply this ink on a porous support or membrane using various methods. In this case, the CL always contains some inactive catalyst sites not available for fuel cell reactions because the electrochemical reaction is located only at the interface between the polymer electrolyte and the Pt catalyst where there is reactant access.

According to Wilson’s 1993 patent [78], the procedure for fabricating a thin-film catalyst layer on the membrane is as follows:

1. Mix a 5% solution of solubilized perfluorosulfonate ionomer (such as Nafion) and 20 wt% Pt/C support catalyst in a ratio of 1:3 (Nafion:catalyst).


2. Add water and glycerol into this mixture in a weight ratio of 1:5:20 (carbon:water:glycerol).

3. Blend the formed mixture until the catalyst is uniformly distributed and the mixture has adequate viscosity for coating. This step pro-duces a catalyst ink.

4. Apply this catalyst ink onto one side of the membrane (Na+ type). Two coats are typically required for adequate catalyst loading.

5. Dry the catalyst-coated membrane in a vacuum at a temperature of approximately 160 C.

Alternatively, the catalyst layer ink can be (1) applied to a PTFE blank or some other substrate and then decal transferred onto the membrane [6,79], or (2) deposited onto the diffusion layer and then hot pressed to the membrane for MEA fabrication [17,80–83]. Based on the nature of catalyst ink and its application method, several thin-film CL fabrication techniques have been developed, including decal transfer [6,79], brush painting [74,84], spray coating [85,86], doctor blade coating [87], screen printing [88–90], ink-jet printing [91,92], and rolling [93]. Currently, screen printing and spray coating have become standard methods for conventional catalyst layer fab-rication. Inkjet printing is also showing promise for fabricating low Pt load-ing CLs.

2.4.2.1.1 Screen Printing

Screen printing is one of most popular methods of ink-based catalyst layer preparation. In a typical screen printing process, the ink slurry is first cast onto certain substrates, and then the CL is transferred to a Nafion membrane by hot pressing. The membrane is converted into Na+ form to avoid swell-ing during hot pressing and decal transfer because, in this form, the Nafion membrane is mechanically strong and stable for hot pressing from 150 to 160°C. In the direct screen printing [77] process, the ink slurry is applied to a membrane in either Na+ or TBA form to stabilize the catalytic layer, thereby enhancing the membrane’s physical strength.

The coating apparatus consists of a silk screen mesh fixed to a frame with sufficient tension to squeeze the ink through the screen and onto the blank substrate (e.g., polyimide). The substrate is fixed on an XY table with adhe-sion tape, and the silken screen mesh is masked, with an open window in the center for screen printing. The silicon rubber squeeze is a fixed support and can be moved in both X and Y directions. A hot-air or IR ramp is used to dry the coating for solvent removal.

The coating procedure consists of positioning the substrate layer under the silk screen mesh, which is not masked, and using a squeegee on top of the mesh at one side. The schematic of this coating apparatus is shown in Figure 2.13. An appropriate amount of ink is micropipetted near the squeegee, and the slurry is first spread across and then pushed through the mesh to the


substrate layer by rapid movements of the squeegee. Hot air is used to dry the catalyst layer coated on the substrate. The same procedure is repeated until the pipette volume of catalyst ink has been transferred to the substrate layer.

2.4.2.1.2 Spray Coating

Similar to screen printing, the spray coating method [95] is widely used for catalyst fabrication, especially in labs. The major difference between the two is that the viscosity of the ink for spray coating is much lower than that for screen printing. The application apparatus can be a manual spray gun or an auto-spraying system with programmed X-Y axes, movable robotic arm, an ink reservoir and supply loop, ink atomization, and a spray nozzle with adjustable flux and pressure. The catalyst ink can be coated on the gas diffu-sion layer or cast directly on the membrane. To prevent distortion and swell-ing of the membrane, either it is converted into Na+ form or a vacuum table is used to fix the membrane. The catalyst layer is dried in situ or put into an oven to remove the solvent.

2.4.2.1.3 Inkjet Printing

Inkjet printers utilize drop-on-demand technology to deposit various materials or “inks.” This is a popular deposition technique used not only in desktop printers, but also to deposit various other coating materials, such as those required for catalyst layer fabrication. Using inkjet printing tech-nology, a research group in the Pacific Northwest National Lab (PNNL) has successfully fabricated CLs for hydrogen-air PEM fuel cells [91]. In their study, a slightly modified commercial desktop inkjet printer was used to deposit catalyst ink directly from a print cartridge onto the Nafion mem-brane to form a catalyst layer. The ink cartridges were filled with the cata-lyst inks. The membrane was secured to a cellulose acetate sheet and fed through the printer using the original paper feed platen. Computer soft-ware was used to control the print parameters, such as electrode dimen-sions, thickness, and resolution. Fuel cell testing on the fabricated CLs showed power densities up to 155 mW/cm2, with a cathode catalyst loading of 0.20 mg Pt/cm2.

These studies demonstrate some of the advantages of inkjet printing for cat-alyst layer fabrication, such as varied composition layer printing, and suggest

Squeegee

Masked Silk Screen

Substrate Film

FIGURE 2.13Schematic of screen printing apparatus. (Reproduced from Rajalakshmi, N. and Dhathathreyan, K. S. Chemical Engineering Journal 2007; 129:31–40. With permission from Elsevier.)


that inkjet-based fabrication technology might be the way to cost-effective and large-scale fabrication of PEM fuel cells in the future. However, the ques-tion still remains of whether or not inkjet fabrication can offer any perfor-mance advantage and still lead to more efficient utilization of the Pt catalyst.

Similarly, Taylor et al. [92] used inkjet printing to deposit catalyst materi-als onto GDLs. Their inkjet-printed catalyst layers with a catalyst loading of 0.020 mg Pt/cm2 showed high Pt utilizations. This research also demon-strated the capacity of the inkjet printing technique to control ink volume precisely down to picoliters for ultralow catalyst loading, the flexibility of using different carbon substrates, and the functionality of gradient catalyst structure fabrication using these techniques.

2.4.2.2 In Situ Catalyst Layer Fabrication

Numerous efforts have been made to develop in situ catalyst layer fabrica-tion methods to lower Pt loading and increase platinum utilization without sacrificing electrode performance.

2.4.2.2.1 Sputter Deposition

The sputter deposition technique is recognized as having great potential for reducing the Pt loading of PEM fuel cells [96,97]. Hirano, Kim, and Srinivasan [98] reported using sputter deposition for fuel cell catalyst layer preparation in 1997. Via this technique, an ultrathin CL (1 μm) with low Pt loading (0.1 mg/cm2) was produced on a gas diffusion electrode. Since then, much effort [99–107] has been put into applying this sputtering tech-nology to the fabrication of CLs, with the twofold goal of improving both fuel cell performance and catalyst utilization.

Figure 2.14 illustrates a low-pressure plasma sputtering reactor setup. The fabrication procedure for the sputter-deposition technique involves having the platinum target opposite the carbon GDL in a vacuum reactor and then sputtering to form a thin CL [108]. The catalyst loading is adjusted by alter-ing the sputtering duration. After this approach has been used, the GDL has been shown by scanning electron microscopy (SEM) analysis to be covered by Pt nanoclusters, which form a relatively dense clustering of Pt films on the surface.

The advantages of sputter deposition include:

1. Precise Pt loading and thickness, as well as controlled microstruc-ture morphology;

2. Much smaller Pt particle size; 3. Homogeneous distribution of the Pt particles on the support and

extremely low metal loadings (down to 10 g/cm2); 4. A simple preparation process that is easy to scale up; and 5. Adaptiveness to various substrates, such as GDL and membrane.


Although the sputter deposition technique can provide a cheap and directly controlled deposition method, the performance of PEM fuel cells with sputtered CLs is still inferior to that of conventional ink-based fuel cells. In addition, other issues arise related to the physical properties of sput-tered catalyst layers, such as low lateral electrical conductivity of the thin metallic films [96,108]. Furthermore, the smaller particle size of sputter-de-posited Pt can hinder water transport because of the high resistance to water transport in a thick, dense, sputtered Pt layer [108]. Currently, the sputter deposition method is not considered an economically viable alternative for large-scale electrode fabrication [82] and further research is underway to improve methods.

2.4.2.2.2 Dual Ion-Beam-Assisted Deposition

Saha et al. [109] have proposed an improved ion deposition methodology based on a dual ion-beam assisted deposition (dual IBAD) method. Dual IBAD combines physical vapor deposition (PVD) with ion-beam bombard-ment. The unique feature of dual IBAD is that the ion bombardment can impart substantial energy to the coating and coating/substrate interface, which could be employed to control film properties such as uniformity, density, and morphology. Using the dual IABD method, an ultralow, pure Pt-based catalyst layer (0.04–0.12 mg Pt/cm2) can be prepared on the surface of a GDL substrate, with film thicknesses in the range of 250–750 Å. The main drawback is that the fuel cell performance of such a CL is much lower than that of conventional ink-based catalyst layers. Further improvement

FIGURE 2.14Schematic of the plasma sputtering reactor. (Reproduced from Kadjo, A. J. J. et al. Journal of Power Sources 2007; 172:613–622. With permission from Elsevier.)


in catalyst utilization and the gas/liquid diffusivity of the CLs prepared by IBAD is necessary.

2.4.2.2.3 Electrodeposition Method

In general, catalyst fabrication always yields inactive catalyst sites in the cat-alyst layer because they do not meet the three-phase reaction requirements. In order to mitigate this drawback and increase Pt utilization, a pulse elec-trodeposition method was developed by Taylor, Anderson, and Vilambi [82] and exemplified by Lee et al. [110]. Electrodeposition was conducted in a Pt plating bath containing K2PtCl4 and NaCl on a Nafion-bonded carbon layer at room temperature. The catalyzed electrode was first rinsed thoroughly with ultrapure water to remove any Pt salt residue, followed by heat treatment at 250°C in a H2 (10%)/N2 (90%) environment for 30 minutes. The ionomer pro-tonation process was carried out by immersing this electrode in 0.1 M H2SO4 solution for 30 minutes with mild heating (80°C). The electrodes were then thoroughly rinsed with water.

Such electrodeposited electrodes (0.025 mg Pt/cm2 on the anode and 0.3 mg Pt/cm2 on the cathode) have demonstrated higher fuel cell performance than a conventional CL (0.3 mg Pt/cm2 on both electrodes). A noticeable increase in catalyst utilization resulted when Pt deposition took place only in the three-phase reaction zone. Because Pt electrodeposition in aqueous solution only occurs in the region of the ionic and electronic pathways, it should be possible to reduce Pt loading significantly and increase Pt utilization in the CL. Pulse electrodeposition is promising and could replace conventional methods for fabricating cost-effective, low Pt loading CLs. However, CL durability may be an issue if the active catalyst layer sites change with time.

2.4.2.2.4 Reactive Spray Deposition Technology

Reactive spray deposition technology (RSDT) has been developed for direct catalyst deposition onto substrates, including polymer electrolyte mem-branes, to form CCMs [111]. The process involves dissolving a Pt precursor in appropriate solvents, followed by spraying the solution with an expan-sion gas through a nozzle to produce micron-sized droplets. The droplets are burnt out in a flame, resulting in metal atoms and/or metal oxide molecules in the gas phase. At the same time, a quench gas is used to induce rapid con-densation of Pt vapors into Pt particles with a size of ~5 nm. A mixture of carbon and Nafion ionomer is subsequently introduced into the gas stream. The cooling effect of the quench gas helps to avoid thermal damage to the ionomer and the membrane substrate. As a result, a thin-film catalyst layer forms on the electrolyte membrane.

In the RSDT process, the steps for introducing catalyst, ionomer, and car-bon into the gas mix are decoupled and can be independently controlled in such a manner that the Pt/C and ionomer/C ratios can be continuously modified during the deposition process. Reactive spray deposition technol-ogy has the capacity and flexibility required to produce compositionally and


structurally optimized CLs. The technology can be used to generate sup-ported and unsupported platinum CLs with thicknesses from 10 to 200 nm and with varied morphologies, including dense films, porous films (packed particles), and dendritic and island-growth structures. However, the usage efficiency of the catalyst ink is quite low in the fabrication process, which is a major drawback of this technology for fuel cell CL preparation.

2.4.2.2.5 Pulsed Laser Deposition

Cunningham et al. [112] used the pulsed laser deposition (PLD) method to deposit platinum onto E-TEK gas diffusion electrodes (GDEs) to prepare low catalyst loading electrodes for PEM fuel cells. In the PLD process, the laser beam is focused by a quartz focus lens onto a polycrystalline plati-num target. The Pt target is continuously rastered across the laser beam, via a dual rotation and translation motion, to obtain a uniform ablation over the entire target surface. The chamber is evacuated by a turbomo-lecular pump and filled with helium at a constant pressure throughout the PLD process. The platinum loadings are controlled by the number of pulses during deposition.

This technique yields a catalyst composed entirely of metal nanoparticles or nanocrystalline thin film, and it allows for control of size and distribution while eliminating the need for a dispersing and supporting medium. The obtained electrodes contained as little as 0.017 mg Pt/cm2 and performed as well as standard E-TEK electrodes (Pt loading 0.4 mg/cm2). The PLD tech-nique may be of special interest as an alternative to the sputtering process in the production of micro fuel cells.

2.4.2.3 Other Methods

2.4.2.3.1 DLR Process

The dry powder spraying technique was developed by the German Aerospace Center (acronym DLR, from Deutsches Zentrum für Luft- und Raumfahrt) for the production of CCMs [113–115]. The DLR preparation process is claimed to be a low-cost and effective manufacturing process for PEM fuel cell CCMs. The basic technique for this production of CCMs is to spray a dry catalyst powder directly onto the membrane, thereby avoiding the wait time for evaporation of solvents in a wet process and, at the same time, achiev-ing good contact between the catalytic layer and the electrolyte membrane. The dry powder preparation process for MEAs is divided into two steps: (1) preparation of the electrode powder by mixing the catalyst powder with different additives, and (2) preparation of the CCM by spraying the catalyst/additives powder onto the membrane.

Figure 2.15 shows a schematic representation of the CCM manufacturing process. The first step for CCM preparation comprises mixing the reactive layer materials (e.g., platinum-supported carbon black) with different amounts of PTFE, polymer electrolyte powder (e.g., Nafion), and/or filler materials


in a knife mill. In order to obtain thin, homogeneous reactive catalyst lay-ers, the materials are atomized and sprayed in a nitrogen stream through a slit nozzle directly onto the proton conducting membrane. Although adhe-sion of the catalytic layer on the membrane is good, a hot rolling or pressing process follows in order to improve the electrical and ionic contacts further. Depending on the degree of atomization, completely covered reactive layers with a thickness as low as 3 μm can be prepared.

2.4.2.3.2 Electrospray Technique

Benitez, Soler, and Daza [116] have developed a novel catalyst deposition method, based on an electrospray technique, to prepare catalyst layers for PEM fuel cells. The electrospray technique involves applying a high volt-age (3,300–4,000 V) between a capillary tube, in which the ink is forced to flow using a high-pressure nitrogen stream, and a carbon cloth substrate. This high electric field generates a mist of highly charged droplets. During the spraying process, the droplets are reduced in size by evaporation of the solvent and/or by “coulomb explosion” (droplet subdivision resulting from high charge density). A high-pressure nitrogen stream is used to force the catalytic ink through the capillary tube. The catalytic ink is put in an ultra-sonic bath to keep it a homogeneous mixture.

During spraying, the carbon substrate is moved by means of an X-Y axes coordinated system controlled by computer software. The CLs formed from the electrospray technique show both morphological and structural improve-ments, which could contribute to better catalyst utilization than is achieved

Coating Nozzle

Powder Supporter

Catalyst

Additive

Nitrogen

Membrane

Roller

FIGURE 2.15Schematic representation of the DLR (dry powder spraying) CCM manufacturing process. (Reproduced from Wagner, N., Kaz, T., and Friedrich, K. A. Electrochimica Acta 2008; 53:7475–7482. With permission from Elsevier.)


by conventional methods. It has been claimed that MEAs fabricated by the electrospray technique exhibit a power density three times higher than those prepared by impregnation methods and eight times higher than those made by conventional spray techniques. Electrospraying is considered a low-cost process and is scalable to volume production.

2.4.2.3.3 Electrophoretic Deposition

Morikawa et al. [117], Louh, Huang, and Tsai [118], and Lough et al. [119] have proposed an electrophoretic deposition (EPD) method to fabricate CLs for PEM fuel cells. In the EPD process, a suspension consisting of ethanol, carbon powder-supported Pt catalyst, and Nafion ionomer is used to obtain a stable dispersed solution. A working electrode (usually a carbon textiles substrate for Pt deposition) and a counterelectrode (Pt) are connected to a high-voltage DC power supply. Voltage (300 V/cm) is applied between the platinum counterelectrode and the working electrode, using a computer-controlled waveform programming controller. The thickness of the prepared CL is controlled by EPD duration and/or suspension concentration.

A well-distributed deposition of Pt/C nanocatalyst and Nafion ionomer on both hydrophilic and hydrophobic carbon-based electrodes has been suc-cessfully obtained using a Pt/C concentration of 1.0 g/L, an electrical field of 300 V/cm, and a deposition time of 5 minutes [118]. The deposition of Pt/C nanocatalysts and Nafion solution via the electrophoretic process gives rise to higher deposition efficiency and a uniform distribution of catalyst and Nafion ionomer on the PEMFC electrodes.

2.4.2.3.4 Sol-Gel Pt Application

Khan and Lin [120] prepared novel Pt-based catalyst layers with controllable Pt loadings by direct deposition of Pt sols. They added a mixture of 0.1 M H2PtCl6(aq) and 0.1 M sodium citrate (J. T. Baker, 99.9%) to a methanol solution under reflux and stirring at 353 K. The reaction was stopped by quenching to room temperature as soon as the solution turned black. The prepared Pt sols were dispersed evenly on a Nafion 117 membrane (DuPont), and the sol-vent was then allowed to evaporate at room temperature. Pt loading was con-trolled by the amount of sol added for deposition. The Pt-deposited membrane showed high specific activity, as well as single-cell performance comparable to conventional CLs prepared from Pt/C catalysts. This method has the dual advantage of easy preparation and good control over Pt loading reduction.

2.5 Catalyst Layer Optimization

An effective catalyst layer must serve multiple functions simultaneously: electron and proton conduction, oxygen or hydrogen supply, and water man-agement. The composition and structure of a CL can affect all these functions


to different extents. Catalyst layer optimization aims to satisfy these require-ments, as well as to maximize Pt utilization, enhance durability, and improve fuel cell performance.

Because the reaction in a CL requires three-phase boundaries (or inter-faces) among Nafion (for proton transfer), platinum (for catalysis), and carbon (for electron transfer), as well as reactant, an optimized CL structure should balance electrochemical activity, gas transport capability, and effective water management. These goals are achieved through modeling simulations and experimental investigations, as well as the interplay between modeling and experimental validation.

2.5.1 Catalyst Layer Composition Optimization

The catalyst layer is composed of multiple components, primarily Nafion ion-omer and carbon-supported catalyst particles. The composition governs the macro- and mesostructures of the CL, which in turn have a significant influence on the effective properties of the CL and consequently the overall fuel cell per-formance. There is a trade-off between ionomer and catalyst loadings for opti-mum performance. For example, increased Nafion ionomer content can improve proton conduction, but the porous channels for reactant gas transfer and water removal are reduced. On the other hand, increased Pt loading can enhance the electrochemical reaction rate, and also increase the catalyst layer thickness.

How to balance Nafion ionomer content and Pt/C loading is a challenge for optimizing CL performance, due to the complexity induced by proton and electron conduction, reactant and product mass transport, as well as electrochemical reactions within the CL. The optimization of such a complex system is mainly implemented through multiple components and scale mod-eling, in combination with experimental validation.

2.5.1.1 Modeling and Simulation to Optimize Catalyst Layers

Optimization of CL composition has been carried out extensively through modeling and simulation:

Using numerical modeling, Wei et al. [121] found a ratio threshold for Pt/C catalyst loading and Nafion ionomer content in the catalyst layer. Beyond this threshold, catalyst utilization could drop dramati-cally. A CL with higher Pt loading could allow a larger range in the ratio of Pt/C catalyst to Nafion loading. Optimal catalyst utilization was reached around 1:1.Song et al. [122] modeled optimal performance as a function of Nafion content as well as Pt loading.Kamarajugadda and Mazumder [123] used numerical modeling and investigated the effects of ionomer (Nafion) loading, catalyst (Pt)


loading, Pt/carbon ratio, agglomerate size, and CL thickness on fuel cell performance. Optimal performance was a function of Nafion content and its distribution around and within the agglomerates. Nafion distribution around the agglomerates had a stronger effect than distribution within the agglomerates, and these researchers found that moderate Pt loading optimized fuel cell performance.Wang, Mukherjee, and Wang [124] investigated the effects of catalyst layer electrolyte and void phase fractions on fuel cell performance using a random microstructure. The model predicted volume frac-tions of 0.4 and 0.26 for void and electrolyte phases, respectively, as the optimal CL compositions.Secanell et al. [125] presented a gradient-based optimization of fuel cell performance. They found that a significant increase in perfor-mance could be achieved by increasing Pt loading and reaching a Nafion mass fraction around 20–30 wt% in the CL.Jain, Biegler, and Jhon [126] optimized Pt distribution along the width of the CL and found that a significant improvement in current density could be obtained by placing higher amounts of Pt adjacent to the catalyst layer/membrane interface.Thepkaew, Therdthianwong, and Therdthianwong [127] conducted a full factorial analysis, including carbon types (Vulcan XC72R and Black Pearls 2000), Pt loading (0.1 and 0.5 mg/cm2), and Nafion iono-mer content (10 and 60%); they found that the key factor affecting the exchange current density or activation loss was Pt loading, and the key factors controlling the ohmic resistance were Nafion content and carbon type. Interactions between these parameters, in particu-lar the interaction of carbon type and Nafion content, controlled the performance of the thin-film CL.

In summary, modeling offers powerful tools and guidance for performance optimization. With advancements in new techniques for micro- and nano-fabrication, it will be possible to engineer fuel cell CLs (electrodes) according to the compositions and structures predicted by modeling and simulation.

2.5.1.2 Experimental Studies on Optimization of CLs

The experimental optimization of Nafion ionomer loading within a catalyst layer has attracted widespread attention in the fuel cell community, mainly due to its critical role in dictating the reaction sites and mass transport of reactants and products [15,128–134]. Nafion ionomer is a key component in the CL, helping to increase the three-phase reaction sites and platinum utili-zation to retain moisture, as well as to prevent membrane dehydration, espe-cially at low current densities. Optimal Nafion content in the electrode is necessary to achieve high performance.


Gode et al. [135] investigated the effect of Nafion content (10–70 wt%) on catalyst layer structure and electrochemical performance; they found that a cathode CL containing 36–43 wt% Nafion displayed optimal performance. Passalacqua et al. [136] examined the effect of Nafion content (14–66 wt%) in a low Pt loading CL (0.1 mg/cm2), and found an optimal ionomer con-tent to be ~33 wt%. Pagainin, Ticianelli, and Gonzalez [83] confirmed that when the Nafion loading was increased from 0.87 to 1.75 mg/cm2, perfor-mance could be significantly improved. If the Nafion loading was increased beyond 2.2 mg/cm2 (equivalent to 33 wt%), the fuel cell performance began to deteriorate at higher current densities. This Nafion loading value of 33% has also been observed in several other recent studies [137–139].

The distribution optimization of Nafion ionomer inside the catalyst layer also appears to be important. Lee and Hwang [133] found that if the Nafion ionomer was impregnated on the CL surface, better performance could be obtained than if the Nafion ionomer was distributed inside the CL. Kim et al. [29] also investigated the Nafion distribution effect using a Nafion dual-gradient CL. They found that when the Nafion content was higher near the electrolyte membrane and lower near the gas diffusion layer, bet-ter performance could be obtained in the high current density region. Xie et al. [32] prepared a three-sublayer CL containing a gradient distribu-tion of Nafion ionomer; they observed that cathode performance could be improved when Nafion content was higher at locations near the catalyst layer/membrane interface and lower at locations near the catalyst layer/carbon paper interface. They argued that more ionomer is needed for pro-ton transport at the catalyst layer/membrane interface, whereas at locations near the GDL interface, less ionomer is better to obtain more porous chan-nels for gas flux.

Apart from the optimization of a single component (Nafion), optimiza-tion of two components, Pt/C and Nafion, has been explored. Sasikumar, Ihm, and Ryu [131] examined the correlation between Pt loading and opti-mum Nafion content in the CL, finding that optimum Nafion content was dependent on Pt loading. If the Nafion content was increased, the Pt load-ing decreased. For electrodes with Pt loadings of 0.5, 0.25, and 0.1 mg/cm2, the best performance was obtained at about 20, 40, and 50% Nafion ionomer loadings, respectively. Qi and Kaufman [138] observed that the best perfor-mance was achieved with Pt loadings of 0.20 0.05 and 0.35 0.05 mg/cm2 for 20 and 40% Pt/C, respectively, but the best performance using 40% Pt/C was only slightly better than that using 20% Pt/C.

Cho et al. [140] examined the performance of PEM fuel cells fabricated using different catalyst loadings (20, 40, and 60 wt% on a carbon support). The best performance—742 mA/cm2 at a cell voltage of 0.6 V—was achieved using 40 wt% Pt/C in both anode and cathode. Antonie et al. [28] studied the effect of catalyst gradients on CL performances using both experimental and modeling approaches. Optimal catalyst utilization could also be achieved when a preferential location of Pt nanoparticles was close to the PEM side;


this demonstrated the possibility of reducing Pt loading using catalyst gra-dients in the CL.

In summary, ionomer content and its distribution inside the CL have a strong effect on fuel cell performance. An optimal ionomer content exists with respect to fuel cell performance. Increasing ionomer loading in the CL can effectively improve the electrochemical active area and overall ionic con-ductivity. However, if the ionomer loading is too high, transport of reactant gas to the reactive sites will be impeded, leading to significant mass trans-port loss. The optimal ionomer loading for a CL depends on what materials and fabrication methods are used. It is generally agreed that a Nafion iono-mer loading in the range of 30–40 wt% is optimal for conventional thin-film catalyst layers.

2.5.2 Catalyst Layer Microstructure Optimization

The microstructure of a catalyst layer is mainly determined by its composi-tion and the fabrication method. Many attempts have been made to optimize pore size, pore distribution, and pore structure for better mass transport. Liu and Wang [141] found that a CL structure with a higher porosity near the GDL was beneficial for O2 transport and water removal. A CL with a stepwise porosity distribution, a higher porosity near the GDL, and a lower porosity near the membrane could perform better than one with a uniform porosity distribution. This pore structure led to better O2 distribution in the CL and extended the reaction zone toward the GDL side. The position of macropores also played an important role in proton conduction and oxygen transport within the CL, due to favorable proton and oxygen concentration conduction profiles.

During catalyst layer fabrication, to enhance mass transport, a porous structure can be created by adding pore formers into the catalyst ink for-mula. Yoon et al. [142] introduced ethylene glycol into the catalyst slurry formulation in order to improve the performance of the catalytic layer. Ethylene glycol acts as a pore former, thereby increasing the secondary pores within agglomerates and assisting gas transport through the catalyst layer. Song et al. [143] used ammonium carbonate as a pore-forming agent in the CL to minimize mass-transport limitations. Fischer et al. [23] increased the porosity by adding pore formers such as volatile filler, ammonium car-bonate, ammonium oxalate, or soluble lithium carbonate to the catalyst slurry. Zhao et al. [144] used NH4HCO3, (NH4)2SO4, and (NH4)2C2O4 as pore formers to prepare CLs. Adding NH4HCO3 made the catalyst dispersion more uniform and the surface more porous, leading to low gas diffusion resistance.

Other efforts have also been made to modify the CL microstructure by controlling the agglomerate size in the catalyst ink. Uchida et al. [145] pro-posed a colloidal ink fabrication procedure using low-dielectric-constant solvents to generate a good network and a uniformity of perfluorinated


sulfone ionomers (PFSIs) on Pt particles in the CL. Wang et al. [146] opti-mized CL microstructure by adding NaOH to suppress Nafion aggregation and achieved a smaller agglomerate particle size distribution in the catalyst ink. The cathode CL made from this catalyst ink showed a high electrochem-ically active surface area (a 48% increase over conventional ink). Fernandez, Ferreira-Aparicio, and Daza [147] also showed that catalyst microstructure could be controlled by solvent composition and evaporation rate.

2.6 Prospects and Conclusion

The high cost of Pt and Pt-group metals (PGMs), along with low Pt utilization in the catalyst layer, is a major barrier to the commercialization of fuel cell technol-ogy. The strategies and methodologies to increase catalyst layer Pt utilization and reduce Pt loading constitute a major effort in fuel cell research and develop-ment. Thus far, researchers have found that these goals can be achieved through optimizing existing CLs with respect to composition and structure, developing novel fabrication technologies, and introducing innovative CL approaches.

In order to make catalyst layers with high platinum utilization and bet-ter performance, we need to determine how various factors affect Pt utiliza-tion. Although this objective has been receiving more attention, we have not achieved a fundamental understanding of the relationships of composition, structure, effective properties, and fuel cell performance—a fact that may limit the optimal design and fabrication of CLs.

Currently, optimization of catalyst layer composition and structure takes place through both experimental and modeling approaches. The experimen-tal approach has some limitations: It is tremendously time consuming and expensive, and trial-and-error experimental optimization is constrained to a narrow parameter range within a reasonable time frame. In addition, the optimal results obtained by experiments often vary from case to case and are difficult to compare with each other. With the rapid development of compu-tation technology, computational modeling can provide a powerful comple-mentary means of quickly and efficiently addressing some issues on a wider scale and in larger dimensions than experiments can handle. Pursuing the interplay between experimentation and modeling will be an effective way to approach fuel cell CL optimization.

New catalyst layer architectures and their corresponding fabrication technologies are highly sought to achieve a breakthrough in PEM fuel cell technology. The 3M-whisker CL structure is considered a promising new approach. The CNT-based catalyst layer structure is promising, but it still suffers from drawbacks such as carbon corrosion, inadequate fuel cell per-formance, and limited durability. To make the CNT-type CL commercially available, fabrication costs and scale-up must also be taken into account. In order to overcome catalyst carbon-support corrosion, noncarbon supports,


including electronically conducting ceramic oxides, should be explored. In addition, ultrathin CLs with controlled or ordered structures, in combina-tion with highly efficient catalyst-coating techniques, could be the avenue to the next generation of PEM fuel cells.

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107

3Proton Exchange Membranes

Timothy J. Peckham, Yunsong Yang, and Steven Holdcroft

CONTENTS

3.1 Introduction ................................................................................................ 1083.2 PEM Properties and Structure–Property Relationships ...................... 108

3.2.1 Proton Conduction ......................................................................... 1083.2.2 Oxygen Permeability and Methanol Crossover ........................ 119

3.2.2.1 Oxygen Permeability ...................................................... 1193.2.2.2 Methanol Crossover ........................................................ 122

3.2.3 Water Transport ............................................................................. 1273.2.4 Mechanical Properties and Chemical Stability of PEMs ......... 129

3.2.4.1 Mechanical Properties .................................................... 1293.2.4.2 Chemical Stability ........................................................... 131

3.3 Brief Overview of Existing PEM Materials ............................................ 1373.3.1 Statistical Copolymer PEMs ......................................................... 137

3.3.1.1 Perfluorinated and Partially Fluorinated .................... 1373.3.1.2 Polyarylenes ..................................................................... 1423.3.1.3 Miscellaneous Monolithic, Statistically

Sulfonated Copolymer PEMs ........................................ 1493.3.2 Block and Graft Copolymer PEMs .............................................. 150

3.3.2.1 Block Copolymers ........................................................... 1513.3.2.2 Graft Copolymers ............................................................ 155

3.3.3 Polymer Blends and Composite PEMs........................................ 1593.3.3.1 Polymer Blends ................................................................ 1613.3.3.2 Ionomer-Filled Porous Substrates and

Reinforced PEMs ............................................................. 1653.3.3.3 Composite PEMs for High-Temperature

Operation and Alternative Proton Conductors .......... 1663.4 Future Directions ....................................................................................... 170References ............................................................................................................. 171


3.1 Introduction

Proton exchange membranes (PEMs) are a key component in PEM fuel cells (PEMFCs) and an area of active research in commercial, government, and aca-demic institutions. In this chapter, the review of PEM materials is divided into two sections. The first will cover the most important properties of a membrane in order for it to perform adequately within a PEMFC. The latter part of this chapter will then provide an overview of existing PEM materials from both aca-demic and industrial research facilities. Wherever possible, the membranes will also be discussed with respect to known structure–property relationships.

3.2 PEM Properties and Structure–Property Relationships

In order to perform effectively within a PEMFC, a membrane should:

possess high proton conductivity (Section 3.2.1);be impermeable to gases (specifically oxygen) and/or fuel (e.g., methanol) (Section 3.2.2);achieve balanced water transport (Section 3.2.3);possess high thermomechanical and chemical stability to fuel cell conditions (Section 3.2.4); andbe an electrical insulator (not discussed).

Ideally, a membrane will have excellent performance in all of these areas. However, it is often found that PEMs will generally perform well in some of these areas while performing only adequately or even poorly in others. This section will present separate overviews for each of these properties as well as observed relationships between the chemical and morphological structures of the membranes.

3.2.1 Proton Conduction

One of the key parameters of a PEM is its ability to conduct protons through the membrane. This parameter is intimately connected with both acid and water content of the membrane and is also affected by the strength of the acid, the chemical structure and morphology of the mem-brane, and temperature. Understanding how all of these properties affect proton conductivity is crucial not only to an understanding of PEMs in general but also to obtaining more effective methods for developing new PEM materials.

Proton Exchange Membranes 109

Transport of cations in solution is usually thought to consist of a solvated cation diffusing through a solution. In addition to the vehicular transport of larger solvated cations (e.g., Na ), protons also move through solution via structural diffusion. This can be visualized as a chain mechanism in which protons are passed from one water molecule to its neighbor so that it appears the protons are migrating through the solution. This structural diffusion has also been called the Grotthuss mechanism. The actual process has been demonstrated by simulations1,2 and nuclear magnetic resonance (NMR) data interpretation3 to occur as shown in Figure 3.1, wherein the proton defect fol-lows the center of symmetry of the hydrogen-bond pattern. This “diffusion” is achieved by the formation and breaking of O-H bonds.4

In the case of PEMs, the situation is more complicated because the sul-fonate counter-ions (in the case of a PEM such as Nafion ) are bound to the polymer chain and are thus relatively immobile, in contrast to the free coun-ter-ion in a small molecule acid such as sulfuric or acetic acid. Tethering of the sulfonate group can be considered to be an impediment to the mobility of the proton as it traverses the membrane. Proton mobility is also affected by the effective mean-free path of connectivity of the conduction pathway as shown in Figure 3.2. In situation (a), the increased number of dead ends and tortuosity of the aqueous domains through which proton transport occurs over the situation in (b) leads to lower overall mobility. This has been demon-strated by Kreuer5 and will be discussed later in this section.

Formation ofHydrogen Bond

(Zundel-ion) (Eigen-ion)

H5O2+

H5O2+

H9O4+

(Zundel-ion)

+

+Breaking ofHydrogen Bond

+

FIGURE 3.1Proton conduction in water. (From Kreuer, K. D. et al. 2004. Chemical Reviews 104:4637–4678.)


In addition, the distance between acid groups may also have an effect upon proton mobility. With protons mediated via positively charge species between sulfonate groups, it is expected that larger distances between these tethered counter-ions will require greater energy in comparison to shorter distances, thus leading to lower mobility in the former case. This is sche-matically represented in Figure 3.3. Furthermore, the mobility of the proton is also affected by the degree of dissociation and its relationship to water content as described in the paragraphs that follow.

Proton conductivity, sH , can be related through the Nernst–Einstein rela-tionship to the activity of protons (aH ) in the membrane as well as to the mobility (mH ) of those protons:

H H HFa (3.1)

= –SO3H

(b)(a)

–

– –– – –

– – – – – – ––

–

–

– –– –

––

FIGURE 3.3Proximity of neighboring acid groups within an aqueous channel. Distance between acid groups in (a) is greater than in (b). (From Peckham, T. J. et al. 2007. Journal of Materials Chemistry 17:3255–3268.)

(a) (b)

FIGURE 3.2Connectivity of aqueous domains in PEMs (white aqueous domains) where the degree of tortuosity of the proton conduction pathway is greater in (a) than in (b). (From Peckham, T. J. et al. 2007. Journal of Materials Chemistry 17:3255–3268.)


where F Faraday’s constant. Both proton activity (and hence concentration) and mobility are dependent upon the strength of the acid groups as well as the water content of the membrane.

Acid strength dictates how easily the proton dissociates from the anionic counter-ion, which, in the case of PEMs, is bound to the polymer and is therefore immobile. In the majority of existing PEMs, the source of protons comes from a sulfonic acid group bound to either a perfluorinated ether (e.g., Nafion and similar materials) or aryl (e.g., polystyrene- or polyarylene-based systems) moiety. Calculations on model compounds, triflic acid, and p-toluenesulfonic acid (with acid ionization constant [pKa] values of –6 and –2, respectively)7 suggest that the proton dissociates itself from the sulfonate group when the number of water molecules per sulfonic acid group, l, equals three.8 It has also been calculated that in order for complete separation to be achieved between the proton and the sulfonate group, the water con-tent must be higher (l 6). Furthermore, greater separation exists between the proton and the sulfonate group in triflic acid than in p-toluenesulfonic acid.8

The effect of acid content upon conductivity is frequently displayed as shown in Figure 3.4, where conductivity is plotted as a function of the ion-exchange capacity (IEC) of the membrane. As can be seen in the figure, pro-ton conductivity is strongly dependent upon acid content for all the PEMs. This is not surprising, given that conductivity is dependent upon proton

Nafion0.25

0.2

0.15

0.1

0.05

Co

nd

uct

ivit

y (S

/cm

)

00.00 0.50 1.00 1.50 2.00

IEC (meq/g)

2.50 3.00 3.50

SPEEKBAM membraneETFE-g-PSSA

FIGURE 3.4Proton conductivity as a function of IEC for ETFE-g-PSSA polyethylenetetrafluoroethylene-graft-polystyrene sulfonic acid, BAM membrane substituted poly(trifluorostyrene) sulfo-nic acid, SPEEK sulfonated poly(ether ether ketone) and Nafion. (From Peckham, T. J. et al. 2007. Journal of Materials Chemistry 17:3255–3268, and Dolye, M. et al. 2001. Journal of Physical Chemistry B 105:9387–9394.)


concentration (Equation 3.1), which itself is based in part upon acid content (i.e., IEC) of the PEM. Also, as generally expected based on Equation 3.1, increasing acid content leads to higher conductivity values. This behavior is exhibited for ETFE-g-PSSA and SPEEK (sulfonated poly (ether ether ketone)). However, in the case of both Nafion and BAM membranes, a maximum conductivity is reached at IEC ~ 1.0 meq/g and 2.0 meq/g, respectively.

In order to understand the underlying causes for these differences in behavior, it is necessary to examine how conductivity is affected by water content. This can be seen in Figure 3.5.

In the region of l ~ 10–20, water has a beneficial effect upon conductiv-ity. This can be understood in terms of proton mobility, whereby increasing water content leads to greater separation between the protons and the bound sulfonate groups and hence the protons are more easily able to migrate from anode to cathode via the Grothus mechanism as described earlier. Proton mobility has also been shown to continue to increase as a function of increas-ing water content for ETFE-g-PSSA, BAM membrane, and SPEEK.6 This was also shown for Nafion and sulfonated poly(ether ether ketone ketone) (SPEEKK) by Kreuer, as seen in Figure 3.6.

At low water contents, DH2O decreases more rapidly for SPEEKK than for Nafion. Proton mobility, Ds, behaves in a similar fashion. When water content is high, Ds is higher than DH2O as a result of the influence of inter-molecular proton transfer. However, at low water contents, the degree of dis-sociation of the proton from the sulfonate groups decreases, thus leading

FIGURE 3.5Proton conductivity as a function of water content (l) for ETFE-g-PSSA, BAM membrane, SPEEK, and Nafion. (From Peckham, T. J. et al. 2007. Journal of Materials Chemistry 17:3255–3268, and Dolye, M. et al. 2001. Journal of Physical Chemistry B 105:9387–9394.)


to increased localization of the proton. Because aryl-bound sulfonic acid groups have lower pKa values than those for the significantly more acidic groups in Nafion, this decrease in Ds is greater for SPEEKK.5

Based on proton mobility alone, it would thus be expected that proton conductivity would increase as a function of increasing water content. However, as already observed in the case of BAM membrane and Nafion (Figure 3.5), this is not always the case. Given that conductivity is also dependent upon proton concentration, it might be expected that some of the decrease in conductivity may be due to changes in proton concentra-tion. As can be seen in Figure 3.7, all of the PEMs (suitable data for Nafion were unavailable) exhibit a decrease in proton concentration with increas-ing water content. However, in the case of BAM membrane, the increase in water content is much higher for similar acid contents versus either ETFE-g-PSSA or SPEEK, so the effective proton concentration is much lower. The result is lower conductivity values, even though it has been shown that mobility values for BAM membrane are generally higher than those for Nafion, ETFE-g-PSSA, or SPEEK.6

In the case of Nafion, a similar situation occurs. There is a sharp increase in proton conductivity and proton concentration as a function of water con-tent followed by a decrease at l > 20. At these higher water contents, Nafion undergoes a similar dilution of proton concentration per BAM membrane in conjunction with a lower mobility value versus ETFE-g-PSSA. However,

1e–3

1e–4

T = 300 K

NAFION Pure Water

PEEKK(70% sulfonated)

Dσ

DH2O

Dσ

DH2O

Dσ

DH2O

1e–5

1e–6

D/c

m2s–

1

1e–7

1e–8

1e–3

1e–4

1e–5

1e–6

1e–7

1e–8

0.1

Water Volume fraction XV

1

FIGURE 3.6Proton mobility (Ds) and water self-diffusion coefficient (DH2O) as a function of the water vol-ume fraction (Xv) in Nafion and SPEEKK, where Xv volume of water in membrane divided by volume of wet membrane. (From Kreuer, K. D. 2001. Journal of Membrane Science 185:29–39.)


at l ~ 15, the conductivity of Nafion undergoes an order of magnitude increase relative to its conductivity at lower water contents. The explanation for this is connected to how Nafion takes up water with increasing acid content. Over the IEC range of 0.7–1.0 meq/g, water content remains consistent at l ~ 15. There is no relative increase in water content, yet acid content is increasing; thus, pro-ton concentration effectively increases, thereby leading to higher conductivity values. At IEC > 1.0 meq/g, water uptake begins to increase significantly and proton concentration effectively decreases—and thus so does conductivity.

From these examples, it can be seen that water content has a strong effect upon proton conductivity. Thus, it is clear that water management is an important factor for efficient PEMFC operation. It will be discussed in Section 3.2.3.

Two other important factors that control the conductivity of PEMs are polymer microstructure and morphology. Within this section, Nafion will serve as the prime example to describe how the formation of hydrophobic and hydrophilic domains relates to proton transport. The microstructures of a few PEMs will then be described to highlight the importance of this area upon proton conductivity.

In Nafion, the hydrophobic perfluorinated segments of the polymer are incompatible with the hydrophilic sulfonic acid groups and thus phase sep-aration occurs. When exposed to water, the hydrophilic domains swell to provide channels for proton transport, whereas the hydrophobic domains provide mechanical integrity and, at least in the case of lower IEC samples,

FIGURE 3.7Proton concentration as a function of water content (l) for ETFE-g-PSSA, BAM membrane, SPEEK, and Nafion. (From Peckham, T. J. et al. 2007. Journal of Materials Chemistry 17:3255–3268, and Dolye, M. et al. 2001. Journal of Physical Chemistry B 105:9387–9394.)


resistance against dissolution. Studies by small-angle x-ray scattering (SAXS) and small-angle neutron scattering (SANS)—with many of the primary stud-ies conducted by Gebel,10 Gebel, Aldebert, and Piner,11 Gebel and Labard,12 and Rubatat et al.13—have been the main methods for elucidating the micro-structure of the polymer with a range of polymer-to-solvent ratios in addi-tion to using different solvents. The work by Hsu and Gierke demonstrated that ionic clusters form in Nafion membranes and theorized that inverted micelles are also formed wherein the composition consists of hydrated clus-ters (40–50 Å diameter) of acid groups within a fluorocarbon phase.14 Studies by Gebel10 have subsequently led to the proposed structural evolution of the microstructure of Nafion from the dry to dissolved state, as shown in Figure 3.8.

Based on Gebel’s calculations for Nafion (where IEC 0.91 meq/g),10 iso-lated spheres of ionic clusters in the dry state have diameters of 15 Å and an intercluster spacing of 27 Å. Because the spheres are isolated, proton transport through the membrane is severely impeded and thus low levels of conductivity are observed for a dry membrane. As water content increases, the isolated ionic clusters begin to swell until, at Xv > 0.2, the percolation threshold is reached. This significant point represents the point at which connections or channels are now formed between the previously isolated ionic clusters and leads to a concomitant sharp increase in the observed level of proton conductivity.

With increasing water content, the ionic domains swell from 40 to 50 Å in diameter and the structure of the membrane is thought to consist of spherical ionic domains joined by cylinders of water dispersed in the polymer matrix. Within this region of water content, proton conductivity steadily increases. At Xv > 0.5, a morphological inversion occurs in which a connected network of aggregated polymer “rods” is now surrounded by water. This network continues to swell for Xv 0.5 0.9 and the conductivity of the membrane approaches the values observed for Nafion solutions.

SAXS analyses have also been performed on sulfonated aromatic poly(arylenes) such as sulfonated poly(ether ketone). From these studies, it has been determined that there is a smaller characteristic separation length with a wider distribution and a larger internal interface between the sepa-rated hydrophilic and hydrophobic microdomains, corresponding to a larger average separation between neighboring acid sites. A schematic illustra-tion as described by Kreuer for the phase-separated structure of sulfonated poly(ether ether ketone ketone) (SPEEKK) can be seen in Figure 3.9 with Nafion for comparison. In general, it can be seen that the water-filled chan-nels through which proton transport is thought to occur are narrower than those in Nafion, as well as exhibiting greater degrees of branching and dead ends and less separation between the channels. This leads to a more tortuous pathway for proton conduction in SPEEKK versus Nafion.

Direct visualization of the microstructure of Nafion has been accomplished using transmission electron microscopy (TEM). Images of dehydrated


Dry Membrane

Swollen Membrane

Percolation

*Structure Inversion*

Connected Network

of Polymers Rods

Colloidal Dispersion of

Rod Like Particles

Solution

0.75

0.50

0.25

0

Water VolumeFraction

Perfluorinated Matrix

Ionic Domains

FIGURE 3.8Structural evolution of Nafion microstructure as a function of water content. (From Gebel, G. 2000. Polymer 41:5829–5838.)


membrane show the presence of roughly spherical ionic clusters (3–10 nm in diameter).15–17 Additional information was obtained by examining very thin films of cast Nafion. Micrometer-sized regions of PTFE were seen to be scat-tered randomly throughout the film with ionic clusters approximately 5 nm in diameter. Examination of these films by electron diffraction showed that Nafion possesses a similar structure to that of polyethylene, suggesting that the backbone of Nafion has a linear, zigzag pattern like that of polyethylene but unlike the twisted chain encountered in poly(tetrafluoroethylene).17

Given that TEM conditions result in dehydration of the membrane, atomic force microscopy (AFM) has also been used to examine the morphology of Nafion because this technique can be used with samples where humidity is varied. Although AFM is limited to surface studies, interesting information has nonetheless been obtained. Low-energy phase images of Nafion N117-K were taken for samples exposed to ambient relative humidity as well as liquid water.18 In the former case, ionic clusters exhibited a uniform distribution of ionic clusters

Wide channels

: H2O

: -SO3–

–

+

1 nm

O

Sulfonated Polyetherketone (PEEKK)Nafion

O

O

O

SO3H

: Protonic

charge

carrier

Narrow channelsLess separatedHighly branchedDead-end channels

Large –SO3–/-SO3

–

separationpKa ~ –1

More separatedLess branchedGood connectivitySmall –SO3

–/-SO3

–

separationpKa ~ –6

— CF2—CF2—CF—CF2—n

OCF2CF—O(CF2)2SO3Hm

CF3

FIGURE 3.9Schematic representation of microphase separation in Nafion and SPEEKK. (From Kreuer, K. D. 2001. Journal of Membrane Science 185:29–39.)


(4–10 nm). For the samples exposed to liquid water, the ionic regions appeared to coalesce into channels; the narrowest dimensions were on the order of 7–15 nm. It was speculated that when the ionic clusters swell in water, the channel-like morphology results from the constraints placed upon the ionic domains by the crystalline regions. It was also noted that because the sulfonyl fluoride analogue has no ionic domains, it does not yield any image contrast.18

More recently, electrochemical AFM has been used to study further the mor-phology of Nafion while also examining proton conductive regions.19 This is accomplished by applying anode catalyst to one side of the membrane and then measuring proton conductivity via an AFM tip that has been modified with a platinum electrode to act as the cathode catalyst for the fuel cell reaction (Figure 3.10). Because protons will only be transported when the circuit is com-pleted, imaging of the surface is coupled not only with its topology but also with regions where proton transport may occur (i.e., ionic domains). In addition, it is possible to measure the current as a function of the applied voltage through the AFM tip. The apparatus is further contained within an ethat allows for changes in the water content of the PEM by variation in relative humidity (RH) values. At low RH (35%), ionic channel size decreases and fewer ionically active areas are observed at the surface, suggesting that some channels have collapsed. At high RH (80%), the ionic domains are presumably swollen and the surface hydro-philic regions are considerably larger. In common with the prior AFM studies, however, the technique is still basically limited to surface studies.

Studies on morphology and conclusions about observed levels of proton con-ductivity have also been carried out on PEMs other than Nafion and sulfonated poly(ether ketone). These include studies in which phenomenological exami-nations of relationships between conductivity and observed microstructure were carried out upon polymer systems where acid content was varied but the basic chemical structure was kept constant. In addition, other systems allowed

Pt-AFM Tip

Membrane

Electrode

A

U

Cathode: O2 + 4H+ + 4e– 2H2O

Anode 2H2O O2 + 4H+ + 4e–

FIGURE 3.10In situ method for the measurement of proton conductivity using electrochemical AFM. (From Aleksandrova, E. et al. 2007. ChemPhysChem 8:519–522.)


comparisons between random and regular distributions of acid groups. These studies will be described in Section 3.3 as part of the review of existing PEMs.

Finally, temperature also has a direct effect upon observed levels of proton conductivity. In part, this is due to the expected relationship between the kinetic energy of the protons and their transport. However, as discussed previously, water content also has a strong effect upon PEMs such as Nafion in which pro-ton transport is mediated by sulfonic acid groups. Because proton transport for this type of acid occurs mainly when water is in the liquid phase, this imposes restrictions upon the effective temperature range for proton transport. This can be seen in Figure 3.11 for Nafion in comparison to other functional groups with different effective operating temperature ranges for proton conduction.

3.2.2 Oxygen Permeability and Methanol Crossover

3.2.2.1 Oxygen Permeability

The efficiency of fuel cells is dependent on the degree of reactant crossover through the membrane. The permeation processes through membranes are governed by the solubility and diffusion of the permeating solute in the mem-brane at any given temperatures and conditions. The solubility of the solute in polymeric membranes is dependent on the chemical nature of the solute and of the corresponding membrane, while the diffusion is determined largely by the morphology of the membrane and the properties of the solute. The

FIGURE 3.11Conductivity of Nafion in comparison to some intermediate-temperature proton conductors and the oxide ion conductivity of YSZ (yttria-stabilized zirconia). (From Kreuer, K. D. et al. 2004. Chemical Reviews 104:4637–4678.)


resulting permeation rate is then determined by a complex interplay between the properties of the system, including the morphology of the polymer, restrictions on the ability of the polymer to swell, and the chemical properties of the solute and polymer, such as hydrophilicity or hydrophobicity.

It has already been recognized that gases dissolve in both the hydrophobic and the swollen hydrophilic domain, whereas most of the gas transport occurs within the swollen hydrophilic domain.20–22 The barrier properties of different PEMs for different gases have been widely studied, mostly with regard to the membrane hydrogen and oxygen permeability for application in a hydrogen fuel cell and methanol permeability for application in a direct methanol fuel cell (DMFC). In general, the barrier properties of different membrane types for a hydrogen fuel cell have been satisfactory despite a minor concern for the side reactions on the electrodes due to permeated fuel or oxygen.23–25 In consideration of space constraints, oxygen permeability will be the focus of this section.

In a PEMFC, the power density and efficiency are limited by three major factors: (1) the ohmic overpotential mainly due to the membrane resistance, (2) the activation overpotential due to slow oxygen reduction reaction at the electrode/membrane interface, and (3) the concentration overpotential due to mass-transport limitations of oxygen to the electrode surface.26 Studies of the solubility and concentration of oxygen in different perfluorinated membrane materials show that the oxygen solubility is enhanced in the fluorocarbon (hydrophobic)-rich zones and hence increases with the hydrophobicity of the membrane. The diffusion coefficient is directly related to the water content of the membrane and is thereby enhanced in membranes containing high water content; the result indicates that the aqueous phase is predominantly involved in the diffusion pathway.26,27

Oxygen reduction and transport characteristics were studied for Nafion 117, BAM, sulfonated polystyrene-b-poly(ethylene-r-butylene)-b-styrene (S-SEBS), and ethylene-tetrafluoroethylene copolymer (EFTE)-g-PSSA membranes by Basura et al.;28 Chuy et al.;29 Basura, Beattie, and Holdcroft;30 and Beattie, Basura, and Holdcroft,31 and disulfonated biphenol poly(arylene ether sulfone) (BPSH) membranes were studied by Zhang et al.,32–34 using microelectrode techniques in an environment that mimics PEMFCs. For these membranes, the effect of pressure on oxygen mass transport properties was found to obey Henry’s law to a first approximation: As the pressure increased, the oxygen solubility (C) increased. O2 diffusion coefficients (D) were found to increase with increasing temperature; although the opposite trend is true for C, the increase in D with temperature is greater than the corresponding decrease in C. Therefore, the overall permeability (D C) increases linearly with temperature.

The effects of equivalent weight (EW g polymer/mol ~ SO3H) and water content on diffusion coefficient, solubility, and permeability of oxygen for fully hydrated BAM, S-SEBS, ETFE-g-PSSA, Nafion 117, and BPSH mem-branes have been studied. It has been found that the diffusion coefficients of all the studied membranes decrease with increasing EW, while the solu-bility correspondingly increases. These trends are the same as found in


previous research for perfluorinated membranes.26,27 The oxygen diffusion coefficient is affected to a greater extent by increasing EW than is solu-bility, possibly because oxygen diffusion in high EW/low water content membranes is significantly affected when connection between aqueous domains is compromised. As a consequence, the permeability of oxygen falls with increasing EW.

It has also been found that, generally, D decreases, C increases, and DC decreases with decreasing water content. In addition, there is reasonable overlap of the data for all membranes and a clear correlation with water con-tent. However, it must be clarified that the water contents were measured at 298 K and represent “wet” (i.e., fully saturated) membranes. Thus, the cor-relation is with the maximum solubility of oxygen within the membrane, whereas the electrochemical mass transport data were taken at 323 K under conditions of 100% RH.

O2 diffusion through the membrane seems to be limited by the percolation network of the diffusion path, which is not only defined by the amount of water in the membrane, but also by the different chemical structure of the membranes. It is difficult to make comparisons of gaseous diffusion behavior among polymers with different structures because polymer morphology can change drastically without appreciable changes in density, and the presence of water and the hydrogen bonds formed between polymer–water moieties also has major effects on system properties.35 However, some points can be made from these particular studies.

It was found that the oxygen diffusion coefficient of BAM membranes is much more dependent on water content than that of EFTE-g-PSSA, S-SEBS, and BPSH membranes. Nafion 117 has a higher diffusion coefficient than BPSH with EW 833 g mol–1, even though Nafion 117 has a much lower water content, with different states of water possibly one of the reasons (see Section 3.2.3).36 BPSH has a lower oxygen diffusion coefficient than other membranes with the same water content, and this might also be associated with its rigid backbone structure compared to other polymers. Under low water content, Nafion 117, BAM, and BPSH membranes exhibit much higher oxygen solubil-ity than S-SEBS and ETFE-g-PSSA do. However, at high water content, oxygen solubility is less affected by the chemical structure of membranes because oxy-gen solubility appears to approach a constant value at high water content.

It was also found that oxygen permeability of Nafion 117 is higher than that of other membranes with the same water content, even though oxygen diffusion of Nafion 117 is relatively slow. This is offset by the correspond-ing high oxygen solubility. For the oxygen permeability of ETFE-g-PSSA and BPSH membranes, the values are surprisingly low, given that their water contents are large (up to 74 and 78 vol%, respectively). Closer inspec-tion of the mass transport data suggests that low permeability appears to be due to poor oxygen diffusion rather than poor oxygen solubility. This is unexpected, given that high water contents are shown to facilitate oxygen diffusion.


However, calculation of the ratio of moles of water to ions reveals that the l values for ETFE-g-PSSA and BPSH membranes are much lower than those obtained for PTFSSA and S-SEBS membranes for similar water contents. ETFE-g-PSSA and BPSH membranes contain a large volume of water, but because the large concentration of ions in the membrane, much of the water is associated with the solvation of ions. Furthermore, it has been seen that the amount of free water in Nafion is higher than in BPSH in spite of the fact that the total water uptake is lower. In the absence of a significant fraction of “free” water, it appears that oxygen diffusion is impeded. BPSH membranes exhibit lower oxygen permeability than other membranes, despite the higher oxygen solubility. This may be associated with its shortage of free water in membrane36 and its rigid backbone. In a recent report, low oxygen perme-ability was also found for SPEEK (at 30 C, 8.7 10–12 mol cm–1 s–1 for SPEEK with IEC 1.88 mmol/g).37

Oxygen37,38 and hydrogen37 permeabilities of recast Nafion films have also been studied. Heat treatment results in morphological changes for recast Nafion film, and the oxygen permeability properties changed significantly around the Tg of Nafion film.38 The oxygen diffusivity increases and the oxy-gen solubility decreases with the decrease in the recast temperature when the heat treatment temperature is below the Tg of Nafion film. The oxygen permeability also exhibits the same tendency as the oxygen diffusivity. However, changes in these properties are not observed when the heat treat-ment is above the Tg of Nafion film.

3.2.2.2 Methanol Crossover

Sluggish methanol (MeOH) oxidation reaction kinetics and MeOH cross-over through the membrane are the two major technical hurdles for DMFC technology. Generally, MeOH crossover in the fuel cell may be defined as the phenomenon of MeOH permeating from the anode com-partment through the membrane to the cathode compartment. The issue of MeOH crossover in DMFC would result not only in fuel loss, but also in an increase in air demand and in a decrease of the cell efficiency due to the reactions and depolarization of permeated MeOH with oxygen at the cathode. Also, the excessive permeation of water to the cathode in liq-uid feed DMFCs associated with MeOH crossover leads to serious water accumulation on the cathode, necessitating high air flows to alleviate flooding effects.

MeOH is transported through the membrane by two modes: diffusion and electro-osmotic drag.39,40 When MeOH comes into contact with the mem-brane, it diffuses through the membrane from anode to cathode and is also dragged along with the hydrated protons under the influence of current flowing across the cell. Therefore, a correlation between the MeOH diffusion coefficient and proton conductivity is observed. The diffusive mode of MeOH transport dominates when the cell is idle, whereas the electro-osmotic drag


dominates when the cell is operating, which means that current is flowing across the cell. MeOH transport through the membrane is accomplished by moving through the ion-cluster pores and the connecting ion channels; in the hydrophobic PTFE region, MeOH has negligible solubility.

The MeOH crossover rate is closely related to several factors, including membrane structure and morphology, membrane thickness, membrane acid content, and the cell operating parameters, such as temperature and MeOH feed concentration. The MeOH crossover rate decreases with an increase in the thickness and can be reduced greatly by using a membrane with suffi-ciently high EW.41–43 Therefore, it may be advantageous to use either thicker or higher EW membranes to reduce the MeOH crossover rate.

However, the disadvantage is the penalty of higher voltage losses due to higher specific resistance as a result of thicker and higher EW membranes. The voltage loss is particularly severe at higher current density operation. Moreover, a thicker membrane means increasing material cost. Ideal mem-branes for DMFC would have no MeOH crossover, but would have some water transport to the cathode to prevent drying of the cathode catalyst layer. Due to the similarity of the MeOH and water molecules, it is difficult (if not impossible) to decouple MeOH from water transport in a DMFC.

Nafion absorbs MeOH more selectively than water,44 and the MeOH dif-fusion flow is higher than the osmotic water flow in Nafion membranes.45 Diffusion coefficients of Nafion 117 determined by different techniques have been reported. Ren et al.42 measured MeOH diffusion coefficients in Nafion 117 membranes exposed to 1.0 M MeOH solutions using pulsed field gradi-ent (PFG) NMR techniques. The MeOH self-diffusion coefficient was 6 10–6 cm2 S–1 and roughly independent of concentration over the range of 0.5–8.0 M at 30 C. A similar diffusion coefficient was obtained for Nafion 117 at 22 C by Hietala, Maunu, and Sundholm with the same technique.46 Kauranen and Skou determined the MeOH diffusion coefficient of 4.9 10–6 cm2 S–1 for Nafion 117 in the presence of 2.0 M H2SO4 solutions at 60 C by using an electrochemi-cal technique.47

MeOH permeability of Nafion 117 membranes was investigated by Tricoli, Carretta, and Bartolozzi, who used a two-compartment glass cell.48 In this setup, one compartment (VA) was filled with a solution of MeOH (8 vol%) and 1-butanol (0.2 vol%) in deionized water. The other (VB) was filled with a 1-bu-tanol (0.2 vol%) solution in deionized water. The membrane was clamped between the two compartments and the solutions were stirred during the experiment. A MeOH flux was established across the membrane owing to the concentration difference between the two compartments. MeOH perme-ability ranges from 1.0 10–6 cm2 S–1 to 1.49 10–6 cm2 S–1 were determined for Nafion 117 membranes at 22 C depending on different pretreatment pro-tocol. High crossover in Nafion membranes is one of the main driving forces to find an alternative membrane for DMFC.

Kim, Kim, and Jung measured proton conductivity and MeOH permeabil-ity for a series of S-SEBS membranes.49 Both proton conductivity and MeOH


permeability increase with increasing degree of sulfonation. A percolation threshold at ~15 mol% sulfonation degree was found for both. However, the increase in permeability does not slow down after reaching the percolation limit. The ratios of proton conductivity to MeOH permeability, which are normally considered as an efficiency indicator, were calculated. MeOH per-meability for Nafion 117 was measured for comparison reasons and found to be 2.6 10–6 cm2 s–1. When comparing this value with the MeOH perme-ability (1.2 10–6 cm2 s–1) of 34 mol% S-SEBS, whose proton conductivity is slightly higher than that of Nafion 117, it can be concluded that the MeOH permeability of S-SEBS is lower by more than a half of Nafion 117. Figure 3.12 shows the ratios of proton conductivity to MeOH permeability, Δ, for various S-SEBS as a function of sulfonation degree.

The efficiency of 34% S-SEBS is twice as high as that of Nafion 117. MeOH permeability decreases with increasing of sulfonation degree, and this indi-cates that MeOH permeability increases faster with increasing sulfonation degree than proton conductivity does. It can be anticipated that the efficiency indicator of S-SEBS with sulfonation degree > ~50% is lower than that of Nafion 117. The water and MeOH uptakes of S-SEBS membranes were directly evaluated, and the results showed that the membranes favor MeOH much more than water, as does Nafion 117 (Figure 3.13); this is in agreement with the result that Δ decreases with increasing degree of sulfonation. Therefore, it would not be an easy task to develop a promising proton exchange mem-brane from S-SEBS that could be utilized for DMFC.

Sulfonated polystyrene-b-(isobutylene)-b-sulfonated polystyrene (S-SIBS) with IEC between 0.5 and 1.0 mmol/g also exhibits approximately 5–10 times higher selectivity than Nafion 117.50 The increased selectivity is thought to be due to lower MeOH and water solubility in polyisobutylene, the major

0 10 20

Sulfonation Degree (mol%)

Nafion 117

30 40 50

0

2

4

Δ ×

10

6 (

Ω–

1cm

–3s)

6

FIGURE 3.12Ratios of proton conductivity to MeOH permeability, Δ, for various S-SEBS membranes as a func-tion of sulfonation degree. (From Kim, J. et al. 2002. Journal of Membrane Science 207:129–137.)


component of the triblock copolymer ionomer examined in this study, and the ordered structure of the block copolymer (lamellar).50

PSSA-grafted polymers based on poly(vinylidene fluoride) (PVDF), ETFE, and low-density polyethylene (LDPE) substrates with <52% grafting degree exhibit lower MeOH diffusion coefficient than Nafion 117 at 25 C.51,52 Hatanaka et al. investigated the MeOH and water uptake and MeOH perme-ability of ETFE-g-PSSA membranes in comparison to Nafion membranes.53 The MeOH and water uptakes at room temperature were proportional to the degree of grafting (IEC). The MeOH uptake of ETFE-g-PSSA was half of those of Nafion membranes, although the water uptakes of both membranes were almost the same. MeOH permeability of ETFE-g-PSSA membranes was about half of that of Nafion in all temperature ranges, and the activation energies of MeOH permeability in these membranes were 16.7–19.3 kJ/mol for Nafion and 18.5–18.9 kJ/mol for the grafted membranes.

For ETFE-g-PSSA membranes with the same IEC, water uptake is higher than MeOH uptake of the membrane, but for Nafion and S-SEBS membranes, MeOH uptake of membrane is always higher than water uptake. Chemical structure and morphology of membranes affect the solvent absorption. Nafion is considered to consist of ionic clusters that are separated from the polymer phase. For grafted polymers, heterogeneity exists to some extent due to the hydrophobic base polymer; however, a regular clustered structure, as in the case of Nafion, has not been proposed for these materials.

Every et al. determined diffusion coefficients of Nafion 117 and BPSH 40 using a modified PFG NMR method.54 Nafion 117 and BPSH 40 were immersed in a MeOH solution of known concentration, and diffusion coef-ficients were measured as a function of MeOH concentration. For a Nafion

0 10 20

Sulfonation Degree (mol%)

30 40

0

3

6

Met

han

ol

to W

ater

Up

tak

e R

atio 9

FIGURE 3.13Ratios of MeOH to water uptakes for S-SEBS membranes as a function of sulfonation degree. (From Kim, J. et al. 2002. Journal of Membrane Science 207:129–137.)


117 membrane immersed in a 0.5 M MeOH solution, the MeOH diffusion coefficient in the membrane is 2.9 10–6 cm2 s–1. Increasing the MeOH con-centration results in a slight increase in the MeOH diffusion coefficient within the membrane, reaching a value of 4.0 10–6 cm2 s–1 at a concentra-tion of 4 M.

The MeOH diffusion coefficients in BPSH 40 exhibit remarkably different behavior. At low MeOH concentrations, the diffusion coefficient is 7.7 10–6 cm2 s–1, but it drops significantly to 2.5 10–6 cm2 s–1 at 4 M. In the range of 4–8 M, MeOH diffusion coefficients of both Nafion 117 and BPSH 40 are roughly independent of MeOH concentration. This trend of decreasing MeOH diffu-sion coefficients upon increasing MeOH concentration is, on the face of it, ideal for DMFCs. The necessary corollary proviso is that the concentration in the membrane increases to a lesser extent than the diffusion coefficient decreases, yielding a net permeability decrease.

In comparing the diffusion behavior of these two membranes at low MeOH concentrations, MeOH in BPSH 40 membranes exhibits significantly higher diffusion coefficients than those in Nafion 117. This may be the result of dif-ferences in morphology. Tapping mode AFM measurements found that, for dry membranes, the domains for BPSH 40 appeared to be 10–25 nm in diam-eter; for Nafion 117, the domains were smaller, approximately 4–10 nm.55 Although one would expect restricted diffusion in both cases, it is possible that the smaller domains limit diffusion to a greater extent.

As the concentration of MeOH increases, the divergent diffusion behavior between the two membrane types is a reflection of the difference in MeOH solubility and its concentration dependence within each membrane. This was verified by solvent uptake measurements. Upon increasing MeOH con-centration, Nafion 117 showed a steady increase in mass, while a sharp drop in total solution uptake was observed for BPSH 40. The lower viscosity of MeOH also affects the fluidity of the solution within the pores. The constant solvent uptake and the increased fluidity of the more concentrated MeOH solutions accounted for the slight increase in diffusion coefficient of Nafion 117. For BPSH 40, increasing the MeOH concentration resulted in a decrease in MeOH diffusion. The solvent uptake measurements showed very similar behavior, indicating that the membrane excludes the solvent upon exposure to higher MeOH concentrations.

In contrast, diffusion of MeOH measured via permeability measurements (assuming a partition coefficient of 1) was lower (1.3 10–6 and 6.4 10–7 cm2 s–1 for Nafion 117 and BPSH 40, respectively) and showed no concentration depen-dence. The differences observed between the two techniques are related to the length scale over which diffusion is monitored and the partition coef-ficient, or solubility, of MeOH in the membranes as a function of concentra-tion. For the permeability measurements, this length is equal to the thickness of the membrane (178 and 132 μm for Nafion 117 and BPSH 40, respectively), whereas the NMR method observes diffusion over a length of approximately 4–8 μm.


3.2.3 Water Transport

The transport of water through a PEM consists of a number of different modes. These are schematically illustrated in Figure 3.14.

The net water flux through the PEM is a combination of electro-osmotic drag and diffusion. Electro-osmotic drag (EOD) is due to the waters of hydration that are transported through membrane as the protons move from anode to cathode, and it is defined by the EOD coefficient, hdrag. In addition to water that can diffuse in the same direction, back diffusion of water to the anode is also possible due to the water gradient that occurs with accumulation of water at the cathode and dehydration at the anode due to EOD. As these processes affect the water content gradient of the membrane, they also have a strong effect upon proton conductiv-ity, which is highly dependent upon water content in sulfonic acid-based PEMs.

The EOD coefficient, hdrag, is the ratio of the water flux through the mem-brane to the proton flux in the absence of a water concentration gradient.57 As hdrag increases with increasing current density during PEMFC operation, the level of dehydration increases at the anode and normally exceeds the abil-ity of the PEM to use back diffusion to the anode to achieve balanced water content in the membrane. In addition, accumulation of water at the cathode leads to flooding and concomitant mass transport losses in the PEMFC due to the reduced diffusion rate of O2 reaching the cathode.

A number of different methods have been developed to measure hdrag. These include:

streaming potential measurements58,59;concentration cells57,60;water flux measurements61;DMFC experiments62,63; andNMR spectroscopy.64

FIGURE 3.14Modes of water transport in an operating H2/O2 PEMFC. (From Siv, A. 2007. Ph. D. Dissertation, Simon Fraser University.)


Unfortunately, the value for hdrag depends upon the type of measurement and also varies according to the chemical nature of the membrane, its acid content (IEC), temperature, and water content. Increasing temperature leads to higher hdrag values, whereas reducing membrane water content via lower RH values results in decreasing hdrag values. With regard to membrane-based comparisons, Nafion membranes exhibit higher hdrag values than aromatic-based polymers, such as SPEEKK, for similar IEC values and water content. The reason behind this is thought to be the smaller channels in polymers such as SPEEKK. However, even when channel sizes are allowed to be simi-lar, Nafion still has a higher value, which indicates that polymer–water inter-actions also have a significant effect upon EOD.

Water self-diffusion coefficients (DH2O) have been determined by pulse field gradient NMR for Nafion (N117)61, BPSH,65 and SPEEKK4 and have been shown to increase with increasing acid and water content. At high water contents, the DH2O values for PEMs approach the corresponding value for pure water. This is due to the increase in volume fraction of free water. At low water content, however, Nafion exhibits greater DH2O values than either BPSH or SPEEKK. In common with the observations for EOD, this has been attributed to the smaller channels in aromatic-based polymers, leading to a considerably lower dielectric constant for the water in the channels.

The nature of the water present within a PEM can also have an effect upon its performance during PEMFC operation. At l > 6, water exists in the three forms previously mentioned in Section 3.2.1: free water, loosely bound water, and nonfreezable water. This has been established by Fourier transform infrared (FTIR) studies66 and more recently by calorimetry and gravimetric analysis.67 Pulse 1H NMR has also been used in conjunction with differential scanning calorimetry (DSC) to analyze the contributions of these three dif-ferent types of water in Nafion and BPSH membranes.36 These values can be seen in Table 3.1.

TABLE 3.1

Contributions to Total l from Free, Loosely Bound, and Nonfreezable Water for Nafion and BPSH Using a Combination of Pulse 1H NMR and DSC

l

PEMsIEC

(meq/g)Water

uptake (%) Total FreeLoosely bound Nonfreezable s (S/cm)

Nafion N1135 0.9 33 20.2 4.9 13.1 2.2 0.11BPSH 0.5 7 8.1 0 4.9 3.2 —

0.9 18 10.8 0 7.2 3.6 0.011.3 29 12.0 0 7.9 4.1 0.051.7 56 18.0 2.2 10.8 5.1 0.082.0 124 33.1 3.7 24.5 4.9 0.11

Source: Kim, Y. S. et al. 2003. Macromolecules 36:6281–6285.


The effect that the amount of free water has upon proton conductivity can be seen from these data. For BPSH membranes where IEC < 1.3 meq/g, there is no free water and conductivity is low, even though at least two samples have IEC values equivalent to or higher than Nafion. In the case of BPSH where IEC 2.0 meq/g, it has a higher l value and more loosely bound water than Nafion. However, the amount of free water possessed by Nafion is higher than for BPSH (IEC 2.0 meq/g) and conductivity is equivalent for the two membranes, despite the considerably higher acid content for the latter.

3.2.4 Mechanical Properties and Chemical Stability of PEMs

The stability of membranes against thermomechanical and chemical stresses is an important factor in determining both their short- and long-term perfor-mance. Transport and mechanical properties of membranes affect the fuel cell performance, while the lifetime of a fuel cell is mostly dependent on the ther-momechanical and chemical stability of the membrane. Thermomechanical and chemical degradation of a membrane will result in a loss of conductivity, as well as mixing of anode and cathode reactant gases.

3.2.4.1 Mechanical Properties

Thin membranes have the advantage of low area specific conductivities and more favorable back diffusion of water in comparison with thicker mem-branes. In the former case, this means that membranes with lower conduc-tivity values could be tolerated. Analysis of voltage loss versus membrane thickness and specific conductivity has revealed that, if a membrane voltage loss of 25 mV at a current density 1 A cm–2 can be tolerated, then existing materials with conductivity values similar to Nafion (0.1 S cm–1) could be pre-pared as 20–30 μm thick membranes. However, thinner membranes also typ-ically exhibit lower mechanical strength than their thicker counterparts and can therefore fail earlier. Therefore, future materials might be suitable with just half the specific conductivity if they can be prepared into membranes of half the thickness and still possess sufficient mechanical strength.68

Liu, Ruth, and Rusch used PTFE-reinforced and nonreinforced perfluori-nated sulfonic acid membranes in accelerated fuel cell life tests to study the factors that may control the durability of membranes, and it was observed that the most significant effect influencing membrane durability in the accel-erated fuel cell life tests was mechanical strength.69 However, PEMs are potentially prone to a range of degradation mechanisms, such as hydrolysis, oxidative attack, depolymerization, and radical induced structural changes. Most of these will lead to chain scission and deterioration of the mechanical and chemical integrity of the membrane.70,71

The molecular weight of Nafion 117 has been reported as 250,000 g mol–1.72 The stress–strain behavior of Nafion 117 shows that it is a tough (~4 107 N m–2 stress at maximum strain) material with high elongation (~190%) at


room temperature.73 With its perfluorinated carbon backbone, it might be expected that the large fluorine atoms in Nafion would cause conformational rigidity due to restricted bond rotations. However, studies indicate that the fluoropolymer chains are rather flexible in the sense of torsion.74,75

Bauer, Denneler, and Wilert-Porada also studied the influence of temperature (30–120 C) and humidity (0 ~ 100%) on the mechanical properties of Nafion 117 membrane via dynamic mechanical analysis (DMA).76 The mechanical behavior of Nafion membranes in a humid atmosphere was observed to dif-fer significantly from that in dry atmosphere, and the influence of water on the mechanical properties of the acid form of Nafion was found to be com-plex. The maximum of the storage modulus (E ) as a function of humidity was shifted to higher humidity values with increasing temperature.

Therefore, at low temperatures, water mainly acts as a plasticizer; at high temperatures, it increases the modulus of the membrane by stabilizing the network of hydrophilic clusters with the result that the glass transition of the ionic regions is shifted to higher temperatures with increasing water activ-ity. An intermediate increase of mechanical strength at very low humidity levels (between 1.5 and 3% RH) was also found and was attributed to the formation of hydrates and hydrogen bridge bonds between vicinal sulfonic acid groups.

The mechanical properties of BPSH copolymers (Mn up to 50,000 g mol–1) have also been studied.77 Dry BPSH copolymer exhibits the characteristics of a tough and ductile high-performance thermoplastic, whereas dried Nafion membrane behaves as an elastomer; it has significantly lower strength but much greater elongation. In the wet state, on the other hand, BPSH exhib-its behavior comparable to dry Nafion 117. The change in the mechanical properties of Nafion 117 under hydrated conditions is less substantial than for BPSH, possibly due to retained crystallinity; however, the same general behavior is still displayed. The greater mechanical property change with hydration may be because the BPSH copolymers are amorphous and have a greater degree of phase mixing in hydrophobic and hydrophilic domains. In spite of the lesser degree of elongation for BPSH and greater property change of the BPSH copolymers under hydrated conditions, it should be noted that the toughness for both dried and wet BPSH 40 membranes was comparable to Nafion 117.

For radiation-grafted membranes, it can be expected that the base polymer will have a large influence upon the mechanical properties of PEMs, as will the degrees of grafting and sulfonation. The pre-irradiation approach requires a much higher irradiation dose to achieve the same level of grafting as compared to the simultaneous method. The reason for this is the gradual but inevitable loss of activated centers during irradiation and storage prior to grafting. A high irradiation dose will lead to degradation of base polymer and thus its mechan-ical properties may also potentially suffer.78 Horsfall and Lovell reported that after preparation of radiation-grafted membranes, samples derived from per-fluoroalkoxy polymer (PFA) and fluorinated ethylene-propylene copolymer


(FEP) base substrates failed to form satisfactory membrane electrode assem-blies (MEAs) due to lack of mechanical integrity.79

In contrast, grafted membranes derived from EFTE, PVDF, and LDPE base substrates can be applied to form MEAs.79 Using PTFE, PFA, FEP, and ETFE as base substrates and styrene as monomer (with addition of divinylbenzene [DVB] as cross-linker) to prepare radiation-grafted membranes, Nezu et al. found that only ETFE-based membranes gave satisfactory mechanical prop-erties while the other materials suffered from poor mechanical strength.80 In all cases, the base materials were reported to exhibit appreciable decreases in strength following irradiation, as well as increased brittleness after the graft-ing reaction. For FEP-based radiation-grafted membranes, it has been dem-onstrated that, upon grafting of styrene/10% DVB onto FEP 25 μm film, the tensile strength increased as a function of graft level, whereas the elongation at break decreased—that is, the grafted membranes became more brittle.81 However, for FEP 75 μm film grafted with acrylic acid82 or styrene/acrylic acid83 followed by sulfonation, it was reported that the tensile strength and elongation at break decreased with the degree of grafting and sulfonation.

Results of DMA measurements on FEP-g-acrylic acid membranes and their sulfonated derivatives showed that the grafted copolymer has a higher storage modulus than the pristine polymer and that the storage modulus increases with increasing of grafting degree. On the other hand, the sulfonated grafted copolymer has a lower storage modulus than its unsulfonated analogue, but the sulfonated grafting polymers with >6% degree of grafting exhibited a higher storage modulus than the pristine polymer.84 Again, the temperature at tan d(max) (where dmax is the phase lag to the applied force) increased after grafting, and the temperature at tan d(max) of the sulfonated graft copolymer was lower than that of its analogue but still higher than that of the pristine polymer.84

3.2.4.2 Chemical Stability

It is usually accepted that end groups have no significant influence on mac-roscopic properties of high molecular weight polymers because their weight is negligible as compared with the whole mass of the polymer and because energy values for bonds in end groups and in the constitutive units are prac-tically equal. However, for perfluoropolymers, hydrogen-containing end groups do have a definite influence on their thermal and chemical stability, as can be expected by comparing the bond strength of C-H with C-F (about 410 and 460 kJ/mol, respectively). The end groups identified in thermoplastic fluoropolymers can be generated during the polymerization process (initia-tor, transfer agent, solvent, contaminants, etc.) or by handling of the polymer (aging, heating, extrusion, chemical reactions, etc.). End groups in fluoropo-lymers that have been identified include carboxylic acid (-CF2-COOH), amide (-CF2-CONH2), perfluorovinyl (-CF2-CF CF2), acyl fluoride (-CF2-COF), diflu-oromethyl (-CF2-CF2H), and ethyl (-CF2-CH2-CH3).85


It is generally believed that the extremely reactive peroxyl (HOO·) and hydroxyl (HO·) radicals are responsible for the chemical degradation of the membrane.86,87 Water in the PEM provides a pathway for hydrogen and oxygen crossover from opposite sides of the membrane. Diffusion rates are slow and generally represent only a 1–3% loss in fuel cell efficiency at a rea-sonable current density. However, oxygen crossover does provide a means for the formation of peroxyl and hydroxyl radicals, which can slowly dete-riorate the membrane.88 Pozio et al. have also suggested that H2O2 may form due to incomplete reduction of oxygen on the surface of the catalyst.89 In common with the anode site formation of H2O2, it is thought that subse-quent reaction of this species with metal contaminants in the membrane (e.g., Fe2 , Cu2 ) leads to the formation of radicals that attack the PEM.

Evidence for the formation of H2O2 and hydroxyl radical during the cata-lytic process in fuel cell operation has been examined.90,91 It is thought that, after their formation, hydroxyl or peroxyl radicals then attack the polymer at the end group sites and initiate decomposition. A possible reaction mecha-nism for attack on an end group such as CF2COOH includes: (1) abstrac-tion of hydrogen from an acid end group to give a perfluorocarbon radical, carbon dioxide, and water; (2) reaction of the perfluorocarbon radical with the hydroxyl radical to form an intermediate that rearranges to an acid fluo-ride and one equivalent of hydrogen fluoride; and (3) hydrolysis of the acid fluoride generating a second equivalent of HF and another acid end group (Scheme 3.1).

The evolution of fluoride ions in actual fuel cell effluent and during labora-tory accelerated life studies has been reported.69,89 One common example of radical generation from peroxide decomposition is in the Fenton test,69,89,93 where peroxyl or hydroxyl radicals can be formed through the reaction of hydrogen peroxide with Fe(II) (Scheme 3.2).

Rf – CF2COOH + .OH Rf – CF2. + CO2 + H2O (1)

Rf – CF2 + .OH Rf – CF2OH

Rf – COOH + HF (3)Rf – COF + H2O

Rf – COF + HF (2)

SCHEME 3.1Reaction of a hydroxyl radical with a CF2COOH end group in Nafion. (From Curtin, D. E. et al. 2004. Journal of Power Sources 131:41–48.)

H2O2 + Fe2+ Fe3+ + .OH + OH–

SCHEME 3.2Formation of hydroxyl radicals. (From Walling, C. 1975. Accounts of Chemical Research 8:125–131.)


Fenton reagents based on Ti3 instead of Fe2 have also been used.94 Bosnjakovic and Schlick detected and identified radicals (HOO·, O2

–, TiOO·, and fluorinated alkyl radical) using ESR spectroscopy in Nafion membranes exposed to the Fenton reagent based on Ti3 .95 TiOO· and O2

– are formed from the initial HOO· radicals. TiOO· was detected in swollen Nafion, while O2

– radicals were detected only in dry Nafion. Healy et al. studied the chemi-cal degradation of Nafion membranes in situ (during fuel cell operation) and ex situ (by the Fenton test, 29% H2O2 solution and either 4 or 16 ppm Fe(II)).96 During fuel cell operation, the degradation rate was quantified by monitor-ing the rate of fluoride release. The proton concentration tracks the fluoride concentration closely for the entire run, indicating that the fluoride collected exists primarily as HF.

Analytical studies using NMR and mass spectrometry have clearly identi-fied a product derived from the pendant chain as the principal organic prod-uct of the degradation reaction. The same product was identified in both the Fenton test water and as a residue extracted from MEAs that were heav-ily degraded during fuel cell operation. These features clearly demonstrate similarities between the in situ and ex situ degradation mechanisms and are consistent with the mechanism that involves degradation along a perfluo-rinated ionomer backbone.92 Even though the Fenton test may be a relevant test for evaluation of oxidative stability of fuel cell membranes, it should only be used as a reference test, rather than as a screening tool for fuel cell mem-brane materials because the concentration of H2O2 in a fuel cell is related to gas crossover. Gas permeability of membranes is also a significant factor that is not taken into account in Fenton tests and mechanical durability and chemical durability are so closely related that any attempt to screen material based on either one of them is biased.97

Membranes containing nonfluorinated components are more susceptible to chemical degradation due to the lower stability of the C-H bond compared to the C-F bond. For sulfonated polystyrene-based PEMs, it has been shown that the benzylic a-hydrogen atoms are susceptible to abstraction reactions and subsequent oxidative degradation of polymer chains88,98,99 Assink, Arnold, and Hollandsworth found much higher stability for the sulfonated tert-bu-tylbenzenesulfonic acid with no benzylic a-hydrogen atom than for the sul-fonated isopropylbenzene with one a-hydrogen atom.98 They also reported that styrene-grafted membranes are much more susceptible to weight loss than the membranes grafted with a-methylstyrene when exposed to oxidiz-ing solutions.

In a study aimed at the identification of products of free radical reactions with polystyrene- and aromatic-based PEMs using model compounds, Hübner and Roduner observed the addition of free radicals to the aromatic rings, preferentially in the ortho position to alkyl- and RO-substituents (in polystyrene- and aromatic-based PEMs, the para position is blocked by the presence of the sulfonic acid group).99 This study demonstrated the com-bined ortho-activation by these substituents and the meta-directing effect


of the sulfonic acid group in an electrophilic addition reaction. The possible bond breaking within the C-O-C connection in aromatic-based PEMs was also demonstrated. Figure 3.15 highlights positions of high sensitivity to free radical attack.

Fenton test (3% H2O2 solution containing 2 ppm FeSO4) results for sul-fonated aromatic polymers have also been reported.100–105 The stability of sulfonated aromatic polymers depends on backbone structure, sulfonation degree, and testing temperature. For polyimides, the time before the polymer is totally dissolved ranges from several minutes to ~9 hours at 80 C100–102 or as long as 24 hours at room temperature.103 A time range from less than 1 hour to 24 hours was reported for a poly(arylene ether).104,105 Compared to Nafion 117,96 aromatic-based PEMs appear to have a much lower oxidation stability.

Very recently, Mittal, Kunz, and Fenton106 and Fenton, Mittal, and Kunz107 carried out studies upon the degradation mechanism of Nafion under PEMFC operating conditions. H2O2 was detected in the effluent of the PEMFC when operating with an MEA in which either cathode or anode was present. When both catalyst surfaces were present, no H2O2 was detected in the effluent. It was proposed that this was because it was very unlikely that any H2O2 would be able to pass through the catalyst layer without undergoing decom-position. PEMFC tests were then conducted in which N2 with low concentra-tions of H2O2 was fed to the cell in which the MEA was present in one of three different ways: Nafion 117 only, Nafion plus anode only, or Nafion plus cathode only.

The degradation rate was then measured using the fluoride evolution rate (FER). The first of these modes was carried out in order to test the theory that Fenton-type contaminants were responsible for helping to initiate chemical

CH

SO3H

CH

SO3H SO3H SO3H SO3H SO3H SO3H SO3H

CH CH CH CH CH CH

O C

O

O

HO3S

n

O C

n

CH3

CH3

S

O

O

HO3S

SPEEK

SPSU

SPS

FIGURE 3.15Possible sites for radical attack in various PEMs. (From Hübner, G., and Roduner, E. 1999. Journal of Materials Chemistry 9:409–418.)


degradation (i.e., a significant FER should be observed due to the reaction between the added H2O2 and the contaminant metal species). Increased temperatures (up to 95 C) and/or H2O2 concentrations did not lead to an increased FER. However, the FER in this case was low, as was the case when cathode- or anode-only MEAs were used. Interestingly, significant FER was only observed when flows of H2 and O2 were combined with the presence of one of the catalyst layers.

Furthermore, when the catalyst layer was contained within the membrane rather than at the surface, H2 and O2 were present to observe a significant FER; N2/H2O2 was not sufficient. Additional tests also established that the FER for a cell under load was dependent upon the H2 crossover rate through the PEM, whereas H2O2 formation as an intermediate in the oxygen reduc-tion reaction at the electrodes was independent of the H2 crossover rate. With all the combined evidence, the authors suggested that an alternative degra-dation mechanism other than harmful species formed via decomposition of H2O2 in the presence of impurities is responsible.

Another concern for polystyrene- and some aromatic-based PEMs is hydrolysis of the sulfonic acid group from aromatic rings as well as hydrolytic cleavage of polymer backbone under fuel cell conditions for aromatic polymers including polyimides, poly(arylene ethers), poly(ether ketones), and poly(ether sulfones). It is well known that the sulfonation of aromatic rings is a reversible process especially at low pH and at elevated temperature (Scheme 3.3). The reversibility of sulfonation, for example, is used in the preparation of trinitrotoluene or picric acid. For the simplest membrane of the class of arylsulfonic acids (i.e., benzenesulfonic acid), the reaction occurs upon treatment with a stream of superheated steam at 180 C.108

Vogel et al. tested the stability of selected polymeric arylsulfonic acids and a variety of low molecular weight arylsulfonic acid derivatives with different additional functionalities, sulfonated arenecarboxylic amides, N-arylphthalimides, and arenesulfonamides on heating in water and in dilute aqueous acid to the upper typical temperature limit for DMFC con-ditions (135 C).109 Poly(styrene sulfonic acid) was found to be stable at 100, 135, and 150 C, with desulfonation taking place on heating to 200 C. For selected sulfonated aramides polymer, desulfonation could not be detected. However, the cleavage of amide bonds resulting in chain scission was found

H2O +

SO3H

H2SO4 +

H

Ar Ar

SCHEME 3.3Sulfonation–desulfonation of aromatic rings.


after heating to 135 C in dilute acid as well as in water. With low molecular weight model compounds, the benzamide linkage, as well as the phthalim-ide linkage, was found to be unstable upon heating to 135 C at pH 7, as well as for the respective derivatives sulfonated at the carbonyl subunit or at the amine subunit.

In addition, desulfonation generally has to be taken into account with benzamides and phthalimides that are sulfonated at the amine subunit. The sulfonamide linkage proved to be stable with sulfonated and nonsul-fonated sulfonamides. Compounds with electron-withdrawing substitu-ents at the sulfonic acid-bearing ring were more stable with respect to the hydrolysis of sulfonic acid groups.

The low stability of five-membered imide rings has been reported by Genies et al.110 They also reported that six-membered imides, obtained from the reaction of naphthalene-1,4,5,8-tetracarboxylic dianhydride with amines, were more resistant to hydrolytic degradation than the five-mem-bered imides. For sulfonated aromatic polymers containing an ether bond, a mechanism (Scheme 3.4) for the hydrolytic cleavage of the ether bond has been proposed in which protonation of the ether lone pair occurs in a rapid step and is followed by a nucleophilic attack of the chlorine anion in the alpha position during the rate-determining step.111 Harsh hydrolytic testing for sulfonated polyimides and poly(arylene ethers) has revealed water stability time ranges from several hundred hours to >4,000 hours.103,104,112

O

A–

H+

O

OH A

H

+

SCHEME 3.4Proposed mechanism of the hydrolysis of ether bond in aromatic-based PEMs. (From Iojoiu, C. et al. 2005. Fuel Cells 5:344–354.)


3.3 Brief Overview of Existing PEM Materials

In this section, highlights from the large array of existing PEM materials will be presented, with examples from both academic and industrial research lab-oratories. The section is divided based on the chemical structure of the PEM:

statistical copolymers;block and graft copolymers; andpolymer blends and other composite PEMs.

Each of these categories is further subdivided based on additional criteria found within each subdivision.

3.3.1 Statistical Copolymer PEMs

The large majority of existing PEMs are based on statistical copolymers. This is due to the in-depth knowledge on this type of copolymer and the vast array of different moieties that may be generated. This section will cover polymers composed of random sequences of the constituent monomer units and joined in a linear fashion. Sulfonic acid groups for these polymers are also statistically distributed along the polymer main chain. However, as will be seen, functionalization with acid groups may be achieved via direct polymerization methods as well as the more commonly employed postsul-fonation techniques.

3.3.1.1 Perfluorinated and Partially Fluorinated

Perhaps the most successful systems to date, PEMs based on perfluorcarbon polymers have been used in portable, stationary, and automotive commer-cial applications of PEMFC technology. In addition to providing an attractive combination of performance and reliability, perfluorosulfonic acid (PFSA)-based polymers demonstrated high durability. As stated in Section 3.2.4.3, current thinking is that radical species generated under PEMFC operating conditions are responsible for chemical degradation of the membrane. With their strong C-F bonds, PFSA-based membranes are thus more resistant to radically induced chemical degradation than hydrocarbon and partially flu-orinated membranes.

Perfluorinated PEMs are based on random copolymers derived from tetra-fluoroethylene and a PFSA-based monomer. A range of different structures has been generated and is shown in Figure 3.16.

Nafion membranes (EW ~ 1,100 g/mol) (1a) are the most widely used and studied of all the perfluorinated PEMs. By comparison, little information is publicly available on structure–property relationships for non-Nafion


perfluorinated membranes. Membranes with chemical structures similar to Nafion have been developed, including Flemion, which has a very simi-lar structure (1a) to that of Nafion but with lower EW (~1,000 g/mol) and Aciplex (1b) (EW ~ 1,000–1,200). Both alternative membranes are said to exhibit similar FC performance to that of Nafion. A less complex variation was developed by Dow Chemical (2a)113 and exhibited higher performance than Nafion.114 Calculations by Paddison and Elliott have also suggested that a shorter chain might be more favorable for proton transport under low-hu-midity conditions.115 Due to the higher cost of the Dow monomer in com-parison to its Nafion analogue, work on this particular system was halted.

More recently, however, new and cheaper routes have been developed for the synthesis of the Dow monomer by Solvay Solexsis.116 This will likely lead to a resurgence in investigating the use of this membrane for PEMFC applications; however, little information is currently available. Newer mate-rials derived from the basic DOW structure have also been developed by 3M Company; the CF2 side chain has been extended to three (2b) carbons.117 Asahi Kasei also appears to be working on similar materials (i.e., 2a–b)118 as well as those with longer carbon chains (e.g., 2c).119 In addition, Asahi Glass has pro-duced a material (3) in which a CF2 spacer group is between the ethereal oxy-gen and the remainder of the sulfonic acid-bearing side chain120 as well those with quinone compounds for improved durability.121 Asahi Glass has also

* CF2 CF2 CF2 CF *

O

CF2

CF

CF3

O CF2CF2 SO3H

1a (x = 1)

1b (x = 2)

m n

* CF2 CF2 CF2 CF *m n

* CF2 CF2 CF2 CF *m n

* CF2 CF2 CF2 CF *

O

m n

x

CF2

2a (x = 2)

2b (x = 3)

2c (x = 4)

SO3Hx

CF2

CF2 CF2 SO3H

O

3

O

CF2

CF

CF3

O CF2CF2 SO2NHSO2CF3

4

FIGURE 3.16Perfluorinated PEMs with sulfonic acid (1a Nafion, Flemion; 1b Aciplex; 2a DOW, Hyflon-ion; 2b 3M, 2c Asahi Kasei; 3 Asahi Glass) and bis[(perfluoroalkyl sulfonyl] (4) groups for proton conduction.


made claims regarding the development of a composite membrane based on its PFSA-based polymers that is capable of sustained 120 C PEMFC opera-tion without significant degradation.122 However, it is not clear whether these claims are based on 3, its original Flemion ionomer, or another system.

Another perfluorocarbon PEM (4) that has been recently developed makes use of sulfonamide groups for proton conduction rather than sul-fonic acid groups.123 These materials have been reported to possess greater thermal stability (up to 400 C in the acid form) than Nafion, and excel-lent chemical and electrochemical stability, as well as less susceptibility to dehydration and oxidative degradation than fluorosulfonic acid-based PEMs.124 These systems are presumed to have nanophase-separated struc-tures similar to their sulfonic acid counterparts. Conductivity tests for samples of 4 with EW values of ~1,200 and ~1,400 g/mol showed signifi-cantly higher values than Nafion (EW ~ 1,100 g/mol) at RH < 70%,125 sug-gesting that the sulfonamide group is less sensitive to water content for proton conduction. Furthermore, MEAs based on 4 exhibited significantly higher performance than Nafion MEAs with membranes of comparable EW values.124

Given the difficulty in synthesizing suitable perfluorcarbon-based mono-mers, partially fluorinated systems have been developed in order to reduce cost while still attempting to retain most of the advantages of the perfluori-nated system (e.g., durability and performance). Two examples of this group are shown in Figure 3.17.

General Electric developed a,b,b-trifluorostyrene-based PEMs in the 1960s.126 It was thought that a fluorinated main chain would significantly

CF

5

CF2CF2 *CF*x y

A SO3H

O O*

F

FF

FF

F

E *

n

6

e.g., E = SO2 NH SO2 (CF2)n SO2 NH SO2

FIGURE 3.17Partially fluorinated PEMs based on trifluorostyrene and a trifluorovinyl ether derivative.


increase the durability of these membranes in comparison to the styrenic ana-logue, while retaining the ease of functionality (i.e., sulfonation) of polysty-rene. These materials, however, are quite brittle and therefore it was difficult to make MEAs.127 However, improvements (via the use of selected substituents on the phenyl rings) were made in the mid- to late 1990s by Ballard Advanced Materials. This led to the development of 5, the so-called BAM membrane (also known as BAM 3G).128 Performance of this membrane is said to be at least on par with, if not exceeding, a Nafion membrane of comparable thickness, although no durability data have been reported in the literature.

Studies on these materials have not shown strong evidence for significant nanophase separation, although SAXS analyses on samples with intermedi-ate IEC values (i.e., ~1.9–2.0 meq/g) appear to show greater ionic aggregation than for samples with higher IEC values (i.e., ~2.2–2.5 meq/g).129 SANS analy-ses on 5 have also been carried out by Gebel and Diat127 and suggest a similar structure with increasing water content, as has been previously proposed for Nafion (see Figure 3.8).10

Trifluorostyrene-based monomers and their derivatives are known to exhibit dimerization preferentially over polymerization in contrast to the hydrocarbon analogue styrene.130–132 Ford, DesMarteau, and Smith,130 Smith and Babb,131 and Smith et al.132 have advantageously used this behavior to produce 6 (where E can be a large number of different spacer groups but also typically be sulfonamide-based) via cyclopolymerization of multifunctional monomers bearing at least two trifluorovinyl ether units. The polymers them-selves have perfluorocyclobutane (PFCB) rings as part of the main chain.

In contrast to the more common sulfonic acid-based PEMs, these polymer sys-tems are based on perfluoro sulfonamide chemistry. The acidity of these groups is said to be relatively similar to that of the acid group in Nafion.133 Clear, free-standing films can be made from these polymers with Tg values of 118–141 C, depending upon the composition of the E spacer groups and with conductiv-ity values approaching that of Nafion (IEC 0.91 meq/g). Other PFCB-based systems, including some copolymers with arylene ether sulfone moieties in the backbone, have also been made in which traditional sulfonic acid groups are used as proton-conducting groups.124,134 Although some conductivity and FC data have been reported suggesting at least similar performance to Nafion, little characterization is available in the literature that would enable an understand-ing of the structure–property relationships for these systems.

An alternative approach to the use of partially fluorinated systems to reduce the cost of fluorinated PEMs has been developed by DeSimone et al.; a perfluo-rinated vinyl ether is copolymerized with a hydrocarbon monomer (styrene), sulfonated, and then subsequently fluorinated to replace existing C-H bonds with C-F bonds (Figure 3.18). Thus yields the perfluorinated, cross-linked sul-fonyl fluoride membrane that can then be hydrolyzed to give the PEM (7).135 Because the membranes are cross-linked, considerably higher acid contents (up to 1.82 meq/g) are possible for these materials in comparison to Nafion, leading also to higher proton conductivity values.


Mechanical properties, however, are still reasonable, which may be because l values for even the highest IEC sample do not exceed ~15 H2O molecules per sulfonic acid group. This is likely a result of the large degree of cross-linking in these materials. FC performances were compared for 7 (IEC 1.67 meq/g) and Nafion 117 at 50 C, 75% RH, and it was found that 7 showed significantly higher performance at current densities > 400 mA/cm2, even though 7 was thicker (190 vs. 175 μm, respectively). The achievable power density for 7 was 150% higher than for Nafion under the given testing conditions.135

Perhaps the most interesting property for these materials is their ability to form high surface area PEMs using micromolding/imprint lithography tech-niques with liquid precursor materials. The process consists of casting the liquid precursor materials onto patterned sacrificial templates and then sub-sequently curing the membrane before removing the template.135 Figure 3.19 shows scanning electron microscopy (SEM) pictures of some of these pat-terned membranes. MEAs based on these materials exhibited considerably higher performances than their flat analogues. It was surmised that this is

* CF2 CF CF2 CF *

F4 F4

SO3HO PFPE O

CF2

F4

* CF2 CF CF2CF *

SO3H

F4

7

where PFPE =

CF2CF2 O CF2OCF2CF2 O CF2CF2m n

FIGURE 3.18DeSimone’s perfluorinated PEM.


due to the fact that the higher surface area of the patterned membranes leads to a higher effective interface between the membrane and the catalyst layer, thereby leading to higher performance. It has also recently been reported that pilot-scale manufacture of micromolded 7 has been demonstrated.136

3.3.1.2 Polyarylenes

Of all the hydrocarbon-based PEMs, this group most likely has the larg-est variety of different systems. This is probably due to the wealth of prior knowledge of the nonsulfonated analogues that have been developed over the last several decades as well as the general expectation of higher ther-mal stability, better mechanical properties, and increased oxidative stability over polystyrene-based systems.137 Within the context of this section, poly-arylenes are systems in which an aryl or heteroaryl ring is part of the main chain of the polymer. This section will, therefore, include polymers such as sulfonated poly (ether ether ketone) and sulfonated poly(imides) but will not include systems such as sulfonated polystyrene, which will be covered in Section 3.3.1.3.

The first example in this class of PEMs to be examined was sulfonated phenol-formaldehyde resins (8).138 These materials were studied for use in the space program, given the relative ease of the synthesis and sulfonation of the base polymer as well as low cost. However, in common with polystyrene, these systems exhibit low oxidative stability98 and thus little work has been carried out since the original studies.

Sulfonated poly(arylene ether)s (SPAEKs) have also been developed for application in PEMs, with sulfonated poly(ether ether ketone) (SPEEK) (9a) as the archetypical example of this group.139 The base polymer of SPEEK is com-mercially available and relatively cheap, and sulfonation is a straightforward procedure using concentrated sulfuric acid. At sufficient levels of sulfona-tion, proton conductivity values for SPEEK are comparable to or higher than those of Nafion. However, this does lead to random copolymers where there

5 μm

(a)

5 μm 5 μm

(b) (c)

FIGURE 3.19SEM pictures of patterned membranes of 7 with feature dimensions: (a) 3 3 1.4 μm; (b) 3 3 1.9 μm; (c) 3 3 3.7 μm. (From Zhou, Z. et al. 2006. Journal of the American Chemical Society 128:12963–12972.)


is a statistical distribution of sulfonic acid groups along the main chain. As was previously discussed in Section 3.2.1, these materials exhibit less phase separation than Nafion and have narrower channels for proton conduction, as well as increased numbers of dead ends. This has been used to explain the greater dependence upon water content for proton conduction in these materials in comparison with Nafion.5

Furthermore, long-term testing of these materials under typical FC operating conditions (e.g., 80 C, 100% RH) has shown that they are unstable, generally losing mechanical strength due to very high water uptakes and hence excessive swelling.137 Other SPAEK systems with statistical distributions of sulfonic acid groups have also been developed, including sulfonated derivatives of poly(ether ketone) (9b),139 poly(phenylene oxide)s (10),140–143 poly(aryl ether ketones) (9c),139,144 polybenzoimidazoles (11),145,146 and polyimides (12) (Figure 3.20).147 However, in common with SPEEK, most of these systems exhibit lower degrees of micro-phase separation than Nafion, as well as a loss of mechanical stability at high degrees of sulfonation or FC operating temperatures due to excessive water uptake.137 Also, as previously stated in Section 3.2.4.2, hydrolytic instability and desulfonation are also potential concerns for these systems.

Another class of polyarylene-based PEMs that has been extensively stud-ied, primarily by Hickner et al.,144 is the sulfonated poly(aryl ether sulfone)s

OH

*

R

*

SO3HSO3H

OH

n

8

* O *

Ph

PhHO3S

n

10

N N

O

OO

O

*

SO3H

SO3H

n

N N

O

OO

O

Ar *

m

12

OO* Ar

O

*

nx

9a (x = 1, Ar = phenyl)

9b (x = 2, Ar = phenyl)

9c (x = 1)

N

N N

N

**

H

H SO3H

11

FIGURE 3.20Examples of sulfonated polyarylenes explored for use as PEMs.


(SuPAES). In common with the previously mentioned PEMs, initial SuPAES materials (see Section 3.3.2.1 for later work on block copolymer derivatives of SuPAES) had a statistical distribution of sulfonic acid groups along the polymer backbone. However, instead of using postsulfonation techniques, sulfonic acid groups were introduced via direction copolymerization; that is, suitable sulfonic acid precursor groups were introduced into one of the monomers (13). The advantages of this method are threefold:

The degree of sulfonation is much more controllable using this method in contrast to postsulfonation.

Sulfonic acid groups may be introduced onto moieties with electron-withdrawing groups, thereby making it unlikely that desulfonation will take place during PEM operation. In contrast, the addition of sulfonic acid groups to polyarylenes via postsulfonation will occur almost exclusively on aryl rings bearing electron-donating groups, which promote both sulfonation and the reverse reaction.

As will be described in Section 3.3.2.1, direct sulfonation also makes it con-siderably easier to generate sulfonated block copolymer polyarylenes.

Direct copolymerization techniques have also been employed in the syn-thesis of sulfonated poly(aryl ether ketones),148,149 polyimides,103,110,150–157 and poly(benzoimidazoles).158,159 The synthesis of random disulfonated biphe-nol poly(arylene ether sulfone) copolymers (BPSH x where x represents the percentage of disulfonated diphenylsulfone moieties in the polymer versus unsulfonated diphenylsulfone moities) (14) is shown in Scheme 3.5.

OH Ar OH + m X S

O

O

S

O

O

X X X+ n

NaO3S

SO3Na

O Ar O* S

O

O

O Ar O S

O

O

*

SO3H

HO3S

m n

1. NMP, Toluene, K2CO3

130 °C, 5 h190 °C, 25 hX = Halide

2. Hydrolysis

13

14

SCHEME 3.5Synthesis of BPSH statistical copolymer (14) via direct polymerization.


Increasing the degree of disulfonated moieties in BPSH leads to higher lev-els of conductivity as well as water uptake. However, although water uptake for BPSH 20 to BPSH 50 increases gradually (~20–50 wt%), swelling increases dramatically for BPSH 60, with a water uptake of ~190%.144 Glass transition temperatures increase from BPSH 20 to BPSH 40, wherein a single transition is observed. In the case of BPSH 50 and BPSH 60, an additional transition is seen. The two transitions in these materials have been attributed to the non-ionic matrix and the ionic clusters, thus suggesting that some degree of microphase separation exists for these higher acid content BPSH samples.55

Further evidence for microphase separation has been seen by AFM.55 As expected, BPSH 00, with no ionic regions, displays no significant features in its AFM image. For BPSH 20, isolated ionic clusters have dimensions of 10–25 nm. These clusters are even more readily discerned from the non-ionic matrix in BPSH 40, but the domains appear to remain relatively segregated from each other. In the case of BPSH 50 and 60, connections between domains are clearly visible, especially in the case of the latter sample. It also should be noted, however, that these samples were in a dehydrated state. Therefore, it might be expected that even in the case of the lower acid content samples, it is likely that some channel formation between ionic domains will still occur upon the uptake of water. This can be clearly seen in its linear conductivity behavior as a function of disulfonated monomer (i.e., the percolation thresh-old has been reached by at least 20–30% content of disulfonated monomer).

Very recently, direct copolymerization has been used to synthesize sul-fonated poly(phenylene sulfones), sPSO2 (15) (Figure 3.21). Given the highly electron-withdrawing nature of the sulfone moiety, postsulfonation routes to these materials have never been previously reported. In the case of 15, this was achieved by copolymerization of sulfonated and unsulfonated 4,4 -difluorodiphenylsulfone monomers with 4,4 -thiobisbenzenethiol. The intermediate poly(phenylene sulfone sulfide), sPSS, was then treated with hydrogen peroxide to oxidize the sulfide moieties to sulfones. IEC values ranging from 1.28 to 2.78 meq/g were made, depending upon the ratio of sulfonated to unsulfonated diphenylsulfone monomers.

Interestingly, the high IEC samples of 15 were found to be insoluble in water, whereas the corresponding sPSS (i.e., 15 prior to oxidation) were solu-ble. Furthermore, the water uptake of 15 was found to be relatively constant

* SO2 SO2

HO3S SO3H

SO2 SO2 *

24 nm

15

FIGURE 3.21Sulfonated poly(phenylene sulfone) via direct polymerization.


up to 140 C, whereas the corresponding sPSS samples, in common with SPEEK,5 were found to exhibit detrimental swelling. This reduced swelling was theorized to be due to the presence of aggregated structures, although the possibility of a small quantity of chemical cross-links induced by the oxi-dation step could not be completely excluded. Membranes based on 15 also exhibited considerably higher resistance to oxidative degradation in com-parison to their sulfide analogues.

Proton conductivity behavior for these membranes in comparison to Nafion can be seen in Figure 3.22. The low solubility of 15—due at least in part to the aggregation or perhaps even some degree of crystallization (known to occur for unsulfonated sPSO2)—enables this system to maintain a high acid con-tent without dissolution in water and thus exhibit higher conductivity val-ues than Nafion for a given water content. Similar to BAM membranes (see Figure 3.5), the insolubility of the sPSO2 membranes allows for the highest IEC sample (sPSO2-360) to reach very high water contents (l ~ 80). However, as observed in the case of BAM membranes, dilution of acid groups leads to overall lower conductivity values.160

It is interesting to note, however, that even though the SPEEKK sample shown in Figure 3.22 has a higher IEC content (1.46 meq/g) than sPSO2-781 (1.28 meq/g), it exhibits overall lower conductivity for a given water content. Based on an examination of the proton mobility values for these polymers, it was suggested that this may be due to some microstructural differences between the two different systems.160 Proton conductivity as a function of

FIGURE 3.22Proton conductivity of 15, Nafion 117, and SPEEKK at T 30 C as a function of water content. IEC values: sPSO2-360 2.78 meq/g; sPSO2-430 2.32 meq/g; sPSO2-781 1.28 meq/g; Nafion 117 0.91 meq/g; SPEEKK-685 1.46 meq/g. (From Schuster, M. et al. 2007. Macromolecules 40:598–607.)


temperature in an all-water atmosphere was also examined, with values for 15 increasing as a function of acid content and the conductivity of sPSO2-430 nearly coinciding with that of Nafion. The observed trends were explained as increasing water volume fraction as a function of increasing IEC (i.e., increasing water content leads to improved percolation within the hydro-philic domains).160

All of the previous examples of polyarylenes have been composed of main chains in which moieties (e.g., sulfone, ketone, ether) other than aryl rings have been present. Having a purely aromatic system where the polymer backbone and side groups are exclusively composed of aryl rings might potentially lead to a system with improved thermochemical stability. Recently, a purely aromatic, rigid-rod PEM was reported by Sandia National Laboratories.161 SDAPPx (where x number of sulfonic acid groups per repeat unit) (16) was synthesized via Diels–Alder polymerization to give polymers with high molecular weights (Mw > 100,000, PDI ~ 2) (Figure 3.23). Subsequent postsul-fonation yielded polymers where up to six sulfonic acid groups per repeat unit could be present with presumably para substitution of the phenyl side groups being dominant due to the directing influence of the neighboring phenyl group as well as steric considerations.161

Samples with IEC values of 0.98–2.2 meq/g (x 1–5) were prepared and found to be insoluble in water. In excess of IEC 2.2 meq/g (x 6), however, the polymer was found to form a hydrogel, thereby eliminating the possibil-ity of forming a suitable film. Conductivity of water-saturated SDAPP was

*

*

(SO3H)x/6 (SO3H)x/6

(SO3H)x/6 (SO3H)x/6 (SO3H)x/6 (SO3H)x/6

n

16

(where x = target sulfonation level)

FIGURE 3.23Sulfonated Diels–Alder poly(phenylene) (SDAPP).


found to range from 0.0013 S/cm (x 1) to 0.123 S/cm (x 5). In compari-son to BPSH and sulfonated polyimide membranes with similar IEC values, SDAPP membranes were found to show slightly lower conductivity values. This was theorized to be due possibly to the nature of the polymer backbone, the concentration density of ionic groups within the polymer, and how water is bound within its microstructure.

In the sulfonated poly(arylene) systems described so far, the sulfonic acid groups have been statistically distributed along the polymer main chain. Poly(arylenes) in which the sulfonic acid sites are separated from the main chain by means of a spacer group have also been developed. Examples of sys-tems in which this has been attempted include poly(p-phenylenes) (17),162–164 poly(p-phenylene)-poly(aryl ether ketone) copolymers (18),164 and polyimides (19, 20).165,166 These are shown in Figure 3.24.

Little characterization of 17 and 18 is available in the open literature. In the case of 19 and 20, both polyimides displayed microphase separation in con-trast to statistically, main-chain sulfonated polyimides for which little, if any, microphase separation is observed. For 19, hydrophilic domains were seen to be on the order of 5 nm in size.165 Similarly, 20 (where x 3, y 10, n 50) with a

* *

OO

O O

m n

SO3H

17

*

O

O

m n

SO3H

18

CO*

N N

O

OO

O

*

O(CH2)xSO3H

O(CH2)xSO3H

n

19 (x = 3, m = 0)

20 (x = 3, 4; y = 6, 10; m = 1 – n)

N N

O

OO

O

(CH2)y *

m

FIGURE 3.24Poly(arylenes) with sulfonic acid-bearing side groups.


low IEC (1.84 meq/g) also exhibited spherical ionic clusters 5–8 nm in size. For the higher IEC (2.31 meq/g) sample of 20 (x 3, y 10, n 70), larger ionic clus-ters (~12 nm) were observed in addition to the smaller ionic clusters. In contrast, for a chemically branched analogue of 20 (x 3, y 10, n 10), even smaller ionic clusters (<5 nm) were observed as the minority, with what appeared to be mainly a relatively uniform distribution of hydrophilic domains. This occurred even though both samples had the same IEC value, and it was theorized that branching prevented the ionic groups from aggregating sufficiently enough to result in large degrees of nanophase separation.166 Interestingly, conductivity studies for 20 as a function of RH over T 80–120 C revealed that branched 20 exhibited higher conductivity than its linear analogue.

One of the main disadvantages of main-chain, statistically sulfonated polyimides is their instability toward hydrolysis. In the case of 19 and 20, by placing the sulfonic acid functionality on a side chain, the water-sen-sitive polyimide backbone phase separates into the hydrophobic domain, thereby protecting it from hydrolysis. For both side-chain sulfonated poly-imides, large improvements were observed over their main-chain sul-fonated counterparts. Open-circuit voltage testing was also carried out on 20 (x 3, y 10, n 50) and it was found to operate without any sign of voltage degradation for over 5,000 hours at both RH 60% (T 68 C) and RH 90% (T 77 C).

3.3.1.3 Miscellaneous Monolithic, Statistically Sulfonated Copolymer PEMs

One of the earliest proton exchange membranes was based on sulfonated polystyrene where divinylbenzene was used as a cross-linking unit for extra stability. Developed by General Electric,167 this membrane (21) was cheap and easy to manufacture,168 and it was used for fuel cells in the Gemini space pro-gram.169 However, due to the sensitivity of the benzylic hydrogen to radical attack, lifetimes for these membranes under FC operating conditions were quite low.170 Thus, little work has been carried out on these systems since their inception.

Inorganic polymers have also been investigated for use as PEMs. Polphosphazenes (22) are composed of phosphorus and nitrogen repeat units. In addition to its relatively high thermal, chemical, and oxidative sta-bility, the polymer backbone in polyphosphazenes is also easy to function-alize with a wide variety of different groups (from the precursor polymer poly(dichlorophosphazene), -[PCl2 N-]-).171 Given these characteristics, it is not surprising that they have been explored for potential use as PEMs. Postsulfonation techniques were found to lead not only to sulfonic acid groups on the phenyl rings but also to the formation of acid–base complexes through the basic nitrogen sites along the polymer backbone, as well as potential cleavage of the backbone itself.172–176 Later, it was found to be more effective to use small molecules containing either protected sulfonic acid or


sulfonamide groups, which were then reacted with -[PCl2 N-]- to form the desired polymer.177,178

Cross-linked, sulfonic-acid-substituted, polyphosphazene-based PEMs have primarily been examined for potential use in DMFC applications due to their low MeOH crossover with reported values 2.5 times lower than that of Nafion.179 These materials have also been shown to display good thermomechanical and chemical stability (in a Fenton test).180 Sulfonamide-substituted polyphosphazenes have exhibited very high power densities that are comparable with Nafion and may be suitable for use in PEMFC applications.181

Another inorganic polymer system, sulfonated polysiloxanes, has also been investigated for potential use as PEMs due to the potential high thermal stabil-ity of the Si-O backbone.182 An early example of this type of PEM is 23a. With the sensitivity of the siloxane backbone to very acidic conditions, it was neces-sary to build the acidic functionality into the monomeric units prior to polym-erization, using mercaptan groups that could be subsequently oxidized to sulfonic acid. Use of trialkoxy silane monomers also permitted the formation of cross-links, thereby increasing the stability of the membrane. IEC values were relatively low, ranging from 0.27–0.37 meq/g. No other data to suggest suitability for FC applications (e.g., conductivity) were provided, however.183

A more recent example from this group is 23b. Conductivity measurements under fully humidified conditions for 23b were on average about an order of magnitude lower than for Nafion. However, under dry conditions, the values were only slightly lower for 23b in comparison to Nafion. Interestingly, the MeOH diffusion coefficients for 23 were 20–40 times lower than for Nafion. This might make these membranes potentially suitable for use in DMFC applications (see Figure 3.25).184

3.3.2 Block and Graft Copolymer PEMs

Block and graft copolymers are composed of significant sequences of differ-ent monomer units, normally in a nonstatistical fashion. For block copoly-mers, these sequences are assembled in a linear fashion whereas in the case of graft copolymers, blocks are grown from or attached to the backbone of another block as branches. In the case of block and graft copolymer PEMs, these sequences may be composed of significantly different chemical units or units that are chemically identical except that one block is sulfonated and the other is not.

As typically observed in the case of non-ionic block and graft copolymers, the immiscibility of the constituent blocks within the copolymers can induce microphase separation beyond even that which normally occurs due to hydrophobic and hydrophilic sites within statistical copolymer PEMs such as Nafion. A relatively recent area of PEM research, ionic block and graft copolymers are interesting from the point of view of providing fundamental understanding about the influence of morphology upon proton conduction


and other PEM properties, as well as for the possibility of developing new membranes suitable for use in PEMFC and DMFC applications.

3.3.2.1 Block Copolymers

An early example of a block copolymer PEM is sulfonated polystyrene-b-poly(ethylene-r-butylene)-b-polystyrene (S-SEBS) (Figure 3.26) (24). Developed by Dais-Analytic Corp., they have been investigated for use as low-cost PEMs in hydrogen PEMFCs operating at ambient temperatures and low current den-sities.185 It has been found that proton conductivity values for S-SEBS for sam-ples with IEC 0.94–1.71 meq/g appear to be relatively independent of IEC. However, water content (l) showed a dramatic change over 1.13–1.71 meq/g, increasing by 100 H2O molecules per ~SO3H.129 It may be that, in common with BAM 3G, S-SEBS PEMs undergo a significant dilution of acid concentration over this IEC range, thereby removing any positive influence of an increase in proton mobility upon increases in water content and overall resulting in relatively little change in conductivity as a function of IEC.

**

**

SO3HHO3S

m n

n

21

P N P*

O

O

N

O

O

*

R

RR

R

A

A

22

(R = H, alkyl)

(A = SO3H, SO2NHSO2CF3)

* Si O

O

ESO3H

*

Si O

OMe

OMe

*

23a (E = CH2)

23b (E = Ph)

m n

FIGURE 3.25Miscellaneous examples of monolithic, statistically sulfonated copolymer PEMs.


Analysis of S-SEBS by SAXS has revealed the presence of cylindrical mor-phologies for a degree of sulfonation of <34%. Interestingly, different mor-phologies can also be observed when membranes are cast from different solvents.186,187 Membranes (27 mol% degree of sulfonation) cast from THF form lamellar morphologies, as seen in Figure 3.27, while those cast from MeOH/THF (tetrahydrofuran) (20/80 v/v) exhibit a diffusive phase boundary with disorderly interconnections between domains. This is due to the differences

(a) (b)

200 nm200 nm

FIGURE 3.27TEM images for S-SEBS (27 mol%) cast from (a) THF; (b) MeOH/THF (20/80 v/v). (From Kim, B. et al. 2005.Journal of Membrane Science 250:175–182.)

CHCH2 CH2 CH2*

SO3H SO3H SO3H

CH2 CH

CH2

*

m n

CH3

x y

24

CH2 CH

CH2

CHCH2 *

mn

CH3

CHCH2*

m

SO3H

CHCH2

m

25

SO

O

O

ArOHO O Ar O Ar'Ar' F

O

HO3S

nx y

26

(Ar = 4,4'-biphenyl; Ar' = 4,4'-benzophenone)

FIGURE 3.26Examples of hydrocarbon block copolymer PEMs.


in affinities of the blocks for the different solvents. Furthermore, it has also been found that S-SEBS with IEC 0.5–1.0 meq/g exhibits a lamellar morphol-ogy, although morphology can also be affected by the casting solvent.188

More recently, Elabd and Napadensky189 and Elabd et al.190 examined a similar triblock copolymer (25) for potential use in DMFC applications. Sulfonated poly(styrene-b-isobutylene-b-styrene), S-SIBS, was also found to adopt different morphologies that were dependent upon the casting solvent. Thus, although THF-cast membranes exhibit a lamellar morphology, lower degrees of order are seen in SAXS analyses when cast from other solvents (e.g., CHCl3 and benzene). Also, changes in proton conductivity (in plane) can also be seen as a result of the different morphologies (e.g., conductiv-ity in CHCl3-cast S-SIBS is 0.023 S/cm, whereas THF-cast S-SIBS shows sig-nificantly higher conductivity at 0.52 S/cm). Also, similar to S-SEBS, S-SIBS exhibits a change in morphology with increasing IEC values, with lamellar microdomains at low ion contents and co-continuous microdomains at high ion content.190

Poly(arylene ether sulfone)-co-poly(arylene ether ketone)-based copoly-mers containing blocks of sulfonated poly(p-phenylene) (PPP) (26) have also been synthesized by Ghassemi, Ndip, and McGrath.191 Telechelic segments of PPP were generated and capped in situ with 4-chloro-4 -fluorobenzophenone. These oligomers were sulfonated and then used in subsequent polymerization reactions to generate block copolymer 26. Proton conductivity values ranged from 0.024 S/cm (IEC 0.70 meq/g) to 0.036 S/cm (IEC 1.20 meq/g). These values are low in comparison to both Nafion and BPSH PEMs.191

More recently, Ghassemi, McGrath, and Zawodzinski investigated the influence of heavily fluorinated blocks upon BPSH-like systems. Fluoro-containing blocks are of special interest due to the greater thermo-oxidative stability and hydrophobicity that they impart upon a polymer.192 Similar to 26, these PEMs have been generated via coupling of telechelic oligomers. The surface morphology by AFM under partially hydrated conditions for 27 was compared with Nafion and BPSH. Whereas Nafion and BPSH appeared to suggest the presence of mainly isolated hydrophilic domains with some local connections, 27 showed a well-defined phase separation with a distinctly different morphology than its random BPSH analogue. This was also used to explain the higher conductivity of 27 versus BPSH at similar IEC values. Furthermore, 27 displayed higher conductivity as a function of decreasing RH in comparison to Nafion N112, possibly due to the morphological differ-ences between the two systems (Figure 3.28).192

Yang, Shi, and Holdcroft,193 Shi and Holdcroft,194,195 Norsten et al.,196 and Rubatat et al.197 have also investigated the effects of introducing fluoro-blocks into a copolymer. Sulfonated polysulfone-b-PVDF (28) was gener-ated via a Williamson ether synthesis of hydroxy-terminated polysulfone and bromide-terminated VDF oligomer.193,198 These PEMs with varying acid content were analyzed by TEM (Figure 3.29).198 Aggregate size was seen to


decrease as a function of increasing acid content for 28 (Figure 3.29b, d). At IEC 0.78 meq/g (Figure 3.29d), there were very large regions of aggrega-tion (50–200 nm), which may have been due to large-scale phase separation of the ionic and non-ionic domains. It should be noted, of course, that these samples are dehydrated and thus may or may not accurately reflect the situation inside a hydrated membrane.

Poly(vinylidene difluoride-co-hexafluoropropylene) (P(VDF-co-HFP)), with terminal CCl3 groups, was generated via the radical polymerization of VDF and HFP in the presence of CHCl3 as chain-transfer agent. These macroini-tiators were then used in the atom-transfer radical polymerization (ATRP) of styrene, which, after subsequent sulfonation, generated the partially fluo-rinated block copolymer 29.194,195 Morphological analysis of 29 with varying acid contents was carried out using TEM (Figure 3.30).195 At low degrees of sulfonation (Figure 3.30a), no distinct domain formation was observed. As ion content increases (20–40 mol%), distinct domain formation can be seen (Figure 3.30b–d) with the presence of ordered, connected ionic channels (8–15 nm in width) with 20–40 nm spacings between domains. At higher sulfonation levels (~50% or greater), the disruption in the ordered domains is such that, at 100% sulfonation (Figure 3.30f), the interfacial region between the hydrophilic and hydrophobic domains is less distinct and the ionic aggregates tend to be disordered.

SO

O

O

O Ar O Ar' O

CF3

CF3

O **x

n

SO3H

HO3S27

(Ar = 4,4'-biphenyl; Ar' = 4,4'-perfluorobiphenyl)

SO

O

O

O CH2CF2 CH2CF2 *

CH3

CH3

*

SO3H

nx

28

CH2CF2 CH2CF CH2x

CF3

Hw

CH CH2 CH Cl

m ny z

SO3H29

FIGURE 3.28Examples of partially fluorinated block copolymer PEMs.


3.3.2.2 Graft Copolymers

Graft copolymer-based PEMs have generally been made via irradiation grafting reactions of a suitable monomer (e.g., styrene) in and onto a dense film as base substrate, with subsequent sulfonation of the reactive units. The base substrate is typically a sheet of a commercially available polymer film. Irradiation can be in situ (i.e., irradiation takes place when both monomer and base substrate are present) or the base substrate can be irradiated prior to introduction of the monomer. In both cases, the monomer reacts with radi-cals generated in the base substrate, leading to the growth of polymer chains affixed to the irradiated film. The advantages of this technique are the poten-tial use of commercially available films, thereby reducing cost as well as pro-viding a material in which the base substrate can be chosen on the basis of its mechanical properties, and the grafted chains on their proton conduction ability. Some examples of radiation-grafted PEMs are shown in Figure 3.31.

Radiation-grafted PEMs based on perfluorinated substrates (32) have been developed by Rouilly et al., who used FEP as the base polymer

(a)

100 nm 100 nm

100 nm 100 nm

(b)

(c) (d)

FIGURE 3.29TEM micrographs of (a) sulfonated polysulfone (IEC 1.55 meq/g); (b) 28 (IEC 1.62 meq/g); (c) sulfonated polysulfone (IEC 0.83 meq/g); (d) 28 (IEC 0.78 meq/g). (From Yang, Y. S. et al. 2004. Macromolecules 37:1678–1681.)


film.199 In situ FC tests have shown that these membranes are capable of performances at least comparable to Nafion under fully humidified, 60 C operation conditions.200 Lifetimes were typically short; however, with the introduction of DVB as a cross-linker during the graft copolymerization with styrene, lifetimes could be improved to around 1,000 hours.86 This could be further improved to over 6,000 hours using a combination of DVB and triallyl cyanuarate as cross-linkers.201

ETFE and PVDF have also been used as base substrates by Shen et al.51 and Horsfall and Lovell.79,202 PEMs (30, 31) made using these materials have been shown to demonstrate comparable FC performance in both PEMFC and DMFC to that of 32. It was concluded from these studies that the primary controlling

A B

C

100 nm 100 nm

100 nm 100 nm

100 nm 100 nm

D

E F

FIGURE 3.30TEM micrographs of 29: (a) degree of sulfonation (DS) 12%, IEC 0.23 meq/g; (b) DS 22%, IEC 0.62 meq/g; (c) DS 32%, IEC 0.89 meq/g; (d) DS 40%, IEC 1.08 meq/g; (e) DS 49%, IEC 1.31 meq/g; (f) DS 100%. (From Shi, Z. Q. and Holdcroft, S. 2005. Macromolecules 38:4193–4201.)


factors for FC performance are the grafting conditions and overall sulfonic acid content, rather than the actual composition of the base substrate.79

The majority of radiation-grafted PEMs are based on PSSA as the proton-conducting component. Less investigated examples include PEMs based on grafted a-methyl styrene98 and glycidyl methacrylate.203 Ballard Power Systems Inc. has also reported that substituted trifluorostyrene monomers may be grafted onto base substrates and subsequently sulfonated to provide radiation-grafted PEMs.204,205 The use of a partially fluorinated polymer as the proton-conducting segment presumably was done to improve resistance to radically induced degradation in these systems (see Section 3.2.5). In com-mon with BAM 3G, however, the exact chemical composition of this system, including the identity of the base substrate or substrates, is not clear.

Few studies have been carried out upon the microstructures of radiation-grafted PEMs. In the case of PVDF-g-PSSA, WAXS (wide-angle x-ray scat-tering) and SAXS measurements suggested the presence of ionic aggregates embedded in a non-ionic matrix.206 Additional analyses using high-angle annular dark-field scanning transmission electron microscopy (HAADF STEM) showed the presence of three phases: (1) dark regions corresponding

SO3H

z

CH *CH2CF2CF2 CF2CF2 CF2CF2*x y

30

SO3H

z

CH *CF2CF2CH2y

*x

31

SO3H

z

*CF2CFCF2CFCF2CFCF2CF*x y

32

CF3 CF3

m mn n

FIGURE 3.31Examples of radiation-grafted PEMs: poly(ethylene-alt-tetrafluoroethylene)-graft-poly (sty-rene sulfonic acid) (ETFE-g-PSSA), 30; PVDF-g-PSSA, 31; and poly(tetrafluoroethylene-co-hexafluoropropylene)-g-PSSA (FEP-g-PSSA), 32.


to PVDF, (2) gray regions corresponding to aggregated sulfonated polysty-rene, and (3) bright dots due to cluster-like sulfonated aggregates dispersed in the aggregated polystyrene regions (Figure 3.32). Nafion, on the other hand, exhibited comparatively ordered and uniform arrays of ionic aggre-gates. Interestingly, when both Nafion and PVDF-g-PSSA samples were treated with a 50% MeOH solution, the ionic regions for Nafion doubled in size, whereas the ionic regions in PVDF-g-PSSA only changed slightly.207

Recently, modern radical polymerization techniques have been used to generate graft copolymer-based PEMs in which the lengths of both the main chains and the graft chains are controlled. Ding, Chuy, and Holdcroft have reported the synthesis of “living” chains of poly(styrene sulfonate) (PSSNa) by TEMPO-initiated polymerization (Scheme 3.6).208,209 The addition of DVB at the conclusion of the polymerization led to the formation of PSSNa mac-roinitiators (macPSSNa), which could then be copolymerized with styrene to form PS-g-macPSSNa graft copolymers. TEM analyses of these graft copoly-mers showed ionic domains (5–10 nm) that are visibly connected, resulting overall in a continuous network. With increasing ion content, the ionic net-work develops to a greater extent. By contrast, TEM analyses of random copo-lymers of PS and PSSNa showed little evidence for microphase separation.

The length of the ionic graft also was found to have an effect upon mor-phology.210 Membranes containing longer ionic side chains (degrees of polymerization 17–102) were seen to microphase separate to a greater extent than those with shorter side chains for a set ionic content. At low ionic con-tent, only isolated domains are observed. With increasing ionic content, a continuous ionic network is observed.

A partially fluorinated polyarylene with grafted PS side chains (33) (Figure 3.33) has also been made via a macroinitiator route.196 TEM analysis of this system showed that increasing the number of side chains resulted in different morphologies (Figure 3.34). For 33a, a wormlike morphology was observed, whereas in the highest graft chain content example (33c), cylindrical domains were observed. In the case of 33b, there were structures suggestive

(a)

100 nm 100 nm

(b)

FIGURE 3.32HAADF-STEM images (Ag stained) of (a) Nafion; (b) PVDF-g-PSSA. (From Huang, H. S. et al. 2006. Applied Surface Science 253:2685–2689.)


of an intermediate morphology composed of both wormlike and cylindri-cal domains. SAXS data were also relatively consistent with this assessment. Proton conductivity data also showed an increase in values with increasing side chain content. Glass transition temperatures (Tg) also increased, ranging from 185 C for 33a to 192 C for 33c.

3.3.3 Polymer Blends and Composite PEMs

All of the examples of PEMs discussed within Section 3.3 until now have been composed of only one polymer system without any other compounds present—be they small molecules, polymers, or solid-state materials. A wide variety of different polymer blend and composite PEMs has been made. However, in this section, only a brief overview highlighting some of the more interesting examples that have been reported in the literature will be presented. For discussion, these types of PEMs have been divided into three categories: polymer blends, ionomer-filled porous substrates and reinforced PEMs, and composite PEMs for high-temperature operation and alternative proton conductors.

SO3Na SO3Na

TEMPO

SO3Na

TEMPO

TEMPO DVB

Living Polymer

SO3Na

m

n

n

n n

k

SCHEME 3.6Synthesis of PS-g-macPSSNa graft copolymers.


O*

CF3

CF3

O

O

CH3

O*

F4

F4

2

22

6

1–y

y

33a-c

HO3S

HO3S

FIGURE 3.33A highly fluorinated, comb-shaped copolymer PEM (for 33a, y 0.19; for 33b, y 0.25; for 33c, y 0.38).

50 nm

(a)

50 nm 50 nm

(b) (c)

FIGURE 3.34TEM images (Pb2 -stained) of 33 with increasing graft chain density: (a) 33a; (b) 33b; (c) 33c. (From Norsten, T. B. et al. 2006. Advanced Functional Materials 16:1814–1822.)


3.3.3.1 Polymer Blends

The blending of two or more polymers is frequently used to try to com-bine the separate desirable properties of each system rather than trying to develop one system with all the properties. In the case of PEMs, this has led to the blending of proton-conducting polymers with non-ionic polymers, low IEC polymers, or polymer-containing basic moieties, particularly for DMFC applications in order to decrease MeOH crossover. These different types of blends will be briefly discussed next.

Attempts to blend Nafion with another polymer have mostly focused on the use of PVDF (as well as its copolymer, PVDF-co-HFP). This is probably due to its chemical and thermal stability as well as to its semicrystalline nature, thus introducing increased mechanical strength to the PEM.211–213 However, these blends have often shown significantly lower conductivity (two orders of magnitude) in comparison to unmodified Nafion samples,214,215 presum-ably due to the hydrophobicity of PVDF resulting in lower water uptake.216 This occurs even when a comparatively low weight fraction of PVDF is pres-ent (20 wt%), although MeOH crossover is considerably reduced.214

Treatment of PVDF by dehydrofluorination and doping with sulfuric acid prior to blending have been shown to improve the hydrophilicity of a Nafion/PVDF blend. Such blends were demonstrated to show comparable conductivity and FC performance to unmodified Nafion and significantly improved over blends in which the PVDF had not been treated. MeOH cross-over rates, however, were not reported.217 PEMs composed of “sandwiches” of Nafion plus Nafion/PVDF blends have also been used as PEMs in order to reduce MeOH crossover and improve DMFC performance.218,219 Other non-ionic polymers that have been blended with Nafion include PPO220 and polypyrrole.221

Ionic polymers other than Nafion have also been included in ionic/non-ionic PEM blends. Poly(ether sulfone) (PES) has been used to strengthen SPEEK as well as sulfonated poly(ether sulfone) (SPES) with contents rang-ing from 20 to 60 wt%.222 The conductivity of the SPEEK component was rela-tively the same as unmodified SPEEK up to about 40 wt%. A similar effect was seen for PES/SPES blends, although the drop in MeOH permeability was more dramatic for PES/SPES from unmodified SPES than for PES/SPEEK from unmodified SPEEK. PVDF has also been used as a blending material to reinforce SPEEK.223 The strength of the PEM was increased over unmodified SPEEK. Although conductivity levels decreased as a function of increasing PVDF content, the selectivity (ratio of proton conductivity to MeOH perme-ability) of the blended PEMs was increased over that of unmodified SPEEK and Nafion.

The grafting of poly(ethylene oxide) chains onto PES and a poly(aryl ether) (PAE) has been used to increase the hydrophilicity of these non-ionic poly-mers.224 These systems were then blended with SPES. SEM analysis showed


the presence of spherical hydrophilic domains (50–100 nm). The hydrophilic-ity of the non-ionic polymer actually improved the water uptake of the blend as a function of increasing non-ionic content for the blend; thus, conductivity was also seen to increase.

Non-ionic polymers have also been blended with ionic block copolymers. Poly(vinyl phosphanate)-b-polystyrene225 and PS-b-SPS226 have been blended with PPO. In both cases, improvements were seen in MeOH permeabil-ity over that of the unmodified block copolymers and conductivity values dropped as a function of increasing PPO content. PVDF has been blended with SEBS in order to improve its mechanical and chemical stability, but aggregation was found to be a problem due to incompatibility between com-ponents. However, it was found that a small amount (2 wt%) of a methyl methacrylate-butyl acrylate-methyl methacrylate block copolymer as com-patibilizer not only led to greater homogeneity but also improved mechani-cal resistance, water management, and conductivity.227,228

In order to improve component compatibility, blends of sulfonated poly-mers have also been investigated. Blends of SPS and sulfonated PPO (SPPO), with IEC values of ~2.5 and ~2.6 meq/g, respectively, have been made.229 Interestingly, the blend showed higher conductivity (albeit higher MeOH permeability) than either of the separate components. DMFC performance was comparable to Nafion. PEMs made from sulfonated poly(aryl ether sul-fone) (SPAES) blended with Nafion exhibited a phase separated morphol-ogy (as determined by SEM) wherein the morphology was controlled by both the blend ratio and the degree of sulfonation of the SPAES component. Conductivity and MeOH permeability were lower than for Nafion.230

Studies have also been carried out on blends of sulfonated poly(ether ketone ketone) (SPEKK) wherein one component has a high IEC value (2.0 meq/g) and the other has a lower IEC value (0.8–1.5 meq/g).231,232 Lower IEC SPEKK as a blend component offers the advantage of improved mechanical strength ver-sus high IEC SPEKK, as would the introduction of PEKK. However, in contrast to using PEKK, lower IEC SPEKK offers the dual advantages of better compat-ibility as well as still providing some degree of proton conductivity.

For SPEKK2.0/SPEKK1.5 blends (where the number denotes the IEC value of the blend component), a relatively homogeneous PEM is found. However, as the IEC value of the second component is lowered, a dispersed particle morphology is observed. Morphology could also be altered by changing the casting conditions (e.g., temperature or solvent). For example, whereas dispersed particles were found for SPEKK2.0/SPEKK1.2 cast from NMP, samples cast from DMAc exhibited a co-continuous morphology. FC performance and mechanical stability of the blends were improved over SPEKK itself.232

Blends of sulfonated polymers and polymers containing basic moieties have also been made. Represented schematically in Figure 3.35, ionic cross-linking between acidic and basic sites generally leads to improved mechani-cal and thermal stabilities.233 Strong interactions within these blends results


from hydrogen bonding as well as from electrostatic forces resulting from proton donation from acidic to basic sites.137

One of the first examples of this type of blend was composed of SPEEK or SPES as the acidic component and diaminated PES, poly(4-vinylpyridine) (P4VP), poly(benzimidazole) (PBI), or poly(ethyleneimine) (PEI) as the basic component.234–236 For blend IEC values of 1.0 meq/g, conductivity values were reported to be good, as was H2/O2 FC performance. Thermal stabili-ties for these blends was also demonstrated to be high (>270 C).235 Other examples of acid-base PEMs include blends of SPPO and PBI,237 sulfonated poly(phthalazinone ether ketone) and aminated SPES,238 SPIs and aminated PIs,239 and SPEEK with PES bearing benzimidazole side groups,240–242 as well as an unusual example in which the blend is composed of sulfonated, hyper-branched polyether and pyridine-functionalized polysulfone.243

Properties of these ionically cross-linked blends have been extensively studied by Kerres et al.244 It was found that the effective degree of ionic cross-linking was dependent upon the size and density of functional groups. In comparing the blend of SPES (IEC 1.6 meq/g) with either PBI (base capacity 6.5 meq/g) or PEI (base capacity 23.2 meq/g), it was found that the effective IEC value of the blend was as expected for SPES/PBI but much higher than expected for SPES/PEI (i.e., the degree of cross-linking was much lower than expected in the case of the latter). It was theorized that this was due to steric restrictions whereby not every amino group in PEI is able to react with a cor-responding sulfonic acid site in SPES.235 Furthermore, acid-base linkages in SPES/PBI were also found to be stronger due to the stronger basicity of the nitrogen centers in PBI versus PEI.

The degree of ionic cross-link density was also found to have an effect upon the degree of swelling in the PEMs (higher cross-link density less swelling) and thus also upon water uptake (inverse relationship between cross-link density and water uptake level). Also, as expected, the greater the number of sulfonate versus sulfonic acid sites in these blends, the lower the observed level of conductivity was.244 It has been noted, however, that these ionic bonds are generally unstable at temperatures higher than 70–90 C and thus they may lose the advantages of the cross-linking (e.g., water insolubility

SO3H

SO3H SO3H

SO3H SO3H

SO3–

NR2H+

SO3–

SO3–

NR2H+ NR2H+

FIGURE 3.35Schematic representation of ionically cross-linked acid–base blend membranes. (From Kerres, J. A. 2005. Fuel Cells 5:230–247.)


and improved mechanical properties) during FC operation at or above this temperature range.137 In order to overcome this disadvantage, more recent work from this group has focused on the use of both covalent and ionic cross-links.137,233,245,246

The effect of blend miscibility on transport properties of SPEKK/PEI has been examined by Gasa, Weiss, and Shaw.247 Miscibility in the SPEKK/PEI system was found to be highly dependent upon the IEC of SPEKK. Thus, mis-cible blends were found for SPEKK at low ionic content (IEC 0.8 meq/g), whereas the blends were only partially miscible at IEC 1.1 meq/g and immis-cible at IEC 1.4 meq/g. In contrast, a blend composed of SPEKK with the non-ionic PES was miscible across a wide range of IEC values (0.8–2.2 meq/g). Conductivity of the high IEC SPEKK/PEI sample (IEC 2.2 meq/g) was found to vary as a function of SPEKK content, as shown in Figure 3.36.

The step transition between 0.4 and 0.5 wt fraction of SPEKK has been attributed to a change in morphology as well as the percolation threshold within this region. At 0.4 wt fraction of SPEKK, the blend consists of SPEKK droplets surrounded by a relatively insulating matrix of PEI. At 0.5 wt frac-tion, interconnectivity of some droplets was observed. As the weight frac-tion of SPEKK increases, the morphology changes from an interconnected, bicontinuous structure to droplets of PEI dispersed in a continuous SPEKK phase. MeOH permeability of these blends was found to be significantly lower than that for Nafion. Interestingly, it was also found that the perme-ability of SPEKK/PEI blends was, in general, lower than that of SPEKK/PES blends. Again, this was attributed to the miscibility and hence morphologi-cal differences between the two different blends.247

0.20

0.15

0.10 ( )/(1 )

Co

nd

uct

ivit

y, S

/cm

0.05

0.000.0 0.2 0.4

SPEKK wt. Fraction

0.6 0.8 1.0

FIGURE 3.36Proton conductivity (23 C, 98% RH) of SPEKK/PEI blends as a function of SPEKK weight frac-tion (SPEKK IEC 2.2 meq/g). Solid line is derived from an effective-medium theory (equation given). (From Gasa, J. V. et al. 2006. Journal of Polymer Science Part B 44:2253–2266.)


3.3.3.2 Ionomer-Filled Porous Substrates and Reinforced PEMs

Thinner membranes generally offer the advantages of lower resistance, lower cost, and improved water management in PEMFC applications. With perhaps the exception of radiation-grafted PEMs, it is difficult to synthesize thinner monolithic membranes due to their reduced mechanical strength. By using reinforcing materials, the fabrication of thinner membranes with acceptable mechanical strength can be potentially achieved. Indeed, PFSA-based PEMs as thin as 5 μm with good conductivity and mechani-cal properties have been made.69,248–250 Generally, Nafion has been the iono-mer of choice used, although PSSA251 and SPEEK252 have also been used. Examples of reinforcing materials include porous PTFE sheets,253 porous polypropylene,248,254 polycarbonate,255 expanded PTFE,256 polysulfone, and glass microfiber fleece.257 The use of some of these supports in PFSA-based PEMs will be described later.

Use of polycarbonate as a support yielded 42.3 μm membranes with a conductivity value of 0.06 S/cm,38 whereas 25 μm PTFE and 28 μm Celgard (microporous polypropylene) gave conductivity values of 0.10 and 0.021 S/cm, respectively.248 However, it has been found that the properties of rein-forced PEMs are dependent not only upon the nature of the support and its thickness but also upon the manner in which they are prepared. For example, Ramya et al. made 35 μm PTFE-based reinforced PEMs and found that conductivity of the PEMs was highly dependent upon the solubility parameter of the solvent.258 Specifically, solvents for which the solubility parameter was more compatible with the acid-containing side chains of Nafion led to reinforced PEMs with improved conductivity due to reduced association of the Nafion with the PTFE substrate, presumably allowing the acid groups to have greater mobility. Lin et al. found that surfactant also had an effect upon proton conductivity of Nafion-impregnated porous PTFE PEMs. It was demonstrated that using 1–3 wt% surfactant in the Nafion solutions facilitated penetration of the ionomer into the pores of the substrate with greater overall conductivity levels versus the situation in which surfactant was not used.259

Water uptake and swelling of reinforced PEMs have also been studied. It has been found that water uptake and water flux of Nafion/porous PTFE PEMs increases with increased ionomer loading and that the rate of water uptake as a function of temperature is actually higher for the reinforced membrane in comparison to Nafion 112.256 However, it was also seen that the actual water uptake of a similar reinforced membrane was about half as much as a Nafion 1135 sample with a commensurate lower degree of swelling.258

FC performances for reinforced PEMs have also been studied.252,259–261 A 20 μm Nafion/porous PTFE membrane was used in a DMFC at 70 C with 2 M MeOH as fuel.260 It was able to achieve 0.2 V at a current density of 390 mA cm–2, which is superior to Nafion 117 and 112, for which a potential of 0.2 V was achieved at


current densities of 310 and 275 mA cm–2, respectively. This performance was attributed to lower proton resistance and/or lower MeOH crossover for the rein-forced PEM.

Porosity and thickness of the porous PTFE substrate have also been found to affect the properties of the reinforced Nafion PEM.262 Under H2/O2 FC operation at 80 C and a gas pressure of 0.2 MPa, the current densities at 0.6 V were found to increase with increasing pore size: 500–850 mA cm–2 for 0.3–0.5 μm pore sizes. Under the same conditions, Nafion 115, Nafion/PTFE–45 μm, and Nafion/PTFE–25 μm were found to be 800, 800, and 950 mA cm–2, respectively.

3.3.3.3 Composite PEMs for High-Temperature Operation and Alternative Proton Conductors

The development of materials suitable for PEMFCs operating above 100 C has attracted a great deal of attention within the last few years. This has been primarily due to the high-temperature PEMFC initiative set out by the U.S. Department of Energy.144 High-temperature operation (>120 C) is potentially desirable for PEMFCs due to the theoretical advantages offered by this approach: (1) decrease in the effect of CO poisoning of the anode cata-lyst, (2) improved fuel oxidation kinetics; (3) increase in cell efficiency; and (4) improved gas transport (no liquid water present to cause interference). However, in addition to all the challenges that high-temperature operation presents to other PEMFC components (e.g., seals, gas diffusion layers, cata-lysts), existing commercial membranes such as Nafion are unsuitable for use at high temperature without expensive and load-parasitical engineering. This is due to the propensity of these membranes to undergo dehydration (with a consummate loss in proton conductivity), loss of mechanical strength, and higher levels of gas permeation.263

A number of different approaches have been used to try to overcome some of these disadvantages of existing membranes. One such approach is to try to prevent water loss from the proton transport pathways, thus maintain-ing proton conductivity above the boiling point of water. Typically, this is attempted by adding hydrophilic inorganic species into the membrane. Furthermore, these particles in themselves may also be capable of proton conduction.

Introduction of inorganic particles may be accomplished by mixing the particles with an ionomer solution prior to casting264 or impregnation of the particles inside a preformed PEM.265 The use of preexisting particles prior to casting frequently leads to an undesirable agglomeration and inhomoge-neous dispersion of particles inside the PEM.266 It is also possible to grow inorganic nanoparticles within the PEM itself and thereby utilize the hydro-philic regions as a template.

A number of different particles for composite PEMs have been used, including silica (SiO2), titania (TiO2), zirconia (ZrO2), clays (primarily


montmorillonite), alumina (Al2O3), zeolites, zirconium phosphate, and het-eropoly acids (HPAs).267 Nafion has been the PEM primarily investigated for use in composite PEMs but other systems (e.g., SPEEK) have also been used. Within the context of this chapter, only Nafion/silica PEM composite PEMs will be covered. For more details on other composite PEMs, Herring pro-vides an excellent review.267

An early Nafion/silica composite was generated in situ using sol-gel chemistry with tetraethylorthosilicate (TEOS) as the silica precursor.265 Characterization of these composite PEMs by SAXS has suggested that silica grows within the confines of the hydrophilic regions of Nafion.268 FTIR and 29Si solid-state NMR spectroscopy have been used to study the formation of the inorganic network inside the PEM.269,270 These studies show that the silica formed within the confines of the PEM is not as highly cross-linked compared with silica formed outside the PEM via the sol-gel process. An examination of mechanical properties has shown that, with increasing silica content, there is an initial strengthening followed by a ductile-brittle transition, suggesting the presence of a SiO2 phase percola-tion threshold.265

A number of conductivity and FC measurements have been carried out upon Nafion/silica composite PEMs.265,271–276 Interestingly, different results are obtained for the two measurements. These composite PEMs have been found to have generally higher water contents than the unmodified PEMs; however, proton conductivity values have been found to be generally lower, with decreasing conductivity as a function of increasing silica content.273,275 A small increase, however, appears to be seen for silica contents < 10 wt%, although the increase is not particularly significant.273

FC measurements, on the other hand, do appear to show a beneficial effect upon both PEMFC and DMFC performance at high temperature. Using Nafion/silica composite PEMs with 6 wt% silica, Adjemian et al. demonstrated current densities of 400 and 75 mA cm–2 at 0.6 V (at 130 C and 3 atm pressure) for the composite and unmodified PEMs, respectively. Furthermore, the performance of the unmodified PEM decreased dramati-cally over 50 hours at 0.65 V, whereas the performance of the composite PEM remained unchanged. These observations were attributed to improved water management for the composite PEMs at higher temperatures. At T > 100 C, the pore structure in unmodified Nafion is thought to be disrupted—a pro-cess of temperature-related structural changes prevented by the presence of silica particles.276 Similar improvements in performance were also observed in DMFC and attributed to both enhanced water uptake and reduced MeOH crossover.271

A number of explanations have been proposed for this seeming discrepancy of the effect of silica upon Nafion in conductivity versus FC measurements:

replacement of unassociated bulk water with hydrophilic silica nanoparticles;


capillary condensation effects as a result of the smaller dimensions of the free spaces in silica-filled pores277; and

improved mechanical properties of the composite PEMs over the unmodified membranes.278

More work must be carried out to establish the major effects leading to this phenomenon.

The proton conduction ability of sulfonic acid groups is highly sensitive to water content. With the demand for PEMFCs capable of operating at temper-atures in excess of 100 C, interest is increasing in the development of PEMs in which proton conduction is enabled by functional groups that can work at low water contents or in the absence of water. This is made possible when the group is capable of functioning not only as the source of excess protons but also as a solvent capable of stabilizing the excess protons and permit-ting efficient proton transport through the membrane. These groups must be amphoteric and thus function as both proton donor and acceptor.279,280 Two functional groups in particular have garnered the most interest: free or bound (i.e., phosphonic) phosphoric acid and free or bound nitrogen-con-taining heterocycles (e.g., imidazole).

One of the first high-temperature PEMs was based on phosphoric acid-doped to poly(2,5-benzimidazole) (PBI).281 This is generally accomplished by soaking a poly(2,5-benzimidazole) film in a solution of phosphoric acid. These PEMs are capable of conducting protons (at high [>90%] acid concen-trations) at temperatures well in excess of 100 C under low humidity condi-tions, as generally are phosphonic acid groups bound to small molecules (e.g., 1,3-ethane-diphosphonic acid) (as can be seen in Figure 3.11).281–283 However, the presence of water in a fuel cell during operation leads to the leaching of the acid and consummate severe drop in proton conductivity.284 Attaching of the phosphonic acid group to a polymer backbone has been used as an attempt to prevent leaching.

Examples of polymers that have been phosphonated include grafted polystyrene,285 polyaryleneethers,286,287 poly(4-phenoxybenzoyl-1,4-phenylenes),288 polyethersulfones,285,288–291 polyphosphazenes,292–294 and polysiloxanes.295 However, the comparative IEC values (based on phospho-nic rather than sulfonic acid) for these phosphonated polymers are relatively low and water is still required for proton conduction. It thus appears that tethering of phosphonic groups via spacer groups to a polymer backbone is not an effective method for generating materials capable of high proton conductivity at high temperature in the absence of water. Very recently, Steininger et al. studied a series of different phosphonic acid-functionalized, small-molecule compounds and suggested that aggregation of a very high concentration of phosphonic acid groups is necessary for high proton conduc-tivity. However, based on their studies, it appears that tethering via spacer groups does not allow for the necessary degree of aggregation. The authors


therefore suggested that alternative immobilization techniques would thus be required in order for these materials to approach the conductivity val-ues of phosphoric acid or PBI doped with high concentrations of phosphoric acid.284

Recently, a sol-gel process has been used to fabricate phosphoric acid/PBI PEMs.296,297 The method consists of using polyphosphoric acid as polymer-ization and casting solvent for PBI. Residual polyphosphoric acid is subse-quently hydrolyzed into phosphoric acid after film formation. Significant improvements in mechanical properties were seen for these PEMs in com-parison to those produced via the usual imbibing process. Conductivity, too, improved, with a value of ~0.25 S/cm at approximately 150 C versus ~0.12 S/cm or less for imbibed membranes. Long-term durability FC tests have also been carried out using these membranes and appear to show little degradation up to 8,000 hours at 120 C without any humidification. Shutdown–start-up tests (262 cycles over 2,000 hours) also showed low per-formance degradation rates, suggesting that the formation of liquid water does not have a decidedly adverse effect upon FC performance and thus that perhaps little phosphoric acid is leached from the membrane dur-ing these cycles.297 However, ex situ leaching tests to determine the ease with which the acid may be leached from these membranes have not been reported to date.

Nitrogen-containing heterocycles—primarily imidazole (34) (Figure 3.37)—have been investigated for their proton-conducting ability. In order to pre-vent leaching, attempts have also been made to bind phosphonic acid sites to a polymer such as polystyrene298 and polysiloxane.299 Conductivity was also demonstrated for these PEMs in the complete absence of water. Interestingly and in contrast to the free versus bound phosphonic acid, there is only about one order of magnitude of difference in conductivity between free and bound imidazole. Furthermore, it has also been found that proton mobility values increase with increasing flexible spacer length (e.g., PEO or alkane) in con-trast to phosphonic acid. This has suggested that proton conduction in these materials is exclusively due to structural diffusion rather than a vehicular mechanism.284 Unfortunately, studies have suggested that the electrochemi-cal stability of the imidazole ring is insufficient for the conditions within an operating fuel cell.300–302

More recently, triazole (35) has been investigated as a potential anhy-drous proton conductor (Figure 3.37).303–305 Studies using cyclic voltam-metry suggest that triazole is electrochemically stable enough for use in fuel cells.304 Furthermore, in a comparison between poly(vinyl triaz-ole) and poly(vinyl imidazole), the former displayed conductivity values almost six orders of magnitude greater at 120 C than the latter.305 Studies based on materials wherein flexible spacer units are between the triazole unit and the polymer backbone have not yet been reported but are pre-sumably underway.


3.4 Future Directions

As can be seen from this chapter, a wide array of different membrane systems has been developed for improved performance in a variety of different FC applications. It is likely that, given the vast amount of knowledge of synthetic polymer chemistry that has been accumulated over almost the last 100 years, many more systems will be generated and tested as potential PEMs. However, it would appear that four areas will be of continued or increasing interest:

development of hydrocarbon-based systems due to a need for lower cost PEMs as well as environmental concerns about fluorocarbon-based materials;synthesis of more organized PEMs through using systems such as block and graft copolymers in order to provide improved proton conduction pathways;development of PEMs capable of high-temperature PEMFC operation (T > 100 C) in order to circumvent some of the existing difficulties (e.g., CO poisoning of catalyst) of FC operation at less than 100 C; andinvestigations into PEM structure–property relationships.

Although all of these areas are important, it is perhaps the last one that may eventually hold the greatest attraction. To date, most PEM development has been carried out without a great degree of understanding about the rela-tionship between the chemical and morphological structures of a PEM and its properties. As knowledge of structure–property relationships for PEMs increases, no doubt more methodical attempts will be made in the area of synthetic PEM chemistry and the need for an iterative approach to the devel-opment of new membranes will be reduced, thus leading to the develop-ment of materials that meet the performance, durability, reliability, and cost requirements necessary for PEMFC and DMFC applications to achieve sig-nificant levels of commercialization.

N

N

N

N

H

35

N

H

34

FIGURE 3.37Molecular structures of imidazole (34) and 1H-1,2,4-triazole (35).


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272. Adjemian, K. T., Srinivasan, S., Benzieger, J. and Bocarsly, A. B. 2002. Investigation of PEMFC operation above 100 C employing perfluorosul-fonic acid silicon oxide composite membranes. Journal of Power Sources 109:356–364.

273. Kim, Y. J., Choi, W. C., Woo, S. I. and Hong, W. H. 2004. Proton conductivity and methanol permeation in Nafion/ORMOSIL prepared with various organic silanes. Journal of Membrane Science 238:213–222.

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191

4Diffusion Layers

Mauricio Blanco and David P. Wilkinson

CONTENTS

4.1 Introduction ................................................................................................ 1924.2 Different Types of Diffusion Layers ........................................................ 196

4.2.1 Carbon Fiber Paper Fabrication ................................................... 1964.2.1.1 Carbon Fibers ................................................................... 1974.2.1.2 Paper Manufacturing ..................................................... 2044.2.1.3 Carbon Fiber Paper with Aerogels ............................... 206

4.2.2 Carbon Cloth Fabrication .............................................................. 2074.2.3 Metal Diffusion Layers ................................................................. 209

4.2.3.1 Metal Meshes ................................................................... 2114.2.3.2 Sintered Metals ................................................................ 2134.2.3.3 Micromachined Meshes ................................................. 2144.2.3.4 Metal Foams ..................................................................... 215

4.2.4 Engineered Diffusion Layers ....................................................... 2154.2.5 Other Diffusion Layers ................................................................. 221

4.2.5.1 Silicon-Based Materials .................................................. 2214.2.5.2 Other Materials ...............................................................223

4.2.6 Performance Comparison between Diffusion Layers .............. 2244.2.6.1 Performance Comparison in PEM Fuel Cells ............. 2244.2.6.2 Performance Comparison in Direct Methanol

Fuel Cells .......................................................................... 2264.3 Treatments and Coatings ..........................................................................227

4.3.1 Hydrophobic Treatments ..............................................................2274.3.1.1 Fabrication Processes and Procedures .........................2274.3.1.2 Effect of Hydrophobic Treatment .................................2294.3.1.3 Hydrophobic Treatment in Direct Liquid Fuel

Cells ............................................................................ 2324.3.2 Hydrophilic Treatments ................................................................233

4.3.2.1 Hydrophilic Treatments for PEM Fuel Cells ...............2334.3.2.2 Hydrophilic Treatments for Direct Methanol

Fuel Cells ..........................................................................234


4.1 Introduction

Figure 4.1 shows a schematic of a typical polymer electrolyte membrane fuel cell (PEMFC). A typical membrane electrode assembly (MEA) consists of a proton exchange membrane that is in contact with a cathode catalyst layer (CL) on one side and an anode CL on the other side; they are sandwiched together between two diffusion layers (DLs). These layers are usually treated (coated) with a hydrophobic agent such as polytetrafluoroethylene (PTFE) in order to improve the water removal within the DL and the fuel cell. It is also common to have a catalyst-backing layer or microporous layer (MPL) between the CL and DL. Usually, bipolar plates with flow field (FF) channels are located on each side of the MEA in order to transport reactants to the

4.3.3 Microporous Layers .......................................................................2344.3.3.1 Fabrication Processes ......................................................2364.3.3.2 Effect of Microporous Layer on Fuel Cell

Performance ..................................................................... 2374.3.3.3 Parameters Affecting Microporous Layers ................. 2394.3.3.4 Multilayered Microporous Layers ................................ 2444.3.3.5 Microporous Layers in Direct Liquid Fuel Cells ........ 246

4.4 Properties and Measurements for Diffusion Layers ............................ 2484.4.1 Thickness ........................................................................................ 2494.4.2 Hydrophobicity and Hydrophilicity ........................................... 251

4.4.2.1 Sessile Drop Method ...................................................... 2514.4.2.2 Wilhelmy Method: Capillary Rise ................................ 2524.4.2.3 Contact Angle of Moving Droplet ................................2544.4.2.4 Internal Contact Angle ...................................................254

4.4.3 Transport .........................................................................................2554.4.3.1 Characterization of General Transport Properties .....2554.4.3.2 Characterization of Gas Transport Properties ............ 2604.4.3.3 Characterization of Liquid Transport Properties ....... 267

4.4.4 Electrical and Thermal Conductivity ......................................... 2724.4.4.1 Electrical Conductivity ................................................... 2734.4.4.2 Thermal Conductivity .................................................... 274

4.4.5 Mechanical Properties .................................................................. 2764.4.6 Corrosion Stability and Degradation .......................................... 278

4.4.6.1 Corrosion Studies ............................................................ 2794.4.6.2 Cold and Freezing Temperatures .................................280

4.4.7 Flow Field Interaction.................................................................... 2824.4.7.1 Pressure Drop Tests ........................................................ 2824.4.7.2 Visualization Techniques ...............................................284

4.5 Future Direction of Diffusion Layers ...................................................... 286References .............................................................................................................288

Diffusion Layers 193

active CL and remove reaction products through the DL. In most fuel cells, coolant plates or flow channels are formed adjacent to each bipolar plate in order to maintain the overall temperature of the cell.

The MPL is normally formed with carbon black and hydrophobic particles (PTFE). The diffusion layer is usually made out of carbon fiber paper (CFP) or carbon cloth (CC) and is a vital component of the MEA and fuel cell because it provides the following functions and properties:

It helps to distribute the reactant gases or liquids evenly from the FF channels of the bipolar plates to the CL so that most of the active zones (and catalyst particles) are used effectively. Thus, the DL has to be porous enough for all the gases or liquids (e.g., liquid fuel cells) to flow without major problems.It helps to remove the water produced and accumulated in the CLs toward the FF channels. The DL must have large enough pores so that the condensed water can leave the CL, MPL, and DL without blocking any pores that may affect the transport of reactant gases or liquids.It provides mechanical support to the CL and the membrane in order for these two components to be unaffected by the pressure that the landings or ribs of the bipolar plate put on them. Therefore, the DL

AnodeCatalyst

Layer

Flow Field Landing

e–

e–

Air

Anode Substrate 250 μm

30 μm5 μm

5 μm250 μm

Cathode Substrate

Membrane

Air Channel

Cathode Flow Field Plate

Fuel Gas Channel

CathodeCatalyst

Layer

FIGURE 4.1Schematic of a polymer electrolyte membrane fuel cell (not to scale).


has to be made out of a material that does not deform substantially after long hours of operation so that it is still able to provide mechan-ical support.It helps to conduct electron flow from the bipolar plates to the CL and vice versa with low resistance between them. In order for the DL to be able to do this successfully, it has to be made of a material that is a good electronic conductor.It helps to transfer the heat produced from the CL to the bipolar plates in order to keep the cell at the desired temperature of operation. Thus, the DL should be made out of a material that has a high level of ther-mal conductivity so that removal of heat is as efficient as possible.

Another important parameter that has to be taken into account when choos-ing the appropriate diffusion layer is the overall cost of the material. In the last few years, a number of cost analysis studies have been performed in order to determine fuel cell system costs now and in the future, depending on the power output, size of the system, and number of units. Carlson et al. [1] reported that in 2005 the manufacturing costs of diffusion layers (for both anode and cathode sides) corresponded to 5% of the total cost for an 80 kW direct hydrogen fuel cell stack (assuming 500,000 units) used in the automotive sector. The total value for the DLs was US$18.40 m–2, which included two carbon cloths (E-TEK GDL LT 1200-W) with 27 wt% PTFE, an MPL with PTFE, and Cabot carbon black. Capital, manufacturing, tooling, and labor costs were included in the total.

In 2007, the same consulting company published another report in which the cost of the DLs had increased slightly to 6% of the overall cost of the stack, compared to the 5% previously estimated [2]. One issue with these analyses and predictions was that they were based on carbon cloth as the diffusion layer, but this material does not reflect what most of the fuel cell companies use (i.e., carbon fiber paper) [3].

In another report, James and Kalinoski [4] performed an estimation of the costs for a direct hydrogen fuel cell system used in automotive applications. The assumed system consisted of an 80 kW system with four fuel cell stacks, each with 93 active cells; this represents around 400 MEAs (i.e., 800 DLs) per system. The study was performed assuming that the DL material used for both the anode and cathode sides would be carbon fiber paper with an MPL. In fact, the cost estimate was based on SGL Carbon prices for its DLs with an approximate CFP value of around US$12 m–2 for 500,000 systems per year. Based on this report, the overall value of the DLs (with MPL) is around US$42.98 per kilowatt (for current technology and 1,000 systems per year) and $3.27 per kilowatt (for 2015 technology and 500,000 systems per year). Figure 4.2 shows the cost component distribution for this 80 kW fuel cell sys-tem. In conclusion, the diffusion layer materials used for fuel cells not only have to comply with all the technical requirements that different fuel cell systems require, but also have to be cost effective.


The task of choosing the correct DL design and material can be quite com-plex because it is important to understand how each parameter and design of the DL affects the overall performance of the cell. Many of the param-eters are interconnected with each other, making the overall design of the DL even more complicated: If one of the properties is changed, then another one will be affected. The following sections will discuss in detail the different materials considered as potential candidates for DLs, the various treatments and coatings used, the various methods used to measure the properties and characteristics of the DLs, and the future direction of DLs. A tremendous

4%

11%

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48%

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8%

11%

37%

14%

13%

4%

4%

3%

Flow plates (stamping)

GDLs

Catalyst Ink

Membrane & catalyzation

MEA frame/gaskets

Coolant gaskets

Endplates & currentcollectorsStack assembly

Other

2%

2%1%1%1%2%

32%

Membr.

36%

Catalyst

Ink

23%

GDL

3%

3%2%2%

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28%

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Sy

ste

ms/

yr

2010 2015

FIGURE 4.2Cost component distribution for a direct hydrogen PEM fuel cell system. (From B. D. James and J. A. Kalinoski. Annual progress report, DOE Hydrogen Program, 700–704. Washington, D.C.: U.S. Department of Energy, 2007.)


amount of work has been published in the area of DLs, particularly in the last few years. This chapter attempts to capture the most relevant aspects of this work.

In this chapter, diffusion layers will be considered as the porous media that help the transport of the reactant fluids and products from one surface to another. In addition, the MPL will be defined as the additional layer or lay-ers (made out of carbon black and water-repellent particles) located between the CL and the DL. It is important to note that although “microporous layer” and “diffusion layer” are the common names for these components, as well as the ones used in this chapter, a number of different names can be found in the literature.

For example, if the DL is used on the side of the cell where the fuel or oxi-dant is in gas phase, then this part can be referred to as gas diffusion layer (GDL). When both the CL and the DL are mentioned as one component, then the name “diffusion electrode” is commonly used. Because the DL is of a porous nature, it has also been called “diffusion medium” (DM) or “porous transport layer” (PTL). Sometimes the DL is also referred to as the compo-nent formed by an MPL and a backing layer. The MPL has also been called the “water management layer” (WML) because one of its main purposes is to improve the water removal inside the fuel cell. In this chapter, we will refer to these components as MPL and DL because these names are widely used in the fuel cell industry.

4.2 Different Types of Diffusion Layers

As discussed previously, a number of different materials have been consid-ered as potential candidates to be used as diffusion layers in PEMFCs and direct liquid fuel cells (DLFCs). The two materials used the most so far in fuel cell research and products are carbon fiber papers and carbon cloths, also known as carbon woven fabrics. Both materials are made from carbon fibers. Although these materials have been quite popular for fuel cells, they have a number of drawbacks—particularly with respect to their design and model complexity—that have led to the study of other possible materials. The following sections discuss in detail the main materials that have been used as diffusion layers, providing an insight into how these materials are fabricated and how they affect fuel cell performance.

4.2.1 Carbon Fiber Paper Fabrication

Carbon fibers have been used as filaments for lamps for nearly a century, since Edison first used them. In the early 1960s, Shindo developed the first modern carbon fiber when he pyrolyzed polyacrylonitrile (PAN) fibers [5].


Carbon fiber introduction into commercial markets happened in the mid-1960s and since then their application has increased substantially. Some of these applications include airplanes, spacecraft parts, compressed gas tanks, automotive parts, bridges, reinforced concrete, structural reinforcement, rec-reational sports equipment, and electrochemical systems [6,7].

Table 4.1 shows some of the parameters and properties of the most com-mon carbon fiber paper materials being produced commercially for use in PEM and DLFCs. Figure 4.3 shows an SEM (scanning electron microscope) picture of a carbon fiber paper without any coating. In the following sub-section we will briefly discuss the fabrication process of carbon fibers and carbon fiber papers.

4.2.1.1 Carbon Fibers

Most carbon fibers use PAN as their precursor; however, other polymer pre-cursors, such as rayon [8], pitch (a by-product of petroleum or coal-coking industries), phenolic resins, and polyacetylenes [6,7], are available. Each com-pany usually uses different precursor compositions for its products and thus it is difficult to know the exact composition used in most commercially avail-able carbon fiber products.

Although there are many variations on how carbon fibers are made, the typical process starts with the formation of PAN fibers from a conventional suspension or solution polymerization process between a mixture of acrylo-nitrile plastic powder with another plastic, such as methyl acrylate or methyl methacrylate, and a catalyst. The product is then spun into fibers, with the use of different methods, in order to be able to achieve the internal atomic struc-ture of the fiber. After this, the fibers are washed and stretched to the desired fiber diameter. This step is sometimes called “spinning” and is also vital in order to align the molecules inside the fiber and thus provide a good basis for the formation of firmly bonded carbon crystals after carbonization [7].

Once the PAN fibers are made, they are heated in an air oven to a tempera-ture between 200 and 300 C for around 30–120 minutes (see Figure 4.4 for a summary of the typical stages used in the manufacturing of carbon fibers from PAN fibers). The exact temperature for this step changes depending on the manufacturer, with some stating that 230 C is the ideal temperature [9] and 220 C preferred by others [6,8]. This stage is sometimes referred to as the “stabilizing” stage because it is used to change the linear atomic bonding of these fibers to a more thermally stable ladder bonding [7] (i.e., thermoset material).

Basically, the fibers grab oxygen molecules from the air so that their atomic binding pattern is reorganized. In the industry, this stage is performed with the use of a number of different techniques. In some cases, the oven is divided into a number of chambers, each with different temperatures; the speed at which the fibers enter and exit the chambers can also be controlled depending on the desired characteristics of the fibers [10]. In other methods,

198 Proton Exchange Membrane Fuel CellsTA

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SE 070816 WD15.7 mm 20.0 kV ×60 500 um

(a)

SE 070817 WD15.7 mm 20.0 kV ×500 100 um

(b)

FIGURE 4.3Scanning electron microscope picture of typical carbon fiber paper sheets used in fuel cells: (a) Toray TGPH-060 CFP with no PTFE (reference bar indicates 500 μm); (b) close-up view of the TGPH-060 CFP with no PTFE (reference bar indicates 100 μm); (c) Toray TGPH-060 CFP with 20% PTFE (reference bar indicates 500 μm); (d) close-up view of the TGPH-060 CFP with 20% PTFE (reference bar indicates 100 μm).


SE 070818 WD16.0 mm 20.0 kV ×60 500 um

(c)

SE 070819 WD16.0 mm 20.0 kV ×500 100 um

(d)

FIGURE 4.3(Continued)


the fibers are heated with air mixed with different gases that can accelerate the stabilization stage.

In the next stage, called “carbonization” or “graphitization” (temperatures > 2,500 C), the stabilized fibers are heated to a temperature between 1,000 and 3,000 C in an oven with an inert gas. Depending on the temperature used, the overall characteristics of the final carbon fibers will change. For example, if the temperature is around 1,000 C, then type A carbon fibers, which have low modulus and low tensile strength, are produced. If the temperature is 1,400 C, then type II carbon fibers with high tensile strength and medium modulus are produced. Type I carbon fibers can be created if the tempera-ture used is higher than 2,500 C (graphitization), which helps to increase the modulus of these fibers but also decreases their tensile strength [5,6].

4.2.1.2 Paper Manufacturing

For fabricating carbon fiber paper, the carbonization process changes slightly (see Figure 4.5 for a summary of the typical stages used in the manufac-turing of CFP). First, the fibers are heated to temperatures between 1,000 and 1,350 C [8,9,11–14] and then the fibers are chopped to predetermined lengths depending on the desired pore size of the paper (i.e., smaller lengths will result in smaller pore sizes) [11]. The diameter of these fibers is usually between 5 and 15 μm [6,9,15,16]. For large-scale production of carbon fiber paper, the chopped fibers are coated with water and a binder (e.g., polyvinyl

PANFibers

Stabilization200–300 °C

in Air

Carbonization> 1000 °C

StabilizedFiber

Type ACarbon Fibers

Carbonization> 1400 °C

Graphitization> 2500 °C

Type ICarbon Fibers

Type IICarbon Fibers

FIGURE 4.4Schematic of the typical steps in carbon fiber manufacturing process from PAN fibers. Dotted borders indicate the steps that depend on the desired final product. (Modified from T. H. Ko. Journal of Applied Polymer Science 43 (1991) 589–600. With permission from John Wiley & Sons, Inc.)


alcohol) [9,16], and the wet material is rolled inside a chamber in order to dry it. Once this process is completed, the mixture is impregnated with a resin.

Phenolic resin is the most common resin because it is highly sticky with a hard carbon structure and low cost [9,12,13]. The paper is then put inside an oven at around 150–175 C in order to cure the resin and to remove the sol-vent [9,12]. The amount of phenolic resin that is impregnated also depends on the desired final characteristics of the CFP. For example, high concentra-tions of the phenolic resin reduce the thickness of the paper, but the surface resistivity decreases significantly [13].

Other types of resins and approaches have also been used. Ji et al. [17] used carbonizable acrylic pulp fibers instead of conventional phenolic resin as a binder material. They claim that by using these acrylic fibers, which are mixed with the carbon fibers during the paper-making step, the total manu-facturing cost for the CFP is reduced because the conventional impregnation step (for the phenolic resin) and subsequent solvent removal are eliminated. In order to reduce the costs of producing CFP, the same research group developed a method in which the step of impregnating the resin to the paper is eliminated [18]. In the process, in addition to adding water and a binder material to the carbon fibers as explained earlier, resin powder is added to the mixture. Then the wet material (which already has the resin) is formed into a mat on a support ready to be dried [18].

PANFibers

Stabilization200–300 °C

in Air

Carbonization1000–1350 °C

ResinImpregnation

Curing150–175 °C

Post-curing150–200 °C

CompressionMolding of Paper

into Sheets

Carbonization andGraphitization

1300–2500 °C in N2

DryingRolling of Fibers

into a Mat

Paper Making

Water andBinder Coating

Choppingof Fibers

FIGURE 4.5Schematic of the typical and general steps in carbon fiber paper fabrication with wet-laid fibers.


After the impregnation of the resin the paper is molded into sheets by the compression molding technique. If a thicker material and better pore distri-bution is desired, then multiple plies of this material are bonded together. Mathur et al. [12] claim that binding more than one ply together improves fuel cell performance compared to single-ply CFPs used as DLs. This is due to the fact that, with multiple plies, the pore size is reduced but the total number of pores is increased; this improves the diffusion of the reactant gases through these papers. After the molding, the sheets are postcured by putting them in an oven at 150–200 C for 2 hours or more.

Once the paper is cured, the carbonization and graphitization steps are performed under an inert environment. The exact temperatures for these two steps depend on the desired characteristics of the final products. In general, it is difficult to determine the best temperatures for these steps. In addition, the rate at which the sample is heated inside the furnace is an important parameter that determines the characteristics of the final product. For example, Mathur et al. [11,12] determined that a carbonization tempera-ture of around 2,000 C (with a 900 C/hour heat rate) and a graphitization temperature of 2,500 C produced a CFP that performed the best in a PEMFC because it improved the hydrophobic properties of the material.

In contrast with this, Liu, Ko, and Liao [13] and Liu et al. [14] reported the fabrication of CFPs that were carbonized at temperatures between 1,300 and 1,400 C. Carbon black particles or graphite powder can also be added to the resin-based solution that is impregnated in the paper in order to improve the electrical conductivity (and decrease contact resistance) of the CFP. By add-ing these particles, it is not necessary to perform the final carbonization or graphitization step in order to achieve high conductivity in the paper [9,13].

It is important to note that the steps explained here were based on wet-laid fibers; however, other procedures can be used to fabricate CFPs based on dry-laid, prestabilized PAN fibers [9]. Basically, the fibers are laid down to form a thin sheet or mat that is exposed to a hydro-entangling step in which a curtain of very fine water jets is impinged onto the moving sheet or mat. This step helps to change the orientation of some of the fibers, thus creating a mechanically bonded CFP. After this, the paper goes through the stabiliza-tion and carbonization stages. These materials are then filled with carbon or graphite particles and resin binders, followed by more carbonization and/or graphitization stages [9,19].

For more information and details regarding the fabrication processes of PAN fibers and carbon fiber papers, please refer to Kinoshita [6], Decrecente, Layden, and Pike [8], and Mathias et al. [9].

4.2.1.3 Carbon Fiber Paper with Aerogels

Another style of DLs using PAN fibers was presented by Glora et al. [20]. They developed carbon aerogel sheets, which were then used as diffusion layers in PEMFCs. In order to fabricate these DLs, resorcinol-formaldehyde (RF) aerogels


were combined with PAN fibers for reinforcement. Once the mixture was stored between two glass plates, it was cured for 1 day at room temperature, 1 day at 50 C, and 1 day at 90 C. Then a solvent exchange against acetone was performed, and after this the samples were dried subcritically. All samples were transformed into carbon aerogels via pyrolysis in an inert environment at 1,000 C.

However, the performance of a fuel cell with these carbon aerogels as DLs was around a factor of six lower than the performance of commercial elec-trodes. This was due mainly to the fact that the authors did not use addi-tional electrolyte when depositing the catalytically active layer, thus causing reduced ionic conductivity between the catalyst (Pt particles) and the mem-brane. In addition, the MEAs with carbon aerogels performed poorly at high current densities because the Pt particles used were 10 times larger than the ones normally used [20].

4.2.2 Carbon Cloth Fabrication

Along with CFPs, carbon cloths have also been widely used materials for dif-fusion layers in fuel cells. Figure 4.6 shows SEM pictures of typical carbon cloth materials used in fuel cells. The majority of these fabrics are made from PAN fibers that are twisted together in rolls. For details regarding how normal PAN fibers and carbon fibers are fabricated, please refer to Section 4.2.1.1. In this section, we will briefly discuss the fabrication process of carbon cloths.

The most typical precursor for carbon cloth is woven PAN fiber, although other precursors such as highly oxidized pitch fibers have also been reported [21]. The first step is to stabilize the PAN fibers (as explained in Section 4.2.1.1) and then produce a spun yarn from the material. Following this, the yarns of stabilized fibers are put into a stretch-breaking machine in order to break all the filaments randomly, thus keeping the tow in a continuous form but with fibers of certain lengths. The resulting yarns are then blended and homoge-nized with the use of different machines in order to improve the yield strength of the material. Then the yarns go through a series of spinning stages in which they are twisted and/or turned in order to hold them together [9].

The yarns then go through the process of weaving or knitting by plain weave or eight harness satin [9]. Once the cloths are ready, they go through a carbonization (or graphitization) process at temperatures between 1,000 and 2,500 C under an inert atmosphere. The exact temperature depends on the desired properties for the final product. Ko, Liao, and Liu [22] presented a study in which it was determined that PAN-based carbon fiber cloths per-formed the best in a fuel cell environment if they were heat-treated at 2,500 C during the graphitization process. The reason for this is that the electrical resistivity declined, the stacking height of carbon layers increased (directly related to conductivity), and the spacing between carbon layers within the cloth decreased with higher heat-treatment temperatures.

This same study also determined that the weaving process used to fabricate the carbon cloths plays an important role in the final thickness of the material,


SE 070813 WD15.8 mm 20.0 kV ×60 500 um

(a)

SE 070814 WD15.8 mm 20.0 kV ×60 500 um

(b)

FIGURE 4.6Scanning electron microscope pictures of typical carbon fiber cloths used in fuel cells: (a) E-Tek carbon cloth “A” with no PTFE (reference bar indicates 500 μm); (b) E-Tek carbon cloth “A” with 20% PTFE (reference bar indicates 500 μm); (c) close-up view of the E-Tek carbon cloth “A” with 20% PTFE (reference bar indicates 50 μm).


which is a critical parameter in fuel cells. Ko [23] proposed carbon cloths of high density and high conductivity made out of oxidized fibers of polypropylene.

Table 4.2 shows the properties of carbon cloth materials that are commer-cially available and have been used as diffusion layers in fuel cells.

4.2.3 Metal Diffusion Layers

Due to their high electrical and thermal conductivity, materials made out of metal have been considered for fuel cells, especially for components such as current collectors, flow field bipolar plates, and diffusion layers. Only a very small amount of work has been presented on the use of metal materials as diffusion layers in PEM and DLFCs because most of the research has been focused on using metal plates as bipolar plates [24] and current collectors. The diffusion layers have to be thin and porous and have high thermal and electri-cal conductivity. They also have to be strong enough to be able to support the catalyst layers and the membrane. In addition, the fibers of these metal materi-als cannot puncture the thin proton electrolyte membrane. Thus, any possible metal materials to be considered for use as DLs must have an advantage over other conventional materials.

Of the most common DL materials, carbon fiber papers are widely known for being mechanically weak because their microstructure is destroyed when excessive compression forces are applied to them (i.e., when compressing a fuel cell). This destruction of the materials affects the porosity, which has a

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direct impact on gas and liquid transport mechanisms within the DL and the fuel cell [25]. Carbon cloths have better mechanical strength due to their com-pressibility. Compression forces on DLs can also affect the overall electrical conductivity of the cell (see Section 4.4.5). Thus, metal meshes (e.g., expanded metals or screens), perforated sheets, felts, and foams have all been consid-ered as possible DLs to overcome some of the conventional DL limitations.

4.2.3.1 Metal Meshes

In electrochemical systems, metal meshes have been widely used as the back-ing layers for catalyst layers (or electrodes) [26–29] and as separators [30]. In fuel cells where an aqueous electrolyte is employed, metal screens or sheets have been used as the diffusion layers with catalyst layers coated on them [31]. In direct liquid fuel cells, such as the direct methanol fuel cell (DMFC), there has been research with metal meshes as DLs in order to replace the typical CFPs and CCs because they are considered unsuitable for the trans-port and release of carbon dioxide gas from the anode side of the cell [32].

Scott et al. [33] designed a DMFC with stainless steel mesh as the anode FF plate that was able to remove the carbon dioxide gas effectively. Later, the same research group was able to demonstrate that using similar meshes as DLs in the anode side also improved the overall gas removal [26,34] (wet-proofed CFP was used as the DL on the cathode side). These meshes were used on the anode side and were made out of catalyzed Ti because similar meshes have been used extensively as catalyzed electrodes in other indus-tries, such as the chlor-alkali industry [26].

Shao et al. [35] not only used a similar Ti mesh to the one presented by Scott’s group but also used a Ti mesh as the cathode DL in a DMFC. The main dif-ference between both meshes was that the one used on the cathode side was coated on both sides with carbon black (Vulcan XC-72) and PTFE (i.e., with MPLs on each side). It was shown that this novel cathode DL performed simi-larly to conventional CFP DLs under comparable conditions. Chetty and Scott [36] also used a catalyzed Ti mesh as the anode DL, but in a direct ethanol fuel cell (DEFC); it performed better compared to a cell with a standard DL (CFP).

Similar metal sheets have also been used as DLs in the cathode of PEMFCs. Wilkinson et al. [37,38] presented the idea of using fluid distribution lay-ers made out of metal meshes with electrically conductive fillers inside the holes of the meshes. A very similar idea was also presented by Hamada and Nakato [39]. Losfeld and Lieven [40] presented another example of fuel cells that use metal meshes as diffusion layers along with metal FF plates.

In most of these studies, the meshes used were expanded metals that are readily available because they are used in many industries. Expanded met-als are usually made with a precision die that slits and stretches the material in a single operation. The material then goes through a set of rollers so that the desired thickness can be obtained [41]. Figure 4.7 shows pictures of an expanded metal mesh similar to those used as DLs in fuel cells.


SE 070826 WD16.3 mm 20.0 kV ×60 500 um

(a)

SE 070827 WD16.8 mm 20.0 kV ×60 500 um

(b)

FIGURE 4.7Scanning electron microscope pictures of a typical expanded metal (reference bars indi-cate 500 μm): (a) picture taken with material sitting flat; (b) picture taken with the material tilted 45 .


4.2.3.2 Sintered Metals

Sintered metals have also been considered as possible DLs in fuel cells since these materials exhibit great mechanical strength and electrical conductivity. In DMFCs, Liu et al. [42] used a sintered stainless steel fiber felt (316L) as the DL (with MPL) on the anode side (see Figure 4.8). In this study it was claimed that due to the high electrical conductivity and hydrophilicity of the sintered material, the power density of the DMFC was improved (compared to a cell with a CFP with MPL as the anode DL). Oedegaard et al. [43] came up with similar conclusions using a metal wire cloth as the anode DL. The metal cloth gave better performance than carbon cloths (E-TEK Type A CC) with different PTFE contents (0, 15, and 30 wt%), and this was attributed mostly to the lower resistance that the metal DL had compared to the wet-proofed CC.

Hottinen et al. [44,45] used titanium sintered meshes as DLs on the cathode side of a PEMFC because the porosity of these metal sheets does not reduce when in compression. It was demonstrated that in order for the cell to achieve the required performance, the sintered meshes had to be coated with platinum. However, the results showed that a cell with CFP (SIGRACET GDL10-BB) as the DL still performed slightly better (especially at high current densities) than the cell with the Pt-coated sintered Ti mesh. Cisar et al. [46] presented another exam-ple in which a DL consisting of sintered metal fibers was used on the cathode side of a PEMFC. Once put together, these fibers were unified or bonded to the FF plate (made out of metal) in order to combine the two components into one.

It is important to mention that sintered metal meshes are widely used as the diffusion layers in unitized regenerative polymer fuel cells (URFCs) and

100 um

FIGURE 4.8Micrograph of sintered stainless steel fiber felt used as DL (reference bar indicates 100 μm). (Reprinted from J. Liu et al. Journal of Power Sources 133 (2004) 175–180. With permission from Elsevier.)


in electrolysis cells. This is because carbon-based materials tend to corrode at high potentials on the oxygen electrode side during the water electrolysis operation [47,48].

4.2.3.3 Micromachined Meshes

Micromachining has also been used in order to fabricate perforated metal meshes as DLs in micro fuel cells. Fushinobu, Takahashi, and Okazaki [49] developed thin, perforated Ti films and sheets using microfabrication tech-nology (see Figure 4.9). However, the performance of the PEMFC using these DLs was substantially lower compared to a cell with CFP as a DL. Wan, Wang, and Mao [50] recently published a study in which a thin Ti substrate with micro flow channels machined on it was used in a micro-PEMFC. This Ti DL was also used as the cathode current collector, flow field plate, and DL. Acceptable performance was achieved with these DLs when they were coated on both sides with a microporous layer.

Zhang, Advani, and Prasad [51,52] also used microfabrication techniques in order to develop a thin, perforated copper foil and use it as a cathode DL in a PEMFC. In addition to the metal DL, an “enhancement” layer was used that con-sisted of a porous material located between the perforated copper foil and the FF plate (CFP was used in this study). This layer improved the overall short-term performance and water management of the cell. However, the authors did not discuss any possible long-term issues related to contamination of the membrane due to the use of a copper DL.

FIGURE 4.9Scanning electron microscope image of a Ti DL with 25 μm diameter holes (reference bar indi-cates 100 μm). (Reprinted from K. D. Fushinobu et al. Journal of Power Sources 158 (2006) 1240–1245. With permission from Elsevier.)


4.2.3.4 Metal Foams

Metal foams have been used in the past in the development of FF plates. However, Gamburzev and Appleby [53] used Ni foams as both a DL and a flow field plate with an MPL layer on one of its surfaces. They observed that such a design had high contact resistance between the nickel foam and the MPL and also increased gas diffusion resistance due to the required MPL thickness. Arisetty, Prasad, and Advani [54] were able to demonstrate that these materials can also be used as potential anode diffusion layers in DMFCs (see Figure 4.10). In fact, the nickel foam used in this study performed better than a carbon cloth (Avcarb 1071HCB) and a stainless steel mesh. However, it was recognized that a major drawback for these foams is their susceptibility to corrosion.

In another study, Chen and Zhao [55] demonstrated that by using a Ni-Cr alloy metal foam as the cathode DL (and current collector), instead of a CFP or CC, the performance of a DMFC can be enhanced significantly due to the improvement of the mass transfer of oxygen and overall water removal on the cathode side. Fly and Brady [56] designed a fuel cell stack in which the distribution layers were made out of metal foams (open cell foams). In addition, more than one foam (with different porosity) could be sandwiched together in order to form a DL with variable porosity.

4.2.4 Engineered Diffusion Layers

One of the main issues with the CFPs and CCs used as DLs is the noncon-trolled variation in porosity (and other localized properties)—for example,

1 mm

FIGURE 4.10Microscope image of a Ni metal foam used as a DL in direct methanol fuel cells (reference bar indicates 1 mm). (Reprinted from S. Arisetty et al. Journal of Power Sources 165 (2007) 49–57. With permission from Elsevier.)


the porosity characteristics between carbon papers are not repeatable [57]. These materials are hard to improve because only average pore sizes and volume densities can be measured and most of the development has been based on empirical parameters (see Figures 4.3 and 4.6). In addition, water management and mass transport limitations, which are some of the major issues for the fuel cell, could be improved considerably if the DLs could be carefully designed, taking into account all their parameters [49]. Therefore, researchers have begun to consider engineered materials in which the porosity and perforations used for liquid and gas transport are specifically designed to achieve the desired performance as possible DLs. It is important to note that CC materials are more ordered and property controlled than CFPs; thus, they could be considered as partially engineered DLs.

One of the first fuel cell designs that used these types of engineered materials was introduced by Wilkinson et al. [37]. They presented the idea of using a fluid, impermeable sheet material (made out of expanded metals, metal screens, or graphite foil) as the DL for the anode and the cathode. These sheets had perfora-tions in the regions corresponding to the active area of the cell and were filled with electrically conductive filler or with catalyst, thus increasing the overall active zone (ideal for liquid fuel cells). In addition, these materials were designed so that, between them, they were able to seal completely any leaks, thus omitting the need of extra seals or gaskets and reducing the overall costs.

Later, the same research group presented a similar design of this perfo-rated material in which more advanced ideas related to the development of the perforations were introduced [38]. They showed that the perforations could increase in size in a graded or banded manner or that their area den-sity could change also in a graded or banded manner (see Figure 4.11a–d). Also, the perforations could have different cross-sectional shapes—for example, straight, tapered, or any other combination (Figure 4.11e); they could also be connected to each other in different ways with passages or grooves (Figure 4.11f). In fact, the idea of integrating both the DL and the FF plate into one component and thus reducing the overall size and cost of the cell even further was also explained in detail. All of these designs can also be tailored to improve the overall gas permeability of the specific DL; thus, the mass transport limitations within the cell can be reduced [38].

Mercuri [58] of Graftech Inc. developed graphite sheet materials made out of compressed expanded graphite particles. In general, these materials consist of a monolithic graphite structure that is flexible, conformable, and electrically and thermally conductive and is not deformed easily when compressed. These graphite products are manufactured from crystalline natural graphite flakes that have been intercalated with acid while exposed to oxygen or air. Then they are thermally shocked to generate exfoliated graphite with a vermiform structure. This graphite is pressed without binders until a flexible continuous graphite sheet is formed. The sheets are then reinforced with resins [59]. To be used as diffusion layers in electrochemical cells, the sheets or foils were per-forated by mechanical impact so that the reactant gases and liquids could flow


through them. This process was performed with a pressing roller, which had protrusion elements that were forced through one side of the graphite sheet.

Similar to the earlier work of Wilkinson et al. [37,38], the perforations could also be designed using different sizes, shapes, and placements in order to reduce water flooding in fuel cells [60]. Thus, a series of different design

(a)

(b)

FIGURE 4.11Perforated pores: (a) different sizes in graded manner; (b) different sizes in banded manner; (c) different density of perforations in a graded manner; (d) different density of perforations in a banded manner; (e) perforations with different cross sections; (f) plan view of perforations with extra passages and grooves. (From D. Knights et al. (2003) US Patent 2003039876.)


configurations could be fabricated to meet certain property requirements. For example, the pores corresponding to the outlet of the cathode flow field plate could be made larger (or with different shapes) so that the liquid water could be removed effectively. Such pores can also be connected directly to the FF channels if both the sheet materials and the bipolar plates are integrated into one component [61]. In addition, these materials can be made hydro-phobic (waterproof), the perforations can be filled with catalyst (important

(c)

(d)



in liquid fuel cells), and the overall electrical and thermal conductivity can be enhanced [62–64]. Figure 4.12 shows pictures of some expanded graphite foils fabricated by Graftech Inc. with 2,500 tips per square inch (Figure 4.12a) and 1,200 tips per square inch (Figure 4.12b).

Unfortunately, few experimental data have been published regarding these types of diffusion layers. Yazici [65] presented a study in which the graphite foils made by Graftech Inc. were used as cathode diffusion layers in DMFCs. Two foils were used; one was made out of 80% expanded graphite and 20% PTFE coated carbon particles to form a porous sheet, and the other was identical to the first except that it was perforated for more permeability with 2,500 tips per square inch (15% open area).

The second foil showed better performance due to more effective transport of gas and liquid within the DL. In fact, when compared to a CC (ELAT with MPL), this second DL did better, especially at dry conditions. Once the cell temperature was increased from room temperature to 80 C, the difference between the CC and the graphite foil DL was not as obvious. In another study, an expanded graphite foil (100% graphite) was used with an MPL on top of one of its surfaces as the cathode DL in a cylindrical PEMFC [66,67].

The research group led by Zhou [68,69] has also investigated the use of per-forated materials as cathode DLs in PEMFCs through modeling. However, the group’s model results were produced based only on conservation of mass and momentum (i.e., energy and electrochemical equations were not considered). This model studied only how different shapes of perforations affected the liquid water removal through the DL.

(e)

(f)



In the previous section, it was mentioned that Zhang et al. [51,52] devel-oped a technique to micromachine a thin metal film made out of copper. This material had a number of perforations that followed a predetermined pattern and such a design can be changed fairly easily by simply changing one of the masks used in the fabrication process. Although copper is not an ideal material to be used as a DL in a fuel cell due to the contamination issues, the development of fabrication techniques similar to those mentioned

0.5 mm

(a)

0.5 mm

(b)

FIGURE 4.12Grafcell flexible graphite diffusion layers by Graftech Inc. with 21% open area: (a) 2,500 tips per square inch; (b) 1,200 tips per square inch.


previously (but with more acceptable materials) is critical to and important for the design and use of engineered materials as DLs.

Aravamudhan, Rahman, and Bhansali. [70] developed a micro direct ethanol fuel cell with silicon diffusion layers. Each silicon substrate had a number of straight micropores or holes that were formed using microelec-tromechanical system (MEMS) fabrication techniques. The pores acted both as microcapillaries/wicking structures and as built-in fuel reservoirs. The capillary action of the microperforations pumps the fuel toward the reaction sites located at the CL. Again, the size and pattern of these perforations could be modified depending on the desired properties or parameters. Lee and Chuang [71] also used a silicon substrate and machined microperforations and microchannels on it in order to use it as the cathode diffusion layer and FF channel plate in a micro-PEMFC.

4.2.5 Other Diffusion Layers

The need for different and novel materials as possible DLs has increased sub-stantially in the last few years—especially with the development of new and more complex fuel cell designs. Furthermore, the interest in small-scale fuel cells to be used as battery replacements in portable electronic devices such as PDAs, laptops, cell phones, music players, etc. has pushed the research for innovative, inexpensive, and efficient fuel cells further [72,73]. Therefore, it is not surprising that most of the recent new DL materials are being used in micro fuel cells.

4.2.5.1 Silicon-Based Materials

Modroukas, Modi, and Chette [74] fabricated mesh structures through etching into silicon wafers followed by gold plating in order to use them as diffusion layers and current collectors on both the anode and cathode sides of PEMFCs. These DLs performed similarly in a fuel cell compared to conventional DLs. Yeom et al. [75] also fabricated silicon meshes that were used as DLs, current collectors (after being coated with a conductive layer), and mechanical supports for the overall cell. This cell was used successfully with three different fuels consisting of hydrogen, formic acid, and methanol. Figure 4.13 shows a schematic of how their silicon meshes were fabricated. These research groups demonstrated similar design ideas to those presented by Aravamudhan et al. [70] and Lee and Chuang [70,71] (see Section 4.2.4).

Xiao et al. [76] designed a complex cell in which a membrane is sandwiched between two silicon substrates. A glass wafer is bonded to the silicon sub-strates to seal the cell and to create reservoirs between the glass and the silicon DLs. Each DL had a number of perforations for the reactants to flow through and rough textures on its surface on which the CL was impreg-nated. These textures on the surfaces were fabricated using micromachin-ing techniques. Both hydrogen/oxygen and methanol/oxygen performances were evaluated using this novel cell design.


(a)

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(g)

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Silicon Gold Polyimide Pt Black Nafion 112

FIGURE 4.13Schematic of the Si-μMEA fabrication process: (a) sputter Au layer on double-side polished wafer; (b) pattern Au layer with liftoff process; (c) spincoat and cure a polyimide layer; (d) perform the double-sided photolithography to pattern etch pits; (e) etch Si in ICP-DRIE to form Au/Si electrode; (f) dice the wafer into a single die; (g) RIE etch the polyimide layer with a shadow mask to expose current collecting region; (h) electroplate Pt black on Au layer; (i) sand-wich both electrodes with Nafion 112 in a hot-press bonder. (Reprinted from J. Yeom et al. Sensors Actuators B 107 (2005) 882–891. With permission from Elsevier.)


Another way to use silicon wafers as DLs was presented by Meyers and Maynard [77]. They developed a micro-PEMFC based on a bilayer design in which both the anode and the cathode current collectors were made out of conductive silicon wafers. Each of these components had a series of micro-channels formed on one of their surfaces, allowing the hydrogen and oxygen to flow through them. Before the channels were machined, a layer of porous silicon was formed on top of the Si wafers and then the silicon material beneath the porous layer was electropolished away to form the channels. After the wafers were machined, the CLs were added to the surfaces. In this cell, the actual diffusion layers were the porous silicon layers located on top of the channels because they let the gases diffuse through them toward the active sites near the membrane.

In a similar design, D’Arrigo et al. [78] also used porous silicon as the DL in a fuel cell. The porous silicon was deposited by chemical vapor deposition (CVD) on top of a silicon wafer that already had microgrooves machined on it. Then, catalytic particles were deposited on top of the porous silicon layer. Unfortunately, no performance-related data indicating whether the cell was acceptable or not were published for this design.

4.2.5.2 Other Materials

Besides silicon, other materials have also been used in micro fuel cells. Cha et al. [79] made micro-FF channels on SU8 sheets—a photosensitive polymer that is flexible, easy to fabricate, thin, and cheaper than silicon wafers. On top of the flow channels, for both the anode and cathode, a paste of carbon black and PTFE is deposited in order to form the actual diffusion layers of the fuel cell. Mitrovski, Elliott, and Nuzzo [80] used a gas-permeable elastomer, such as poly(dimethylsiloxane) (PDMS), as a diffusion layer (with platinum elec-trodes embedded in it) for liquid–electrolyte-based micro-PEM fuel cells.

Other diffusion layer approaches can also be found in the literature. Chen-Yang et al. [81] made DLs for PEMFCs out of carbon black and unsintered PTFE comprising PTFE powder resin in a colloidal dispersion. The mixture of these materials was then heated and compressed at temperature between 75 and 85 C under a low pressure (70–80 kg/cm2). After this, the DLs were obtained by heating the mixture once more at 130 C for around 2–3 hours. Eventually, the amount of resin had a direct influence on determining the properties of the DL. The fuel cell performance of this novel DL was shown to be around a half of that for a CFP standard DL. However, because the manufacturing process of these carbon black/PTFE DLs is inexpensive, they can still be considered as potential candidates.

Diffusion layers have been developed that combine more than one type of material on them. For example, Koschany [82] proposed the use of lay-ered DLs made out of two materials with different gas permeabilities and manufactured one on top of the other. Normally, the materials with the lowest permeability were made out of expanded graphite or metal, and the


other materials were made out of carbon fibers. Another example of com-bining more than one material is the one published by Voss, Kollmann, and Kollmann [83], in which a new material called POLYMET is described. This is a composite polymer with a porous three-dimensional polymeric struc-ture metalized with a surface coating of different types of metals and alloys. The end result is a woven/felt-based material that can be used as a DL in fuel cells and that helps to reduce the internal resistance of a PEMFC with a membrane.

Campbell et al. [84] developed DLs made out of glass fiber webs filled with carbon and PTFE particles. The same research group later designed special DLs made with different carbons claiming to improve the overall fluid diffu-sion toward the catalyst layer [85].

4.2.6 Performance Comparison between Diffusion Layers

It is important to understand how each type of DL performs under different operating conditions in order to determine which is best for the appropriate application. As expected, most of the studies in which the performances of more than one DL are compared deal with CFP and CC DLs. Publications in which other DL materials are compared to conventional designs are scarce, and only a very few have been mentioned previously (please refer to Sections 4.2.3 and 4.2.4). Therefore, performance comparison of the conventional DL materials, CFPs and CCs, will be the emphasis in this section.

4.2.6.1 Performance Comparison in PEM Fuel Cells

In PEMFCs, Ralph et al. [86] tested a Ballard Mark V single cell with two different DLs: a carbon cloth (Zoltek PWB-3) and a carbon fiber paper (Toray TGP-090); all the other operating conditions stayed the same for both cases. It was observed that the carbon cloth demonstrated a distinct advantage over the CFP at high current densities (>600 mA/cm2), while at low current densi-ties both DLs performed similarly. It was claimed that this was because the CC material enhanced mass transport properties and improved the water management within the cell due to its porosity and hydrophobicity.

It is important to note that this single cell was tested at 100% relative humidity and the cell temperature was 80 C. Spernjak, Prasad, and Advani [87] used a transparent fuel cell (the end plate was made out of polycarbon-ate) to visualize the water accumulation inside the FF channels with dif-ferent DL materials. It was observed that, at humidified conditions, the CC (untreated—no PTFE—AvCarbon TM 1071 HCB) was able to perform and remove the water better than a CFP (untreated—no PTFE—TGP-H-60). In fact, it was concluded that water was trapped inside the DL when the CFP was used, resulting in flooding of the CL.

Similar results were also reported by Moreira et al. [88] after they tested a small single cell with both CFP and CC as the DLs. However, the difference


between the performances of each DL was not as large; in fact, the CFP did better at lower current densities (due to lower ohmic losses). This slight dif-ference in observation may be caused by the fact that these tests were per-formed with dry gases (i.e., no humidification) and a cell temperature of 25 C. Williams, Kunz, and Fenton [89] and Williams et al. [90] performed an exten-sive analysis with different materials as DLs over a wide range of humidity conditions. The CFP DL (with an MPL) showed greater performance at low relative humidity conditions, while the more hydrophobic CC (E-TEK CC with MPL) was superior at saturated inlet conditions. Antolini [91] had simi-lar observations and experimental results.

In order to improve the performance of fuel cells, Wilkinson and St-Pierre [92] and Johnson et al. [93] compared typical CFP cathode DLs with modi-fied DLs (from CFP or CC) that improved the mass transport at high current densities. Figure 4.14 shows the different DLs that were used to improve the cell’s performance. Similar strategies can be implemented in other types of DLs, such as metallic or engineered.

Through mathematical modeling based on experimental results, Wang, Wang, and Chen [94] were also able to conclude that, at high humidity condi-tions, CC performs better than CFP because of the highly tortuous structure of the CFP, which leads to an increase in mass transport limitations under these conditions. In addition, the detachment of water droplets from the CFP smooth surface is difficult, thus increasing the amount of water films blocking the pores. On the other hand, at dry conditions, the CFP is able to retain water within the MEA and maintain the necessary hydration for the membrane.

In addition, Akyalçin and Kaytakoglu [95] were able to draw similar con-clusions experimentally. They also concluded that CFP performs better at

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FIGURE 4.14Polarization curves showing the effect of modified diffusion layers (air/H2, 2/1.5 stoichiom-etry, 3.1/3.1 bara, 100/100% relative humidity, 80 C). (Reprinted from D. P. Wilkinson and J. St-Pierre. Journal of Power Sources 113 (2003) 101–108. With permission from Elsevier.)


dry conditions because CC has a highly coarse fiber network, which leads to the active catalyst particles entering the pores of the CC, thus causing high ohmic resistance due to low contact with the CL. However, it is impor-tant to note that Frey and Linardi [96] stated that the carbon cloth, which is more flexible with a structure that does not change during pressing, had better electrical contact with the CL and the DL. On the other hand, due to its smooth surface, the CFP had less impact on the electrode surface, creating an inferior electrical contact with the CL.

It is important to note that most of these studies concentrated on the cath-ode side of the PEMFC because it has been shown that the difference between materials on the anode side does not represent a major performance loss due to flooding at high- or low-humidity conditions [97]. Therefore, the type of DL used on the anode side does not result in major differences over a wide range of operating conditions.

4.2.6.2 Performance Comparison in Direct Methanol Fuel Cells

In DMFCs, methanol crossover and carbon dioxide gas management are critical issues that have be dealt with. Argyropoulos, Scott, and Taama [98] used a transparent fuel cell (the anode end plate was made out of acrylic) to visualize the CO2 evolution and management on the anode side. Both CFPs and CCs were used as anode DLs and it was observed that CFP (Toray carbon paper) was not a suitable material due to its poor gas removal properties.

On the other hand, with the same amount of PTFE as the CFP, carbon cloth (E-TEK type A carbon cloth) performed better and was able to eject the gases within the DL more effectively, thus giving more access to the methanol.

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FIGURE 4.15Polarization curves for Toray carbon paper 260 and carbon cloth A. (Modified from A. Oedegaard et al. Journal of Power Sources 127 (2004) 187–196. With permission from Elsevier.)


Oedegaard et al. [43] were also able to confirm that a fuel cell using CFP (TGP-H-090) as the anode DL achieved lower current densities than when the cell used CC (E-TEK type A CC) (see Figure 4.15). In addition, the cur-rents were very unstable with the CFP because CO2 bubbles were blocking access of methanol to the CL.

4.3 Treatments and Coatings

After the diffusion materials are fabricated, a number of treatments and coatings are still necessary in order to tailor the final properties for these materials based on the specific fuel cell application and the associated oper-ating conditions. The following sections will explain in detail the different treatments that are normally used on diffusion layers for fuel cells. Brief examples showing how these treatments change the performance of the DLs will also be discussed.

4.3.1 Hydrophobic Treatments

Diffusion layers used in PEM and DLFCs are normally treated with a hydrophobic agent such as PTFE or fluoroethylenepropylene (FEP). This treatment increases the hydrophobicity of the materials because most of the CFPs and CCs are not hydrophobic enough after fabrication. In addi-tion, it is important to note that by coating the DL with these agents, we mean that the whole material (including the fibers) is coated, rather than just the top surface of the material. For cathode DLs, this coating is extremely important and vital because most of the water produced and accumulated inside the cell exits through the cathode side. For the anode DL, this coating is not as critical but still important (especially when dealing with back diffusion of water) and it can provide some structural strength to the DL.

All the examples discussed in the following subsections are meant to be taken simply as an indication of how important it is to determine the correct hydro-phobic content for a specific DL. It is also critical to take into account all the possible operating conditions in which the fuel cell will function because these parameters will have a direct impact on the overall performance of the cell.

4.3.1.1 Fabrication Processes and Procedures

4.3.1.1.1 PTFE Treatments

The PTFE can be coated onto the diffusion layer in many ways. The most typical way is to dip the material slowly into a solution containing a pre-determined amount of PTFE, at a rate slow enough that the solution can adsorb onto the paper [99,100]. After the paper is left inside the solution for a


determined time (depending on the desired amount and the material used), the DL is put inside an oven at 80 C on top of a number of pins attached to a plate so that uniform distribution of the solution is achieved [101]. These steps are repeated until the weight of PTFE desired is achieved.

The material is further heated at 270 C in order to heat-treat and remove any extra surfactant or moisture originating from the PTFE solution. Following this step, the material is sintered at a temperature in the range of 340–350 C [100,101]. It is important to note that the literature contains a number of varia-tions with respect to these process steps. However, the overall process and order of steps are very similar in most published reports. The common steps that may change are the dipping time, the drying temperature, and the sin-tering temperature. In some cases, the drying and sintering steps are com-bined, with the temperature increased stepwise from room temperature to around 350 C [102].

4.3.1.1.2 Other Hydrophobic Treatments

Other ways of making a hydrophobic layer on the DLs for PEM fuel cells can also be found in the literature. For example, Chiu and Wang [103] devel-oped a method to coat the DL with a hydrophobic layer through sputter deposition of the PTFE. This technique forms fluorine-containing func-tional groups on the DL surface, consequently creating a hydrophobic layer. In this method, a PTFE sheet (target) and diffusion layer (substrate) are used in the process. A radio frequency power supply is connected to the target for plasma generation and the substrate is kept electrically floating. Ar, N2, or dry air can be used as the sputtering gases. The advantage of this coating technique is that it produces coatings that are more uniform than those produced by other methods.

Pai et al. [104] implemented a method in which the carbon fibers used to fabricate carbon cloths were treated by a CF4 plasma method, thus improving the hydrophobic property of the final material. The CF4 method is adapted from the radio frequency (rf) plasma treatment that is applied to polymers to modify their surface properties. Through this method, the whole surface of the substrate is hydrophobic without having residual materials that block or seal the pores within the CC. Therefore, the step of coating the DL with PTFE is avoided. A fuel cell that used plasma-treated diffusion layers performed better than a similar cell with PTFE-coated DLs. Taniguchi and Yasuda [105] presented another method using plasma polymerization to coat a CFP with hexafluoropropylene (HFP). It was observed that this treatment increased the wetting contact angle of the CFP surface and the cell’s performance improved significantly after just 5 minutes of the plasma treatment. Although, this pro-cess is still at an early stage, it is very reproducible and reliable.

Another method to improve the hydrophobic treatment of DLs was pre-sented by Wang, Shi, and Du [106]. They coated the CFP fibers (TGP-H-090) with a sucrose aqueous solution for around 6 hours. After that the material was heated at 400 C in order to carbonize the sucrose solution. This process


was repeated until the desired amount of carbon was achieved. Through this method, the surface of the carbon fibers was significantly rougher and thus the PTFE could be homogeneously distributed over the whole surface of the DL. This improved the hydrophobicity of the material under lower PTFE loadings (compared to CFPs without the extra carbonization procedure). A carbonized CFP (used as the cathode DL) with 10 wt% PTFE had the best performance compared to normal CFPs with the same and higher amounts of PTFE.

Ji and Kumar [107] presented an invention in which, after treatment with PTFE, the cathode diffusion layers are coated with a silicon solution in order to enhance the hydrophobic properties of the DLs. This silicon does not have to be coated over the whole surface of the DL, but could be coated in just cer-tain areas, depending on the design of the cell, the location of the cell within a stack, and the desired hydrophobic properties. It was demonstrated how a DL with the silicon coating improved the performance of a single fuel cell when operating at high relative humilities [107].

4.3.1.2 Effect of Hydrophobic Treatment

In PEMFCs, it is critical to determine the appropriate amount of PTFE con-tent on both the anode and cathode DLs because that can change the perfor-mance of the cell substantially. The most common loadings of PTFE and FEP are from 5 to 30 wt%. Bevers, Rogers, and von Bradke [100] studied the effect of PTFE content and sinter temperatures on the performance of a DL (Sigri PE 704 CFP). It was demonstrated that high PTFE content and high sinter temperatures led to high hydrophobicity as well as to low electrical conduc-tivities, which decreased the performance of the fuel cell.

On the other hand, with high PTFE content (>40 wt%), the diffusion prop-erties (i.e., gas permeability) of the CFP decreased because the high quan-tities of PTFE particles partially filled some of the pores, thus decreasing the overall permeability of the DL. However, with high sinter temperatures (from 400 to 420 C), the diffusion actually increased because, at higher sin-ter temperatures, the available PTFE is allowed to coat paper fibers more thoroughly, thus moving the PTFE from spaces between fibers to the fibers themselves [100]. This allows better gas flow diffusion through these areas.

Similar conclusions regarding how high PTFE content decreases the gas permeability and porosity of CFPs (TGP-H-060 and TGP-H-090) were pre-sented by Park et al. [102]. It was also observed and concluded that the main transport mechanism of the water through the DL was shear force or evap-oration, instead of capillary forces (main force in CLs). They also showed that a CFP (TGP-H-060) with 15 wt% PTFE had the best performance with relatively dry conditions compared to a thicker paper with the same hydro-phobic content.

Lin and Nguyen [108] did an extensive study on the effect of the PTFE content on the performance of Toray and SGL Carbon SIGRACET carbon


fiber papers. It was observed that increasing the hydrophobicity of the DL enhanced both the gas and water transport when the fuel cell operated with high levels of humidity. However, excessive amounts of PTFE reduced the amount of hydrophilic pores, thus deteriorating the water flow out of the CL and the DL. Velayutham et al. [109] used carbon fiber papers (Avcarb P50) with different PTFE contents (from 7 to 30 wt%) and with an MPL having 20 wt% PTFE (the hydrophobic content of the MPL was not changed). It was observed that when the cell was operated at 55 C and ambient pressure con-ditions, the DL with 23 wt% PTFE performed very well in both the low and high current density regions (see Figure 4.16).

In addition, it was shown that the performance of the DLs improves at first when the content of PTFE is increased, but that once a maximum point is reached, the performance decreases rapidly. In this study, the cell was operated with anode and cathode relative humilities of >100% and <100%, respectively.

Shimpalee, Beuscher, and Van Zee [110] used a mathematical approach based on experimental data to determine how the water flooding in PEMFCs is affected by the PTFE content in DLs. The main goal was to use the mentioned model as

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FIGURE 4.16Effect of PTFE in the DL of PEM fuel cells. Operating conditions: fuel cell temperature of 55 C; anode and cathode dew point temperatures of 65/40 C; hydrogen and air stoichiometries of 1.5/2.5; ambient pressure for air and hydrogen; serpentine flow field for anode and cathode. Anode DL was an AvCarb P50 CFP (PTFE content not specified); cathode DL was an Avcarbon P50 CFP. Electrodes were fabricated using 20% Pt/C. Anode catalyst loading was of 0.25 mg Pt cm–2 and 0.5 mg cm–2 of Nafion (5% solution); cathode catalyst loading was of 0.50 mg Pt cm–2 and 1.0 mg cm–2 of Nafion (5% solution). Type of membrane was not specified. (Reprinted from G. Velayutam et al. Fuel Cells 7 (2007) 314–318. With permission from John Wiley & Sons, Inc.)


the basis of a simple approach that could help fuel cell designers determine the correct hydrophobic content based on the desired properties and parameters.

Tüber, Pocza, and Hebling [111] used a transparent PEM fuel cell in order to visualize the water buildup inside the FF channels and on the surface of different DLs (treated and not treated) while operating the fuel cell at room temperature. In the case of the hydrophobic CFP (TGP-H-090 with 25 wt% PTFE), it was observed that the water appeared randomly distributed along the flow channels. The water produced at the cathode side seemed unable to penetrate the CFP until enough pressure was built up and then small drop-lets were formed on the surface of the DL.

Similar observations were presented by Spernjak et al. [87], who also devel-oped a transparent fuel cell to visualize the different behavior of treated and untreated DLs. This cell gave the indication that with treated DLs the water produced at the cathode side emerged as droplets on the surface of the material over the entire visible area. However, with the untreated DLs, water preferred to be in contact with the side walls of the channels; with time, the water accumulated and formed films and slugs near the flow field walls. This behavior caused greater water management issues and lower gas transport toward the active catalyst areas.

Regarding the effect of the PTFE content on the thermal conductivity of the DL material, Khandelwal and Mench [112] observed that as the PTFE percentage increased, the thermal conductivity decreased. The greatest dif-ference occurred between 0 and 5 wt% PTFE, in which a 35% reduction of thermal conductivity was recorded. However, between 5 and 20 wt% PTFE, the reduction in this property was just 29%. Thus, the authors concluded that that the thermal conductivity would approach an asymptotic value with further increase of PTFE content in DLs.

In air-breathing PEM fuel cells, Jeong et al. [113] were able to demonstrate that with high PTFE content, the fuel cell performed poorly at high current densities because the high amount of PTFE lowered the porosity of the DL, as discussed previously. They concluded that cathode DLs (Toray CFPs) with 5–10 wt% PTFE performed the best (the PTFE content of the anode DL was kept constant at 20 wt% PTFE).

Staiti et al. [114] coated TGP-090 carbon fiber papers with different amounts of FEP. The FEP content ranged from 0 (untreated) up to 60 wt%. The DLs with 20–30 wt% FEP gave the best performance overall compared to untreated DLs. Another study in which FEP was used as the hydrophobic agent was presented by Lim and Wang [101]. They reported that a CFP with 10 wt% FEP performed substantially better than a CFP with 30 wt% (both CFPs were TGP-H-090) at high humidity conditions. This was attributed to the surface modification of the DL material with excessive amounts of FEP resulting in significant obstruc-tion of surface pore and thus limiting the transport of gases and liquids.

When no humidification is used on the cathode side of a fuel cell, Song, Uchida, and Watanabe [115] showed that the low current density performance between treated (10 and 30 wt% FEP) and untreated CFPs (TGP-H-120) was


practically the same. At high current densities, the untreated DL performs bet-ter than the other DLs. However, it is important to note that in this study the air was dry and all the DLs were also coated with an MPL, which affects the overall behavior of each component. More information on how MPLs affect the fuel cell performance is discussed in Section 4.3.3.

Although in most cases it is necessary to have an even coat of the hydro-phobic layer within the diffusion layer, sometimes it is desirable to have dif-ferent regions within the surfaces that are more (or less) hydrophobic than others. Mathias et al. [116] developed a process in which the PTFE-coated DLs are placed over a vacuum fixture. Then the substrate surface (the one away from the draw) undergoes grinding while still exposed to the vacuum draw. This procedure creates PTFE dust that is pulled through the substrate by the vacuum draw. Thus, it is possible to create high and low particle (and hydrophobic) density regions.

4.3.1.3 Hydrophobic Treatment in Direct Liquid Fuel Cells

In DMFCs, Scott, Taama, and Argyropoulos [117] changed the PTFE content (from 0 to 40 wt%) of the anode DL (E-TEK type A CC) in order to observe how this affected the methanol and carbon dioxide transport through the DL. At very high levels of PTFE, the performance of the cell decreases due to an increase in resistance losses. On the other hand, when an untreated CC was used, the observed performance was the lowest of all the materials investigated. In this study it was concluded that the ideal amount of hydro-phobic agent for the anode DL is around 13–20 wt% (see Figure 4.17).

Gogel et al. [118] compared two CFPs, one untreated and one treated (25 wt% PTFE) as the anode DL (both were TGP-H-120). The fuel cell was operated at a cell temperature of 110 C, and it was observed that the DL without any treatment performed better. However, the difference between both materials was very small and the methanol permeation was actually reduced (increased Faradaic efficiency) with the treated DL. A possible explanation for this is that methanol is oxidized more effectively at the anode due to the formation and stabilization of carbon dioxide bubbles in the active area. As a consequence, the methanol concentration gradient across the membrane is reduced.

Xu, Zhao, and Ye [119] reported similar results in which it was shown that untreated CFPs (TGP-H-090) performed better than CFPs that were treated with PTFE. They also stated that the limiting current density decreased when the DLs were coated with the hydrophobic agent, possibly due to the lower permeability and higher mass transfer resistance within the material. Thus, from their point of view, without taking into account the methanol crossover, they recommended not using PTFE-treated CFPs as anode DLs. The same research group also investigated the influence of hydrophobic treatments on the cathode DL (TGP-H-090 CFPs) of a DMFC [120]. They concluded that high PTFE loading decreased the cell’s performance and caused fluctuations in the current when the cell was discharged at a given voltage. This was


attributed to the decrease in porosity in the CFP due to the blockage of pores with PTFE particles. Thus, it was recommended to use low PTFE loadings in order not to decrease the cell’s performance.

4.3.2 Hydrophilic Treatments

As mentioned previously, hydrophobic treatments are one of the most com-mon ones for DLs in most cases and operating conditions. However, to a lesser extent, hydrophilic properties in DLs may also be desired when deal-ing with specific conditions (such as dry). It is important to note that the amount of work found in the literature regarding hydrophilic treatments in DLs is very limited.

4.3.2.1 Hydrophilic Treatments for PEM Fuel Cells

Campbell, Chisham, and Wilkinson [121] found that the catalyst utilization in the electrode and fuel cell performance could be improved by making the carbon-supported catalyst hydrophilic. This was done by treating the carbon-supported catalyst with a suitable acid such as nitric acid in order to introduce the surface oxide group on the carbon. In principle, this same approach could be applied to the carbon components of the DL and MPL.

Tüber et al. [111] used a transparent fuel cell (at 30 C) to visualize both hydrophobic and hydrophilic diffusion layers. They discussed the idea of

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FIGURE 4.17Influence of PTFE content in the anode DL of a DMFC. Operating conditions: 90 C cell tempera-ture; anode at ambient pressure; cathode at 2 bar pressure; methanol concentration of 2 mol dm–3; methanol flow rate of 0.84 cm3 min–1. The air flow rate was not specified; there was a parallel flow field for both sides. The anode catalyst layer had 13 wt% PTFE, Pt 20 wt%, Ru 10 wt% on Vulcan XC-73R carbon; TGP-H-090 with 10 wt% PTFE as cathode DL. The cathode catalyst layer had 13 wt% PTFE, Pt 10 wt% on carbon catalyst with a loading 1 mg cm–2 Pt black with 10 wt% Nafion. The membrane was a Nafion 117. (Reprinted from K. Scott et al. Journal of Applied Electrochemistry 28 (1998) 1389–1397. With permission from Springer.)


hydrophilizing a CFP to absorb and distribute the product water inside the DL, thus preventing the water clogging of gas channels. In their approach, the carbon fiber paper was dipped into a solution of sodium dodecyl sulfate (SDS) and distilled water for around 15 minutes. After this, the paper was dried in an oven at 60 C for around 30 minutes. A fuel cell was operated at a constant voltage of 0.5 V with three different DLs: a standard, hydrophobic, and hydro-philic CFP. It was observed that in all three cases the current density decreased at the beginning of the tests; however, the hydrophilic DL was the one that performed the best because the drop in current density was not as sudden as with the other two DLs. This result was attributed to the water inside the DL being distributed homogeneously as a result of the hydrophilic coating.

One downside to using this hydrophilic layer is that it degrades with time; cations (Na+) emerge from the sodium dodecyl sulfate solution and displace the H+ ions inside the polymer electrolyte membrane, thus decreasing the conductivity. It was also stated that hydrophilic treatment for PEMFCs that operate at high temperatures (or high operating conditions) may not be appropriate; however, this group did not present any results with hydro-philic DLs at high temperatures. Jung et al. [122] added SiO2 particles, which are hydrophilic, to the CL of both the anode and cathode to improve the cell’s performance at low relative humidity. This hydrophilic coating could also be applied to the DL (or MPL) instead.

4.3.2.2 Hydrophilic Treatments for Direct Methanol Fuel Cells

In direct methanol fuel cells, a hydrophilic treatment with Nafion can be applied to the anode diffusion layer to facilitate wetting and methanol mass transport. Zhang, Colbow, and Wilkinson [123] presented a number of exam-ples in which ionomer coatings were added to different CFP substrates used as anode DLs. The ionomeric coating of the anode in the vicinity of the CL provides good proton conductivity and improves wetting that enhances the access of aqueous methanol. At high current densities, the cell’s performance improved significantly when this hydrophilic layer was added. This treat-ment can also be applied to the DL and the MPL [124–127]; that is, the PTFE particles that are usually added to the MPL are replaced with the ionomer.

Colbow, Zhang, and Wilkinson [128] showed that the performance of liq-uid feed fuel cells could be increased by oxidizing the carbon diffusion layer. The DL was electrochemically oxidized in acidic aqueous solution (impreg-nated in some cases with proton-conducting ionomer) prior to application of the electrocatalyst.

4.3.3 Microporous Layers

A layer of carbon black and PTFE is usually deposited on top of one of the DL surfaces (forming a diffusion double layer) as shown in Figure 4.18. This catalyst backing layer or MPL forms smaller pores than the DL (20–200 nm


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FIGURE 4.18Scanning electron microscope pictures of SGL Carbon GDL 25 BC (CFP with MPL): (a) top view of the MPL (reference bar indicates 500 μm); (b) cross section of the DL with MPL on top (reference bar indicates 100 μm).


pores for MPLs [129] and 0.05–100 μm pores for typical CFP DLs [130,131]) and acts as another mechanism to reject water, which is critical—especially when the fuel cell is operated at high humidity levels [130]. The MPL also provides support for the CL, which is located on top of it or on the surface of the proton exchange membrane.

The catalyst layer usually consists of carbon-supported catalyst or carbon black mixed with PTFE and/or proton-conducting ionomer (e.g., Nafion iono-mer). Because the sizes of the pores in a typical DL are in the range of 1–100 μm and the average pore size of the CL is just a few hundred nanometers, the risk of having low electrical contact between both layers is high [129]. Thus, the MPL is also used to block the catalyst particles and does not let them clog the pores within the diffusion layer [57,90,132,133].

It is also important to take into account that one of the main issues with the carbon fiber paper or cloth used as the DL is the uncontrolled variation in porosity (and other localized properties) of these manufactured conventional diffusion layers; that is, the porosity characteristics between carbon papers are not repeatable [57]. These materials are difficult to improve because only average pore sizes and volume densities can be measured and much of the development has been based on empirical parameters. Thus, extensive work has focused on optimizing the MPL in order to reduce the differences within carbon paper fiber and cloth diffusion layers.

4.3.3.1 Fabrication Processes

Prior to coating one of the surfaces of the diffusion layer with the MPL, a homogeneous suspension is prepared by mixing and stirring in an ultra-sonic bath at room temperature carbon black particles and PTFE in an alco-hol solution (e.g., isopropyl alcohol) for a set period of time. It is important to note that instead of PTFE, polyvinylidene fluoride (PVDF) can also be used [134]. After the slurry or ink is completed, it can be coated on top of one of the surfaces of the diffusion layer in different ways, such as spraying, rolling, screen printing, brushing, using the doctor-blade technique, etc. The most common application method is to spray the mixture onto the DL and then let it dry in air at a temperature in the range of 80–120 C for about 1 hour. After this, the material is heat-treated at 280 C for 30 minutes to remove the dispersion agent contained in PTFE and then at 350 C for another 30 min-utes to sinter the layer [135,136]. Temperatures may change depending on the materials used and on the desired properties.

Jian-hua et al. [137] studied how the MPL sintering temperature and time affected the performance of a fuel cell. It was observed that the best perfor-mance was obtained when the MPL was sintered at 350 C for 30 minutes. Lee et al. [138] demonstrated that DLs with an MPL that was screen printed or sprayed performed substantially better than when the MPL was roll coated. In general, the spraying and screen-printing methods are the most widely used.


Yu et al. [139] developed a dry-deposition technique for coating the MPL onto a diffusion layer. This method consisted of forcing a mixture of carbon and PTFE powder through a mesh with the help of a vacuum pump located underneath the DL material. Once the mixture passed through the mesh, it was deposited on the surface of the substrate (still with the help of the vacuum pump). After this, the DL, with the MPL, was sintered at 350 C in order to melt the PTFE particles and bind all the particles together. Once the thermal stage was completed, the MPL was subjected to a rolling step in order to adjust the total thickness of the layer (MPL and DL).

All of the previously mentioned techniques involve coating the MPL on top of a substrate, such as the DL or the membrane [125]. However, the MPL can also be made out of a carbon-based polymer porous sheet that is simply placed between the CL and DL when assembling the fuel cell [129,140–142]. The sheet-based MPL approach is not as common and is not widely used. In fact, in previous years a commercially based MPL film was available (Carbel MP, W. L. Gore & Associates, Inc.) but is now no longer manufactured. Shi et al. [140] prepared an MPL sheet using a two-roll-shaft roller to roll a mixture of carbon black and PTFE repeatedly. Once the desired toughness in the film was achieved, it was sintered. This MPL sheet performed well, especially under high-humidification conditions.

Another important point regarding the fabrication process of MPLs is that, typically, when carbon fiber paper is used as the DL, the MPL is coated on only one surface of the CFP. However, when a carbon cloth DL is used, it is normally coated on both sides with MPLs. Section 4.3.3.4 will discuss these DLs with multiple microporous layers in more detail.

4.3.3.2 Effect of Microporous Layer on Fuel Cell Performance

Microporous layers are now commonly used to improve the overall perfor-mance of a fuel cell. It is believed that they play a key role in overall water management within the fuel cell. However, it is still unclear exactly how the MPL affects the water transport mechanism inside the DL and the MEA. Therefore, it is important to discuss briefly some of the different studies that have tried to tackle this issue.

Passalacqua et al. [143] were able to prove that when an MPL is interposed between the DL and the CL, the performance of the cell improves substan-tially. They concluded that the MPL reduced the size of the water droplets, thus enhancing the oxygen diffusion. This layer also prevented the catalyst particles from entering too far into the DL. Park et al. [102] concluded that with the addition of an MPL, both water management and electrical conduc-tivity improved.

Similar observations were also presented by Song et al. [115] and Holmstrom et al. [97], especially when the fuel cell’s performance at high current den-sities was investigated. In fact, it was shown that DLs without an MPL at the cathode side experienced major mass transport losses (and resistance) at


high-humidity conditions due to water flooding [97]. The MPL has also been shown to improve the start-up performance of a PEMFC effectively by sup-pressing water accumulation at the electrode, which is especially important when dealing with subzero conditions [144].

In addition to improving water management in fuel cells, MPLs have also been studied in order to understand how they can affect other aspects of the fuel cell performance. Mirzazadeh, Saievar-Iranizad, and Nahavandi [145] used a three-electrode electrochemical cell to study the oxygen reduction reaction (ORR) and determine whether this was the main parameter that improved the overall performance of the cell. It was concluded that the use of the MPL improves performance at high current densities, but at low cur-rent densities a DL without MPL shows better performance. Williams, Kunz, and Fenton [146] observed that a DL without MPL had a higher limiting cur-rent density compared to a DL with MPL. However, the MPL was consid-ered still to be critical because it improves the current collection and reduces resistance. Unfortunately, none of these studies provides an explanation or mechanism of how the MPL helps fuel cell performance; rather, they just show the final results.

After testing a number of DLs with and without MPLs, Lin and Nguyen [108] postulated that the MPL seemed to push more liquid water back to the anode through the membrane. Basically, the small hydrophobic pores in the MPL result in low liquid water permeability and reduce the water transport from the CL toward the DL. Therefore, more liquid water accumulated in the CL is forced toward the anode (back diffusion). This reduces the amount of water removed through the cathode DL, decreases the number of blocked pores within the cathode diffusion layer, and improves the overall gas trans-port from the DL toward the active zones.

One drawback of using the MPL is that the water saturation in the CL increases and causes more flooding [108]. Through neutron radiography imaging, Owejan et al. [147] were able to observe that MEAs that had cathode DLs with MPLs had better distribution of water over the active area at high current densities. DLs without MPLs tended to have more water accumulated in one location of the active area (closer to the outlet). One issue with this work was that the water accumulation observed was for the whole MEA and the water quantities were not separated between the anode and cathode sides.

Through the use of a transparent fuel cell, Spernjak et al. [87] were able to visualize the anode FF plate (and DL without MPL) while operating the fuel cell with a cathode that had MPL on the DL. It was observed that liquid water was present in the anode flow field only when an MPL on the cathode side was used. Again, this is an indication that the cathode side creates a pressure barrier that pushes the water toward the anode. These observations agree with the ones presented mathematically by Weber and Newman [148]. Although they did not do any experimental work, their two-phase fuel cell model concluded that the MPL acts as a valve that pushes water away from the DL toward the anode though the membrane.


Pasaogullari, Wang, and Chen [149] also presented a two-phase fuel cell model in which the effect of MPL was studied. They concluded that the water flux toward the anode is enhanced when the following MPL character-istics are used: smaller pore size, lower porosity, larger thickness, and higher hydrophobicity. It is important to note that similar conclusions have been presented in studies related to MPLs used in direct methanol fuel cells (see Section 4.3.3.5 for more information).

One contradictory point regarding how the MPL works is related to the water saturation in the CL of the cathode. Nam and Kaviany [150] stated that using an MPL near the CL means that the water condensed in the DL cannot enter the CL, thus reducing the overall saturation of the active catalyst zones. This idea was also presented by Pasaogullari and Wang [151], who concluded that in the presence of an MPL, the liquid saturation in the CL is reduced substantially. These concepts contradict those presented earlier because it is not clear whether the liquid saturation does in fact increase in the cathode catalyst layer. This may depend directly on the rate at which the water goes back (or is forced) to the anode.

To shed some light on these issues and to be able to have a better under-standing of the water transport when using MPLs, Atiyeh et al. [152] presented an experimental method designed to investigate the net water drag coefficient in order to have a better indication of the amount of water flowing from the cathode to the anode. They observed that the performance of the fuel cell improved when the anode, the cathode, or both had microporous layers.

However, after implementing the water balance measurements, they were not able to observe a significant difference on the net water drag coefficient for a fuel cell with a cathode MPL and an anode without an MPL compared to a cell without any MPLs. It is important to note that they were able to observe that the MPL does in fact improve the fuel cell performance and sta-bility when operating at constant conditions (i.e., the voltage fluctuations are significantly reduced when the cathode DL has an MPL). These results do not correlate with the observations presented earlier; thus, more experimental work is necessary to investigate the process behind how the MPL helps the performance of the fuel cell.

4.3.3.3 Parameters Affecting Microporous Layers

In addition to the way in which the MPL is manufactured, other MPL param-eters directly affect fuel cell performance. These include thickness of the MPL, carbon loading, PTFE content, type of carbon particles, etc. The follow-ing subsection will briefly discuss them.

4.3.3.3.1 Thickness and Carbon Loading

The thickness of the MPL is related to the total distance that both the gases and the water have to flow in order to reach or leave the active catalyst zone. Paganin, Ticianelli, and Gonzalez [153] analyzed the influence of different


DL thicknesses on overall fuel cell performance and concluded that the per-formance of DLs with MPL increased considerably when the MPL thickness was increased from 7.5 to 17.5 m. Their explanation was that very thin MPLs provide poor electrical contact between the CL and the current-collecting FF plate because the electrical resistance is increased due to the roughness of the carbon cloth DL. If the MPL is too thin, the amount of carbon/PTFE is insufficient to provide good electrical contact for the collection of the current generated in the three-phase reaction zone of the catalyst layer.

In addition, the performance between the 17.5 and 25 m MPLs was very sim-ilar, but once the thickness was further increased, the cell’s performance at high current densities dropped. This voltage loss can be attributed to an increase in the diffusion or “travel” distance for the reactants, as well as an increase in electrical resistance. Thus, an adequate diffusion distance, resistance, and mechanical strength are necessary for the MPL (and DL) to be effective.

Another way to control the thickness of the MPL is through the amount or loading of the carbon particles in the layer. Jordan et al. [154] showed similar results to Paganin and colleagues’, demonstrating that out of four loadings used, the lowest and the highest loadings showed the lowest performances—once more indicating that intermediate thicknesses (loadings) are more favor-able. One interesting observation was that the best loading when using air in the fuel cell was slightly lower than for oxygen. This was attributed to the lower partial pressure of oxygen in air resulting in a lower gas flux to the DL. Zhigang and Arthur [57] also investigated MPL carbon loading in detail and showed that the performance of the cell was very similar for loadings between 2 and 4 mg cm–2.

Figure 4.19 shows results by Park, Lee, and Popov [155] for the effect of carbon loading in the MPL on fuel cell performance. It can also be seen that a fuel cell without an MPL performs very poorly compared to the others. Jian-hua et al. [137] recently presented very similar results.

4.3.3.3.2 Hydrophobic Content

In Section 4.3.1 the importance of using the correct amount of hydrophobic agents in the DL was discussed because of how it affects fuel cell performance. Similar conclusions can be drawn for the effect of the PTFE content on MPLs. Paganin et al. [153], Zhigang and Arthur [57], Song et al. [133], and Lufrano et al. [156] demonstrated in their studies that the performance of a fuel cell increased when the PTFE content in the MPL was between 15 and 35 wt%. It was explained that the water removal within the DL and MPL deteriorates with very high PTFE content due to the blockage of pores. On the other hand, with very low PTFE loading, the water cannot be removed efficiently, even though the conductivity of the carbon and CL is improved [156].

Various research groups have been able to demonstrate that the best PTFE loading in the MPL is around 20 wt% when a fuel cell is operated at fairly high humidity conditions [109,136,137,155,157]. It is important to note that in most cases, at low current densities (<0.2 A cm–2), differences due to PTFE


1.2

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FIGURE 4.19(a) Polarization curves of PEMFCs with different carbon loadings in the microporous layer and (b) current densities at 0.45 and 0.60 V as a function of carbon loading in the MPL. Operating conditions: 75 C cell temperature; anode/cathode dew temperatures of 77/75 C; air stoichiom-etry of 2.0. The hydrogen stoichiometry was not specified; serpentine flow fields were both sides. Cathode DL was a 10 wt% PTFE SGL-10CA CFP; anode DL was an E-TEK CFP. PTFE content was not specified. Cathode and anode catalyst layers had loadings of Pt 0.4 mg cm–2; membrane was Nafion 112. (Reprinted from S. Park et al. Journal of Power Sources 163 (2006) 357–363. With permission from Elsevier.)


content are minimal because the amount of water produced is low and the gases can flow through the MPL without major issues.

Antolini, Passos, and Ticianelli [158] used carbon cloth DLs with two MPLs (one in each side) and performed an extensive study of the optimum PTFE content for each MPL (see Section 4.3.3.4 for more detail on DLs with multiple MPLs). For the MPL next to the CL, the PTFE content that resulted with the best performance was 15 wt%; it gave support to the CL and kept the mem-brane hydrated. For the MPL next to the FF channels, the amount of PTFE that gave the best results was 30 wt%; it improved the cell’s water manage-ment without affecting the gas flow through the DL.

4.3.3.3.3 Types of Carbon Particles

Another important parameter that has an effect on the overall performance of the fuel cell is the type of carbon particle used in the MPL. Two of the most common carbon particle types used in this layer are Vulcan XC-72R and acetylene black (AB). Jordan et al. [154,159] were able to show that micropo-rous layers with AB (1.25 mg cm–2 with 10 wt% PTFE) performed better than MPLs with Vulcan XC-72R carbon black. They suggested that the reason for this result was the lower porosity of the acetylene black, which made it bet-ter at removing water from the MEA, thus leading to improved gas flow and diffusion toward the catalyst layer.

Passalacqua et al. [160] studied four different types of carbon/graphite particles for the MPL: Vulcan XC-72, Shawinigan acetylene black (SAB), Mogul L, and Asbury 850 graphite. At a loading of 2.5 mg cm–2 (same loading for all the samples), the carbon blacks performed the best. This was attrib-uted to their homogeneous micropores, which enhance the mass transport characteristics and reduce the possibility of flooding compared to the big flakes in Asbury 850 and the large agglomerates in Mogul L. After further tests, it was concluded that SAB was a better candidate than Vulcan XC-72 because AB had the largest pore volume and small average size; this helped to reduce the mass transport issues and improved the water removal inside the fuel cell.

Antolini et al. [161] also compared SAB and Vulcan XC-72 as possible can-didates in MPLs, but in this case they used carbon cloth DLs with two MPLs. From their results, it was concluded that the Vulcan XC-72 gave slightly higher electrocatalytic activity for the ORR. On the other hand, MPLs near the FF that used SAB had better performance. Thus, it was suggested that for improved fuel cell performance at high pressures (around 3 atm), the ideal cathode MPL compositions would use the Vulcan XC-72 in the MPL next to the CL and SAB in the MPL next to the flow fields.

Other types of carbon blacks, such as Ketjenblack EC-600JD and Denka, have also been used in MPLs for carbon fiber paper DLs [139,162]. Compared to Vulcan XC-72R, both carbon black powders performed better due to their lower resistance and their excellent water transport capabilities, which limit water flooding at higher current densities [162].


Chen, Matsuura, and Hori [163] reported that a modified microporous layer (called a “water management layer”—WML) was implemented to improve the water management in a fuel cell. The novel component of this layer, which had Ketjenblack EC-600JD as the carbon powder, was that its proper-ties (composition, thickness, porosity, and hydrophobicity) were gradually changed along with the gas flow direction, so that effective water manage-ment in the MEA could be achieved. Fuel cell testing demonstrated that a DL with this novel MPL performed better than a DL without an MPL, especially at low humidity conditions.

Wang et al. [164] developed an MPL that consisted of a composite car-bon black mixture that had 90 wt% AB carbon and 10 wt% Black Pearls 2000 (BP) carbon. The AB gave a homogeneous surface rich with pores, which had a great number of hydrophobic macro- (7–70 μm) and mesopo-res (0.05–7 μm) that were beneficial to the gas diffusion. In contrast, BP presented a stronger compactness of carbon and some larger cracks due to its smaller pore size. In addition, BP has more hydrophilic micropores (<0.05 μm) that act as water flow paths, which increase the possibility of flooding. Therefore, combining these two carbon blacks into one layer with PTFE creates an MPL that performs substantially better than other MPLs that have just one carbon black. The weaknesses of each carbon are covered by the other carbon.

Kannan, Menghal, and Barsukov [165] and Kannan and Munukutla [166] used a new form of partially ordered graphitized nanocarbon black, called Pureblack carbon, as part of the MPL for a carbon fiber paper DL. In addition, the nanocarbon black was mixed with nanofibrous carbon (Showa Denko) in order to improve mechanical strength. It was demonstrated that this com-posite MPL with Pureblack and the nanofibrous carbon performed better than an MPL with Vulcan CX-72R under fully humidified and ambient pres-sure conditions, especially at higher current densities.

In another study done by the same research group, three different MPLs were studied: 100% Pureblack, 100% single-walled carbon nanotubes (SWCNTs), and 50% Pureblack and 50% SWCNTs [167]. The MPL with 50% SWCNTs and 50% Pureblack gave the highest fuel cell performance (100% relative humid-ity) due to the SWCNTs forming more favorable pores that improve the gas transport as well as the water management characteristics at high current densities. Owejan et al. [168] also used a graphitized-based carbon black to mitigate corrosion issues in the MPL during fuel cell start-up and shutdown.

Carbon nanofibers (CNFs) and carbon nanotubes (CNTs) have also been considered as raw materials for the microporous layers in order to improve electrical conductivity and permeability. Park et al. [169] tested a number of MPLs with different carbon materials: CNFs, Vulcan XC72, and Black Pearl 2000. It was found that a DL (TGP-H-060 CFP) with an MPL containing a mix-ture of 75 wt% Vulcan and 25 wt% CNF gave the best performance, perform-ing slightly better than a commercial DL (SGL Carbon SGL31BC) with MPL. It was suggested that the reason for these results was that carbon nanofibers


enhanced the gas permeability and electric conductivity while not affecting the water management of the cell.

4.3.3.3.4 Pore Formers in Microporous Layers

The pore distribution in the DL and the MPL is a key parameter that has a direct effect on the mass-transport processes within the MEA and fuel cell. Kong et al. [170] used a pore former, Li2CO3, which was mixed with the carbon black and PTFE prior to coating the DL to modify the pore distribution in the MPL. The DL with MPL (and the pore former) is treated with 1 M H2SO4 and then washed several times with distilled water. In this step, Li2CO3 is removed by dissolving the salt in aqueous sulfuric acid. Various loadings of the pore former were used to investigate the effect of the pore distribution on the overall fuel cell performance. The optimum pore former content was found to be directly related to the operating conditions of the cell because the pore distribution affects the mass transport properties of the MPL. It was observed that an MPL with a pore former performed better than a conven-tional MPL. It was concluded that, to reduce mass-transport issues, it is best to tailor the morphology of the DL with control of the pore-size distribution through the use of pore formers.

A similar study was performed by Jian-hua et al. [137], who used (NH4)2SO4 as the pore former due to its high solubility in water. After fuel cell testing was performed, it was observed that the pore former improved the perfor-mance of the cell at higher current densities (>0.9 A cm–2), indicating that control of the pore distribution in the MPL and DL was critical to enhancing the efficiency of the fuel cell system.

4.3.3.4 Multilayered Microporous Layers

Another important point regarding the fabrication process of MPLs is the fact that, typically, when carbon fiber paper is used as the DL, the MPL is coated just on one surface of the CFP. However, when a carbon cloth is used, a homogeneous water suspension of carbon powder and PTFE is filtered under vacuum onto both faces of the carbon cloth mate-rial to form the MPLs [153,158,161,171], followed by drying and sintering as mentioned earlier. Antolini et al. [161] were able to demonstrate that carbon cloth with double MPLs, for both the anode and the cathode sides, showed better performance than when a CFP was used as the cathode DL with one MPL. At low current densities, the difference between the two DLs was not as obvious, but it became more evident at higher current densities because the limiting current densities for each case are quite different (~1.6 A cm–2 for CFP vs. ~2.7 A cm–2 for CC) (see Figure 4.20 for more details).

Wang et al. [131] used carbon composite MPLs (AB and Black Pearls 2000), similar to those presented in Reference 164, and coated them on both sides of a carbon fiber paper (similar to that for carbon cloths; see Section 4.3.3.1). In this


study, CFP with a single MPL and CFPs with two MPLs were compared while maintaining the same total carbon loading on the MPLs (1.0 mg cm–2). It was shown that a single MPL DL had a lower performance than a dual-MPL DL (with carbon loadings of 0.7 mg cm–2 for the side near the CL and 0.3 mg cm–2 for the side near the flow fields). These results were attributed to the fairly thick and compact single MPL, resulting in limited gas diffusion through this layer. In addition, the absence of the MPL on the side near the FF increased the contact resistance between the DL and the bipolar plate. The DL with a dual MPL pro-vided a change in porosity within the DL, thus creating a gradient that helped the gas and water transport within it.

Wang et al. [140] presented another example of multilayered MPLs that consisted of a basic MPL (Vulcan XC-72 and PTFE) on top of a DL and a mixture of carbon black (Black Pearl 2000) and Nafion sprayed onto the peripheral region of the MPL (i.e., around the active area). This was done on both the anode and cathode sides. The extra MPL, also called a water transfer region (WTR), was used in order for the MEA to be self-humidified. The WTR would receive the excess water at the cathode and then the water

Carbon paper/Vulcan

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FIGURE 4.20Polarization curve for PEMFCs with two different cathode diffusion layers: carbon fiber paper with one MPL and carbon fiber cloth with two MPLs. Operating conditions: cell temperature of 85 C, O2/H2 dew point temperatures of 90/100 C; gas pressures of 2 atm. CFP DL was a TGP-H-090 with 20 wt% PTFE in the MPL. CCs were PWB-3 from Stackpole; cathode CC had 15 wt% PTFE in the MPL near the CL and 30 wt% PTFE in the MPL near the flow field. The anode CC had 15 wt% PTFE in both MPLs; carbon loading on the MPL was not specified. The catalyst Pt loading was 0.4 mg cm–2 and the Nafion loading was 1.1 mg cm–2 for all catalyst layers; the membrane was a Nafion 115. (Modified from E. Antolini et al. Journal of Power Sources 163 (2006) 357–363. With permission from Elsevier.)


would diffuse through the membrane toward the anode, thus humidifying the membrane and reducing the ohmic losses.

Fuel cell testing showed that this WTR improved the fuel cell perfor-mance significantly for operation at a temperature of 40 C (with dry gases). Unfortunately, the fuel cell was not tested at higher temperatures. Another issue is that it is unclear how the fuel cell was modified in order to be able to have this extra layer (on each side) because it is out of the active area and may be affected by the use of seals. In addition, these MPLs represent several addi-tional manufacturing steps that have to be added for each MEA produced.

One final example of multiple layer MPL was presented by Kannan, Cindrella, and Munukutla [172]. A four-layer MPL was fabricated by using nanofibrous carbon, nanochain Pureblack carbon, PTFE, and a hydrophilic inorganic oxide (fumed silica). The first three layers were made out of mixtures of the nanofibrous carbon, Pureblack, carbon, and PTFE. Each of these three layers had different quantities from the three particles used. The fourth layer consisted of Pureblack carbon, PTFE, and fumed silica to retain moisture con-tent to keep the membrane humidified. Therefore, by using these four layers, a porosity gradient was created that significantly improved the gas diffusion through the MEA. In addition, a fuel cell with this novel MPL showed little performance differences when operated at various humidity conditions.

4.3.3.5 Microporous Layers in Direct Liquid Fuel Cells

In direct liquid fuel cells, the use of MPLs is also very popular and most of the details explained earlier also apply to the liquid fuel cells. However, some of the parameters differ from those in PEM fuel cells because there are other mass transfer-based issues in DLFCs, especially on the anode side related to methanol crossover and CO2 production.

For this type of fuel cell, a number of reports studying anode MPLs have been published. Neergat and Shukla [124] used a hydrophobic MPL on the cathode (carbon black and PTFE) and a hydrophilic MPL on the anode (carbon black and Nafion) (see Section 4.3.2). Different types of carbon particles were used (Vulcan XC-72, acetylene black, and Ketjenblack) and it was concluded that Ketjenblack was the carbon that showed the best performance when it was used on both the anode and cathode MPLs with 10 wt% Nafion and 10 wt% PTFE, respectively. A similar design was also used by Ren et al. [173] in a passive DMFC. Improvement of the DMFC performance by using a hydro-philic MPL, as discussed previously, was also demonstrated by Lindermeir et al. [125]. They compared both hydrophilic and hydrophobic MPLs for the anode DL, and it was observed that the former improves the mass transport of the MEA.

Shao et al. [126] studied the effect of Nafion in the anode MPLs and deter-mined that the ideal content was 10 wt%. It was observed that this ionomer reduces the number of pores in the DL and MPL that are in the diameter range of 10–100 nm. Therefore, higher loadings of Nafion would result in


lower porosity and in methanol transport issues. Similar conclusions and observations were presented by Zhang et al. [127], who designed a trans-parent cell in order to observe the CO2 gas bubbles inside the anode flow field plates. It was observed that with hydrophilic MPLs on the anode side, uniform CO2 gas bubbles were formed and the methanol solution flowed through the anode DL to the CL uniformly and easily. On the other hand, with hydrophobic MPLs, larger, nonuniform CO2 bubbles were formed, thus decreasing the performance of the cell.

Mao et al. [174] recently presented research in which Nafion ionomer parti-cles were used as hyperdispersant agents in the MPL of a cathode DL. It was shown that this ionomer helps to decrease the particle size of the PTFE in the MPL. Thus, increasing the Nafion particle content gradually decreased the PTFE size and decreased the hydrophobicity in the layer. In fuel cell testing, an MPL having 1 wt% ionomer showed the best performance; it improved the gas permeability and electronic conductivity.

Park et al. [175] investigated different raw materials such as acetylene black, Vulcan XC-72R, and RuO2 in MPLs for both anode and cathode DLs. In this study, PTFE was used in both sides. After fuel cell testing, it was found that as a raw material for the MPL in the anode DL, RuO2 improves slightly the cell’s performance because it reduces the losses related to metha-nol diffusion. For the cathode MPL, RuO2 improves the cell’s performance significantly because it reduces the water flooding, especially when the cell is operating at low temperatures and with air as the oxidant gas. In addition, a mixture of AB and RuO2 was used in the MPLs for both DLs and a signifi-cant improvement in the cell’s performance, compared to MPLs with just AB, was observed.

The optimum amount of PTFE in the anode MPL depends on the operating conditions and design of the DMFC. Dohle et al. [176] used a 500 W DMFC stack and observed that anode MPLs with 13 wt% PTFE had the best perfor-mance. Peled et al. [177] designed anode MPLs that had around 20–40 wt% PTFE and determined that the layers with lowest PTFE content performed the best.

In contrast with these two studies, Imazato et al. [178] took into account a wider range of PTFE content in the MPL and determined how this affected the overall performance of both anode and cathode MPLs. Surprisingly, it was found that the best performance for both MPLs was achieved when the PTFE content was between 75 and 85 wt%. This PTFE content is signifi-cantly higher than that tested by any other research group. Even though the cell’s resistance increased with the increased PTFE content, the compacted surfaces of the MPLs formed large cracks; thus providing more paths and channels and increasing the permeability of the MPL. Further research is required to confirm these results and determine the effect of high PTFE con-tent on MPL performance.

To reduce methanol crossover and improve water back diffusion through the membrane in passive DMFCs, Kim et al. [179] designed an MEA with


multilayer MPLs for both the anode and cathode diffusion layers. For the anode, a CFP was coated with a mixture of nanosilica particles and PVDF. Then a thin CFP (TGP-H-030) was put on top of the MPL and was used as the catalyst backing layer for the CL that was sprayed on it. On the cathode side, a CFP was coated with carbon black and PTFE to form an MPL on top of which a layer of ordered mesoporous silica (OMS) and PVDF was sprayed. A catalyst backing layer (TGP-H-030 CFP) was then located on top of the multi-layer MPL. This novel MEA showed that the water back diffusion increased significantly compared to an MEA with single MPLs on both the anode and cathode. In fact, in the fuel cell testing, it was observed that the novel MEA was able to achieve a large power density. This was considered to be a result of the lower methanol crossover and less water flooding.

Peled et al. [177] also designed a novel MEA in order to improve the water back diffusion from the cathode to the anode side. They used a liquid-water barrier layer (LWBL), which consisted of a paste, made out of PTFE and car-bon black particles, that was inserted in the pores of the CFP to form a layer inside the paper. Up to seven layers were necessary in order to achieve a uniform layer of 20–50 μm in thickness. Testing showed that the LWBL on the cathode DL creates a hydraulic pressure that forces (or pushes) the water back from the cathode toward the anode, thus improving the cell’s water management at different operating conditions.

Xu and Zhao [180] investigated in detail how a hydrophobic MPL on the cathode DL reduces water crossover through the membrane and thus improves significantly the gas transport and the limiting current densities. In an air-breathing DMFC, a similar study was performed by Song, Lee, and Kim [181] that also demonstrated that a cathode MPL helps to reduce the water transport coefficient and increase the methanol utilization. The use of a hydrophobic MPL on the cathode DL is considered to create a hydraulic pressure buildup in the MPL, which results in a large pressure differential across the membrane; hence, water backflow from the cathode to the anode is achieved [182]. Thus, the MPL significantly reduces water crossover through the membrane and improves methanol consumption.

4.4 Properties and Measurements for Diffusion Layers

To design the optimal diffusion layer for a specific fuel cell system, it is important to be able to measure and understand all the parameters and char-acteristics that have a direct influence on the performance of the diffusion layers. This section will discuss in detail some of the most important prop-erties that affect the diffusion layers, such as thickness, hydrophobicity and hydrophilicity, porosity and permeability (for both gas and liquids), electri-cal and thermal conductivity, mechanical properties, durability, and flow


field interaction. This section will also discuss the most common methods that have been used to measure and investigate these properties.

4.4.1 Thickness

The thickness of the diffusion layer is directly related to the mass transport of gases and liquid within the material because it determines the length of the flow path. The electrical conductivity and resistance of the DL are also affected substantially by the thickness of the material. Therefore, to choose an optimal DL, there has to be a compromise between the thickness of the material and the properties mentioned before.

The thickness of the DL material without any compression pressure can be measured by using a micrometer. However, depending on the material, the thickness can change substantially when there is compression. Thus, another method is to exert a specific compression pressure (which must be reported) on the sample while measuring the thickness [9]. In addition, the thickness measurements should be performed at various points over the sample and at multiple times, in order to be as precise as possible. Chang et al. [183] designed a test stand that can measure the thickness of the sample material with a thickness gage at different compression pressures.

Schmitz et al. [184] tested various carbon fiber papers with different thick-nesses as cathode DLs in PEM fuel cells. It was observed that the cell resistance dropped when the thickness of the DL increased; thus, thicker materials are desired in order to improve the electrical conductivity. It was also mentioned that the optimal thickness for the DL is usually between the thinnest and the thickest materials because the two extremes give the lowest performance. In fact, in thin DLs, the water produced can fill pores within the material, resulting in flooding and the blockage of available flow paths for the oxygen. Similarly, Lin and Nguyen [108] concluded that thinner DLs (without MPLs) were more prone to liquid water accumulation than thicker ones.

Prasanna et al. [185] were also able to observe an optimum thickness of DLs for fuel cells experimentally. They demonstrated that the thicker DLs experi-ence severe flooding at intermediate current densities (i.e., ohmic region) due to low gas permeation and to possible condensation of water in the pores as the thickness of the DL increases. On the other hand, as the thickness of the DL decreases, the mass transport losses, contact resistance, and mechani-cal weakness increase significantly [113,185]. Through the use of mathemati-cal modeling, different research groups have reported similar conclusions regarding the effect of DL thickness on fuel cell performance [186–189].

Cha et al. [190] studied the relationship between DL thickness and the dimensions of the FF channels in micro fuel cells. They concluded that, with micro fuel cells, matching the thickness of the DL to the dimensions of the FF channels improves the overall performance and prevents DL damage.

In DMFCs, Xu et al. [119] tested various carbon fiber papers with different thicknesses (TGP-H-030, TGP-H-060, TGP-H-090, and TGP-H-120) as anode


diffusion layers. It was shown that when the thinnest CFP was used, the overall performance of the cell was the lowest. When the thickest CFP was used as the DL, the performance was reduced, especially at high current densities, thus affecting the limiting current density. These workers claim that this was due to the increased through-plane mass transfer resistance. Once again, it was concluded that the optimum CFP DL had an intermediate thickness (i.e., TGP-H-090).

The porosity of a DL material is also directly related to the thickness of the material. The optimal thickness of the DL also depends on the material porosity. If the porosity is low, then the DL should be quite thin in order for the cell to perform well. If the porosity is high, then a thicker DL should be used [191]. More information about the porosity of the DL is discussed in Section 4.4.3.1.

One critical parameter that affects the thickness of the diffusion layer is the compression force used in the fuel cell in order to avoid any gas leaks and to assure good contact between all the components. However, this compres-sive force can deform the diffusion layer and hence affect the performance of the cell. More information regarding how the compression forces affect the diffusion layer is discussed in Section 4.4.5. Ideally, the material used as the DL should be able to resist this compression force or pressure without affect-ing most of its parameters. Figure 4.21 shows a schematic of the cell voltage (performance) at a given current density, resistance, and DL porosity as a function of the cell’s compression.

In conclusion, at an intermediate optimum thickness, a diffusion layer allows for (1) gas diffusion toward the CL, (2) liquid transport from the CL toward the flow field channels, (3) good contact with both the bipolar plate

Compression

Porosity

Cel

l P

erfo

rman

ce

Res

ista

nce

, Po

rosi

ty

Resistance

Cell

Performance

FIGURE 4.21Schematic of the cell performance, resistance, and DL porosity as a function of the fuel cell’s compression.


and the active zone, and (4) acceptable mechanical support and strength. Thorough studies should be performed in order to find the ideal DL mate-rial, taking into account the specifications and operating conditions of the corresponding fuel cell system.

4.4.2 Hydrophobicity and Hydrophilicity

One of the most common ways to characterize the hydrophobicity (or hydrophilicity) of a material is through measurement of the contact angle, which is the angle between the liquid-gas interface and the solid surface measured at the triple point at which all three phases interconnect. The two most popular techniques to measure contact angles for diffusion lay-ers are the sessile drop method and the capillary rise method (or Wihelmy method) [9,192].

4.4.2.1 Sessile Drop Method

The sessile drop method has been widely used in the past as a way to deter-mine the contact angle of a liquid on a solid surface. In this technique, a drop of liquid is deposited on the surface of the material; once the droplet has settled (or has become sessile), the contact angle between the liquid and the solid is measured. This measurement can be done by using a goniometer or by using a digital system that takes an image of the droplet and calculates the angle, q, of a tangent line at the triple-phase point (see Figure 4.22). If the value of this angle is >90 , then the material is considered hydrophobic; if the value is <90 , then it is considered hydrophilic.

It is important to note that, with materials used as diffusion layers in fuel cells, the size of the water droplet should be fairly small (<1 mm in diameter) to avoid the size and weight of the droplet affecting the mea-surements [9]. It is recommended that various contact angle measurements be performed at different locations on the same sample material to ensure homogeneity. These measurements can also be performed in an environ-mental chamber in order to imitate the desired temperature and operating conditions. However, the contact angle has to be measured before the liquid evaporates. Figure 4.22 shows an example of the contact angle of water on a Grafcell flexible graphite film indicating that the material is hydrophobic. As expected, the contact angle is >90 because this material is coated with a hydrophobic layer.

For the DMFC, Zhang et al. [127] used the sessile drop method to study the wettabilities of liquid methanol solutions on the surface of the anode DLs and MPLs. They were able to observe that the contact angles of the materials were the smallest with low PTFE content. In addition, the effect of Nafion ionomer content on the MPL (to increase hydrophilicity; see Section 4.3.2) was also shown through the contact angle measurements (i.e., smaller con-tact angles compared to MPLs with PTFE).


4.4.2.2 Wilhelmy Method: Capillary Rise

It is important to note that a number of different techniques are based on the Wilhelmy method. Here, we will describe only the capillary rise method (based on the Wilhelmy gravimetric plate technique) presented by Lim and Wang [101] and Wang [192]. For more detail on the other Wilhelmy methods, please refer to Mathias et al. [9].

The capillary rise technique is considered to be very useful for DL materi-als, especially if the angle is less than 90 and/or for measurements that are taken under different temperature conditions [192]. In this method, a sample material is immersed in a container filled with water and the meniscus height is measured with a microscope (see Figure 4.23). Contact angles between the water and DL are calculated using Equation 4.1 and by measuring the

Solid Material

(a)

(b)

FIGURE 4.22Contact angle: (a) schematic of the measurement of a contact angle; (b) example of the contact angle for a Grafcell flexible graphite film fabricated by Graftech Inc.


meniscus heights from the initial water level (prior to immersing the mate-rial) to an interfacial line between the air, liquid, and solid material:

sin .1

2

2gh

(4.1)

Specimen

Microscope withCCD Camera

Vertical PositioningTool

Light Source

Temperature

Control

Vertical Translator X-Y-Z

Translator

PC

(a)

(b)

FIGURE 4.23(a) Schematic of an experimental setup for contact angle measurement using the capillary rise method. (b) Example of a meniscus line on the surface of a wet-proofed CFP in water pool. (Modified from C. Lim and C. Y. Wang. Electrochimica Acta 49 (2004) 4149–4156. With permis-sion from Elsevier.)


whereq is the contact angleΔr is the difference between the water and air densitiesg is the gravitational constanth is the meniscus heights is the liquid-gas surface tension of water

If the meniscus height is negative (i.e., lower than the water level in the container), then the surface of the material is hydrophobic; if the height is positive (i.e., higher than the water level in the container), then the surface of the material is hydrophilic [101]. This method is ideal in determining the contact angle of the DL material at higher temperatures because it is easier to keep the temperature of the water in the container stable. In the sessile drop method, the temperature of the water changes rapidly once it is in contact with the DL material; therefore, it is difficult to measure at high temperatures.

4.4.2.3 Contact Angle of Moving Droplet

The contact angle of a liquid in a DL surface changes depending on the rate at which the gases are flowing through the flow field channels. Kumbur et al. [193] studied how the PTFE content in different DLs changes the difference between the advancing and receding angle (also called contact angle hyster-esis) of a water droplet placed on top of a DL that is being pushed by air. They concluded that high PTFE content increased the deformation of the droplets, resulting in higher contact angle hysteresis. They explained that this was due to the decreased surface interfacial tension of the water molecules on the carbon fibers. These results are important when the optimal PTFE content is determined and when the two-phase flow transport inside the flow field chan-nels and on the surface of the DLs is studied. For more information regarding this, please refer to Section 4.4.3.

4.4.2.4 Internal Contact Angle

Both the sessile drop and Wilhelmy methods are important when determin-ing the hydrophobicity of the surface of a material because they can measure the external contact angle. However, these techniques do not give any infor-mation regarding the hydrophobic or hydrophilic properties of the internal fibers and pores inside a CFP or CC. Determining the internal contact angle of a DL material is critical because it can provide information regarding the water transport mechanism within a specific material. Due to the complex internal structures of CFPs and CCs, measuring these internal contact angles has been quite challenging.

Gurau et al. [194] proposed a method to estimate the internal contact angle to water by combining the Washburn technique with the Owens–Wendt


two-parameter theory and the method of least squares. They tested materi-als with 30 and 70% PTFE and were able to observe that the samples with higher PTFE content have the highest contact angles with water and the low-est surface tensions. They also claimed that their method could be used with catalyst and microporous layers.

4.4.3 Transport

Issues with mass transport resistance, especially at higher current densities, represent an important hurdle that fuel cells need to overcome to achieve the required efficiencies and power densities that different applications require. Diffusion layers represent one of the major fuel cell components that have a direct impact on these mass transport issues; thus, optimization of the DLs is required through the use of different experimental and characterization techniques.

In this section, we will briefly discuss different testing techniques that are widely used to measure most of the important mass transport properties of the diffusion layers. It is important to note that these techniques can also be used with MPLs. The first subsection will explain methods that deal with properties that affect both gas and liquid mass transport, and the other two subsections will discuss only techniques that measure gas and liquid trans-port properties, respectively.

4.4.3.1 Characterization of General Transport Properties

As stated earlier, CFP and CC are the most common materials used in the PEM and direct liquid fuel cell; due to their nature, it is critical to under-stand how their porosity, pore size distribution, and capillary flow (and pressures) affect the cell’s overall performance. In addition to these proper-ties, pressure drop measurements between the inlet and outlet streams of fuel cells are widely used as an indication of the liquid and gas transport within different diffusion layers. In this section, we will discuss the main methods used to measure and determine these properties that play such an important role in the improvement of both gas and liquid transport mechanisms.

4.4.3.1.1 Porosity

Porosity determination is one of the most common characterization steps performed in the early stages of choosing the optimal DL for a specific fuel cell application. Porosity represents the ratio of pore (void) volume to total bulk volume of the material. This property can be calculated from mea-suring the bulk density of the material, which is calculated from the areal weight and the thickness [9]. Then the porosity is determined by using the bulk density and the solid-phase density of the material, which can be mea-sured experimentally with a gas pycnometer. The porosity also depends on


the uncompressed and compressed thicknesses of the material. The follow-ing equation shows the relationship of the average porosity, e, to density and thickness:

1 bulk

s

ott

(4.2)

whererbulk is the bulk densityrs is the solid phase densityto and t are the uncompressed and compressed thicknesses, respectively

However, this porosity takes into account all the open pores—even those that are not connected between each other, which are useless in fuel cell operation. Therefore, the effective porosity, which counts only the intercon-nected pores, is more critical when determining the optimal diffusion layer in a fuel cell. This porosity can be determined by using volume filtration techniques. For example, a porous sample is immersed in a liquid that does not enter inside the pores (e.g., mercury at low pressures) and then the total volume of the material can be determined. Next, the specimen is put inside a container of known volume that contains an inert gas, and the changed pres-sure is recorded. After this, a second evacuated chamber of known volume is connected to the system, and the new pressure is recorded. With these pres-sures and the ideal gas law, the volume of open pores and thus the effective porosity can be determined [195].

Other methods have also been used to measure the hydrophilic and hydro-phobic porosity of diffusion layers in fuel cells. Dohle et al. [176] determined the fraction of hydrophobic and hydrophilic pores by first filling a sample material with water, which fills the hydrophilic pores, followed by weigh-ing and drying of the sample. After this, the material is filled with decane (highly wetting liquid), which fills both the hydrophilic and hydrophobic pores, and then it is weighed. With the two measured weights, the fractions of hydrophilic and hydrophobic pores in the sample can be determined. In another method, Wang et al. [164] designed an apparatus in which the speci-men was placed horizontally on a flat surface at room temperature. Then water vapor carried by air was passed through the material and, with time, liquid water accumulated in the hydrophilic pores of the sample. The differ-ence in weight of the material before and after the test was used to determine the hydrophilic porosity of the DL.

4.4.3.1.2 Pore Size Distribution—Mercury Porosimetry

One of the most popular methods to measure the pore size distribution in diffusion layers is mercury intrusion porosimetry (MIP); this technique is


able to span the measurement of pore sizes ranging from 1 nm up to 1 mm [196]. In this method, a nonwetting liquid (e.g., mercury) fills the chamber where the sample material is under vacuum and the volume is recorded. At this point, the liquid has not yet filled the pores of the sample due to its high surface tension. The pressure applied to the nonwetting liquid is gradually increased, resulting in a change in volume of the liquid in the pores. Both the pressure and the change in volume are accurately measured. Combining these data with data about surface tension and contact angle of the liquid allows pore diameter, pore volume, pore volume distribution, and pore sur-face area to be computed and determined [9,197].

One issue with this method is that, for larger pores shielded by smaller ones, the corresponding pressure necessary for the liquid to intrude them corresponds to the entry pressure for the smaller pores; thus, the volume for the larger pores is incorrectly attributed to smaller ones. In addition, the assumption that the contact angle of the nonwetting liquid is the same on all solid surfaces is not completely correct because diffusion layers with different treatments have pores with different wetting properties. For example, two pores of the same size may have different PTFE content and the entry pres-sure necessary for the liquid to penetrate them will be different for each pore, thus affecting the overall results and the calculated capillary pressures [196].

Due to the issues mentioned here and the fact that, in fuel cells the wet-ting property may change within the DL, it is necessary to have another technique that can differentiate between hydrophobic and hydrophilic pores within a sample. This information would allow researchers to improve the design of diffusion materials.

4.4.3.1.3 Pore Size Distribution—Method of Standard Porosimetry

The method of standard porosimetry (MSP) is another method used to study pore distribution in a porous material. With this technique, any wetting fluid can be used as the working liquid. Therefore, water can be used as the liquid in order to determine the hydrophilic pores. For the distribution of the overall porosity, a liquid with strong wetting proper-ties, such as octane, can be used because it can completely wet graphite, carbon, and PTFE. Another advantage of this method is that the contact angle of the liquid with the solid surface is not necessary to obtain the necessary data [196].

This method is based on the principle of capillary equilibrium, providing a nondestructive technique in which the total porosity and capillary pressure-saturation data are measured [198,199]. The idea is that when two partially saturated porous materials are in contact, the system moves toward an equi-librium state in which the capillary pressure of a liquid in one of the samples is the same for the other material.

Figure 4.24 shows a schematic of the MSP method [199]. The sample DL material is placed between two standard materials that have known capillary


pressure-saturation curves and data and similar porosity to the DL speci-men. One end of the stack assembly, from which the liquid evaporates, is open and the other end is closed. It is important to have the materials fully saturated prior to the test because this method does not work if the wetting phase only moderately wets the porous material [196].

Once the stack has been assembled, a small portion of liquid is evaporated slowly by flowing a dry, inert gas while keeping the three materials in close contact so that capillary pressure equilibrium is maintained. At each point, the mass of each sample is measured to gather the wetting phase saturation, thus allowing determination of the corresponding capillary pressure of the standard material from its known parameters and data. At equilibrium, this pressure is equal to the capillary pressure of the DL sample. This procedure is repeated in short steps until all the liquid is evaporated in the whole stack.

As stated previously, the capillary pressure data, the overall pore distribu-tion (when octane is the working fluid), and the hydrophilic pore distribution (when water is the working fluid) can be obtained through this technique. In addition, these measurements can be used with different compression pres-sures of the sample DL and with a wide range of temperatures inside the system [200,201]. For more information regarding this technique, please refer to the paper by Volfkovich et al. [198].

Temperature

Controller

Standard Porous

Material

Standard

Porous Material

DM SamplePDM = Psample

PDM = Psample

Fluid

Pool

Capillary

Pressure Balance

Compression

Controller

Known

Standard

Porous Material

PC

S

FIGURE 4.24Schematic of the MSP method. (Reprinted from E. C. Kumbur et al. Journal of the Electrochemical Society 154 (2007) B1295–B1304. With permission from ECS—the Electrochemical Society.)


4.4.3.1.4 Pore Size Distribution—Capillary Flow Porometry

Another method to measure pore size distribution is capillary flow porom-etry [202,203], in which a sample material is soaked with a low surface ten-sion liquid that fills all its pores. Then, gas pressure is applied on one side of the sample in order to force the liquid out of the pores. At low pressures, the flow rate is close to zero; however, as the pressure increases, the flow rate also increases and the amount of liquid inside the pores decreases. Thus, the flow rate is determined as a function of pressure and is then used to calculate the desired pore characteristics, such as pore size distribution, largest pore diameter, and mean flow pore diameter.

4.4.3.1.5 Capillary Pressure

The determination of the capillary pressure of a diffusion layer is critical, not only to have a better understanding of the mass transport mechanisms inside DLs but also to improve their design. In addition, the accuracy of mathematical models can be increased with the use of experimental data obtained through reliable techniques. Both Gostick et al. [196] and Kumbur et al. [199] described and used the MSP method in detail to determine the capillary pressures of different carbon fiber paper and carbon cloth DLs as a function of the nonwetting phase saturation. Please refer to the previous subsection and these publications for more information regarding how the capillary pressures were determined.

Fairweather et al. [204] developed a microfluidic device and method to mea-sure the capillary pressure as a function of the liquid water saturation for porous media with heterogeneous wetting properties during liquid and gas intrusions. In addition to being able to produce plots of capillary pressure as a function of liquid water saturation, their technique also allowed them to inves-tigate both hydrophilic and hydrophobic pore volumes. This method is still in its early stages because the compression pressure and the temperatures were not controlled; however, it can become a potential characterization technique that would permit further understanding of mass transport within the DL.

Nguyen et al. [205] designed a volume displacement technique that was used to measure the capillary pressures for both hydrophobic and hydro-philic materials. One requirement for this method is that the sample material must have enough pore volume to be able to measure the respective displaced volume. Basically, while the sample is filled with water and then drained, the volume of water displaced is recorded. In order for the water to be drained from the material, it is vital to keep the liquid pressure higher than the gas pressure (i.e., pressure difference is key). Once the sample is saturated, the liquid pressure can be reduced slightly in order for the water to drain. From these tests, plots of capillary pressure versus water saturation correspond-ing to both imbibitions and drainages can be determined. A similar method was presented by Koido, Furusawa, and Moriyama [206], except they studied only the liquid water imbibition with different diffusion layers.


4.4.3.1.6 Pressure Drop Measurements

When water is accumulated in the DLs (and FF channels), the pressure drop of the fuel cell increases significantly. Therefore, techniques in which the pressure drop is measured carefully have been implemented in order to correlate such measurements with the water flooding inside dif-fusion layers. These methods can be used generally with any FF design. However, due to the nature of its design, the interdigitated flow field can be used to obtain even more information about the characteristics of both the gas and liquid transport in different diffusion layers. With the inter-digitated FF, gas is forced to flow through the DL between inlet and outlet channels [207].

Thus, a number of research groups have implemented cells with this FF design in order to monitor the pressure drop and use these data as a diagnos-tic response to monitor the extent of DL flooding. He, Lin, and Nguyen [208] used a cathode interdigitated flow field plate and monitored the pressure drop between the inlet and outlet to obtain real-time flooding information in the DL depending on the type of materials used. They observed that this method worked well due to the dependency of the gas permeability of the porous diffusion layers on the liquid-water content. Through this method, it was also shown that poor water removal results in liquid-water buildup in the cathode and that increased air flow rate or cell temperature can improve this condition.

Ito et al. [209] also used an interdigitated FF to observe the flooding in DLs implicitly. In addition, the estimation of the water saturation of different DL materials was obtained based on the differential pressure through the interdigitated flow field. This study concluded that the carbon cloth DL was able to remove more water than the carbon fiber paper DL. It was also shown that the water saturation in the DLs increased with the increase of the cur-rent density, gas utilization ratio, and humidification temperature. The over-all permeabilities for a CFP and a CC were determined with three different gases (air, nitrogen, and hydrogen) supplied to the cell (without running electrochemically) at various flow rates while the pressure drop was mea-sured. With these data, Darcy’s law is implemented to calculate the corre-sponding overall permeability for each material. Yamada et al. [210] used a similar technique for the measurement of flooding in the DLs of fuel cells.

4.4.3.2 Characterization of Gas Transport Properties

One of the main parameters that would improve the overall performance of a fuel cell is better mass transport of reactants through the diffusion layer toward the active catalyst zones. In order to quantify and character-ize how well the gas mass transport is in a specific DL material and design, it is important to measure the in-plane and through-plane permeabilities. Most of the published permeability results report the viscous permeability


coefficients based on Darcy’s law, which describes the single-phase flow through a porous medium for low fluid velocities [211]:

u

kP gg

g

gg g( )

(4.3)

where

ug is the superficial gas velocitykg is the gas-phase permeabilitymg is the gas-phase viscosity

Pg is the gas-phase pressure gradientrg is the gas-phase densityg is the gravitational constant

Neglecting gravity and rearranging Equation 4.3 yields

P

kug

g

gg .

(4.4)

In other cases, researchers assume that the inertial resistance to flow in DLs adjacent to conventional flow fields is negligible and they tend to lump both viscous and inertial coefficients together. This may not be a correct assumption, especially when dealing with flow fields like the interdigitated design [129,212], in which higher velocities are experienced in the pores of the DL. The Forchheimer equation is an extension of Darcy’s law and takes into account the inertial resistance at high velocities [211,213]:

P

ku u ug

gg g g g g| |

(4.5)

where the first term on the right-hand side is Darcy’s law and bg is the gas-phase inertial coefficient. Most of the commercially available apparatuses do not measure both coefficients [129]. Thus, in the past few years, research groups have attempted to develop methods in which both permeability coef-ficients are determined separately.

4.4.3.2.1 In-Plane Permeability

Although in-plane permeability is critical in order to understand in detail the transport mechanisms of fluids inside diffusion layers, it has not been as commonly used (and measured) as through-plane permeability. The fol-lowing are a few examples of how in-plane permeability can be determined


for porous materials used as diffusion layers in fuel cells. It is important to note that some engineered DLs (see Section 4.2.4) may not have any in-plane permeabilities because they are based on nonporous materials with through pores or orifices. In addition, because of their structure, carbon cloths may also have limited in-plane permeability.

Gostick et al. [212] designed an apparatus in which the in-plane permeabil-ity was measured as a function of the DL thickness with different compres-sion pressures (see Figure 4.25). The DL specimen was compressed between plates, which had spacers of different thicknesses in order to control the total thickness in each test. The sample was located between two grooves or chan-nels, one of which corresponded to the inlet of the air and the other to the air outlet. Therefore, the air had to flow in plane through the sample in order to

(a)

(c)

(b)

Flow ControlValve

VolumetricFlow Meter

Air InInlet H

eader

Outlet H

eader

ΔP

Sample

Air Out

Face Plate

Rubber Face Seal Gasket

Bolt Hole

Spacer

FIGURE 4.25Experimental apparatus used to measure in-plane permeability: (a) assembled view; (b) exploded view; (c) sectioned view. (Reprinted from J. T. Gostick et al. Journal of Power Sources 162 (2006) 228–238. With permission from Elsevier.)


access the outlet channel. The pressure drop between the inlet and exit was measured as well as the outlet flow rate.

For each thickness, at least 10 different flow rate measurements were obtained in order to cover the range of flow rates that a DL experiences dur-ing normal fuel cell operation. To obtain the corresponding permeability, the pressure drop results were plotted as a function of the mass flow rate. After this, the Forchheimer equation was fitted to the plotted data to determine the viscous and inertial permeabilities. As expected, the in-plane permeabilities of each sample DL material decreased when the compression pressure was increased. It is also important to mention that these tests were performed in two perpendicular directions for each sample in order to determine whether any anisotropy existed due to fiber orientation.

Feser et al. [214] used a radial flow apparatus to determine the viscous in-plane permeability of different DLs at various levels of compression (see Figure 4.26). A stack of round-shaped samples, with each layer of material separated with a brass shim, was placed inside two plates. Thicker shim stock was also used in order to control the total thickness of the stack of samples. Compressed air entered the apparatus through the upper plate and was forced through the samples in the in-plane direction. After this, the air left the system and flowed through a pressure gage and a rotameter in order to measure the pressure drop and the air flow rate. The whole apparatus was compressed using a hydraulic press; for each compression pressure, 10 dif-ferent flow rates were used.

Flow Inlet

Brass

Shim

GDL Sample Layer Separators

O-Ring Seal

Flowmeter

Flow OutletHydraulic Press

Pressure GaugePressureGauge

FIGURE 4.26Radial flow permeability testing apparatus. (Reprinted from J. P. Feser et al. Journal of Power Sources 162 (2006) 1226–1231. With permission from Elsevier.)


With these data and Darcy’s law, the in-plane viscous permeabilities were determined. Only the viscous permeability coefficient was determined because it was claimed that the inertial component was undetectable within the error limits of measurement for these tests. It is important to mention that this technique could also be used to measure the permeability of diffu-sion layers with different fluids, such as liquid water.

Gurau et al. [129] presented another apparatus used to measure the in-plane viscous and inertial permeability coefficients. In their method, an annular DL sample was placed between an upper and lower fixture. The gas entered the upper fixture and was then forced through the DL into the outlet ports (open to the atmosphere). A strain sensor was located in the upper fixture in order to determine the thickness of the DL (i.e., deformation) because the whole assembly was compressed to a determined pressure. In this method, the flow rate, temperatures in both fixtures, and pressures were monitored in each test. Once the data were collected, the in-plane permeability was deter-mined from the Forchheimer equation by application of the least squares fit analysis method.

Chang et al. [183] presented a similar design in which two discs (with ori-fices in the middle) were used to compress the sample material. Pressurized air (without any moisture) was then passed through the orifices of the discs toward the sample DL, which then flowed peripherally to the atmosphere. The two discs were compressed in order to see how the permeability of the DL changed as a function of the clamping pressure. The permeability coefficient was solved using Darcy’s law; thus, only the viscous in-plane permeability was taken into account. Other, similar techniques can be found in the literature [215–217].

4.4.3.2.2 Through-Plane Permeability

Through-plane permeability is usually one of the most common parame-ters given by manufacturers for carbon fiber papers and carbon cloths, even though it is often not specified as through-plane permeability. It is impor-tant to note that commercial instruments, such as permeameters and Gurley method instruments, are used in the fuel cell industry to measure this per-meability [197,218].

In order to determine the viscous and inert through-plane gas permeabilities of diffusion layers at varied compression pressures, Gostick et al. [212] designed a simple method in which a circular specimen was sandwiched between two plates that have orifices in the middle, aligned with the location of the material. Pressurized air entered the upper plate, flowed through the DL, and exited the lower plate. The pressure drop between the inlet and the outlet was recorded for at least ten different flow rates for each sample. The inert and viscous per-meabilities were then determined by fitting the Forchheimer equation to the pressure drop versus flow rate data as explained earlier.

Gurau et al. [129] designed a device in which the sample DLs were placed between two main components, each with a cylindrical compartment and an annular compartment (see Figure 4.27). The compartments in the upper


component were connected with a number of ports in order to minimize the pressure drop when the gas flowed through this component. The gas entered the upper component and was separated between the compartments; then, it was forced to flow through the DL sample into the compartments of the lower part. The gas in the lower annular compartment exited through a back-pressure regulator and then was released to the atmosphere. The flow rate of the cylindrical compartments was measured in the downstream line (in the lower component), and the pressure drop between the upper and lower parts of the cylindrical compartments was also recorded.

The pressure difference between the lower cylindrical compartment and the lower annular compartment was measured. To make sure that the flow through the sample material had only a through-plane component (z-direc-tion), the back-pressure valve was adjusted until the lower annular pressure was equal to the pressure experienced by the lower cylindrical compartment. The temperatures of the inlet and outlet lines, as well as the flow rate of the gas leaving the lower cylindrical compartment, were monitored. A wide

CylindricalCompartments

AnnularCompartments

Ports

Upstream Fixture

T

T1

GDLSample

ΔP

P

D

Q

Flow Meter

Back PressureRegulator

Downstream FixtureP2

P3 P2 P3

P1 P1 P1

T

T2 P0

P0

P1 – P2

ΔP P3 – P2

FIGURE 4.27Experimental apparatus to measure through-plane permeability. (Reprinted from V. Gurau et al. Journal of Power Sources 165 (2007) 793–802. With permission from Elsevier.)


range of flow rates was used for each sample. Once the data were obtained, the inertial and viscous permeabilities were determined through the use of the Forchheimer equation (as explained earlier).

Other methods to study the through-plane permeabilities were presented by Chang et al. [183] and Williams et al. [90]. However, these methods only determined the viscous permeability coefficient with Darcy’s law and did not take into account the inertial component of the permeability.

4.4.3.2.3 Relative Permeability

Relative permeability is defined as the ratio between the permeability for a phase at a given saturation level to the total (or single-phase) permeability of the studied material. This parameter is important when the two-phase flow inside a diffusion layer is investigated. Darcy’s law (Equation 4.4) can be extended to two-phase flow in porous media [213]:

P

Kkug

g

rgg ,

P

Kkul

l

rll ,

(4.6)

whereK is the single-phase permeabilitykrg is the gas-phase relative permeabilitykrl is the liquid-phase relative permeabilityml is the liquid-phase viscosityul is the superficial liquid velocityPl is the liquid-phase pressure gradient

Nguyen et al. [205] used a technique in which a constant mass flow rate of water-saturated air was forced through a water-saturated sample. It was explained that the shear force of the gas flow dragged water out of the sam-ple. In addition, the saturated air was needed in order to prevent water loss from the sample by evaporation. Once a steady state was achieved, the pres-sure difference between the inlet and outlet of the apparatus was recorded. After the tests were completed, the sample was weighed to obtain its water content. Thus, the relative permeability was calculated from the pressure drop, the water content in the sample, and the mass flow rate [205].

Koido, Furusawa, and Moriyama [206] used a technique based on the steady-state test method for reservoir rock, sandstone, and other porous media. In this method, a DL is sandwiched between similar DLs on the inlet and outlet sides. The material on the inlet is used to guarantee homogeneous distribution of liquid water in the planar direction, while the material at the outlet minimizes the flow in the outlet. Liquid water is introduced first and then a constant flow rate of air is injected. Once it is at steady state, the pressure difference between the inlet and outlet is measured. The sample is then weighed and the permeability is calculated in a way similar to that of Nguyen and colleagues [205].


4.4.3.2.4 Direct Visualization of Gas Bubbles in Direct Methanol Fuel Cells

For DMFCs, a number of studies using transparent fuel cells—mainly to visu-alize carbon dioxide bubbles in the flow channels and diffusion layers—have been conducted. For example, Argyropoulos, Scott, and Taama [32,98] designed transparent direct methanol fuel cells with the anode side made from acrylic so that carbon dioxide gas evolution and flow behavior could be studied. Current was withdrawn from the anode side of the cell using a peripheral stainless steel strip embedded into the acrylic block, which contacted the MEA. With the aid of a high-speed video camera, appropriate computer software, and trans-parent acrylic cells, gas evolution was recorded in a fuel cell working environ-ment. Other examples of transparent DMFCs can be found in the literature [219–221]. Although these studies do not describe the transport mechanism of the gas bubbles inside the diffusion layers, they still can give valuable informa-tion regarding the location and amount of gas bubbles on the DL surface.

4.4.3.3 Characterization of Liquid Transport Properties

Accumulation of water inside the DLs and CLs may cause serious failure modes that can significantly deteriorate the performance and lifetime of a fuel cell. To ensure appropriate water removal, it is vital to understand the water transport mechanism inside a fuel cell, especially in the DLs. Because CFP and CC contain complex structures and porosities, many researchers have developed methods that could facilitate the characterization and design of optimal diffusion layers with proper water removal capabilities. A lot of work has also been performed on mathematical models that attempt to ana-lyze the water flooding and transport inside DLs. A comprehensive review describing these models can be found in Sinha, Mukherjee, and Wang [222]. This section will discuss only examples of the experimental techniques.

4.4.3.3.1 Transparent Fuel Cells

In recent years, the use of transparent fuel cells has increased substantially due to the need for a better understanding of liquid water accumulation on the surface of the DLs and flow through the FF channels. In most transparent cells, either the cathode or the anode (or both) has transparent polycarbonate end plates that act as windows and sit on top of the corresponding FF plates. These plates are normally thin and made out of metal, such as stainless steel or gold-plated brass, and their thickness is equal to the depth of the FF chan-nels (i.e., the channels are machined all the way through the plate). Thus, the transparent end plate also acts as part of the channels.

To avoid condensation on the surface of the windows, the overall temperature is controlled with the use of heaters located in different places within the cell. A digital camera is then placed on top of the transparent end plate and is connected to a computer that gathers and analyzes the images. The performance of these cells compared to that of typical fuel cells is usually lower due to increased ohmic resistance caused by the use of FF plates with machined-through channels.


An example of a transparent PEMFC was presented by Spernjak, Prasad, and Advani [87], who used a 10 cm2 transparent fuel cell to investigate different cath-ode DL materials’ (with and without MPLs) influence on water management. The FF channels had a single-path serpentine design with rectangular channel cross sections 1 mm deep and 0.8 mm wide. In these researchers’ study, the ana-lyzed images corresponded to those in the lower section of the cathode’s active area (closest to the outlet) because most of the water droplets were observed in this area away from the inlet. To observe how different DLs affected the water transport in the anode, this side was also visualized (see Section 4.3.3.2).

In another example, Ous and Arcoumanis [223] visualized just two small sections within the cathode single-path serpentine FF channel, in one of the turns and in the middle section of the plate. With the use of a zoom lens, water droplets on the DL surface were investigated at different current densities and relative humidities at a low cell temperature. The main focus of the study was to observe the water droplets and how they grew in the surface of the diffu-sion layers up to the point at which the neighbor droplets would make contact with each other and form larger droplets or even water slugs. The attachment of the water droplets and slugs to the DLs was also studied with different flow rates. In addition, image processing was used to obtain the average drop-let size at specific time intervals. This allowed these workers to correlate the droplet sizes with the changes in current densities and voltage.

Yang et al. [224] used a transparent cell with straight channels to observe in detail the water formation and accumulation on top of the diffusion lay-ers. A similar approach was developed by Zhang, Yang, and Wang [225], except that they investigated two different modes of water removal between the diffusion layers and the FF channels. One mode was observed at high flow rates and was based on droplet detachment by the shear force of the gas flow followed by a mist flow in the channel. The other mode (at low flow rates) was by capillary wicking onto the more hydrophilic channel walls fol-lowed by annular film and/or liquid slug flow in the channel.

Ge and Wang [226] utilized a transparent fuel cell with two different FF designs—straight and serpentine—in order to study the water accumula-tion on the anode side of a fuel cell. No water accumulated on top of the DL surfaces in either FF design during all the tests. The literature contains many other examples of transparent cells used as tools to better understand the water transport mechanisms in fuel cells [111,227–232].

4.4.3.3.2 Neutron Imaging in Fuel Cells

To date, a number of research groups and institutes have started to use neu-tron imaging as a tool that provides spatial information on water distribution throughout a fuel cell. In fact, this technique can be used to observe the water accumulation and storage inside the FF channels and diffusion layers since the neutrons provide the contrast necessary to image the hydrogen and water in a PEMFC without suffering a significant reduction in the quality of the water visualization while passing through the cell’s metal housing [233,234].


Kowal et al. [235] used this method to compare the liquid water distribu-tion in the fuel cell with CFP and CC as cathode DLs at different operating conditions and with a parallel flow field channel design for both anode and cathode plates. It was observed that the CFP DL experienced more flooding at lower current densities than the CC, and it retained more water near the landing widths than in or under the channels (60 vs. 40%, respectively). In addition to showing better performance and water removal, the CC resulted in more uniform water coverage on the landing widths and in the channels of the FF.

Owejan et al. [236] investigated the water accumulation in diffusion lay-ers when the cathode flow field had an interdigitated design. It was shown that the DL accumulated water at approximately the same rate at which the water was produced at the CL. In addition, the DL operating under saturated conditions accumulated water until it hit a critical point at which point the water was removed from the DL at a rate similar to the one corresponding to the production of water. A similar study was then performed by the same research group but with serpentine flow fields [237]. In this case, the flow behavior of a fuel cell stack was simulated by using a single cell coupled with a bypass flow loop for the cathode flow. This study is a good example to simulate and probe the water accumulation inside a diffusion layer at more realistic conditions. However, the water inside the MEA cannot be differen-tiated between the corresponding components. More examples of this tech-nique can be found in Section 4.4.7.

It is important to mention that another visualization technique based on magnetic resonance imaging (MRI) has been developed in order to observe the water flow inside fuel cells. However, in this technique, all the materi-als used in the fuel cell have to be nonmagnetic. For this reason, the water content in the CL and DL (made from CFP or from CC) would be difficult to visualize with MRI.

4.4.3.3.3 X-Ray Radiography

Manke et al. [238] used synchrotron x-ray radiography as an in situ experi-mental technique to study the liquid water transport mechanisms and dynamics in a DL. This technique gives better resolution than neutron imaging, a few micrometers versus 100 μm, respectively. They used a single fuel cell with graphite plates and metal end plates that had small sealed holes (8 mm in diameter) used as the “windows” so that the cell could be visualized. It is important to note that in this method the obtained images are only in the in plane and not the through plane of the diffusion layers. These images can be used to study the breakthrough of water on the surface of the DL and to observe the accumulation of water in certain specific areas of the DL (but not in the whole active area).

After running the cell at different flow rates and conditions, Manke and colleagues were able to determine how the liquid water droplets inside the DL were ejected into the channel and merged into larger droplets. In addition


to the optical images, they were also able to estimate the amount of water at certain locations as a function of time. These observations allowed them to conclude that the water transport mechanism changed within the DL depend-ing on the location.

4.4.3.3.4 Ex Situ Visualization Techniques

The research group led by Dr. Djilali at the University of Victoria has devel-oped an ex situ experimental technique using fluorescent microscopy to study the liquid water transport mechanisms inside diffusion layers and on their surfaces [239–243]. The diffusion layer is usually placed between two plates (the top plate may or may not have a channel); the liquid water, which is pumped through a syringe pump, flows from the bottom plate through the DL. Fluorescein dye is added to the water for detection with the microscope.

After the tests, Djilali’s group used mathematical assumptions and equa-tions to correlate the intensity of the dye in the image with the depth in the gas diffusion layer. With this method they were able to study the effect of compression on diffusion layers and how that affects water transport. Water removal in a flow channel has also been probed with this technique and it was observed that, with a dry DL slug, formation and flooding in the FF channels followed the appearance and detachment of water droplets from the DL. Even though this is an ex situ technique, it provides important insight into water transport mechanisms with different DLs and locations.

Another ex situ technique used to understand the two-phase flow in dif-fusion layers was developed by Chapuis et al. [244]. In this study, a com-plex, numerically generated, hydrophobic, fibrous medium model was constructed. First, a number of cylinders (1 mm in diameter) were drilled into a polycarbonate plate at different locations (given by the model). Then this plate was used as a mold to fabricate a slice of silicon RTV containing the array of solid cylinders. This slice was then sandwiched between a metal plate and a polycarbonate sheet. Once this porous medium was completed, water at different flow rates was poured from the top toward the bottom. The water flow was then recorded with a digital camera in order to analyze the corresponding images. A mathematical model with the same geometry was used to compare the observations obtained by the visualization tech-nique. Although this technique shows promise, it is very complex and heav-ily dependent on the mathematical model that generates the cross section of the porous medium.

4.4.3.3.5 Liquid Permeability

To measure the water permeation of various diffusion layer samples, Benziger et al. [245] placed the sample DL in a pressurized membrane filtration cell. Then water was slowly added to the cylinder in the filtration cell, and the amount of water that flowed through the DL was measured as a function of time. Once all the water was drained, the sample was weighed to determine


the amount of water retention. The threshold pressure, defined as the hydro-static pressure value at which the water started to penetrate the DL, was also measured. The liquid permeability was then determined through the following equation:

ln

hh

K gA

to

water

cell (4.7)

whereh is the height of the water head in the cellho is the initial height of the water head in the cellK is the permeabilityrwater is the density of the waterg is the gravitational constantAcell is the cross-section area of the filtration cellt is the time

Park, Lee, and Popov [136] used a similar technique to determine the liq-uid permeation in different diffusion layers. Feser, Prasad, and Advani [214] used the same method explained in Section 4.4.3.2 to measure the liquid in-plane permeability of DLs. When water was used, flow was forced from a pressurized tank (0–200 kPa) through the apparatus (and the sample), and the outlet water was then collected with a graduated cylinder.

4.4.3.3.6 Water Balance Analysis

Performing water balance measurements and calculations is important when studying water management and flooding in PEMFCs and DMFCs in detail because it can give an insight on how different components can affect the amount of water that leaves each side of the fuel cell. In recent years, many researchers have used this technique to determine the overall net water drag coefficient, which represents the moles of water per mole of protons dragged from the anode through the membrane to the cathode. This coefficient can determine how different diffusion layers (and MPLs) affect the overall water transport inside the MEA.

As briefly mentioned in Section 4.3.3.2, Atiyeh et al. [152] performed water balance measurements and calculations to determine the effect of using DLs with MPLs (on either or both cathode and anode sides). In their fuel cell test station, water collection systems were added in order to be able to collect and measure accurately the water leaving both anode and cathode sides of the fuel cell. Based on the operating conditions (e.g., pressures, temperatures, relative humidities, etc.) and the total amount of water accumulated at the outlets of the test station, water balance calculations were performed to determine the net water drag coefficient. Janssen and Overvelde [171] used this method to observe how different operating conditions and fuel cell materials affected


the drag coefficient. They noticed that no significant change in the drag coef-ficient was observed due to the high level of hydrophobicity in the DLs that were used.

In another study, Nakajima, Konomi, and Kitahara [144] studied the water accumulation in different components of the fuel cell at simulated start-up cycles. Each component was weighed before and after each test; once a test was completed, water balance analysis was performed. Through this analy-sis, the effect of different diffusion layers was probed in detail, and it was concluded that the DLs with higher gas permeability were able to remove water more efficiently. It was also observed that the MPL was effective in improving start-up performance of the fuel cell by suppressing water accu-mulation at the CL and within the DL.

In DMFCs, the water balance analytical method has been used as a tool to study the fuel (methanol) and water crossover from the anode toward the cathode. Xu, Zhao, and He [120] and Xu and Zhao [180] performed a thorough investigation of how different cathode DLs and MPLs affected the total water crossover from the anode side. In order to be able to perform the water balance equations, they also collected the water at both outlets of the cell. This analysis technique was vital for them to be able to observe how different characteris-tics for the cathode DL affect not only the overall performance of the fuel cell but also the net water drag coefficient and water crossover in DMFCs.

Liu, Lu, and Wang [246] tested a number of different MEA designs through water balance analysis in order to observe how each design affected the water crossover. With this technique, they were able to show how the fuel crossover through the membrane decreased significantly when the water crossover was also limited by the new MEA designs. Song, Lee, and Kim [181] used this measurement method to improve the microporous layer design of the cath-ode DLs based on how the MPL affected water crossover and water balance.

In conclusion, performing a water balance analysis in a fuel cell not only is critical in order to determine the effectiveness of membranes water crossover (or drag), but it is also a very effective tool when investigating how different diffusion layers can manipulate water flow and water management within the MEA.

4.4.4 Electrical and Thermal Conductivity

Electrical and thermal conductivity are important diffusion layer properties that affect the fuel cell’s overall performance. The material chosen to be the DL in a fuel cell must have a good electrical conductivity in order for the electron flow from the FF plates to the CLs (and vice versa) to have the least possible resistance. Similarly, the DL material must have good thermal prop-erties so that heat generated in the active zones can be removed efficiently. Therefore, in order to choose an optimal material it is critical to be able to measure the electrical and thermal conductivity. In this section, a number of procedures used to measure these parameters will be discussed.


4.4.4.1 Electrical Conductivity

The following subsection will briefly discuss the main methods used to mea-sure in-plane and through-plane electrical conductivity for diffusion layer materials. This parameter is critical for optimal fuel cell performance.

4.4.4.1.1 In-Plane Conductivity

The most typical way to measure the in-plane electrical conductivity of a diffusion layer is through the use of the four-point probe method. A small current is applied across the sample material; a separate set of voltage mea-suring probes that are in touch with the material are used to measure the resulting voltage drop in order to calculate the resistance. With these val-ues, the in-plane resistivity, r, can be calculated with the following equa-tion [9,233]:

RwxL

(4.8)

whereR is the resistanceL is the distance between the voltage measuring probesw and x are the width and thickness of the sample, respectively

Nitta et al. [216] designed a method to measure the in-plane conductivity of a DL as a function of the compressed thickness. The sample material was placed on a plate and compressed on both ends by graphite current collec-tors. Steel gages were located between the graphite blocks and the plate in order to maintain constant thickness of the DL while tests were conducted. A specific current range was applied to the apparatus and the voltage drop was measured in order to calculate the total resistance of the system.

Through assumptions and the use of values for known resistances of the materials used in the apparatus, the actual bulk resistance of the DL material could be calculated. This resistance was then used so that the electrical con-ductivity could be solved. Nitta and colleagues noted that the in-plane conduc-tivities of the DL materials were a linear function of the compressed thickness (i.e., the conductivity increased when the thickness decreased with increased compression pressure). This resulted from a decrease in thickness that led to a loss of porosity in the DL materials and higher contact between fibers.

4.4.4.1.2 Through-Plane Conductivity

The most typical way of measuring this parameter is to place a sample material between two plates (compressed at a defined pressure) and apply a direct current through the DL. The voltage drop between the plates is measured to determine the resistance [9,100,247]. Through this system, the resistance can be measured at different compression pressures. To minimize the contact resistance between


the plates, they are usually made out of gold-plated copper. The through-plane conductivity can then be determined by implementing an analysis of the total resistance within the system and the separate resistances from each component that is part of the testing apparatus. Once the bulk resistance (or resistance of the DL) is solved, the conductivity can be measured as well. Both Mathias et al. [9] and Nitta et al. [216] show in detail how this analysis can be performed in order to calculate the final through-plane conductivity of a material.

Similarly to in-plane conductivity, through-plane conductivity seems to be a linear function of the compressed thickness of the DL; that is, con-ductivity increases linearly with thickness decrease. Nitta et al. [216] also observed that in-plane conductivity was larger than through-plane con-ductivity; however, the difference was not as large as that found in previ-ous studies [9].

4.4.4.2 Thermal Conductivity

As mentioned by Mathias et al. [9], reliable methods to measure the thermal conductivity of diffusion layers as a function of compression pressures are very scarce in the open literature. Khandelwal and Mench [112] designed an ex situ method to measure accurately the thermal conductivities of differ-ent components used in a fuel cell. In their apparatus, the sample materials were placed between two cylindrical rods made out of aluminum bronze (see Figure 4.28). Three thermocouples were located equidistantly in each of the upper and lower cylinders to monitor the temperatures along these com-ponents. Two plates located at each end compressed both cylinders together. The temperatures of each plate were maintained by flowing coolant fluids at a high flow rate through channels located inside each of the plates. A load cell was located between two plates at one end so that the compression pres-sure could be measured.

Prior to any test, the system was controlled for 8–10 hours to make sure that steady state was achieved. Once steady state was reached, the tempera-tures of all the thermocouples were measured and recorded to determine the temperature drop across the sample material. Because the heat flux is proportional to the temperature difference, the through-plate thermal con-ductivity across the material can be determined using Fourier’s law [141]:

q k

Txi i

i (4.9)

wherei represents x, y, or z dimensionqi is the conduction heat transfer flux per unit area in the i directionki is the thermal conductivity in the i directionT is the temperature


These tests were performed on materials with the same characteristics but with different thicknesses; thus, the intrinsic thermal conductivity could be resolved at different temperatures and compression pressures. Through these tests, the thermal conductivity of TGP-H carbon fiber papers was mea-sured and achieved the same value as that reported by the manufacturer. In addition, it was observed that the thermal conductivity of the CFPs decreased from 1.80 0.27 W m–1 K–1 (at 26 C) to 1.24 0.19 W m–1 K–1 (at 73 C). This result was suggested to be due to the presence of carbonized thermosetting resin on the CFPs. The thermal conductivity of the resin, which is a thermo-setting polymer and acts as a binder, decreases with increasing temperature. For carbon cloth (without any resin), no significant changes in thermal con-ductivity were noted when the temperature was increased.

In a similar method, Ramousse et al. [248] designed a technique wherein the sample material is placed between two copper plates that have thermo-couples located at their centers. Copper plates were chosen due to the high thermal conductivity of copper and to ensure a uniform temperature distri-bution. Fluxmeters to measure the thermal flux between both plates were located beside each copper plate. At each end of the apparatus, end plates

AluminumBronze

FIGURE 4.28Schematic of the equipment used to measure thermal conductivity of thin materials. (Reprinted from M. Khandelwal and M. M. Mench. Journal of Power Sources 161 (2006) 1106–1115. With per-mission from Elsevier.)


were used to keep the system compressed to control the sample thickness. In addition, the top end plate was heated and the bottom plate was kept cooled with a water bath. The whole system was kept well insulated to avoid any temperature losses. The mean value of the heat flux through the sample mate-rial was used to calculate the thermal conductivity with the help of a math-ematical model. Their results resembled those presented by Khandelwal and Mench; compared to other values found in the literature (mostly based on mathematical models), their values are much lower.

In another study, Nitta, Himanen, and Mikkola [249] used a design very similar to those explained previously, except that the two rods used were made out of graphite and the bottom rod was heated with a heating element. The sample materials were placed between both rods. The temperature of the upper rod was kept constant, at a lower value than the lower rod, by a cooling block with coolant fluid flowing through it. A total of four thermo-couples (two per rod) were at different locations of the rods (two very close to the sample material) in order to measure and record the temperatures. The compression pressure could also be controlled from the top end plate. The values did not change significantly over a wide range of pressures; thus, no clear dependence of the thermal conductivity on the compression pressure was found.

It is important to note that Vie and Kjelstrup [250] designed a method of measuring the thermal conductivities of different components of a fuel cell while the cell was running (i.e., in situ tests). They added four thermocouples inside an MEA (i.e., an invasive method): one on each side of the membrane and one on each diffusion layer (on the surface facing the FF channels). The temperature values from the thermocouples near the membrane and in the DL were used to calculate the average thermal conductivity of the DL and CL using Fourier’s law. Unfortunately, the thermal conductivity values pre-sented in their work were given for both the DL and CL combined. Therefore, these values are useful for mathematical models but not to determine the exact thermal characteristics of different DLs.

4.4.5 Mechanical Properties

The membrane and catalyst layers in a fuel cell are thin and delicate compo-nents that require mechanical support in order to prevent rupture or substantial bending when a compression pressure is applied to the whole cell. Therefore, the diffusion layers must provide the necessary mechanical support to those components without affecting the other parameters discussed previously.

The compressive behavior of a DL is a very important mechanical property. Therefore, to study the mechanical properties of various diffusion materials (carbon cloths, carbon fiber papers, and carbon felts), Escribano et al. [251] used a compression cell. The sample diffusion materials were placed between the two plates of the cell, and the thickness and deflection of each sample were measured as a function of the compression pressure. These researchers


performed two successive compressions on the specimens: The first one rep-resented the pressure used to manufacture the MEA and the second repre-sented the compression pressure caused by the bipolar plates when the whole fuel cell is assembled. Both compressions were performed in a pressure range of 0–10 MPa.

It was observed that the CC had the highest compressibility with the great-est compressive strains (70% at 10 MPa and 55% at 2.5 MPa) and the CFP showed two different behaviors depending on the pressure range. At low pressures, the thickness decreased rapidly and the compressive strain had a small slope that increased significantly after 3 MPa. In the second compres-sion, the CFP behaved more like a hard stop; this was also observed in the carbon felt. The compressive strains of the paper were 55% at 10 MPa and around 20% at 2.5 MPa. The residual strains after the first compression for both CFP and CC were very similar (30 vs. 25%).

Among the three materials, the carbon felt had the most stable behavior, with the lowest compressive strain (25% at 10 MPa and 20% at 5 MPa) and almost no residual strain (<5% after first compression). It was also shown that CCs with hydrophobic treatment and CCs with MPLs improved the limit of compressibility by decreasing the slope of compressive strain versus com-pressive stress.

Lee and Merida [215] studied the strain of different CFPs (with PTFE coat-ings) before and after a number of aging tests, which consisted of compress-ing the samples at 200 or 400 psi while heating them up to temperatures in the range of 80–120 C. The strain was calculated by measuring the thickness of each sample material before and after the tests (see Section 4.4.6 for more information). Temperature was found to have a greater effect on maximum strain than the compression pressures that were applied. This result was attributed to the temperature dependence of PTFE. Thus, after aging at high temperatures, the mechanical properties of the PTFE deteriorated, resulting in a lower mechanical stability of the diffusion layer.

Flexural or bending behavior is also a critical mechanical parameter, depending on the DL material that is used, because the FF channels and landing widths can deform the DLs. In addition, it is not desirable for the DL to intrude inside the channels because this causes blockage of the gas and liquid flow, thus increasing the overall pressure drop within the flow field. Mathias et al. [9] mentioned that a number of reasonable tests can be done to determine the bending behavior of a DL. For example, in the ASTM D790 method, a sample material is placed between two supports and a force is applied in the center. From this test, the displacement of the material versus the force response is used to determine the flexural modulus and strength. In other tests, the specimen is clamped on the ends and then the material is bent in the center in order to determine the modulus and strength based on the material’s response.

Nitta et al. [216] studied the intrusion of the DL inside a channel with the use of a simple apparatus. The setup consisted of two aluminum plates with the


DL and two steel gages placed between them. The gages were used to allow precise control of the compressed DL thickness. Different bottom plates were used in order to study a wide range of different flow channels. A dial indicator was located at the bottom of the channels in order to measure the DL intrusion inside the channel. The DL under the channel did not present a significant change in thickness regardless of the channel width. Slight compression of the DL under the channel was observed, but only at very high compression pres-sures. It was also observed that the DL thickness under the landing widths (or ribs) changed significantly; this was considered to be a result of the loss in porosity. Therefore, within the active area, the porosity of the DL can change substantially depending on the FF design and the compression pressures.

It is important to operate the fuel cell at different compression pressures in order to determine the correct compression pressure for a DL material. If the applied compression pressures are too high, the DLs may deform, both the porosity and permeability of the DL decrease, and the probability of failure modes increases significantly. On the other hand, if the pressures are too low, then gas leaks and serious contact resistance between the components of the cell may be present. Various studies have been presented in which the compression pressure of the fuel cell is varied in order to observe how the cell’s performance is affected [25,183,252]. In general, there is an optimal compression pressure range in which the cell’s performance is the highest; however, this depends on the DL material and on the MPL thickness (see Figure 4.21).

4.4.6 Corrosion Stability and Degradation

Fuel cells must be able to operate without major obstacles for long periods of time in a number of applications (e.g., utility-based applications). However, determining methods to measure the lifetime (and durability) of a fuel cell and its components is a complex matter; such methods require a significant amount of testing time and resources [253]. Durability testing is often done over a wide range of conditions (e.g., constant voltage, constant current, dif-ferent temperatures and humidities, power cycles, etc.). During durability testing, a number of in situ measurements are usually performed, such as polarization curves, current distribution, hydrogen crossover, analysis of effluent water, etc.

Once the durability testing of the fuel cells is finalized, the internal com-ponents are then characterized. For diffusion layers, some of these charac-terization techniques include SEM to visualize surface changes, porosimetry measurements to analyze any changes in porosity within the DL and MPL, IGC (inverse gas chromatography) to identify relative humidity effects on the hydrophobic properties of the DLs, contact angle measurements to observe any changes in the hydrophobic/hydrophilic coatings of the DL, etc. [254,255].

Ex situ tests are also performed with most of the components used in a fuel cell in order to understand the individual contributions of the fuel cell


components to the overall degradation of the system. In addition, ex situ methods are useful because they can accelerate testing and aging effects for fuel cell components. However, the amount of data and reports found in the literature regarding these methods for diffusion layer materials and MPLs is quite limited; most of the studies have been focused on performance rather than durability.

4.4.6.1 Corrosion Studies

One of the common ways in which fuel cell components experience degra-dation is through corrosion. Carbon particles in the CL are susceptible to electrochemical (voltage) corrosion and contain Pt particles that catalyze oxi-dation reactions. The carbon fibers in CFPs and CCs and the carbon black in MPLs are not as susceptible to these issues because they are not part of the electrochemical reactions and do not contain Pt particles. However, they can still go through chemical surface (hydrogen peroxide) oxidation by water or even by loss of carbon due to oxidation to carbon monoxide or carbon diox-ide [256,257].

Oxidation of the DL can cause changes in the hydrophobicity of the mate-rial, resulting in flooding in the MEA and in an increase in mass transport limitations within the DL. Frisk et al. [257] at 3M Company demonstrated how the performance of a fuel cell decreases after a long-term test, mainly due to the mass transport loss increase over time. A number of methods have been designed to study the oxidation of different DLs with microporous lay-ers. In the first method, the sample materials were submerged in a bath of 15 wt% H2O2 at 82 C for a set time period. After each test, the weight loss and contact angle (sessile drop method) of the specimens were measured. The weight loss was a result of the oxidation of the carbon articles in the MPL. The contact angle measurements demonstrated how the hydrophobic properties of the materials decreased as the time during which they were submerged in hydrogen peroxide increased.

In the second method, a voltage was applied to the DL samples while they were in contact with an electrolyte. From this method, the corrosion current density (or rate of oxidation) could be determined and analyzed. It was evi-dent that as the voltage increased, the corrosion current density increased substantially. These methods can be used to select appropriate materials to be used as diffusion layers in fuel cells.

Borup et al. [254,255,258] have studied the corrosion of DLs. They aged different types of hydrophobic treated DLs for around 1,000 hours in deion-ized water at 80 C. After these aging tests, the samples were fuel cell tested at different relative humidities. It was observed that the DLs that were aged behaved like hydrophilic DLs; they showed the best performance under dry conditions and the worst under high-humidity conditions due to flooding. On the other hand, hydrophobic DL materials that were not aged showed the lowest performance during dry conditions and the best


performance under high-humidity conditions due to their ability to remove water effectively.

These researchers also demonstrated that increasing the carbon-oxide layers by contacting the carbon-based materials with hydrogen peroxide increased the hydrophilic properties of the carbon surface. This mechanism was used to explain the loss of hydrophobicity on the DL surface after long hours of fuel cell operation. To further study the aging of DLs, different car-bon fiber papers with hydrophobic treatment were aged at different condi-tions. Figure 4.29 shows how the contact angle of the CFPs decreases after exposure to nitrogen or air at 60 or 80 C for long periods of time. After aging, the DL materials substantially lose their hydrophobic properties, indicating that water flooding is more prone to happen after long hours of fuel cell operation. It was also shown that exposing the DL materials to air is more aggressive than exposing them to nitrogen.

Through the use of x-ray-induced photoelectron spectroscopy (XPS), Schulze et al. [259] were able to demonstrate that the PTFE particles, which are coated on the diffusion layers, decomposed after fuel cell testing for more than 1,600 hours. This resulted in a change of the hydrophobic properties of the DL. Unfortunately, the mechanism behind the decomposition of PTFE was not explained.

4.4.6.2 Cold and Freezing Temperatures

To understand how the properties of diffusion layers change when exposed to extreme cold (and freezing) temperatures, Lee and Merida [215] used a number

96

93

90

87

84

Co

nta

ct A

ngl

e (d

eg)

81(a)

(b) (c)

(d)

(e)

(f )

(g)

Hydrophobic

Hydrophilic

78Unaged

TGP-H (Plain)UnagedTGP-H

UnagedTGP-H

TGP-H N2,60 °C, 460 hr

TGP-H N2,80 °C, 460 hr

TGP-H N2,60 °C, 680 hr

TGP-H N2,80 °C, 680 hr

FIGURE 4.29Contact angle measurements of Toray carbon fiber papers: (a) untreated; (b) 17.2 wt% FEP; (c) 17 wt% FEP; (d) 17.2 wt% FEP—aged 460 hours in N2 and 60 C; (e) 16.7 wt% FEP—aged 460 hours in N2 and 80 C; (f) 16.9 wt% FEP—aged 680 hours in air and 60 C; (g) 17.0 wt% FEP—aged 680 hours in air and 80 C. (Reprinted from J. Borup et al. ECS Transactions 3 (2006) 879–886. With permission from ECS—The Electrochemical Society.)


of ex situ tests to study the bending stiffness, electrical resistivity, air permeabil-ity, surface contact angle, porosity, and water vapor diffusion of DLs after aging under freezing conditions. After 50 consecutive freeze–thaw cycles between –35 and 20 C, no obvious issues with most of the mentioned properties were observed. The only two properties that increased were the in-plane and through-plane permeabilities and this was due to material loss during the permeability measurements. It was concluded that the MPL was weakened after exposure to freezing conditions, thus having material loss from air flow through the GDL.

In another, similar study, Mukundan et al. [260] performed 100 freeze–thaw cycles (from –40 to 80 C) with different types of CFPs and CCs. After 100 cycles, no obvious degradation was observed in the carbon cloth DL; in fact, the performance of the fuel cell slightly improved. On the other hand, after 45 cycles, the CFPs showed significant breakage of the carbon fibers at the edges between the flow channels and the landing widths (or ribs). Thus, it was con-cluded that this breakage could potentially become a serious failure mecha-nism in PEM fuel cells when the system was started at subzero temperatures.

After treating different fuel cells to 100 freeze–thaw cycles (from –40 to 70 C), Kim, Ahn, and Mench [261] concluded that stiffer materials used as diffusion layers improved the uniform compression with the CL, resulting in fewer issues after the freeze and thaw cycles. On the other hand, more flex-ible DLs failed to improve the compression; the CL left open spaces for ice films to be formed, resulting in serious issues after the freeze–thaw cycles. However, even with the stiffer materials tested, such ice films were still evi-dent and caused delamination of the DL and CL, surface damage in the CL, and breakage of the carbon fibers. This resulted in increased electrical and mass transport resistances.

Cold-start tests are also important in order to observe how freezing tempera-tures affect the material properties of the diffusion layers. Oszcipok et al. [262] dried a fuel cell through purging with N2 prior to cooling it down to –10 C. At this point, the cell was started and operated at different voltages. After more than 10 tests, the cell was disassembled and it was observed that the FF channel pattern was visible on the cathode DL.

Contact angle tests showed that the hydrophobic properties on those areas corresponding to the channels had changed substantially (more hydrophilic) compared to other parts of the DL (e.g., those close to the landing widths). The same effect was visible on the surface of the MPL, indicating that loss of the PTFE particles was experienced on both surfaces of the DL. Similar observations regarding the hydrophilicity of the areas in which the FF chan-nels are present were presented previously by St-Pierre and Jia [263].

In conclusion, it is evident that most of the degradation issues in diffusion layers are related to the hydrophobic properties of the materials. Therefore, techniques that investigate the different degradation and corrosion modes that may affect the diffusion layers inside fuel cells are vital in order to design and choose optimal diffusion layers with the goal of achieving the best reliability and performance possible.


4.4.7 Flow Field Interaction

An optimum relationship between the DL and the flow field channels is a key factor in the overall improvement of the fuel cell’s performance at both high and low current densities. Currently, flow field designs are typically serpentine, interdigitated, or parallel [207,264]. The FF plate performs sev-eral functions: It is a current collector, provides mechanical support for the electrodes, provides access channels for the reactants to their respective elec-trode surfaces and for the removal of product water, and it prevents mixing of oxidant, fuel, and coolant fluids.

Therefore, it is necessary to have good interaction between the diffu-sion layers and the FF plates—not only from a mass transport standpoint but also to maintain optimal electrical and thermal conductivity between them. Section 4.4.4 explained in detail measurement techniques to deter-mine the electrical resistance in diffusion layers. It is important to note that most of those methods can also be implemented in order to calculate the contact resistance between the DLs and the FF plates. In this subsection, we will focus mostly on mass transport interactions between these two components.

The basic nature of the flow field design dictates many of the diffusion layer requirements. For example, Wilkinson, Vanderleeden, and Zimmerman [265] developed FF features for supporting fluid diffusion layers in fuel cells. Unfortunately, a fairly limited number of studies investigating the interac-tion between the flow fields and the diffusion layers can be found in the literature, compared to the amount of modeling work. Through mathemati-cal modeling, a great number of parameters and operating conditions can be simulated in a shorter period of time and without the costs involved in machining, fabrication, and testing of fuel cells. However, experimental data are still critical in order to be able to compare and improve the modeling work. Details on mathematical modeling of flow fields and diffusion layers are beyond the scope of this chapter.

4.4.7.1 Pressure Drop Tests

Measuring the pressure drop in an FF over a wide range of flow rates is a common step when different FF designs are examined. Similarly, this method is used when how different diffusion layers affect the overall pressure drop of the FF is studied. For the conventional serpentine designs, convection through the diffusion media due to adjacent channel pressure drop must be controlled to prevent gas shorting [9]. Williams et al. [266] analyzed the pressure drop in a single serpentine flow field with and without a DL. They observed that the pressure drop of the FF without the DL was significantly higher than with the DL. It was concluded that the difference in pressure drop between the two cases was caused by gas flow from the higher pressure channel to the adjacent lower pressure channel through the DL.


Park and Li [267] also performed an experimental and numerical study in which a serpentine FF with and without diffusion layers was analyzed (see Figure 4.30). For the case in which a DL was not used, an imperme-able plate was placed between the anode and cathode plates in order to perform the pressure drop tests. It was observed that the pressure differ-ence between the two cases was as large as 80% of the pressure drop in the case without the DL. It was also explained that the reason for the large pressure difference between both cases was due to the cross-flow phe-nomena between adjacent channels through the porous diffusion layer. These researchers also performed the pressure drop tests when the fuel cell was running; it was observed that the pressure drops were higher in this case than when the cell was inactive. It was believed that this pressure difference was a result of liquid water blocking reactant flow in either the channels or the DLs.

Soler, Hontanon, and Daza [268] tested two different FF designs with a num-ber of carbon fiber paper and carbon cloth DLs in order to determine the best combination. They measured the pressure drop of the flow field in a nonactive fuel cell with each DL material with oxygen, air, and nitrogen. The researchers

160Experiment (without GDL)

Experiment (with GDL)

Correlation (without GDL)

3D simulation (without GDL)

3D simulation (with GDL)

Experiment (running cell)

140

120

100

80

ΔP

(k

Pa)

60

40

20

0

0 500 1000 1500

ReDh

2000 2500

FIGURE 4.30Pressure drop as a function of the Reynolds number for experiments and simulations with and without the diffusion layer on the cathode side. Experimental data are from active operating PEMFC at a current density of 0.2 (+) and 0.4 A cm–2 (x); air stoichiometry (humidified) is 3.0; operating temperature is 80 C. (Reprinted from J. Park and X. Li. Journal of Power Sources 163 (2007) 853–863. With permission from Elsevier.)


determined that the effect of DL permeability on PEMFC performance depends strongly upon the FF design. Depending on this design, the effect of the DL’s permeability may cause major performance losses (see Section 4.4.3 for more information regarding permeability of DLs).

Pharoah [269] presented similar results, in which computational fluid dynamics (CFD) analysis was performed on a long serpentine flow channel and a GDL. It was observed that as the flow channel got longer, the maxi-mum pressure difference between channels increased, resulting in larger driving forces for convective flows through the GDL. In brief, it can be con-cluded that the significance of convective flow through the GDL depends on the pressure drop (driving force) between two adjacent channels in each FF and the permeability of the DL.

In the case of interdigitated flow field designs, the diffusion layer permea-bility must be high enough to permit gas passage without excessive pressure drop [9,207]. Yamada et al. [210] performed pressure drop tests with parallel and interdigitated flow fields in order to study the water accumulation in different diffusion layers when running the fuel cell. From these measure-ments, they were able to observe how different DLs affected the performance of the cell as a result of mass transport losses caused by flooding.

He, Yi, and Nguyen [270] modeled a PEM fuel cell with an interdigitated FF and studied the influence of this flow field on the electrode (DL plus CL) performance. It was shown that the performance of the cathode improved as the number of channels increased. Decreasing the landing width (i.e., increasing the number of channels) decreased the flow path, resulting in a lower pressure drop and a higher gas-flow rate. However, a certain mini-mum amount of landing was needed in order to have an acceptable contact area and conductivity.

4.4.7.2 Visualization Techniques

As discussed in Section 4.4.3, there has been a lot of interest in using visu-alization techniques to observe liquid and gas transport through flow field channels and on the surfaces of different diffusion layers. Some of these meth-ods have facilitated the study of the interaction of the flow field and DLs. For example, Owejan et al. [147] used neutron radiography imaging in order to visualize the liquid water inside two different FF designs and a number of dif-fusion layers (with and without MPL). Both FF designs had five path serpentine channels, with the same cross-sectional channel area, but one had rectangu-lar cross-sectional shaped channels and the other one had triangular shaped channels. Through this visualization technique they were able to observe how the accumulation of water in each DL increased with higher current densities and to determine what combination of flow field design and diffusion layer presented the best performance. It was concluded that one of the most impor-tant parameters for a DL to perform well is high in-plane permeability, which yields effective gas transport over the whole area.


Turhan et al. [271] also used neutron imaging and observed that, with a serpentine FF, carbon fiber paper DLs were more saturated with water than carbon cloth DLs. This was due to more effective water removal from the carbon cloth, especially under the landing widths of the FF channels. The same research group later used the same technique to evaluate serpentine flow fields with different landing-to-channel ratios [234]. They concluded that the water accumulation (or storage) in the diffusion layer was a function not only of the diffusion layer itself but also of the FF geometry and surface of the flow field wall characteristics. They also observed that a reduced num-ber of channel–DL interfaces corresponded to fewer corners and turns for a serpentine FF design; this resulted in lower water accumulation in the cell (because corners in the FFs are prime areas for water accumulation).

Kramer et al. [272] used this same technique to compare two different flow field designs—interdigitated and serpentine—and their interactions with the cathode diffusion layer. It was shown that the bottom of the interdigitated chan-nels got plugged with liquid water that was not removed properly. On the other hand, the serpentine FF could transport the water inside the channels more effectively, but inside the cathode DL accumulation of water was still evident.

Another example of neutron imaging is the one presented by Yoshizawa et al. [273], who compared the performance of carbon cloth and carbon fiber paper with a parallel FF design. The CC had a better performance than the CFP at high current densities, but the CFP showed less water content over the whole active area. Thus, it was concluded that the CC diffusion layer was less influenced by the accumulation of water because the transport of oxygen toward the catalyst zones was sufficient while still keeping the membrane humidified.

Transparent fuel cells are also common tools used to visualize and observe the water accumulation inside FFs and on the surfaces of diffusion layers. Liu, Guo, and Ma [227] tested interdigitated and parallel flow fields with CFP DLs. It was observed that the former FF design enhanced the mass trans-fer when the gas flow was forced to pass through the DL. In fact, the water flooding areas in the interdigitated channels were substantially smaller than in the parallel channel.

Another example of a visualization technique used to study the inter-action between FFs and DLs was presented by Lozano et al. [274]. They developed an ex situ method in which the flow in a DL located on top of an FF plate was examined by using acetone vapor planar laser-induced fluorescence. These researchers tested four different FF designs: parallel, diagonal-parallel, parallel-serpentine, and cascade designs. By means of flow analysis, they were able to conclude that the flow through the CFP is mostly controlled by the pressure gradient across itself. Therefore, even a nonuniform velocity distribution of the reactant gas in the plate could result in an acceptable distribution over the CL if the pressure drop was sufficiently low across the DL. On the other hand, a uniform flow in an FF plate may result in a defective gas distribution in the CL if the pressure gradient across the DL is too high.


In conclusion, most of the visualization techniques that have been devel-oped can be used to study in detail the interaction between diffusion layers and flow fields. However, it is important to note that there are still a number of issues with these techniques that may limit their precision.

In addition to the three basic FF designs mentioned, various new FF chan-nel concepts (e.g., biomimetic or fractal flow fields, improved mass trans-fer channels with variable channel cross section, etc.) have been proposed recently [275]. In all cases, the DL requirements and design will depend on the type of FF design. Therefore, it is critical to understand the relationship between any flow field design and the corresponding DL.

4.5 Future Direction of Diffusion Layers

Key milestones in the progress of the PEMFC have been demonstrated using conventional diffusion layer materials such as carbon cloth and carbon fiber paper. These include demonstration of fuel cell stack power density, durabil-ity, dynamic and operational response, ability to operate on multiple fuels, and field-trial and low-volume commercial testing in a number of applica-tions. Significant challenges and opportunities still remain for the diffusion layer in the fuel cell. There is a drive toward more functional diffusion lay-ers, lower cost materials and manufacturing processes, better tolerances and reproducibility, and control of properties over the active fuel cell area, along with more fundamental understanding and predictive performance.

New approaches to the diffusion layer are developing that will likely include integration of functions, predetermined performance design for appli-cations, and low-cost materials and processing for high-volume manufacture. Presently, the diffusion layer and the microporous layer provide two important mass transport functions. However, it is likely in the future that these separate components will be integrated into one component making use of approaches such as graded porosity. Integration of functionality in the diffusion layer is an important future direction. Figure 4.31 shows a cross section of an integrated flow field and diffusion layer first proposed and tested by Wilkinson et al. [37] that gave good performance and improved power density.

Surprisingly, little has been reported on this approach and virtually no experimental results or optimized modeling of the concept has been shown. In another approach, Wilkinson et al. [37] demonstrated the concept of introducing integration of sealing and the flow field into a perforated fluid impermeable sheet material acting as the diffusion layer as a fully inte-grated, low-cost approach to the diffusion layer and other cell functions. In this and later work [38], it was shown that the perforation size, shape, and spatial orientation could be changed to improve performance.

Graftech Inc. was the first company to look at developing perforated diffusion layers from graphite sheet materials made out of compressed expanded graphite


particles [58], but the material is still not widely adopted. A tremendous oppor-tunity remains in the perforated approach to provide an engineered approach to mass transport in the diffusion layer to optimize performance for different oper-ating conditions (e.g., dry and flooded conditions) and for gradients in operating conditions across the active area. In addition, such rational design approaches provide the opportunity to make connections between high-level performance effects and subscale properties and can greatly simplify modeling.

Another important future area for diffusion layers is the use of three-di-mensional catalyzed diffusion layers for liquid-based fuel cells. This allows the three-phase active zone to be extended into the diffusion layer to increase per-formance and utilization and reduce crossover [276,277]. Recent work by Lam, Wilkinson, and Zhang [278] has shown the scaleable use of this concept to create a membraneless direct methanol fuel cell. In other work by Fatih et al. [279], the

Ion ExchangeMembrane

AnodeCathode

Fuel Cell

Separator

LayersSeparator

Layers

Cathode

Groove

Anode

GrooveMembrane

Electrode

Assembly

FIGURE 4.31Cross section of lightweight fuel cell membrane electrode assembly having integral reactant flow passages. (From J. P. Wilkinson et al. (1999) U.S. Patent 5976726.)


CFP cathode diffusion layer itself was used as the catalytic surface for a redox fuel cell.

Such new approaches to the diffusion layer increase its importance in the fuel cell as a critical component along with the catalyst layer, membrane, and flow field.

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305

5Bipolar Plates and Plate Materials

Tim Cheng*

* The author is associated with Automotive Fuel Cell Cooperation (AFCC). AFCC is owned by Daimler, Ford, and Ballard Power Systems. The company was established in February, 2008 and is mainly based on the automotive fuel cell division of the previous Ballard Power Systems.

CONTENTS

5.1 Introduction ................................................................................................3065.2 Functions, Structures, and Performance Requirements

of Bipolar Plates ..........................................................................................3085.2.1 Basic Structure of a Bipolar Plate and Its Role in a Stack .........3095.2.2 Functions and Performance Requirements of Bipolar

Plates and Plate Materials ............................................................. 3115.3 Traditionally Used Materials of Bipolar Plates and Major

Technical Barriers ...................................................................................... 3145.4 Progress and Challenges in the Development of Bipolar Plate

Materials ...................................................................................................... 3155.4.1 Composite Plates ............................................................................ 316

5.4.1.1 Carbon/Carbon Composite Plates ................................ 3175.4.1.2 Thermoset-Based Composite Plates ............................. 3195.4.1.3 Thermoplastic-Based Composite Plates ....................... 321

5.4.2 Metal Plates ....................................................................................... 3255.4.2.1 Types of Metals Used in Metal Plates .......................... 3265.4.2.2 Key Fabrication Processing ............................................ 328

5.5 New Activities and Development Trends ..............................................3335.5.1 Development of New Materials with Better Performance

and Durability ................................................................................3345.5.2 Development of New Manufacturing Processes with

Higher Efficiency and Lower Cost ..............................................3355.5.3 Strong Competition between Different Candidate

Materials and Processing ..............................................................3365.5.4 Scale-up from Lab-Scale to Large-Quantity Production .........336

5.6 Summary ..................................................................................................... 337References .............................................................................................................340


5.1 Introduction

Proton exchange membrane fuel cells (PEMFCs) including direct metha-nol fuel cells (DMFCs) have been recognized as the most promising clean-energy converting devices due to their high efficiency and zero emission. A fuel cell contains several important and cost-dominant components, including electrocatalysts/catalyst layers, membrane, gas diffusion layers (GDLs), and bipolar plates. In this chapter, the focus will be on the bipolar plates.

The major function of a bipolar plate, or simply called “plate,” is to connect each cell electrically and to regulate the reactant gas (typically, hydrogen and air in a hydrogen fuel cell) or reactant liquid (typically, methanol in a DMFC) and liquid or gas coolant supply as well as reaction product removal in desired patterns. This plate must be at least electrically conductive and gas and/or liquid tightened. Considering these important functions and the larger fraction of volume, weight, and cost of the plate in a fuel cell, it is worthwhile to construct this chapter with emphasis on the current status and future trend in bipolar plate research and development, mainly for the plate materials and fabrication process.

Materials are the basis for any device design and manufacture and, in particular, for fuel cells. Therefore, material development is critical for plate design and manufacture of fuel cell development and commercialization. Normally, development or down-selection of materials for a plate is deter-mined by its major engineering requirements, which include:

performance or properties determined by functions;durability or lifetime;manufacturability, particularly the availability for mass produc-tion; andcost.

These requirements are also applicable to the case of fuel cells.On the other hand, fuel cell applications and the market will determine

the priority order of these requirements. The major markets or applications where fuel cells have been applied or have a large potential to be applied include [1]:

transportation;materials handling;backup power, mainly in telecommunication;

Bipolar Plates and Plate Materials 307

portable applications in laptop computers and other small electronic products;

stationary or co-generation of electricity and heat for residential or industrial applications; and

military applications.

The requirements for plate materials in a fuel cell stack for different mar-kets or applications can be quite different due to fuel cell working conditions and specific needs for the power, lifetime, weight, volume, size, and accept-able cost range. For example, in addition to basic requirements of all plate materials for their common functions, the plate material used in transporta-tion fuel cells, such as that used in automotive applications, would be signifi-cantly different from requirements in stationary stacks in terms of working temperature range, density, durability, and lifetime.

One obvious reason for the different requirements is that the fuel cells in a movable vehicle must have light weight, high power density, as well as the capability to work in aggressive conditions with frequent cycling of load, temperatures, and humidity. In addition, the automotive fuel cells must have long durability in extreme weather conditions (very hot or very cold). However, the stack in stationary applications will need a much longer work-ing lifetime with less critical requirements for stack weight, size, and special mechanical and thermal capability due to the slightly compromised work-ing conditions. Similarly, small volume and dimension would be the critical requirement of fuel cells for portable applications.

Therefore, the type of applied market of a stack and its cascaded special requirements for plates should be considered when we work on materials development or make materials selection for a bipolar plate. In addition to the importance of materials development and selection for a plate, the plate design and manufacturing process play important roles in assuring fuel cell performance, durability, and low cost. In addressing this, finite element analysis (FEA) has been popularly used in modeling plate performance once the initial design is completed. The manufacturing tolerance and quality will also affect performance and durability of a plate. Hence, material develop-ment and selection of the plate should consider the requirements and limita-tions of the design and manufacturing processes.

In this chapter, we will pay attention to the basic or common materials requirements of the plate according to its functions in fuel cells. The empha-sis will be put on plate materials used in transportation fuel cells because these applications, more directly for automotive, have potentially the largest market for fuel cells and the related material requirements are most chal-lenging [1]. The various plate materials, fabrication process, and major chal-lenges will be introduced and analyzed. The underlying mechanism and development trends will also be discussed.


5.2 Functions, Structures, and Performance Requirements of Bipolar Plates

The bipolar plate with multiple functions, also called a flow field plate or sepa-ration plate (separator), is one of the core components in fuel cells. In reality, like serially linked batteries, fuel cells are a serial connection or stacking of fuel cell units, or so-called unit cells; this is why fuel cells are normally also called stacks (Figure 5.1) [2]. The complicated large fuel cells or module can consist of a couple of serially connected simple fuel cells or cell rows. Except for the special unit cells at two ends of a simple stack or cell row, all the other unit cells have the same structure, shape, and functions.

The unit cell mainly consists of a membrane electrode assembly (MEA), plates, and seals. The MEA includes polymeric electrolyte (membrane) with catalyst and electrode (normally carbon powders) layers on each side sur-face (Figure 5.1). Due to the weak nature of the membrane and electrode/catalyst layers, they are supported by carbon papers as GDLs (not indicated in Figure 5.1). Depending on the composition and structure, the electrode functions as either anode or cathode in a unit cell. The core electrochemical reaction between the hydrogen (in hydrogen fuel cells) and air is processed in the unit cell to generate the electron flow from the anode to cathode in the external circuit. This forms output electrical voltage, current, and power, as introduced in the previous chapters.

Single Cell

Components

Gas Flow

Bipolar Plate

Gas Flow

Bipolar Plate

Multi-cell Stack

Electrode/CatalystPolymeric

ElectrolyteElectrode/Catalyst

FIGURE 5.1Schematic indication of unit (single) cell components, including bipolar plates, and the rela-tionship between the unit cell and cell row or multiple-cell stack. (Reprinted with permission from Costamagna, P. and S. Srinivasan. 2001. Journal of Power Sources 102:242–252.)


As shown in Figure 5.1, the plates sandwich with unit cell components and play a critical role to assure and support the function of the unit cell and whole fuel cells (explained later). The plate directly contacts the GDL (nor-mally made of carbon paper) in either the anode or cathode side of a unit cell. The plate normally occupies a majority of the weight; a large cost in a stack depends on the type and specific design of the fuel cells. It has been esti-mated that plates account for about 60–80% of the total weight and 30–45% of the cost of a stack [3,4]. Development of new design, material, and fabrication process of the bipolar plates to meet the requirements of performance, dura-bility, manufacturability, and cost would make significant contributions to improving energy density and efficiency and reduce the cost of a stack. This is one of the major challenges in leveraging fuel cells to commercialization.

5.2.1 Basic Structure of a Bipolar Plate and Its Role in a Stack

Figure 5.2 schematically exhibits the structure and reactant flow of a simpli-fied stack designed by Bac 2 Conductive Composites Inc. [5] that contains three unit cells. Some components, such as GDLs, are not shown in the sim-plified diagram. Each unit cell includes an MEA and a plate (the anode plate, cathode plate, and coolant plate are not differentiated). Gas flow channels or fluid fields are on the surface of each plate.

In real fuel cells, the plate normally includes anode plate, cathode plate, coolant plate (if not integrated with anode and cathode plate), and end plate. The anode plate and cathode plate directly contact the anode side and cath-ode side, respectively, of an MEA or a membrane through the corresponding GDL, as explained before. The closed fluid flow channels or fluid flow fields

Membrane Electrode Assembly

Repeat Unit

Gas Flow Channels

H2

H2

O2

O2

ElectroPhenEndplate

ElectroPhenBipolar Plate

FIGURE 5.2Schematic indication of bipolar plates in simplified PEM fuel cells. The bipolar plates and end plates (ElectroPhen ) were designed by Bac 2 Conductive Composites Inc. (http://www.bac2.co.uk/fuel-cell-applications/ (accessed Dec. 2008).)


are formed when the plate is compressed to the GDL on assembling. The reactant (fuel and oxidant) under a certain pressure flows through the flow field or flow channels from inlet port at one end of a plate to the outlet port at the other end of the plate. There is a seal in each port to avoid leaking of fluids. Throughout the process, the reactant diffuses across the GDL to sup-ply the gas to each electrode uniformly and assure the core electrochemical reaction in fuel cells for the production of electrical current.

Normally, an anode plate in a unit cell is bonded or welded together with a cathode plate in neighbor unit cells to form a single plate, as shown in Figure 5.2. This is why the plate is called a “bipolar” plate. In addition to the reactant fluids, the coolant flow is also provided and flows through the plate, either an independent coolant plate or integrated with anode and cathode plate (not shown in Figure 5.2). The integrated approach for coolant transmission is normally used to save volume and cost. The cool-ant channel in the integrated plate is formed on the inside surface of the bonded cathode plate and anode plate [6]. The major components in a cool-ant are a mixture of glycol and deionized water. The coolant flow, which crosses over its flow field, plays a vital role of thermal and water man-agement in fuel cells because large amounts of heat and water are pro-duced from the core electrochemical reactions in the anode and cathode, respectively.

The plate at the two ends of a cell row or stack is called the end plate and has a slightly different structure from that of normal bipolar plates in the stack. The end plate actually is a “single-polar” plate with only the fluid field on the inside surface contacting the anode or the cathode of the unit cell at either end of the stack. The outside surface of the end plate is flat with fluid ports as shown in Figure 5.2. The end plate normally contacts the other cell row or system as electrical and fluid input/output connec-tions. Because the end plate is normally made of the same material through similar processing to that of the bipolar plate in a stack, the bipolar plate and end plate will be called a plate hereafter in this chapter unless their differences are addressed.

Although a fluid always flows from the inlet port through the channels to the outlet port in each bipolar plate, there are various design patterns or styles of flow path or flow field to meet performance and durability require-ments [4]. In general, the flow field design should generate and maintain the stable and homogeneous fluid flow in the fluid field so that the reactant can be appropriately supplied to electrodes through sufficient diffusion or permeation across the GDL. Moreover, the generated heat and water in elec-trochemical reactions in electrodes should be quickly transferred out to keep the required temperature and humidity in unit cells.

For example, due to the inevitable friction and impact of the channel sur-face to fluid flows, particularly at the bending part of the flow channels, how to reduce fluid flow rate changes or fluid pressure drop from inlet to outlet has to be considered in the fluid field design. One key part of the flow field


that influences the pressure drop or fluid flow rate is the flow channels close to each port. The direction and cross-sectional area of the flow channel in this part change greatly and the flow rate variations are difficult to control. The similar channel design in the coolant plate determines the cooling effi-ciency, temperature distribution, and condensed water management in the unit cell. The major styles of fluid fields include parallel, interdigitated, ser-pentine, etc.

5.2.2 Functions and Performance Requirements of Bipolar Plates and Plate Materials

According to the structure, location, and role of the plate in fuel cells men-tioned earlier, it is clear that the full function of the bipolar plate would be very important for the electrochemical reactions, heat and water manage-ment, and electrical current and power transfer in a stack. The specific func-tions of bipolar plates include:

provide uniform and stable reactant (fuel and oxidant) fluid flow to electrodes through GDLs that meets the requirements of efficient and stable electrochemical reactions in unit cells;provide serial electrical connections between two neighboring unit cells and electrical output to another cell row in a module or to the system;deliver the required coolant flow, transfer the generated heat, and assure fast heating at the start stage, particularly the start in a cold environment;carry out the produced water in the cathode, provide and maintain the required humidity for the membrane, and on the dry anode side, to assure good management of both humidity and heat;structurally support and separate each fragile unit cell, including separation of the hydrogen and oxidant, and sustain the compres-sive loading in a cell row or stack; andseal fluids with port seals and MEA seals to avoid fluid leakage.

Although it is difficult to determine the quantitative requirements of plate and plate materials appropriately for various fuel cells and different applica-tions in a development phase, such a target would be helpful to direct the development effort and make necessary trade-offs. The cascaded perfor-mance requirement targets in 2010 and 2015 for bipolar plates of fuel cells in transportation applications were set by the U.S. DoE (Department of Energy) according to functions of the plate mentioned before and overall require-ments of performance, reliability, manufacturability, and cost of a stack, as shown in Table 5.1 [7]. The technical target in the DoE’s multiyear research, development, and demonstration plan has been popularly and worldwide


referred to in many published papers and presentations, development road maps, and funding applications. To a certain degree, the target can be utilized as a criterion in development and selection of the plate and plate materials.

In general, the requirements for plate and plate materials listed in Table 5.1 can be classified as electrical properties, corrosion resistance, mechanical properties, gas permeability or soundness, weight, and cost. Although the output power of a stack is determined by many factors, one simple interpre-tation of the 2010 DoE cost target is ~US$2 per plate, which has often been quoted as a convenient estimation [8]. In addition to the cost requirements, the other requirements are all linked either directly or indirectly to the func-tions of the plate mentioned earlier.

For example, the required lower bulk electrical resistance and surface con-tact resistance are directly related to reducing internal power consumption in fuel cells to achieve maximum power output. The requirements of high flexural strength and flexibility (ultimate strain) are important to assure no distortion of fluid fields and no crack in a plate sustained in the large com-pressive loading when each unit cell is assembled together as a stack. This is particularly important when the thickness of the plate becomes thinner and thinner (can be close to or less than 1 mm [9]) and the dimension of the fluid field becomes smaller and smaller. Whether it is elastic or plastic, the large

TABLE 5.1

DoE 2010/2015 Performance Targets of Bipolar Plates for Transportation Fuel Cellsa

Properties Units 2010 Target 2015 Target

Electrical conductivity (bulk) Siemens per centimeter >100 >100Electrical resistivityb Ohm-centimeter 0.01 0.01Flexural strengthc Megapascals >25 >25Flexural flexibility Percent (deflection at

midspan)3–5 3–5

Corrosion resistance Microamperes per square centimeter

<1 <1

Hydrogen permeation rate cm3 sec–1 cm–2 (80 C, 3 atm = < 0.1 mA/cm2)

<2 × 10–6 <2 × 10–6

Weight Kilogram per kilowatt <0.4 <0.4Costd U.S. dollars per kilowatt 5 3

a Skipped column of “2005 status.”b Includes contact resistance.c Developers have used ASTM C-651-91, “Standard Test Method for Flexural Strength of

Manufactured Carbon and Graphite Articles Using Four Point Loading at Room Temperature.”d Based on 2002 dollars and costs projected to high-volume production (500,000 stacks per

year).Source: U.S. Department of Energy. 2007. Technical targets: Bipolar plates. Multiyear research,

development and demonstration plan. http://www1.eere.energy.gov/hydrogenand-fuelcells/mypp/pdfs/fuel_cells.pdf (accessed Dec. 2008).


deformation of the fluid channels or the plate would cause a pressure drop and fluid sharing in the fluid flow. This directly results in flow rate reduction and less and/or inhomogeneous supply of the reactants or coolant. All these will cause stack performance or durability problems. This is particularly true because the bipolar plate is a highly inhomogeneous and anisotropic component due to its fluid field structure.

One requirement for the plate used in automotive fuel cells that is not listed in Table 5.1 is the small volume required to reach the high power density of the whole stack, which is important for achieving the high energy efficiency of the stacks applied in transportation. As mentioned earlier, for fuel cells applications in transportation or automotive working conditions (such as cycling and frequent stopping and starting), the requirement of the higher working temperature, more severe vibration, exposure to very hot or very cold ambient environments in different regions, and high power density is very aggressive in comparison with that of the fuel cells in the other applica-tions. However, the technical target listed in Table 5.1 is not only a guide for plates used in automotive markets but also is a good reference for the plate in fuel cells used in other markets.

As mentioned at the beginning of the chapter, the material and related manufacturing process are the basis of any component, including the plate. Except for weight and cost targets, which are counted at the plate com-ponent level, the other requirements in Table 5.1 can be directly applied for plate materials as major required properties. The additional and more specific requirements for plate materials have been suggested in many reference papers based on experiences in testing and applications. These include:

high thermal conductivity (>20 W/mK [10]);high purity or low content of volatile components (VOCs) and extractable components (EOCs);chemical and mechanical compatibility with GDLs;no significant property degradation in a stack working environ-ment for a long term (5,000 h in 2010 and 2015, refers to Table 3.4.3 in “Technical Targets for Transportation Fuel Cell Stacks” [7]);good manufacturability or formability to fabricate the plate with the fine fluid field and small tolerance (e.g., smaller than 0.05 mm [10] for mass production); andhigh surface smoothness.

The purpose of these requirements is to assure that the plate or plate mate-rials have the required performance, reliability, and low cost. Among them, the first item is related to thermal management of the fuel cells. The second and third items are important for component compatibility in unit cells. As mentioned in previous chapters, the membrane and catalyst are particularly


sensitive to contamination, including many types of EOCs or VOCs from the other unit cell components. The fourth item is durability requirements cascaded from transportation fuel cells. The fifth and sixth items are require-ments related to cost, manufacturability, and quality in mass production.

For example, the small tolerance and low surface roughness of the plate in manufacture are critical for assuring the high electrical contact conductivity, low fluid flow resistance, and low water holdup to meet performance require-ments of the plates. Moreover, to play the role of removing generated water in the cathode side—particularly to avoid flooding when the current density is high, the surface of the cathode plate may need hydrophobicity [11] so as to better adjust hydrophobic and hydrophilic properties of plate materials in cathode and anode plates. This area needs further study.

5.3 Traditionally Used Materials of Bipolar Plates and Major Technical Barriers

According to the requirements of bipolar plates and plate materials men-tioned in Table 5.1, graphite—particularly high-density graphite, including synthetic or natural graphite—was utilized initially to manufacture the plates [12]. This is mainly because graphite has a desirable combination of excellent corrosion resistance, high chemical stability or lack of poisoning components to affect the catalyst and membrane, and good electrical conduc-tivity in the fuel cell working environment. These advantages made graphite a good initial material with which to fabricate bipolar plates. The fluid flow field in the graphite plate was normally machined.

The typical graphite plate is made by POCO (an acronym for the Pure Oil Company), which is part of Entegris, Inc. [13]. The company is a leading mate-rial developer, specializing in premium graphites and carbides, along with many other advanced materials. POCO produces a variety of high-strength, fine-grained, isotropic graphites with consistent microstructures that can be used for plates. The most popular material is PyroCell, a nonporous graphite that does not allow leaking of the cell stack.

However, the graphite has inherent shortcomings in mechanical properties and formability due to its nature of microstructure. The pure graphite only contains the carbon. The unique properties and shortcomings of the graph-ite, largely different from the normal carbon material or diamond (also only containing the carbon), come from its special microstructure. The graphite has a layered structure of carbon atoms. The atoms bond together in the layer by strong covalent bonding with the short atomic distance; however, the connec-tion between layers is through weak van der Waals bonding with the larger dis-tance. Hence, the graphite has essentially weak flexural strength and is prone to fracture due to the special structure. This is why graphite plate can only be


formed by machining processes with very high costs and long time periods and why defects are easily produced.

Normally, the postmachining process, such as resin impregnation, has to be added to avoid gas permeation [14]. The material cost of the graphite also is high. These make graphite plate as expensive as ~US$100–200 per plate [12], which can account for ~60% of the stack cost [6]. In addition, the brittle and porous nature and lower strength restrict the minimum thickness of the bipolar plates to only about 5–6 mm, which seriously limits efforts to develop mass production and to reduce the volume and weight of fuel cells for many applications, particularly for transportation and portable applica-tions. In addition to mechanical properties, the bulk electrical conductivity of the graphite is not very high. For example, the metal has much higher electrical conductivity than that of the graphite (see Table 5.2).

It is clear that the high cost and poor mechanical properties and formability are major technical barriers for the graphite plate to meet market requirements. Significant efforts have been made to develop alternative bipolar plate materials for overcoming the barrier and reaching the other technical requirements, together with development of plate design and the fabrication process. Nevertheless, the performance of graphite bipolar plates, such as POCO graphite plates, has been taken as a benchmark in the development of alternative materials.

5.4 Progress and Challenges in the Development of Bipolar Plate Materials

Many alternative materials have been investigated to replace graphite in the fabrication of bipolar plates. The major candidate materials with poten-tial to overcome the technical barriers and reach the targets mentioned in

TABLE 5.2

Comparison of Major Properties between SS316 and Graphite

Properties Density ThicknessElectrical

ResistivityCorrosion Current

Permeability for Hydrogen

Modulus of

Elasticity

Unit Gram per centimeter

Millimeter Siemens per centimeter

Milliampere per square centimeter

Cubic centimeter per square centimeter per second

Megapascal

SS316 8.02 0.16 5 × 106 <0.1 <10–12 193,000

Graphite 2.25 2.50–4.00 1,000 <0.01 10–2 ~ 10–6 4,800

Source: Reprinted with permission from A. Kumar and R. G. Reddy. In Fundamentals of advanced materials and energy conversion proceedings, ed. D. Chandra and R. G. Bautista, 41–53. Seattle, WA: TMS.


Section 5.2.2 include composites and metals. From a cost reduction point of view, it is estimated, according to the cost model, that the cost percentage of the plate in a stack can be reduced from ~60 to 15–29% if the graphite plate were replaced by the composite plate or metal plate [15]. However, many uncertain factors are involved in the estimation. The progress and major challenges in development of bipolar plates fabricated by these candidate materials will be introduced in the following part of this section.

5.4.1 Composite Plates

The composite as a unique material normally refers to hybrid or mixed mate-rials between dispersed filler or reinforcement in the form of fiber, powder, flake, etc. and the continuous matrix. The composite applied to manufac-ture the plate mainly belongs to the type of nonmetallic composite with both nonmetallic filler and matrix, although research work on composite plates with metal filler has been carried out. This research will not be specifically introduced here because no promising results have been reported so far. The other sandwiched composite plate consists of layered metals and thermal expanded graphite and was developed by Russian scientists [11]; however, no technical details were released.

In nonmetallic composite, the filler is normally stronger and harder or has special functional properties. The matrix is relatively ductile and tough. The nonmetallic composite used to fabricate the plates introduced here normally contains a large fraction (up to ~70–80%) of graphite fillers (or carbon fill-ers in some cases, such as that in carbon/carbon composite) and a relatively smaller fraction of the nongraphite polymer matrix. The continuously dis-tributed polymer in the composite plate plays the role of a binder to bond the fillers together; this is slightly different from that in the conventional com-posites with the larger fraction of the polymer matrix. Nevertheless, for con-venient description and comparison of different composite plates with the various polymers in this chapter, “matrix” and “filler” as conventional terms in the composite material have been utilized and will be referring to polymer materials and graphite, respectively. According to the definition, the non-metal composite plate will be called polymer (thermoset or thermoplastic)-“based” composite plates.

The function of the graphite (or carbon) filler in the composite plate is mainly to assure the required electrical conductivity, mechanical strength, and chemical stability. To this end, the filler in the composite plate plays the same function as that in graphite plates. However, the composite plate can have different matrix materials. According to the difference of matrix materials, the composite as plate materials mainly includes carbon/carbon composite (actually, carbon/graphite; see the following section), thermoplas-tic-based graphite composite, and thermoset-based graphite composite.

In all composite plates, a certain percentage of the graphite in the traditional graphite plate is replaced by the cheaper nongraphite matrix material so that


the plate cost can be significantly reduced. The ductile and tough matrix (such as thermoplastic and thermoset) in the composite improved the mechanical properties and forming capability of the graphite. Hence, the more robust composite plates with smaller thickness and finer flow fields can be manufac-tured by a more efficient process than expensive and slow machining.

Additionally, the nature of the composite provides the flexibility to tailor the properties of composite plates through selecting different combinations of filler and matrix material and different content ratios between the filler and matrix so as to meet various requirements and priorities for different fuel cell applications. The normal manufacturing process [11] of the poly-mer-based composite plates includes mixing and forming (e.g., molding) of graphite fillers together with polymer powders/flakes or impregnation of graphite plates with a monomer or an oligomer. This is followed by polym-erization to make the plate gas impermeable.

5.4.1.1 Carbon/Carbon Composite Plates

It is easy to be confused about carbon/carbon composite and thermoset-based graphite composites. By conventional definition, carbon/carbon composite refers to the composite with carbon fiber filler (which may contain the graph-ite filler as well) in a carbon or graphite matrix after the manufacturing pro-cess is completed. However, the initial matrix material in the composite is in the form of filler-compatible thermosets (e.g., phelonics, epoxies). The poly-mer matrix will become carbon or graphite through the pyrolysis reaction at a very high temperature in the fabrication process. Because the matrix in the manufactured composite plate is in the form of carbon or graphite, it is better to differentiate the carbon/carbon composite plate from the thermoset-based graphite composite plates (See next section).

In addition to the advantages of the composite plate over the traditional graphite plate mentioned before, the carbon/carbon composite plates have the advantage of lower density (about 30% lower than the thermoset- or ther-moplastic-based composite plates [16]) and higher manufacturing efficiency. This offers the potential of continuous production in comparison with the machining process for graphite plates.

One typical example of carbon/carbon composite plates is that made by Oak Ridge National Laboratory (ORNL) in the United States [12]. The com-posite preform was fabricated by a slurry-molding process from the mixed slurry between short carbon fibers (graphite fibers were also added in some sample plates) and the phenolic resin. The mass ratio between fiber reinforce-ment and phenolic matrix is 4:3. The phenolic matrix improves the mechani-cal properties and dimensional stability of the plate. A subsequent vacuum molding process was utilized to fabricate composite plates and fluid fields with relatively high resolution (Figure 5.3, [11]).

The molded plate with the large volume of pores was further coated by chemical vapor-infiltrated (CVI) graphitic carbon at ~1,500 C with 5 kPa


pressure to form a hermetically sealed and highly conductive surface. The phenolic resin binder was pyrolyzed in the process at high temperature. The graphite coating on the surface and the carbon/graphite matrix of the plate assure a high electrical conductivity (200–300 S cm–1) different from that of the thermoset composite plate (see the following section). The plate also has low density (0.96 g/cm3), thin thickness (1.5–2.5 mm), and high biaxial flex-ural strength (about 175 MPa). The major performances are comparable to or better than that of traditional dense POCO graphite plates.

Porvair Fuel Cells Technology, a division of Porvair Advanced Materials Inc. (PAM), licensed the carbon/carbon composite plate technology in 2001 from ORNL [18]. Porvair worked together with UTC Fuel Cells Inc. and received funding from the DoE Hydrogen Program from 2003 to 2006. The company focused on further development of near-net or net-shape forming processing and scale-up of the technology from the lab level. It has worked on molding process development and optimization, plate microstructures and properties improvement (Figure 5.4, [19]), sealing process improvement, shrinkage control, and cost reduction.

The improved process reduced the cycle time from 3 min to less than 10 sec and showed excellent capability and controllability with the value of Cpk at 2.33. The ex situ testing of the plate exhibited an encouraging performance, which included high electrical conductivity (600 ~ 800 S/cm), high flexural strength (600 ~ 7,000 psi), and low hydrogen permeability (smaller than 2 10–6 cm2/sec)—all of which exceeded or reached DoE 2010 targets (see Table 5.1)

10 mm

FIGURE 5.3The carbon/carbon bipolar plate fabricated by U.S. Oak Ridge National Laboratory. (Besmann, T. M. et al. In U.S. National Laboratory R&D Meeting. http://www.pnl.gov/microcats/ott-review/ottmeeting/14-Besman.pdf (accessed Dec. 2008).)


[16]. The stack level in situ testing in simulated driving cycle conditions for 2,000 h indicated that the plates had good comparable performance with machined graphite plates.

Although carbon/carbon composite normally showed good performance and durability, the major issue for the composite plates is that cost is still high [6]. This is mainly related to its complicated manufacturing process and the use of the larger quantity of the expensive graphite. How to improve the manufacturing process for reduction of the overall cost of carbon/carbon plates is a big challenge.

5.4.1.2 Thermoset-Based Composite Plates

Thermoset-based graphite composite is one of the composite materials often used to fabricate bipolar plates. The major filler or reinforcement in the com-posite is graphite in a form of powder, flake, or fiber, with additions of carbon powder/fiber (mainly to reduce the cost).

The used graphite can be natural graphite, which normally has higher electrical conductivity but may contain various inclusions with different sizes. The natural graphite can also have inhomogeneous crystallinity and properties, depending on the control of mining, pulverizing, and filtering processes. On the other hand, the artificially synthesized graphite has been popularly used as the filler of composite plates. The synthesized graphite is produced from thermal treatment of carbon powders or compounds at extremely high temperatures (about 2,500–3,000 C), which has higher purity but may not be fully graphitized sometimes.

In addition to graphite, carbon filler has also been added in composite plates. The major difference between graphite and carbon fillers in the com-posite plate is that the latter is less electrically conductive but is cheaper than

Sealed Surface

Porous Interior

FIGURE 5.4Cross-section images with different magnifications of carbon/carbon plates made by Porvair Advanced Materials. The image on the left with lower magnification exhibits a landing and two half-channels in one plate. The image on the right with higher magnification shows the porous composite and dense sealed surface. (Haack, D. 2006. Annual merit review report, U.S. DoE Hydrogen Program. http://www.hydrogen.energy.gov/pdfs/review06/fc_33_haack.pdf (accessed Dec. 2008).)


graphite fillers. Hence, the relative content between graphite and carbon fill-ers influences the balance between performance and cost of the thermoset-based composite plates.

According to the functional requirements of the plate, the graphite filler normally has to form a percolation network to facilitate electron transforma-tion. This defines the minimum content of the graphite or that together with the possible carbon addition. Although the range of the ratio between the filler and matrix in the composite plates varies greatly, the typical graphite filler (or together with carbon fillers) content is about 50–80 wt% [6], corre-sponding to higher electrical conductivity larger than 10 S cm–1.

On the other hand, the matrix material in the composite plates should meet many critical requirements, including high chemical stability, high and low temperature stability, good chemical and thermal compatibility with the graphite or carbon filler, low moisture absorption, high purity and low EOC and VOC content, low gas permeability, good manufacturability, and low cost. The thermosets commonly used as the matrix in composite plates include epoxies, phenolics, and vinyl esters. Many companies in this area have expended great effort to develop a special grade of thermoset materials to meet the aggressive requirements of the plate.

GRAFCELL@ is a typical thermoset composite that consists of a large per-centage of the expanded or flexible exfoliated natural graphite filler and epoxy matrix made by GrafTech AET Inc. [20]. Natural graphite has high thermal conductivity and anisotropy characteristics in comparison with artificial graphite. The GRAFCELL composite has been utilized to make a bipolar plate (Figure 5.5); the competitive performance [6,21] is due to its high electrical conductivity and low contact resistance, high corrosion resistance, low gas permeability, high thermal conductivity (500 W/mK) and thermal diffusivity (33 times higher than that of stainless steel), ease in forming flow fields through embossing, and relatively low cost.

The major shortcoming of the GrafTech composite plate is its inherently lower flexural strength, mainly due to its high content of graphite fillers. This restricts plate thickness reduction and the increase of related volume power density. Due to the resulting inhomogeneous gas supply and high-pressure drop, the lower flexural strength can also affect the working stability and durability of the plate.

To overcome the disadvantage and develop the new thermoset based com-posite plates, GrafTech is working together with Case University and Ballard Power Systems, supported by U.S. DoE funding in 2007, to develop a new gen-eration of thermoset composite plate [22]. The major target for the 2-year proj-ect is development of next-generation automotive composite plates based on expanded graphite and resin capable of operation at 120 C—very challenging for the resin matrix. The specific goals include identifying new matrix mate-rials and better graphite, improvement in mechanical and other properties, reduction in plate thickness to 1.6 mm, and cost reduction to US$6 per kilowatt to meet requirements of commercialization and high-volume production.


5.4.1.3 Thermoplastic-Based Composite Plates

Thermoplastic-based composite plate is another major type of composite plate with the thermoplastic matrix. The major advantage of this type of plate is that the well-developed injection molding processing in the thermo-plastic industry can be used to manufacture the plates for saving cost and improving production efficiency.

Normal plastics, such as polypropylene (PP), with low prices and good abil-ity of the injection molding process can be used as matrix materials in com-posite plates. However, the lower working temperature (around 70–80 C), lower

FIGURE 5.5GRAFCELL@ plates (thermoset-based graphite composite plates) made by GrafTech AET. (GrafTech International. 2008. http://www.graftechaet.com/GRAFCELL/GRAFCELL-Products/ Flow-Field-Plates-(FFP).aspx (accessed Dec. 20, 2008).)


chemical resistance, and lower creep strength of the plastics restrict utilization of the plastic composite plates in many types of fuel cells working in critical con-ditions. The other plastics applied as plate matrix materials of fuel cells work-ing in aggressive conditions mainly belong to the category of high-performance thermoplastics. Examples include polyphenylene sulfide (PPS), polyvinylidene-fluoride (PVDF), and liquid crystal polymer (LCP). It was reported that the most notable plastics being investigated and developed for thermoplastic-based com-posite plates are PVDF and LCP, although the cost of PVDF is high [8].

Ticona Engineering Polymers, a business unit of Celanese Corporation, has designed and manufactured thermoplastic-based composite plates, including end plates [23]. The thermoplastics utilized in the plate include Vectra LCP and Fortron PPS. The plastics all have good mechanical proper-ties at relatively high temperature, excellent chemical corrosion resistance, and capability for high loading of graphite fillers to achieve the required electrical conductivity. As shown in Figure 5.6, the thermoplastic-based com-posite plates made by conventional injection molding exhibit nice features with high production efficiency. The plates have been installed into stacks and in situ testing has been conducted.

Bac2 Conductive Composites [5] is a British fuel cell company that designs and manufactures thermoplastic-based composite and composite plates with a trademark of Electrophen™, as schematically indicated in Figure 5.2. The unique characteristic of the composite plate is that its plastic matrix is elec-trically conductive and cured at room temperature. The inherently conduct-ing pathway is produced during the polymerization reaction of the plastics, although no further mechanism is released. The resulting conductivity can be further enhanced by adding conductive fillers. The raw materials used to fabricate the plastic composite are relatively cheaper and the composite can be easily molded into robust plates with enhanced flexibility and lower cost than those of similar alternative plates. The Electrophen technology has

FIGURE 5.6Injection molded thermoplastic bipolar plates made of Vectra LCP (liquid crystal polymer) and Fortron PPS fabricated by Ticona Engineering Polymers. (http://www.ticona.com/redesign/index/markets/innovation/fuel_cell.htm Ticona Engineering Polymers. 2008; accessed March 2008.)


patents pending in several countries and the composite has been said to be in the process of commercialized production; however, the company does not disclose the technical details.

Cho and his colleagues in the Fuel Cell Research Center, Korea Institute of Science Technology, developed polymer-based composites consisting of a very high content (~90%) of graphite powder, ~10% of unsaturated polymer (not named) powders, and a small amount of organic solvent and additive [15]. The 2 mm thick plate was prepared by a thermal compressing process. The single-cell stack was installed with the composite plate. A series of ex situ and in situ testing (80 C, 1 atm pressure), including durability testing for 500 h, has been conducted.

Figure 5.7 compares the performance before durability testing of the single-cell stack installed with a composite plate and a graphite plate, respectively. No significant difference was exhibited in the similar curves after 500 h durability testing [15], which showed a close trend. The results indicated that the composite plate, particularly composite B plate, showed promising ex situ properties, in situ performance, and low degradation in durability testing—all comparable with that of the traditional graphite plate.

In DoE–funded, thermoplastic-based composite plate projects, NanoSonic and Virginia Tech (VT) did not use injection molding processing. Instead, they developed a wet-lay process to fabricate a series of thermoplastic-based composite plates, such as PPS plates and PVDF plates [24]. The key step in the wet-lay process includes mixing thermoplastic fibers, graphite fibers, and carbon fibers together in water to make a homogeneous slurry, which was

1.2

1.0

0.8

0.6

0.4

0.2

0.0

0.8

0.6

0.4

Po

wer

Den

sity

/W c

m–

2

0.2

0.0

0.0 0.4 0.8


Vo

ltag

e/V

1.2 1.8

Graphite

Composite A

Composite B

FIGURE 5.7The initial performance (polarization and power density changes with the current density) of single cell stacks consisted of two composite plates and graphite plates, respectively. (Reprinted with permission from Cho, E. A. et al. 2004. Journal of Power Sources 125:178–182.)


dried to form sheets. The sheets were further hot-compression-molded into plates. No additional machining process was needed due to the high dimen-sional resolution of the process; therefore, the wet-lay process can be applied in continuous production in large scale with high efficiency and low cost.

The other difference compared to many thermoplastic composite plates is that both thermoplastic and reinforced graphite and carbon materials in NanoSonic/VT plates are in the form of fibers instead of powders; this con-tributes to an improvement in the mechanical properties of composite plates. In addition, due to the inherently high strength of PPS, the ratio of carbon and graphite fibers to the thermoplastic powder in PPS-based plates is very high (~76:24) to achieve the high electrical conductivity. Some performances of the plate reached or exceeded the DoE target for 2010. For example, the flexural strength and in-plane conductivity of the composite reached 96 MPa and 271 S/cm, respectively [24].

The excellent corrosion resistance of PPS is also helpful for meeting the requirements of chemical stability of the overall plates. In comparison with conventional graphite plates, the PPS-based plates have better mechanical properties and good corrosion resistance.

In addition to the companies just mentioned, many other companies manufacture commercialized polymer-based composite plates [12], includ-ing DuPont, H2Economy, ICM Plastics, Ned Stack, and SGL Carbon GmbH (SIGRAFLEX@). Identifying the best thermoplastic material, further improv-ing the process, and reducing the cost and conducting stack level testing will be the major tasks of the work in development of thermoplastic composite plates.

Overall, as mentioned previously, the major challenge in composite plates is how to maintain the best balance between electrical conductivity and mechanical properties. This is mainly influenced by the ratio between the fillers and matrix. Regarding the electrical conductivity, well designed and manufactured composite materials show performance comparable to that of graphite. For example, the bulk electrical conductivity of a well designed and fabricated composite can reach 300 S cm–1, close to that of the graphite [21].

However, if the volume ratio between the filler and the matrix is too high, the composite can become brittle, similar to dense graphite, and the advan-tage of composite plates in cost reduction will not exist. Hence, the normal composite plate contains about 50–80% fillers so as to have balanced electri-cal conductivity and mechanical properties with the accepted low cost.

In addition to the amount of filler content, the shape, size and size distribution, surface wettability, interface bonding, and compatibility with the matrix resin of the filler can all influence electrical conductivity, mechanical properties, and other performance characteristics of the composite plates. As mentioned previously, to achieve higher electrical conductivity, the conductive graphite or carbon fillers must form an interconnected or percolated network in the dielectrical matrix like that in GrafTech plates. The interface bonding and compatibility between


the filler and matrix directly affect the overall performance and durability of composite plates. The porosity of the matrix and interface bonding between the filler and matrix will influence the gas permeability, particularly the hydrogen permeability. All these materials-related parameters in the composite plate have to be well controlled to achieve better performance, durability, and lower cost.

In general, due to the simpler manufacturing process and the cheaper polymer matrix, the thermoset- or thermoplastic-based composite plates have lower cost than that of carbon/carbon composite plates. The thermoset-based composite and thermoplastic-based composite plates have their own advantages and limitations, which are largely determined by the properties of the matrix materials and the ratio between the filler and matrix.

One advantage of a thermoset composite plate over a thermoplastic one is its relatively short molding cycle time. Like the normal thermoplastics, the thermoplastic composite plates have to be cooled to a lower temperature before the part can be demolded. The thermoset composite plates normally have a short curing time at the curing temperature. The plate can be removed after curing from the mold even when it is very hot. These advantages are important to reducing cost and enhancing production efficiency, which com-prise one of the major compulsory targets that all fuel cell components must reach to meet requirements of commercialization.

The other difference [21] is that thermoset composite shows higher bal-anced strength and toughness compared to thermoplastic composite, with a little brittleness. Thus, thermoset composite plates can be fabricated rela-tively thinner and have more applications. On the other hand, as mentioned before, a well developed injection molding process has been applied as the fabrication process of thermoplastic-based composite plates to assure high quality and low cost.

5.4.2 Metal Plates

The metal plate has been newly developed and has attracted much atten-tion in recent years. For example, General Motors indicated recently [25] that metal plates had been selected and would be utilized in its newly designed fuel cells to be installed in commercialized vehicles soon.

Obviously, in terms of the base material, the metal plate differs greatly from the conventional graphite plate and improved composite plate, such as carbon composite, thermoset composite, and plastic composite. The major advantage of the metal plates is that the metal normally has much better electrical conductivity and formability in comparison with the graphite and composite materials. The metal does not have pores, so it has good hydro-gen and the other fluid impermeability. Due to the nature of metallic bond-ing, the metal has higher strength, modulus, toughness, and related shock resistance. Owing to special mechanical properties, the metal is more easily formed into the thinner plate and finer fluid field with the higher and more stable quality in mass production than graphite or composite plates.


These advantages are important to enhance the power density and durability of fuel cells applied in the automotive industry. The overall cost of the metal plates, such as SS316L plates, is competitive with that of the graphite and composite plates. Ease of recycling or reuse is another good point for metal plates, particularly when environmental protection is becoming a critical challenge for all countries in the world. Moreover, due to the long development history of various metals and their fabri-cation processing (including forming, welding, and corrosion-resistant coating processes), in the automotive, chemical, and many other indus-tries, various metals and technologies are available for fabrication of the required metal plates.

In addition to many advantages, the major shortcoming of the metals is less corrosion resistance and larger density than that of graphite and composites. Hence, the metal plates must be coated with a corrosion-resistant surface layer after forming and welding. Also, plate thickness must be reduced in the forming process to compensate for its high density so as to compete with conventional graphite and composite plates.

5.4.2.1 Types of Metals Used in Metal Plates

Researchers have tried to fabricate plates using many different metals—mainly, stainless steel, aluminum alloys, titanium alloys, nickel alloys, cop-per alloys, intermetallic alloys, and metal-based composites such as carbon fiber-reinforced aluminum alloys, carbon fiber reinforced copper alloys, etc. [26]. Although Ta, Hf, Nb, Zr, and Ti metals show good corrosion resistance and chemical stability [6], the cost of these metals is too high for them to be used as materials in metal plates. That is why relatively cheaper iron-based alloys, particularly stainless steel, have been popularly studied as plate mate-rial. In the following sections, we will introduce stainless steel (SS) and SS plates, which have been extensively investigated and show promise for the final applications [6,11].

“Stainless steel” normally refers to steel with a high content of nickel and chromium so that the steel has a dense chromium oxide protective layer show-ing strong corrosion resistance. The role of the nickel in stainless steel is mainly for avoiding formation of the brittle sigma phase and extending the austenite phase area to room temperature. The austenite phase is an iron-based solid solution with a face-centered cubic (F.C.C. or FCC) structure and good plastic-ity or formability. In alloy steel, the alloying metal atoms and carbon atoms are solutes with a certain solubility in the iron solid solution. According to requirements of shape formability and corrosion resistance mentioned before, the various austenite stainless steels can be used to make the plates. Among them, stainless steel 316L (SS316L), which contains high chromium (~16–18%), and nickel (~10–14%) have received more attention in recent years [6,11]. This is due to their high resistance to almost all types of corrosion, including pitting,


grain boundary, and contact corrosion in the normal working environment, and to their high strength.

The stable and dense chromium oxide passive layer formed on SS316L is the major reason for its high corrosion resistance and chemical stability. The letter “L” refers to the low carbon content (<0.03%), which is important to keep good capability of being welded and for maintaining high corro-sion resistance in the welding line with the small heat affected zone (HAZ). As shown in Table 5.2, in comparison with benchmark graphite, SS316 (the listed properties of SS316 are very close to those of SS316L) has much higher electrical conductivity and elastic modulus, much lower hydrogen permeability, and makes thinner thickness of the plate. Obviously, SS316 has higher density and lower corrosion resistance. In addition to austenite stainless steel, such as SS316L, the duplex phase (austenite/ferrite) stainless steel (e.g., SS2205) has also shown a good combination of high corrosion resistance and strength, which have been considered the base materials of metal plates [27].

However, the currently used polymer membrane in PEMFCs essentially has a strong acidic nature, as introduced in the previous chapter. The affected fluids circulated in plates, membrane, and other components are the major source resulting in chemical corrosion of the plate. The severe electrochemi-cal corrosion around the plate may link to the inevitable electrical potential difference between the GDL and corresponding plate, the strong flow flu-ids, and the electrochemical reaction environment in each cell. The elevated working temperature and humid atmosphere in the fuel cells accelerate the corrosion. Much testing indicates that even a stainless steel plate, such as the SS316L plate, cannot get rid of surface damages in the superimposed chemi-cal corrosion and electrochemical corrosion in the fuel cell working condi-tions for a long time.

Once the plate starts to corrode, many problems appear to affect perfor-mance and durability, even serious failure, of the fuel cells. For example, the interface contact resistance between the corroded metal plates and GDL will increase to reduce the power output. The corrosion products (mainly various cations) will contaminate the catalyst and membrane and affect their nor-mal functions because the polymer membrane essentially is a strong cation exchanger and the catalyst is susceptible to the ion impurity. Hence, adding a corrosion-resistant coating to the metal plate will almost inevitably assure the performance and long-term durability of a stack.

However, some efforts [27] have been made to develop noncoated duplex-phase stainless steel with high chromium content, such as 2205 steel, for application as metal plate materials and getting rid of the expensive coat-ing. The passive layer in the steel was formed quickly in the fuel cell work-ing environment, but no details of in situ testing were reported. In addition, higher electrical interface contact resistance (ICR) of the passive layer in the duplex-phase stainless steel is an issue for the application.


5.4.2.2 Key Fabrication Processing

Fabrication processing of metal plates is different from that for the graphite and composite plates mentioned in previous sections. The major manufactur-ing process of metal plates includes forming the thin plate with the required fluid fields and coating the corrosion-resistant layer on the plate surface. For the stack with integrated coolant plates, the anode plate and cathode plate need to be welded together and the coolant channels formed between two plates. Hence, the welding process normally is also the processing necessary before coating.

5.4.2.2.1 Forming Process

As mentioned previously, to compensate for the high density and reach the target of the overall weight of the plate and power density of a stack, the metal plate must be very thin: down to ~1 mm or even 0.1 mm [9,36]. To shape the complicated fluid field and fluid port structure with high accuracy and limited tolerance on such thin metal plates is a big challenge. That is why there is a strict plasticity requirement for metals used to make the plate (flex-ural flexibility of 3–5%; see Table 5.1). Depending on the type or plasticity of the metals, the commonly used forming processes include compression molding, stamping, and embossing. [28]. Design of fluid field details, design and manufacture of the forming mold, and selection of forming process parameters such as loading control and deformation ratio can all influence the quality of the formed plate. Warping, distortion, and cracks are common defects that occur during the process.

Figure 5.8 shows schematically a design architecture concept of metal plates [29]. The thin anode and cathode plates were stamped to form a hydro-gen flow field and air/oxygen flow field, respectively. The coolant flow field was formed simultaneously and the closed channel was generated when the anode plate and cathode plate were bonded together. The cross-section shape of the flow field or flow channel varies depending on the required flow sup-ply in the specific fuel cells.

5.4.2.2.2 Coating Materials and Coating ProcessCoating Materials. In addition to the basic requirements of high chemical and electrochemical stability to work in the harsh environment of unit cells, the coating layer or material should meet aggressive requirements to assure its performance and durability. These include thermal and chemical compat-ibility with substrate metal, high bonding or adhesion strength to the sub-strate, low electrical ICR with GDL, high purity, low VOC and EOC, low cost, and ease of manufacture in mass production. According to the men-tioned function of the plate, it is very important that the coating have low ICR. Appropriately selected coating compounds, such as transition metal nitrides, borides, and oxides, could have the ICR with the carbon GDL when the coating layer thickness is very thin.


The underlying mechanism has not been thoroughly understood. The thermal and mechanical properties between the ceramic coating and metal substrate normally have larger differences due to their different chemical bonding nature. Understanding how to keep and maintain the desired ther-mal and mechanical compatibility and strong bonding between the coating layer and substrate is crucial for the durability of the coating. This is par-ticularly so when the plate works in aggressive cycling conditions of both temperature and moisture. Owing to the limit of plate thickness or stack dimension, the coating layer must be very thin—down to micrometers or even thinner; this creates more challenges for a reliable coating.

According to previous general requirements, various coating materials have been studied. They mainly include different ceramic compounds and the inert metal or carbon [3,9,11]. Due to strong covalent chemical bonding, the ceramic compounds, such as nitrides, borides, and oxides, have high chemical stability and corrosion resistance compared with stainless steel. The specific ceramic compounds utilized as coating materials of metal plates will be introduced later, together with introduction of the coating processing.

Because the coating on anode plates must contact the hydrogen at a certain temperature with high humidity, the coating compounds should also have good stability for reduction. The stability of carbides was found to be not very good in the hydrogen reduction atmosphere, so carbides are normally not selected as the coating on metal plates. However, stable or inert graph-ite and diamond-like carbon have been utilized as coating material for the

Liquid

Adhesive

Bonding

H2

O2 or Air

FIGURE 5.8Architecture of stamped metal plate concept with MEAs. (Reprinted with permission from Granier, J. 2004–2005. CLEFS CEA 50/51: 76–78.)


plate [3]. The inert metals, such as gold and silver with good chemical stabil-ity and good formability, have also been investigated. The additional advan-tage for both inert metals and carbon materials is that they all have excellent electrical conductivity and low electrical ICR with carbon GDLs. However, one apparent challenge for the utilization of the inert metals, diamond-like carbon, and graphite is their high cost. In addition, the conductive polymer [3] has also been tested as a metal plate coating material, although no appli-cation report has been released so far.

Coating Processing. A certain pretreatment process, such as surface cleaning activation, should be required to achieve better bonding strength between the coating layer and plate substrate. There are critical require-ments, such as uniform thickness with low tolerance, high surface smooth-ness, no pinholes, and no inclusions, for coating and coating process. The coating quality influences not only corrosion resistance but also electrical contact with the GDL and fluid flows in channels. The commonly utilized coating processing in metal plates mainly includes immersion coating, CVD (chemical vapor deposition), PVD (physical vapor deposition), electroplat-ing, and electroless plating [28].

Sputtering and evaporation in high-vacuum and inert-gas-protective con-ditions are typical examples of PVD processing used for metal plate coat-ing. The advantage of PVD processing includes high dimensional accuracy, fewer coating defects, and good surface finishes. However, the PVD coating process normally involves high temperatures, so it can easily result in resi-due stress and cause distortion or warping of very thin plates. The cost of a PVD coating is also relatively high.

The nitridation and boronization processes are examples of CVD process-ing used for metal plate coating, which will be introduced later. A plating process operated at lower temperature and with lower cost has been uti-lized in the surface coating of metal plates. The major disadvantage of the plating process is that the coated layer is less dense and can easily have more defects, such as pinholes, than that of PVD coatings. In addition to coating material cost, the coating process time is one of the major factors in determining the overall cost of the coating. How to increase the deposition rate or reduce the coating time remains a challenge in the coating process for metal plates.

Nitridation, a surface coating process to coat nitrides, such as TiN and CrN, on metals, has attracted more attention for applications in metal plates in recent years [9]. Either through the conventional chemical heat treatment process or PVD process to coat a hard surface layer on metal components, the nitriding process has been popularly used in molds, tools, and many other products. The nitride layer on the component surface is formed through the reaction between nitrogen-containing gas and a certain alloying element, such as chromium, molebdum, vanadium, on the surface of metal com-ponents in a relatively low temperature range (~400–600 C). Although the nitride layer is normally thinner than 1 mm, it has good corrosion resistance,


high bonding strength with the metal substrate, and high density or fewer voids. This is why the process has been utilized in coating metal plates [30].

Boronization is another similar chemical heat treatment process to form a boride layer on the metal component surface through reaction between the boron-containing gas or powder/liquid medium and metal element (e.g., iron, titanium, chromium) on the component surface. The boride has high corrosion resistance and relatively high hardness, but a brittle interface between the boride layer and steel substrate can appear if control of the pro-cess is not appropriate. In addition, the boronization process temperature is slightly higher (~500–700 C) and, after boronization, the thin metal plate can be distorted or deformed. The boronization processing has also been applied to metal plate surface coating. For example, powder-pack process-ing has been utilized on the metal plate to form a Ni3B layer on a nickel-clad steel plate surface [31]. The boride layer dramatically increased the corrosion resistance of the plate, although the layer thickness was not homogeneous. Further development of the processing is needed.

In addition to applying the coating process on metal plate surfaces to meet the critical requirement of corrosion resistance, a conventional cladding pro-cess has also been utilized in the fabrication of metal plates for the same purpose [8,30]. The cladding is a mature process for fabrication of multilayer metal plates or sheets; it combines a mechanical compression process and metallurgical process to add to thin layers of different metals to the matrix metal plate or sheet. According to the requirement, the thin metal foil may be cladded on one side or each side of the matrix sheet. The very thin top (and bottom) layer in the multiple-layer metals is normally made of a softer metal (e.g., Nb, Au) with very good plasticity and special physical properties, such as corrosion resistance, which the matrix plate does not have. The clad-ding process is conventionally done by co-rolling a thin metal foil or sheet and a matrix plate or sheet. The very high compression force and possibly added heating promote the formation of good bonding—not only van der Waals bonding but also possibly metallurgical bonding—between the layers of metals.

Weil et al. [31] have successfully applied a cladding process to fabricating Nb/SS430/Nb and other metal plates. They down-selected niobium as the cladding metal from many other candidate metals for its corrosion resistance, surface electrical contact resistance, formability, and cost. One of their tested plates was fabricated by the roll cladding process. The single anode or cathode plate was stamped to form the complex fluid field. The anode and cathode plates were brazed together to form coolant channels between two plates. The plate consisted of an SS430 core substrate with a thickness of 450 μm; a niobium cladding layer of only 50 μm thickness was on the top and bottom sides, so the total thickness of the clad plate was 0.1 mm.

These researchers also tried nickel foil cladding on the SS430 to make the plate before the boronization process was applied to add a protective coating to the nickel surface for improvement of the corrosion resistance. The major


tested or estimated properties of Nb-clad SS430 plates and boronized Ni-clad SS430 plates in comparison with the DoE 2010 target are listed in Table 5.3. It can be seen that the properties of both clad metal plates are over or close to DoE 2010 targets. It is interesting to notice that the boronized Ni-clad SS430 plates have higher electrical conductivity and 10% lower cost than that of Nb-clad SS430 plates. More results of in situ testing and durability of the plate would be expected.

5.4.2.2.3 Welding

The other major fabrication process of the metal plate is welding: The anode and cathode plates are normally welded together to form a single plate with the integrated coolant channels between two plates. Due to the very thin thickness of the metal sheets, the overheating and longer heating time in normal welding processing will generate degradation of metal properties, large residue stress, and distortion in the weld line and closed HAZ in the plate. Using spot welding with very small beam size and high welding speed, the solid-state optical fiber laser has been utilized in metal plate welding to mitigate these problems. One example is the YLP-100 single-mode ytterbium fiber laser manufactured by IPG Photonics, which has a narrow beam and

TABLE 5.3

Properties and Estimated Cost of Nb-Clad SS430 Plates and Ni-Clad SS430 Plates with Boronized Layer Made by U.S. Pacific Northwest National Lab to Compare with U.S. DoE Bipolar Plate Technical Targets

Characteristic 2010 Target Nb-clad Steel Ni-clad Steela

Cost ($/kg) 6 5.0c 4.5d

Weight (kg/kW) <1 0.7c 0.72d

H2 permeation rateb <2 × 10-6 0.35 × 10–7 c 0.37 × 10–7 d

Corrosion (mA/cm2) <1 <1e TBDf

Conductivity (S/cm) >100 0.6 × 105 0.8 × 105

Resistivity ( .cm2) 0.01 <0.01g <0.01Flexural strength (MPa) >4 Expected to meetf Expected to meetf

Flexibility (% at midspan) 3–5 Expected to meetf Expected to meetf

Source: Weil, K. S. Annual progress report, U.S. DoE Hydrogen Program. http://www.hydrogen.energy.gov/pdfs/progress06/v_d_3_weil.pdf (accessed Dec. 2008).

a Boronized.b In cm3/sec.cm2 at 80 C and 3 atm.c Assumes a <20 μm thick Nb cladding layer over a 130 μm steel core. Cost informa-

tion is based on discussions with ATI Wah Chang, a leading worldwide producer of Nb sheet product.

d Assumes a 40 μm thick cladding layer over 110 μm steel core.e Anticipated based on preliminary screening tests.f Yet to be determined.g At clamping pressures > 1.5 MPa.


deep penetration. The laser is said to be utilized in welding the thin SS316L plates in fuel cells and in other applications [32].

The good weldability of the metals (such as the previously mentioned 316L) used to fabricate the plates is necessary to achieve a high-quality welding process. The SS316L has good weldability, mainly due to its lower carbon content. The major alloying element, chromium, responsible for cor-rosion resistance in stainless steel can easily form carbides, instead of oxides, during the melting from chromium oxide and resolidification on welding. The process will diminish the corrosion resistance in the welding area due to a reduction in chromium content in the iron solid solution, which was attracted into the carbides. The lower carbon content will reduce the forma-tion of carbide so as to reduce localized corrosion tendencies resulting from welding. The common defects, such as voids, cracks, and distortion, should be avoided in the welding process.

As mentioned earlier, in addition to SS316L, many other metals have been investigated as metal plate candidates. Overall, SS316L still showed better performance and durability and lower cost, according to available reports so far. Andrukaitis recently conducted ex situ and in situ testing on a series of metal plates, including three stainless steels (SS316L, SS420, and SS904L), Al6061, titanium alloy (titanium grade 2), and one unknown alloy (GTI alloy) plates with different types of coatings, the nature of which were not released [8]. He found that the SS904L plate coated by Ineos Chlor FC7 showed the best performance of all the metal plates and comparable performance with that of baseline graphite plates, as well as immeasurably low voltage deg-radation rates. However, the cost of the stainless steel and coating is higher than that of the graphite plate.

Alternatively, the SS316L plates with Acheson Colloids Electrodag EB008 coating have a combination of lower cost than that of graphite plates and an acceptably low voltage degradation rate of 22 μV/h. The information on metal plate dimension, coating thickness, stack configuration, and test con-dition details for these results was not reported in the presentation.

5.5 New Activities and Development Trends

Due to increasing global warming and energy resource shortages, R&D work has increasingly been focused on developing fuel cells, including the plate as one of the key components of the cells. The market is demanding many efforts, the scale-up from such as lab level or pilot production to near-term, small-scale production. Even mass production of fuel cells in the near future has been made, particularly by many automotive manufacturers worldwide. The major new activities and related development trends according to avail-able information are briefly summarized next.


5.5.1 Development of New Materials with Better Performance and Durability

To achieve the goal of required performance, durability, and cost of plate materials, one approach is improvement of the control of the composition and microstructure of materials, particularly the composite, in the material designing and manufacturing process. For example, in the direction of devel-opment of thermoplastics-based composite plate, CEA (Le Ripault Center) and Atofina (Total Group) have jointly worked on an innovative “microcom-posite” material [33]. The small powders of the graphite platelet filler and the PVDF matrix were mixed homogeneously by the dispersion method. The filler and matrix had a certain ratio at the microlevel in the powder accord-ing to the optimized properties requirements. The microcomposite powders were thermocompressed into the composite plate.

In addition, the well-controlled geometry and structure of microcomposite particulates are advantageous for a more efficient forming process, such as an extrusion or extrusion–compression (extrusion and subsequent brief compres-sion) process. With this processing, the graphite filler up to 80% by volume can be uniformly added to the composite plate with the attractive properties, including the high electrical conductivity of 250–300 S/cm, high thermal con-ductivity of 110 W/m/K, high ultimate bending strength of 35 MPa, and low gas permeability.

How to reduce the cost yet achieve even better performance and durability of the plate and its materials is a big challenge. In addition to many activities on composite materials and related plates mentioned earlier, a lot of effort has focused on further developing coating, substrate materials, and coating process-ing of metal plates to face the challenge. Brady et al. [34] at ORNL cooperated with ATI Allegheny Ludlum Corp. (to cast and develop the matrix alloy) and GenCell Corp. (to stamp the alloy foil for subsequent nitridation) to develop Cr nitrided (such as Ni-Cr, Fe-Cr) protective and Fr-Cr-V stamped metallic plates. The DoE agreed to support this financially from 2007 to 2010.

The major nitrides on the plate surface include CrN2 and CrN, which have small grain size and uniform distribution on the SS316L steel matrix. Brady and colleagues successfully fabricated nitrided thin metal plates and con-ducted single-cell drive-cycle durability testing for a time up to 1,160 h. The metal plate did not show performance degradation and corrosion attack on the coated dense layer, but specific test conditions were not described. To reduce nitrided plate cost, they modified the processes and matrix alloy compositions to prepare nitrided Fe-Cr and Fe-Cr-V alloys, replacing the expensive nitrided SS316L in the metal plates. A process of preoxidation and segregation of vanadium to the surface was found favorable to formation of nitrides on the plate surface in Fe-27Cr-6V alloy. The addition of vanadium also helped to improve corrosion resistance of Cr-nitride coating. Particularly, they found that nitriding the Fe-Cr-V alloy plates can significantly reduce the electrical surface resistance.


These researchers plan to optimize the bulk alloy system further and improve nitriding and stamping processes to fabricate nitrided metal plates with a thick-ness not larger than 0.1 mm. The plate is expected to have the equivalent per-formance to that of graphite plates, 5,000 h working lifetime durability, and the DoE required a cost of US$5 per kilowatt. In addition, the concept of the cheaper and lighter metal plates made of aluminum or aluminum–magnesium alloys and coated by graphite-based conductive polymer has been patented [11].

Together with the new design concept, the new special material may make a big jump in the development route of the plate and plate materials. For exam-ple, a foam or porous material with open pores, such as stainless-steel foam, Ni-Cr foam, or carbon cloth, was inserted between the flat SS316L plate and GDL [35,36]. The foam replaced the complicated flow field and showed better functions of reactant gas accessing the electrodes and enhanced the perfor-mance in initial single-cell testing. The three-dimensional foam increased the fluid flow accessible surface area from 60 to 70% of the conventional channel structure to greater than 90%; this can increase the homogeneity of the fluid flow supply and related local current density. The porous material also has the potential to reduce the weight of the plate and to develop catalyst-bearing surfaces. The optimization of the structure, shape, and dimensions of the pores in the foam may further enhance the function of the new concept.

The major issue found in testing is the corrosion of the foam material and resultant contamination of the membrane. The high manufacturing cost of the metal or carbon foam with the required pore shape, size, and distribu-tion also is a challenge. Further study and testing of the corrosion mecha-nism, selection of appropriate coating, a capillary process involved in the tiny pores, and related water retention are necessary to identify whether the new material and concept can be finally applied in the plate.

5.5.2 Development of New Manufacturing Processes with Higher Efficiency and Lower Cost

The manufacture of plate materials and plates is a vital factor in reaching the final goal, particularly the cost goal, for the plate. The forming process of very thin metal foil with a complicated flow field is one such challenge for metal plate development. One example of taking the challenge in this direction is American Trim’s development of a new, efficient, high-velocity electromagnetic forming (HVEF) process for thin metal foils [37]. Together with GM and the Department of Materials Science and Engineering at Ohio State University, in April 2007, the company was awarded US$1 million by the Ohio State Department of Development for further development of processing and applications in manu-facturing the metal plates. In this cooperative project, American Trim planned to develop its new forming process to produce metal plates with the required flow field features: quality and cost. The company expected that achievement in the project could go for manufacture of commercial quantities of plates to support a full production line. In April 2009, the company announced it had


completed requirements for the project supported by the state grant and had demonstrated metal plate production using HVEF technology. No details of the new HVEF process and its cost reduction range have been released.

In addition, the popularly utilized and efficient microelectronic manufac-turing process, including excimer laser lithography processing, was reported to apply in the manufacture of portable fuel cells and flow fields of plates in normal fuel cells [11]. Although the processing is complicated, it is believed that the manufacturing cost will be extremely low when the volume becomes high, according to experience in the electronic product industry. Similarly, the efficient magnetron sputtering process was utilized to fabricate plates with sil-ver or copper coating on a plastic billet, which has the structure of a polymer/metal sandwich.

5.5.3 Strong Competition between Different Candidate Materials and Processing

As introduced in this chapter, there are many candidate materials and cor-responding manufacturing processes to produce plates. The fuel cells and their components, such as the plate, have not reached the stage of commer-cialized production and all in situ testing conducted so far is still within a limited time scale; thus, it is difficult to judge which material and corre-sponding manufacturing process is the best for plates—even for a fuel cell in a known application. However, the potentially huge market for fuel cells is very attractive, so competition among the different materials, manufactur-ing processes, and plate design concepts is very strong.

The competition among different types of plates has mainly focused on composite plates and metal plates, the sign of which can be seen from DoE-awarded plate-related projects in 2007 [38]. That year, the DoE Hydrogen Program awarded three plate-related projects, including one project on metal plates and coating process, and two projects on composite plates (one for thermoplastic-based plates and another for carbon/carbon plates with focus on the manufacturing process). Strong competition on plate materials and manufacturing process would be helpful for accelerated development of materials, customized fluid field design, and manufacturing technology of the plate to reach the goal of required performance, durability, and cost. For example, in their research report, Carlson, Garland, and Sutton [9] pointed out that the available thickness of graphite plates (most likely referring to thermoset composite plates with graphite fillers, although they were not clear about this) is close to that of metal plates and the overall cost is even lower in comparison with SS316-based metal plates.

5.5.4 Scale-up from Lab-Scale to Large-Quantity Production

To match the road map of the overall development of fuel cells and meet the strong market demand for clean energy resources, many fuel cell companies


and automotive manufacturers have made great effort to scale up R&D results or field trials to production with larger volume. In addition to cost and durability, how to make the manufacturing process of fuel cells, including the plates, capable of the mass production is a challenge. The conventional manufacturing process of the plates at the lab level or at small scale mainly consists of several separate and isolated steps, such as molding (composite plates) or forming (metal plates), bonding or welding, or resin encapsulation (composite plates) or coating (metal plates). The process is slow, inefficient, and expensive, and it is difficult to control the overall quality of the plate; this cannot be leveraged to mass production.

To be suitable for the commercialization requirements of fuel cells, more attention must be paid to the study of a continuous manufacturing process, which can be further developed as an automotive production line, if nec-essary. The roll- or coil-based metal foil continuous manufacturing line, including flatting, embossing (or the other forming process), cutting, weld-ing, and coating of metal plates, is one of the examples [28]. The improved manufacturing process will be effective for large reduction of the cost as well as strengthening quality control of the plates for mass production.

5.6 Summary

Bipolar plates in PEMFCs were conventionally made of graphite with excel-lent corrosion resistance, chemical stability, and high thermal conductivity. However, graphite has a high cost, poor mechanical properties, and very little formability due to its microstructural nature. This limits its further applications as plate material and forces a search for alternative solutions. Nevertheless, the performance, durability, and cost of the graphite plate (e.g., POCO graphite and graphite plates) have been taken as benchmark refer-ences to compare with those of alternative materials.

Great progress has been achieved in development of better plate materi-als and plates. The composite, which includes the carbon/carbon composite, thermoset-based composite, and thermoplastic-based composite, is one of the major alternative materials to graphite as plate material. The carbon/carbon composite plate has good performance and durability but higher cost. Strong efforts in improvement to the manufacturing process to reduce the overall cost are necessary. The thermoset-based and thermoplastic-based composite plates reduce use of expansive graphite and improve the poor mechanical properties and formability of graphite, which can be utilized for mass pro-duction of the plates. However, improvement of mechanical properties of the composite plates is limited by the small fraction of polymer materials com-promised by electrical conductivity requirements. This can influence further reduction of the dimension of the plate and flow field and the increase of plate power density, particularly that for applications in automotive industries.


Metal is another major candidate material for the plate. It has large advan-tages in electrical conductivity, mechanical properties, and fabrication pro-cessing to make thinner plates and finer flow fields with high power density and lower cost in mass production. The metal plates also have the benefit of the ability to be recycled and consistency in forming and welding qual-ity. However, the need for a defect-free and electronically conductive surface coating to increase corrosion resistance, as well as achieving an overall opti-mum of high performance, better durability, efficient manufacturing pro-cesses, and lower cost, remains a big challenge.

It is interesting to note that the development of plate materials from graph-ite to composite to metal is a process from application of the expansive pure graphite to less graphite in composites until no graphite is contained in the metals. In a certain sense, the development process may suggest cost reduc-tion as a dominant factor in the development of plate materials and plates when the other requirements are maintained.

The competition between different plate materials or plates has become more severe in recent years; this is beneficial for fuel cell design and allows manufacturing companies to make a better choice. The major competition is focused on polymer-based composite plates and metal plates. As quali-tatively shown in Table 5.4, each material has its advantages and shortcom-ings. To this end, it is difficult and also too early to make a judgment on which of these two plate materials is better. In addition, as mentioned at the beginning of this chapter, with different market applications, the fuel cells,

TABLE 5.4

Qualitative Comparison of Major Properties and Cost of Three Plate Materialsa

Plate Material GraphitePolymer-Based Composite Metal

Corrosion resistance Excellent Good Poor, need coatingElectrical conductivity Good Fair ExcellentMechanical strength Poor Fair ExcellentMechanical flexibility Poor Good ExcellentThermal conductivity Excellent Fair GoodFormability Poor Good ExcellentGas permeability Poor, need seal Poor, need seal ExcellentDensity Lowest Low HighCost High Lower LowerMass production Difficult Capable Capable

a The term used in the table is only for general qualitative comparison because the actual properties and cost are determined by the specific composition and manufacturing processing of each material.


including their plates and plate materials, have some different requirements and priority rankings of the requirements due to their various working con-ditions and market expectations. After further extensive development, the specific composite plate and metal plate may both be good options for fuel cells applied in different markets.

Although great progress in the development of various plates and plate materials has been achieved and different plate materials have competed strongly with each other, the overall development of plates and plate materi-als has been focused on cost reduction and durability improvement. This is dictated by the market requirements of fuel cells as alternative energy resources. It is impossible to go to mass production and sales of fuel cells and related plates without such improvement, given that the relatively large gap of cost and durability between fuel cells and conventional petroleum-based energy resources still exists. To reach the goal, as mentioned previ-ously, the major trend in research and development of new plates as one of the key components of fuel cells includes exploring new materials with better performance and durability but lower cost, the development of more efficient and cheaper fabrication processing, and moving the focus from lab-level research to production with large volumes.

It may be valuable to emphasize again that, although materials and related manufacturing processes are very critical for the performance and durability of plates, the other working chain, such as plate design, also plays an important role in making the required plates. For example, results from Kumar and Reddy [26] that their continuously tested fuel cells did not show performance reduction after 1,000 h of running could be due to the better flow field design, which helped to efficiently remove the dissolved metal ions from the metal plates. In addition, the small electrical ICR of the coating is important for the metal plate; however, due to the high compres-sive rigidity and larger tolerance in thickness of the plate, it is important to design an appropriate compressive load or corresponding compressive displacement of metal plates to achieve good electrical contact between the metal plates and neighboring GDL without damaging the weak com-ponent. It was found that in a gold-clad stainless steel plate with 0.1 mm thickness, the compressive displacement of ~20 mm could better meet the conflict requirements [39].

Overall, the major challenge for different plate materials is how to meet balance and comprehensive property requirements, long durability, mass production capability, and low cost. With the increase of production scale from the lab level, the requirements of quality and consistency control will also become more challenging. Simpler ex situ testing (including accelerated lifetime degradation testing and accelerated stress testing), as well as long-term in situ testing at stack level under strict conditions close to real applications, is critical in accelerating the improvement iteration of plate materials and plates through to final commercialized production and applications.


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12. Besmann, T. M., J. J. Henry, Jr., E. Lara-Curzio, et al. 2003. Optimization of a carbon composite bipolar plate for PEM fuel cells. Materials Research Society Proceedings 756:F7.1.1–F7.1.7.

13. POCO Graphite. 2008. Materials and services. http://www.poco.com/tabid/54/Default.aspx (accessed Dec. 2008).

14. Cunningham, B., and D. G. Baird. 2006. The development of economical bipolar plates for fuel cells. Journal of Materials Chemistry 16:4385–4388.

15. Cho, E. A., U. S. Jeon, H. Y. Ha, et al. 2004. Characteristics of composite bipo-lar plate for polymer electrolyte membrane fuel cells. Journal of Power Sources 125:178–182.

16. Haack, D. 2004. Scale-up of carbon/carbon bipolar composite. Annual prog-ress report, U.S. DOE Hydrogen Program. http://www.hydrogen.energy.gov/pdfs/progress04/ivd1_haack.pdf (accessed Dec. 2008).


17. Besmann, T. M., J. W. Klett, and J. J. Henry, Jr. 2000. Carbon composite bipo-lar plates for PEM fuel cells. Pacific Northwest National Lab. In U.S. National Laboratory R&D Meeting. http://www.pnl.gov/microcats/ottreview/ottmeeting/14-Besman.pdf (accessed Dec. 2008).

18. Haack, D., M. Janney, and E. Sevier. 2006. Scale-up of carbon/carbon bipolar composite. Annual progress report, U.S. DoE Hydrogen Program. http://www.hydrogen.energy.gov/pdfs/progress06/v_d_1_haack.pdf (accessed Dec. 2008).

19. Haack, D. 2006. Scale-up of carbon/carbon bipolar composite. Annual merit review report, U.S. DoE Hydrogen Program. http://www.hydrogen.energy.gov/pdfs/review06/fc_33_haack.pdf (accessed Dec. 2008).

20. GrafTech International. 2008. Flow field plates (FFPs). http://www.graft-echaet.com/GRAFCELL/GRAFCELL-Products/Flow-Field-Plates-(FFP).aspx (accessed Dec. 20, 2008).

21. Mepsted, G. O., and J. M. Moore. 2003. Performance and durability of bipolar plate materials. In Handbook of fuel cells—Fundamentals, technologies, and applica-tions, ed. W. Vielstich, H. A. Gasteiger, and A. Lamm. Vol. 3: Fuel cell technologies and applications, 286–293 (chap. 24). New York: John Wiley & Sons.

22. Adrianowycz, O. 2007. Next-generation bipolar plates for automotive PEM fuel cells. Annual program review, U.S. DoE Hydrogen Program. http://www.hydrogen.energy.gov/pdfs/review07/fcp_10_adrianowycz.pdf (accessed Dec. 2008).

23. Ticona Engineering Polymers. 2008. Fuel cells. http://www.ticona.com/rede-sign/index/markets/innovation/fuel_cell.htm (accessed March 2008).

24. Bortner, M. J. 2006. Economical high-performance thermoplastic composite bipolar plates. Annual merit review, U.S. DoE Hydrogen Program. http://www.hydrogen.energy.gov/pdfs/review06/fcp_38_bortner.pdf (accessed Dec. 2008).

25. Gerard, D. 2007. Materials and materials challenges for advanced automo-tive technologies being developed to reduce global petroleum consumption. MS&T’07 Abstract Book, 161. ASM & TMS.

26. Kumar, A., and R. G. Reddy. 2002. PEM fuel cell bipolar plate—Materials selec-tion, design and integration. In Fundamentals of advanced materials and energy con-version proceedings, ed. D. Chandra and R. G. Bautista, 41–53. Seattle, WA: TMS.

27. Wang, H., and J. A. Turner. 2004. Investigation of a duplex stainless steel as polymer electrolyte membrane fuel cell bipolar plate material. Journal of Power Sources. 128:193–200.

28. Wind, J., A. LaCroix, A. Braeuninger, et al. 2003. Metal bipolar plates and coat-ings. In Handbook of fuel cells—Fundamentals, technologies, and applications, ed. W. Vielstich, H. A. Gasteiger, and A. Lamm. Vol. 3: Fuel cell technologies and applica-tions, 294–307. New York: John Wiley & Sons.

29. Granier, J. 2004–2005. Innovative concepts for bipolar plates. CLEFS CEA 50/51:76–78.

30. Brady, M. P., B. Yang, J. A. Wang, et al. 2006. Formation of protective nitride sur-face for PEM fuel cell bipolar plates. Journal of the Minerals, Metals and Materials Society 8:50–57.

31. Weil, K. S., Z. G. Yang, G. G. Xia, et al. 2006. Development of low-cost, clad metal bipolar plates for PEM fuel cells. Annual progress report, U.S. DoE Hydrogen Program. http://www.hydrogen.energy.gov/pdfs/progress06/v_d_3_weil.pdf (accessed Dec. 2008).


32. IPG Photonics. 2008. YLR-SM series: 100 W to 1.5 kW single mode CW ytterbium fiber lasers. http://www.ipgphotonics.com/apps_mat_single_YLR_SM.htm (accessed March 2008).

33. Salas, J. F. 2004–2005. Molded organic composites. CLEFS CEA 50/51:78. 34. Brady, M. P., P. F. Tortorelli, J. Pihl, et al. 2007. Nitrided metallic bipolar plates.

Annual progress report, U.S. DoE Hydrogen Program. http://www.hydrogen.energy.gov/pdfs/progress07/v_b_2_brady.pdf (accessed 2008).

35. Kumar, A., and R. G. Reddy. 2004. Materials and design development of bipo-lar/end plates in fuel cells. Journal of Power Sources 129:62–67.

36. Riva, R. 2004–2005. A promising concept: Porous materials. CLEFS CEA 50/51:79–80.

37. America Trim. 2008. High-velocity metal forming. http://www.amtrim.com/hvmf (accessed Dec. 2008).

38. U.S. DoE Hydrogen Program. 2007. Annual progress report. http://www.hydrogen.energy.gov/annual_progress07_fuelcells.html#b (accessed Dec. 2008).

39. Matsuura, T., M. Kato, and M. Hori. 2006. Study on metallic bipolar plate for proton exchange membrane fuel cells. Journal of Power Sources 161:74–78.

343

6Physical Modeling of Materials for PEFCs: A Balancing Act of Water and Complex Morphologies

Michael H. Eikerling and Kourosh Malek

CONTENTS

6.1 Introduction ................................................................................................3446.1.1 Energy Conversion in Polymer Electrolyte Fuel Cells .............344

6.2 Challenges for Materials and Operation ................................................3466.3 Physical Theory and Molecular Modeling of Materials .......................347

6.3.1 The Membrane ...............................................................................3486.3.2 The Catalyst Layers .......................................................................348

6.4 Complex Morphology and the Role of Water ........................................3496.4.1 The Role of Water ...........................................................................3496.4.2 The Hierarchy of Scales ................................................................ 351

6.5 Structural Organization and Dynamic Properties of Ionomer Membranes ............................................................................. 3526.5.1 Status of and Directions in Membrane Research ......................3536.5.2 Structure and Dynamics in the Membrane ............................... 3556.5.3 Molecular Modeling of Self-Organization ................................. 359

6.5.3.1 Atomistic Simulations .................................................... 3596.5.3.2 Mesoscale Coarse-Grained Simulations ...................... 362

6.6 Water Sorption in PEMs ............................................................................ 3696.6.1 Structure of Water in PEMs: Classification Schemes ................ 3696.6.2 Phenomenology of Water Sorption ............................................. 3706.6.3 Thermodynamic Model of Water Sorption ................................ 371

6.7 Proton Transport from the Bottom Up ................................................... 3816.7.1 Proton Transport Phenomena in Membranes ........................... 3816.7.2 Pore-Scale Models of Proton Conduction ...................................3836.7.3 Proton Mobility near the Polymer–Water Interface ..................3856.7.4 Network Model of Membrane Conductivity ............................. 3906.7.5 Electro-Osmotic Drag .................................................................... 394

6.8 Membrane in Fuel Cell Modeling............................................................ 397


6.1 Introduction

6.1.1 Energy Conversion in Polymer Electrolyte Fuel Cells

The powerful biological machinery of energy conversion proceeds via redox reactions in aqueous media that involve electron and proton transfer between molecular entities.1,2 Nature devised concerted sequences of these processes that generate electrochemical potential gradients across cell membranes and thereby enable the storage and the release of electrical energy.

In an analogous, yet in principle simpler manner, electrochmical energy con-version proceeds in polymer electrolyte fuel cells (PEFCs), as depicted in Figure 6.1. Polymer electrolyte fuel cells (PEFCs) represent the most versatile member of the family of fuel cells.3,4 Due to unrivaled thermodynamic efficiencies of direct con-version of chemical energy into electrical energy, high energy densities, and ideal compatibility with hydrogen as a fuel, PEFCs are widely considered a primary solution to the global energy challenge.4–7 They could replace internal combustion engines in vehicles and supply energy to a plethora of power-hungry portable and stationary applications. The successful introduction of PEFCs hinges on the devel-opment of advanced materials and the engineering of cells and stacks that opti-mize power densities and voltage efficiencies. Moreover, fuel cells compete with existing energy conversion technologies in terms of their range of viable operating conditions, stability, and overall system cost.

Under normal operation of an H2/O2 fuel cell, anodic oxidation of H2 (or other hydrocarbons or alcoholic fuels)—that is, H2 2H 2e–—produces protons that move through the polymer electrolyte membrane (PEM) to the cathode, where reduction of O2 (i.e., ½ O2 2H 2e– H2O) produces water. The overall redox process is H2 ½ O2 H2O. The electronically insulating PEM forces electrons produced at the anode through an external electric cir-cuit to the cathode to perform work in stationary power units, drive trains

6.9 Ionomer in Catalyst Layers: Structure Formation and Performance ................................................................................................4036.9.1 Challenges for Design and Operation of Catalyst Layers ........4046.9.2 Multiscale Modeling Scheme of Catalyst Layers ......................4066.9.3 Mesoscale Simulations of Self-Organization

in Catalyst Layers ...........................................................................4096.9.4 Main Results of Macrohomogeneous Catalyst

Layer Models .................................................................................. 4126.9.5 Water Management in Catalyst Layers ....................................... 414

6.10 Concluding Remarks ................................................................................. 419Symbols ................................................................................................................423Greek Symbols .....................................................................................................425References .............................................................................................................426

Physical Modeling of Materials for PEFCs 345

of electric vehicles, cell phones, laptops, etc. The maximum obtainable work from the controlled progression of electron and proton transfer processes in PEFCs corresponds to the Gibbs free energy difference, ΔG, between prod-uct state (H2O) and reactant state of the redox couple (H2/O2). Direct alcohol fuel cells follow the same thermodynamic principles; the redox couples are (CH3OH/O2), (C2H5OH/O2), etc. and the product state involves H2O and CO2.

The thermodynamic efficiency of the cell is given by the ratio of ΔG over the reaction enthalpy ΔH,

th

GH

T SH

1 ,

where ΔS is the reaction entropy.The portion ΔQ ΔH – ΔG TΔS of ΔH is transformed into heat. Ideal theo-

retical efficiencies eth are determined by the types and amounts of reactants and by the operating temperature. Fuel cells have an efficiency advantage over combustion engines because the latter are subdued to the Carnot limitation. High thermodynamic efficiencies are possible for typical fuel cell reactions (e.g., eth 0.83 (at 25 C) for H2 1/2O2 H2O(l)). The electrical potential differ-ence between anode and cathode, Eeq –ΔG/ntF, which is also called the elec-tromotive force or open-circuit voltage, drives electrons through the external

Flo

w F

ield

Flo

w F

ield

Cat

alys

t L

ayer

Cat

alys

t L

ayer

Gas

Tra

nsp

ort

Lay

er

Gas

Tra

nsp

ort

Lay

er

Po

lym

er

Ele

ctro

lyte

Mem

bra

ne

Anode

+

O2, Air

Water

H2, Fuel

Electrical Load

Net Reaction: H2 + ½ O2 H2O

H+

H2 H2O2H+ + 2e– ½ O2 + 2H+ + 2e–

e– e–

–

Cathode

FIGURE 6.1Schematic depiction of seven-layer structure and basic processes in polymer electrolyte fuel cells under standard operation with hydrogen and oxygen.


metal wiring. Here, nt is the number of electrons transferred in the overall reaction (nt 2 for the reactions in an H2/O2 fuel cell, as specified previously).

6.2 Challenges for Materials and Operation

Chemical reactions and thermodynamic considerations of PEFCs are seem-ingly straightforward. The challenges for optimizing fuel cell operation are hidden, however, in the fine microscopic details of structures and processes in their materials. The multilayered design of a single PEFC, depicted in Figure 6.1, includes a proton-conducting PEM sandwiched between anode and cathode. Each electrode compartment comprises an active catalyst layer (CL), which accommodates finely dispersed nanoparticles of Pt on a carbo-naceous substrate; a gas diffusion layer (GDL); and a flow field (FF) plate that serves at the same time as a current collector (CC) and a bipolar plate between neighboring cells in a fuel cell stack.

Fuel cell operation entails (1) coupled proton migration and water fluxes in the PEM, (2) circulation and electrochemical conversion of electrons, pro-tons, reactant gases, and water in CLs, and (3) gaseous diffusion and water exchange via vaporization/condensation in pores and channels of CLs, GDLs, and FFs. All components of an operating cell have to cooperate well in order to optimize the highly nonlinear interplay of these processes. It can be estimated that this optimization involves several 10s of parameters.

In general, the fuel cell voltage E is a complex function of the working current density j0; that is, E f(j0). This relationship is sensitive to material structures, cell design, and operation conditions. Irreversible voltage losses due to mass transport by diffusion and convection of reactant gases and product water in diffusion media, migration of protons and electrons in conduction media, non-uniform distributions of reaction rates, and limited kinetics of thermally acti-vated electrocatalytic processes at fuel cell electrodes reduce E relative to Eeq, thereby causing the diminished voltage efficiency eV E/Eeq. The major objec-tives of worldwide research efforts in the development of advanced materials and the engineering of cells and stacks is to maximize fuel cell efficiency,

fc th V eq

nFEnFE T S

,

and power density, P j0E, at given cost and lifetime of the cell.Concentrating on the operation of the so-called membrane electrode

assembly (MEA), E includes irreversible voltage losses due to proton conduc-tion in the PEM and voltage losses due to transport and activation of elec-trocatalytic processes involved in the oxygen reduction reaction (ORR) in the cathode catalyst layer (CCL):

E j E R j j j j( ) ( ) ( ) ( ).0 0 0 0 0eq

PEM CCL other (6.1)


In general, the membrane resistance RPEM should be considered as a func-tion of j0, as indicated in Equation (6.1). Nonuniform distribution of water in PEMs due to improper water balance can lead to nonlinear effects in RPEM(j0).8 Under extreme conditions, PEM dehydration on the anode side can give rise to a limiting current density in E(j0). The term hCCL(j0) h(z 0) h0 accounts for irreversible voltage losses in the CCL. The remaining contribu-tions to E( j0) in Equation (6.1), lumped together in the term hother(j0), account for parasitic losses due to limited diffusion and nonuniform distribution of reactants in porous gas diffusion media and flow channels7,9–12; ohmic losses in electronically conducting media; methanol crossover13–15; contamination of the catalyst by impurities16; and electrocatalytic losses due to CO poison-ing on the anode side17–19 because of adsorbed reaction intermediates of the oxidation reactions in direct alcohol fuel cells.20,21 At the systems level, effi-ciency considerations include further detrimental factors including fuel pre-treatment, fuel consumption, etc.

6.3 Physical Theory and Molecular Modeling of Materials

This chapter gives an overview of the state of affairs in physical theory and molecular modeling of materials for PEFCs. The scope encompasses systems suitable for operation at T < 100 C that contain aqueous-based, proton-conducting polymer membranes and catalyst layers based on nano-particles of Pt.

As discussed in Eikerling, Kornyshev, and Kulikovsky,11 the theoretical framework fulfills an integral function in linking the various disciplines in fuel cell research. At the fundamental level, theory helps to unravel complex relations between chemical and morphological structures and properties, bridging scales from molecular to macroscopic resolutions. Understanding these relations supports the design of novel, tailor-made fuel cell materials. In fuel cell diagnostics, theory relates ex situ properties of materials to their in situ fuel cell performance; it helps in identifying root causes of nonoptimal fuel cell operation, which in many cases are not amenable to direct measurements. Theory provides valuable input for cell and systems optimization. Approaches in engineering lead to uncontrollable results if they are based on oversimpli-fied structural models and unsettled understanding of fundamental physics. For instance, it would be pointless to study water management in a PEFC with-out appropriate structural pictures of PEMs, CLs, GDLs, and CCs. All these structural elements have to cooperate well in a properly balanced cell.

The materials of greatest interest in view of fundamental understanding and design are the polymer electrolyte membrane and the catalyst layers. They fulfill key functions in the cell and at the same time offer the most compelling opportunities for innovation through design and integration of advanced materials.


6.3.1 The Membrane

To a large degree, the PEM determines the operational range of a fuel cell (i.e., the feasible temperatures, pressures, and humidification requirements). Peculiar properties of the PEM—namely, being gas tight, highly proton con-ductive, and electronically insulating—are of fundamental importance for the fuel cell principle (i.e., the spatial separation of partial redox reactions on anode and cathode sides).

The membrane should be a medium with high concentration and mobility of protons. Moreover, membranes should be mechanically and chemically stable over desired operation times. Fuel cell lifetime requirements range from 3,000 to 5,000 h for car applications up to 20,000 h for bus applications and up to 40,000 h for stationary applications. Notwithstanding the wide range of these overall lifetimes, they have to be accomplished under widely varying operating conditions. Whereas fuel cells for automotive applications need to be compatible with operating temperatures from –40 to 100 C and a wide dynamic power range, stationary applications are much less demand-ing in this respect but require the longest lifetimes.

PEMs can contribute to a significant fraction of irreversible voltage losses during fuel cell operation due to their resistance to proton transport and the crossover of reactant gases. The crossover of unreacted fuel from anode to cathode is a major problem of direct alcohol fuel cells because PEMs that contain large volume fractions of liquid water easily dissolve and transport polar alcohol molecules like methanol or ethanol. The voltage losses in PEMs can be particularly harmful if PEFCs are operated outside their benign range of operation—that is, under conditions that are too hot (T > 90 C) or too dry. In addition to voltage losses incurred directly in PEMs, structure and pro-cesses in the polymer electrolyte affect water management in PEFCs in all components and at all scales.

6.3.2 The Catalyst Layers

The catalyst layers (the cathode catalyst layer in particular) are the powerhouses of the cell. They are responsible for the electrocatalytic conversion of reactant fluxes into separate fluxes of electrons and protons (anode) and the recombi-nation of these species with oxygen to form water (cathode). Catalyst layers include all species and all components that are relevant for fuel cell operation. They constitute the most competitive space in a PEFC. Fuel cell reactions are surface processes. A primary requirement is to provide a large, accessible sur-face area of the active catalyst, the so-called electrochemically active surface area (ECSA), with a minimal mass of the catalyst loaded into the structure.

The generation of electric current in modern catalyst layers proceeds at nanoparticles of Pt that are randomly dispersed on a porous carbon sub-strate with pores of nanoscopic dimension (1–10 nm).22,23 A certain fraction of larger pores (10–100 nm) is needed for the supply of gaseous reactants and


the removal of product water. Overall, current catalyst layers thus possess a random composite structure with bifunctional porosity. The distribution of water in these dual porosity networks plays a major role in determining the interplay of transport and reaction.

Electrochemical reactions occur only at Pt particles at which electrons, protons, reactants, and liquid water meet. This demands interpenetrat-ing percolating networks of Pt/C, Nafion–ionomer, and pores. Major con-straints of this design are: (1) statistical limitations of the Pt utilization due to the random three-phase morphology, and (2) nonuniform reaction rate distributions due to the high thickness (~10 mm) in relation to characteristic diffusion lengths of reactant molecules. These conditions cause inefficient utilization of Pt and problematic voltage losses; those due to oxygen reduc-tion in the CCL (~400 mV) diminish cell efficiency by 30–40% at desired operating current densities of j0 1 A cm–2. An increase by a factor of 10 in the ECSA of Pt reduces voltage losses by 60–120 mV.

Pt is, however, an expensive and limited resource. For a 60 kW fuel cell vehicle, the cost of Pt would be over $2,400 at current cost and loading of Pt. Even worse, replacing combustion engines in all existing vehicles by fuel cell drive systems at no penalty in power would exceed the known reserves of Pt. Catalyst layer design, therefore, strives to reduce the Pt loading markedly at no penalty in the fuel cell voltage.

The previous discussion asserts that design, fabrication, and implementa-tion of stable and inexpensive materials for membranes and catalyst layers are the most important technological challenges for PEFC developers. A pro-found insight based on theory and modeling of the pertinent materials will advise us how fuel cell components with optimal specifications can be made and how they can be integrated into operating cells.

6.4 Complex Morphology and the Role of Water

Two common threads will connect the various aspects considered in this chapter: the pivotal yet double-edged role of water for the operation of PEFCs and the hierarchy of scales that has to be considered in theoretical model-ing, physicochemical characterization, and materials design, as illustrated in Figure 6.2.3,5,7

6.4.1 The Role of Water

From the operational point of view, the major challenge is to understand the versatile role of water for structure and processes in fuel cells. As the main prod-uct of the reaction, the presence of water is unavoidable in any type of fuel cell running on hydrogen, methanol, or other hydrocarbon-based fuels. In PEFCs in particular, water molecules determine the interactions between molecular


components, which control the self-organization phenomena during structural formation in PEMs and CLs.22 These phenomena influence the phase segrega-tion, the adhesion properties between nano- and mesoscopic phase domains, and the stability of the self-organized structures in these materials.

In the hydrated ionomer membrane, liquid-like water acts as the pore for-mer, pore filler, and proton shuttle.5,24 The water distribution and the random network morphology of aqueous pathways determine proton conduction at

HydrophobicPolymer

Membrane MembraneCatalystLayer

CatalystLayer

Microporous

Layer MicroporousLayer

Gas TransportLayer

Gas Channel Flow FieldPlate

Fuel CellStack

Single FuelCell

PlatinumParticle

Ionomer IonomerCoating

CarbonParticle

CatalystParticle

Agglomerates of CatalyzedCarbon Particles

MacroscaleMesoscale

Distance Scale, m

Molecular Scale, Nanoscale

Random Compos.,Effective Properties

Fueling

Options

Lifetime,

Aging

Water

Handling

Distribution of

Species and Reactions

Fabrication and Optimization of MEAs. Engineering of Cells & Stacks

Characterization and

Operation of MEAs

10010–110–210–310–410–5

Distance Scale, m

10–610–710–810–9

O2

O2 O2H2 H2

H+ H2O

10–10

Self-organization,Adhesion

Electrocatalysis:Structure & Kinetics

Proton Conductors:Molec. Mechanisms

CarbonParticle

Proton

Water

–

+

SO3–

FIGURE 6.2The hierarchy of scales that determine structural properties, physical processes, and perfor-mance of PEFCs. Relevant phenomena are indicated.


mesoscopic to macroscopic scales in PEMs and CLs. In the porous CLs, the liquid water saturation determines the interplay between electrochemical con-version of reactants, vaporization exchange of water, and transport of reactants and products to and from the catalyst sites via diffusive and convective flux.

A poorly balanced water distribution in the fuel cell can severely impair its performance and cause long-term effects due to structural degradation. If PEMs or CLs are too dry, proton conductivity will be poor, potentially lead-ing to excessive joule heating, which could affect the structural integrity of the cell. Too much water in diffusion media (CLs and GDLs) blocks the gas-eous supply of reactants. As these examples show, all processes in PEFCs are linked to water distribution and the balance of water fluxes.

Establishing the links between microstructure and water balance in fuel cell materials and components depends on fundamental understanding and knowledge of parameters for the following major aspects:

How does the local water content in fuel cell media depend on the materials’ microstructure and the operating conditions?

By which mechanisms does it attain local equilibrium?What are the mechanisms and the transport coefficients of water fluxes

(diffusion, convection, hydraulic permeation, electro-osmotic drag)?What are the mechanisms and the rates of phase changes between

liquid water, water vapor, surface water in porous substrates, and strongly bound interfacial water in PEMs?

What are the rates of transfer of water across interfaces between dis-tinct fuel cell media (e.g., between PEM and CL)?

6.4.2 The Hierarchy of Scales

Figure 6.2 illustrates the hierarchy of scales in PEFCs that encompasses 10 orders of magnitude, from angstroms to meters. This figure highlights the physical phenomena that must be addressed in fuel cell modeling. Obviously, a well-devised hierarchy of approaches in physical theory and molecular modeling is needed in order to understand how structural details and pro-cesses at all resolutions determine the operation of PEMs and CLs in operat-ing PEFCs.

Molecular-level studies of mechanisms of proton and water transport in PEMs require quantum mechanical calculations; these mechanisms deter-mine the conductance of water-filled nanosized pathways in PEMs. Also at molecular to nanoscopic scale, elementary steps of molecular adsorption, sur-face diffusion, charge transfer, recombination, and desorption proceed on the surfaces of nanoscale catalyst particles; these fundamental processes control the electrocatalytic activity of the accessible catalyst surface.23 Studies of sta-ble conformations of supported nanoparticles as well as of the processes on their surface require density functional theory (DFT) calculations, molecular


dynamics studies of reaction pathways, and kinetic modeling of reactivities based on Monte Carlo simulations or mean field approximations.

At the mesoscopic scale, interactions between molecular components in membranes and catalyst layers control the self-organization into nanophase-segregated media, structural correlations, and adhesion properties of phase domains. Such complex processes can be studied by various theoretical tools and simulation techniques (e.g., by coarse-grained molecular dynamics sim-ulations). Complex morphologies of the emerging media can be related to effective physicochemical properties that characterize transport and reac-tion at the macroscopic scale, using concepts from the theory of random het-erogeneous media and percolation theory.

Finally, conditions for operation at the device level can be defined and bal-ance equations for involved species (i.e., electrons, protons, reactant gases, and water) can be formulated on the basis of fundamental conservation laws of continuum mechanics and fluid dynamics. Therefore, full relations among structure, properties, and performance could be established; this in turn would allow predicting architectures of materials and operating condi-tions that optimize fuel cell operation. These relations would permit studies of operational aspects related to voltage efficiency, power density, catalyst utilization, water handling, and lifetime/degradation issues. Some progress along these lines has been made over the last decade. However, due to the structural complexity of materials and the subtlety of interrelated functional requirements, a plethora of formidable challenges remains to be addressed in the future.

In this chapter, we will mainly address the vital topics in theoretical mem-brane research. Specifically, we will consider aqueous-based proton conduc-tors. Our discussion of efforts in catalyst layer modeling will be relatively brief. Several detailed accounts of the state of the art in catalyst layer research have appeared recently.7,25,26 We will only recapitulate the major guidelines of catalyst layer design and performance optimization and discuss in some detail the role of the ionomer as a proton-supplying network in catalyst lay-ers with a conventional design.

6.5 Structural Organization and Dynamic Properties of Ionomer Membranes

In this section, we describe how modeling and computational tools can be used to investigate the self-organization phenomena in hydrated poly-mer electrolyte membranes.22 It is expected that such studies will generate basic knowledge for design and fabrication of novel, functionally optimized materials for PEFCs. The currently used PEMs exhibit random morpholo-gies. Tailoring their physical properties in view of optimized performance of PEFCs is a multiscale problem.


The bulk of experimental data and modeling results on structure and properties of PEMs has shown that PEMs are not represented well as a homogeneous electrolyte solution or as a rigid porous rock.4,5,7,8 This implies that simple volume-averaging techniques or macroscopic models of flux in porous media cannot be applied straightforwardly to PEMs. Water in PEMs fulfills a versatile role as the pore former, the pore filler, and the pro-ton shuttle. The ingredients for building a sound theoretical framework of transport properties and fuel cell performance of PEMs demand an appro-priate structural picture of the phase-segregated membrane morphology and understanding of the mechanism by which water equilibrates with the membrane matrix.

In the main body, this section presents recently employed mesoscale com-putational methods that can be utilized to evaluate structural factors dur-ing fabrication of PEMs. These simulations provide density distributions or maps and structural correlation functions that can be employed to analyze the sizes, shapes, and connectivities of phase domains of water and polymer; the internal porosity and pore size distributions; and the abundance and wetting properties of polymer–water interfaces.

6.5.1 Status of and Directions in Membrane Research

Aqueous-based polymer electrolyte membranes are the archetypal materi-als employed as proton conductors in fuel cells.27 Over the last 30 years, Nafion ionomer, designed at E. I. DuPont de Nemours in the mid-1960s, has survived as the material of choice due to its high durability and good perfor-mance in PEFCs. Other perfluorosulfonic acid (PFSA) ionomer membranes, such as Flemion (Asahi Glass), Aciplex , and Dow (Dow Chemicals), have been tested extensively as well. These PFSA ionomers vary in chemical structure of polymeric backbones and side chains; ion exchange capacity (IEC), which is defined as the total number of chemical equivalents for ion exchange (i.e., acid head groups, -SO3H) per unit weight of the dry resin; and thickness.

For instance, the Dow experimental membrane and the recently intro-duced Hyflon Ion E83 membrane by Solvay–Solexis are “short side chain” (SSC) fluoropolymers, which exhibit increased water uptake, significantly enhanced proton conductivity, and better stability at T > 100 C due to higher glass transition temperatures in comparison to Nafion.28,29 The membrane morphology and the basic mechanisms of proton transport are, however, similar for all PFSA ionomers mentioned. The base polymer of Nafion, depicted schematically in Figure 6.3, consists of a copolymer of tetrafluoro-ethylene, forming the backbone, and randomly attached pendant side chains of perfluorinated vinyl ethers, terminated by sulfonic acid head groups.30–33

The “catch-22” of these systems is their dependence on liquid water as the working fluid. In contact with water, hydrophobic and hydrophilic polymer constituents self-organize into phase-segregated random structures that


bear water-filled pathways for the transport of protons. In state-of-the-art PEMs, the evaporation of weakly bound water at temperatures exceeding 90 C extinguishes the favorable proton transport through bulk-like water. Moreover, in an operating PEFC, parts of the PEM close to the anode side attain a poorly hydrated state at high current densities due to the electro-osmotic drag effect.1,5,7,8,28,34,35 This incurs critical current density effects in the fuel cell voltage.7,8

The excellent prospects of PEFCs as well as the undesirable dependence of current PEMs on bulk-like water for proton conduction motivate the vast research in materials synthesis and experimental characterization of novel PEMs.36–41 A major incentive in this realm is the development of membranes that are suitable for operation at intermediate temperatures (120–200 C). Inevitably, aqueous-based PEMs for operation at higher temperatures (T > 90 C) and low relative humidity have to attain high rates of proton trans-port with a minimal amount of water that is tightly bound to a stable host polymer.33,37,40,42,43 The development of new PEMs thus warrants efforts in understanding of proton and water transport phenomena under such condi-tions. We will address this in Section 6.7.3.

Membrane research is a rather diverse field, exploiting perfluorinated iono-mers, hydrocarbon and aromatic polymers, and acid-base polymer complexes. Polyether and polyketo polymers with statistically sulfonated phenylene groups such as sPEK, sPEEK, and sPEEKK or polymers on the basis of benz-imidazole have been tested as well. Recent reviews on membrane synthesis and experimental characterization can be found in the literature.36,38,39,41,44

0.6 nm 1.5 nm 10 nm

Random Phase

Separation

Assembly into

Fibrils

CF2

CF2

CF2

CF2

FC O CF2 FC O CF2CF2

CF3

SO3

CF2–

H–

FIGURE 6.3Schematic depiction of the structural evolution of polymer electrolyte membranes. The pri-mary chemical structure of the Nafion-type ionomer on the left with hydrophobic backbone, side chains, and acid head groups evolves into polymeric aggregates with complex interfa-cial structure (middle). Randomly interconnected phases of these aggregates and water-filled voids between them form the heterogeneous membrane morphology at the macroscopic scale (right).


Notwithstanding water being a genuine proton shuttle that nature relies on for vital functions in living organisms, tremendous progress has been made in the development of water-free proton conductors, which operate robustly at elevated temperatures where electrocatalytic reactions at electrodes run fast. As reported by Sossina Haile,37 the intermediate temperature “super-protonic” conductor, CsH2PO4, has advanced from a laboratory curiosity to a viable fuel cell electrolyte. The superprotonic phase exists between 230 and 260 C and rapid proton transport there is believed to emerge from the high degree of polyanion rotational disorder. Another promising example of pro-ton conductors for high-temperature operation are polybenzimidazole (PBI) polymer membranes doped with phosphoric acid, which can operate effi-ciently up to 200 C without any humidification.45 Results of Brian Benicewicz at Rensselaer Polytechnic Institute have demonstrated very good perfor-mance and durability in PEFCs and direct methanol fuel cells (DMFCs).

6.5.2 Structure and Dynamics in the Membrane

Numerous experimental techniques have been employed to understand the morphology and dynamics of polymer and water in hydrated ionomer membranes at different time and length scales. Techniques that have been applied include small-angle and wide-angle x-ray scattering (SAXS, WAXS), small-angle neutron scattering (SANS),46–53 quasi-elastic neutron scatter-ing (QENS),54–56 infrared (IR) and Raman spectroscopy,57,58 time-dependent Fourier transform infrared (FTIR) spectroscopy,59 nuclear magnetic resonance (NMR),60,61 electron microscopy,62 positron annihilation spectroscopy,63 scan-ning probe microscopy,64 scanning electrochemical microscopy (SECM),65 and electrochemical impedance spectroscopy.66,67 Recent reviews on the application of experimental techniques have been conducted.33,44,68

Based on the analysis of WAXS data, Longworth and Vaughan69,70 pro-posed a cluster model of ion aggregation in polyethylene ionomers. A ther-modynamic theory of ion aggregation in organic polymers was proposed by Adi Eisenberg in 1970.71 In the early 1980s, T. D. Gierke and colleagues at DuPont de Nemours developed the first morphological models of hydrated Nafion membranes, using information derived from SAXS data.46,72,73 The so-called Gierke model describes the PEM as a random network of inverted spherical micelles with nanoscopic dimension that are confined by anionic head groups of the polymeric side chains. Aqueous pathways consisting of these inverted spherical micelles and aqueous necks are embedded in an inert and structureless polymer host.

This oversimplified random network model proved to be rather useful for understanding water fluxes and proton transport properties of PEMs in fuel cells.7,8,24,30 It helped rationalize the percolation transition in proton conduc-tivity upon water uptake as a continuous reorganization of the cluster net-work due to swelling and merging of individual clusters and the emergence of new necks linking them.24


Later, G. Gebel and his colleagues at the CEA (Centre d’Études) in Grenoble, France, developed a more refined structural picture on the basis of small angle (SANS, SAXS) and ultrasmall angle x-ray scattering data (USAXS). They identified cylindrical or lamellar-like polymer aggregates of elongated hydrophobic polymer backbones as the prevailing structural motif at the nanometer scale that form the membrane skeleton.51,74–79 These aggregates are lined on their surfaces by the anionic side chains, surrounded by water and protonated water clusters. Schmidt-Rohr and Chen’s recent simulations of small-angle scattering data of hydrated Nafion support a tubular-structure model.80 The suggested structure consists of an array of parallel cylindrical but otherwise randomly packed ionic water channels, embedded within a locally aligned polymer matrix. Introducing crystallites of hydrophobic polymer as physical cross-links was found to be crucial for reproducing the scattering data. The authors demonstrated that other struc-tural models failed to explain peculiar details of experimental scattering curves.

Figure 6.3 illustrates the three major levels in the structural evolution of Nafion-type membranes: from the primary chemical architecture of the ionomer to the random heterogeneous membrane morphology at the mac-roscale. Self-organization of polymer backbones upon hydration leads to the formation of a hydrophobic skeleton that consists of interconnected elongated fibrillar aggregates.74–79 Dissociation of hydrophilic acid groups (-SO3H) releases mobile protons into the aqueous subphase that fills the void spaces between aggregates. Polymeric side chains, which bind the hydrated anions (-SO3

–), remain fixed at the surface of hydrophobic polymer aggre-gates, where they form a charged, flexible interfacial layer relative to which protons and water molecules move.

The structure of this interface determines the stability of PEMs, the state of water, the strength of interactions in the polymer/water/ion system, the vibration modes of side chains, and the mobilities of water molecules and protons. The charged polymer side chains contribute elastic (“entropic”) and electrostatic terms to the free energy. This complicated interfacial region thereby largely contributes to differences in performance of membranes with different chemical architectures. Indeed, the picture of a “polyelectro-lyte brush” could be more insightful than the picture of a well-separated hydrophobic or hydrophilic domain structure in order to rationalize such differences.81

The simple water channel models74,75,80 can explain the ionomer peak and the small-angle upturn in the scattering data of the unoriented samples as well as of the oriented films. Interestingly, the helical structure of backbone segments82 is responsible for the stability of the long cylindrical channels.83 The self-diffusion behavior of water and protons in Nafion is well described by the water channel model. The existence of parallel wide channels at high water uptake favors large hydrodynamic contributions to electro-osmotic water transport and hydraulic permeation.


The study of the dynamical behavior of water molecules and protons as a function of the state of hydration is of great importance for understanding the mechanisms of proton and water transport and their coupling.33 Such studies can rationalize the influence of the random self-organized polymer morphol-ogy and water uptake on effective physicochemical properties (i.e., proton conductivity, water permeation rates, and electro-osmotic drag coefficients).

Pulsed field gradient (PFG)-NMR experiments have been employed in the groups of Zawodzinski and Kreuer to measure the self-diffusivity of water in the membrane as a function of the water content.84,85 From QENS,54–56 the typical time and length scales of the molecular motions can be evalu-ated. It was observed that water mobility increases with water content up to almost bulk-like values above l ~ 10, where the water content l nH2O/nSO3H is defined as the ratio of the number of moles of water molecules per moles of acid head groups (-SO3H). In Perrin et al.,86 QENS data for hydrated Nafion were analyzed with a Gaussian model for localized translational dif-fusion. Typical sizes of confining domains and diffusion coefficients, as well as characteristic times for the elementary jump processes, were obtained as functions of l; the results were discussed with respect to membrane struc-ture and sorption characteristics.86

NMR relaxometry is the most suitable technique to investigate the proton motion in the range of 20 ns to 20 ms.87 The NMR longitudinal relaxation rate, R1, measured over three orders of magnitude of Larmor angular fre-quencies, w, is particularly sensitive to host–water interactions and thus well suited to study fluid dynamics in restricted geometries. In polyimide mem-branes, a strong dispersion of R1 was found that closely followed a power law—R1 w a —in the low-frequency range (correlation times from 0.1 to 10 ms). This is indicative of strong attractive interactions of water with “inter-facial” hydrophilic groups of the polymeric matrix (good wetting behavior).

Variations of R1 with l suggest a two-step hydration process: solvation and formation of disconnected water clusters centered on polar head groups, fol-lowed by the formation of a continuous hydrogen-bond network. At low l, R1 depends logarithmically on w, suggesting bidimensional diffusion of pro-tons in the interfacial region between polymer and water.

QENS studies on the dynamics in PEMs suggest that water and protons attain microscopic mobilities that are similar to those in bulk water. The experimental scattering data include two components, which correspond to “fast” (a few picoseconds) and “slow” (>100 ps) motions. The latter indicates the existence of the hydronium ion as a long-lived entity in Nafion.87 The local and long-range diffusion coefficients of water, probed by QENS, range from Dt 0.5 10–5 cm2 s–1 to Dt 2.0 10–5 cm2 s–1 and Dlr 0.1 10–5 cm2 s–1 to Dlr 0.6 10–5 cm2 s–1 for water contents in the range of l 3–18.55 The diffusivity of protons in hydrated Nafion 117, obtained from QENS data, is DH3O 0.5 10–5 cm2 s–1 at l 15 and 300 K.55 The corresponding values of water and proton diffusivities in bulk water are 2.69 10–5 cm2 s–1 and 0.62 10–5 cm2 s–1, respectively.68


The reduction of the long-range diffusivity, Dlr, by a factor of four with respect to bulk water can be attributed to the random morphology of the nanoporous network (i.e., effects of connectivity and tortuosity of nanop-ores). For comparison, the water self-diffusion coefficient in Nafion measured by PFG-NMR is Ds 0.58 10–5 cm2 s–1 at l 15.84,85 Notice that PFG-NMR probes mobilities over length scales > 0.1 mm. Comparison of QENS and PFG-NMR studies thus reveals that the local mobility of water in Nafion is almost bulk-like within the confined domains at the nanometer scale and that the effective water diffusivity decreases due to the channeling of water molecules through the network of randomly interconnected and tortuous water-filled domains.55

Figure 6.4 shows that long-range diffusivities of water in Nafion membranes measured by QENS, Dlr, are equal to self-diffusivities determined by PFG-NMR, Ds, at l > 10. In well-hydrated membranes, the major geometric constraints for water mobility due to the phase-segregated, random network morphology of

Bulk Water 25 °C

FIGURE 6.4Local (Dt) and long-range (Dlr) diffusion coefficients of water in Nafion membrane probed by QENS, illustrating the enhanced water dynamics with increasing hydration level of the membrane. Self-diffusivity of water in Nafion probed by PFGNMR and self-diffusivity of bulk water (dashed horizontal line) are given for the sake of comparison. [We gratefully acknowledge permission of the author to adopt this figure from Jean-Christophe Perrin, PhD thesis: “Etude Experimentale Multi-Echelles de la Dynamique de L’eau Dans Les Membranes Ionomeres Utilisees en Piles a Combustible,” 2006, Université Joseph Fourier, Grenoble I, France. Université Joseph Fourier, http://tel.archives-ouvertes.fr/docs/00/11/54/18/PDF/TheseJCPerrin2006.PDF]


aqueous domains unfold at the scales probed by QENS, which is up to several tens of nanometers. This tells us that there are no restrictions of mobility aris-ing between the nanometric and the micrometric scales when the membrane is sufficiently hydrated. At low hydration, the difference between the two dif-fusion coefficients, Dlr and Ds, could be attributed to the lamellar structure of Nafion at the nanometric scale, as revealed by NMR relaxometry.87

The important impact of these experimental insights for molecular model-ing is that the development of structure versus property relations of PEMs does not require multiscale approaches going all the way to the macroscopic scale. Rather, the main job is done if one arrives at the scale of several 10s of nanometers. Notably, operation at low hydration emphasizes even more the importance of (sub)nanoscale phenomena controlled by explicit interactions in the polymer–water–proton system.

6.5.3 Molecular Modeling of Self-Organization

Transport properties of hydrated PFSA membranes strongly depend on nanophase-segregated morphology, water content, and state of water. In an operational fuel cell, these characteristics are indirectly determined by the humidity level of the reactant streams and Faradaic current densities gener-ated in electrodes, as well as the transport properties of catalyst layers, gas diffusion layers, and flow fields.4,5,7,8,11,12,30,34

Despite satisfactory experimental evidence for cylindrical or ribbon-like shapes of backbone aggregates in Nafion membranes, understanding of how transport properties are correlated with the microstructure is still lacking. Computer simulations at meso- to microscale and down to atomistic levels can complement experiments in understanding the processes of structural formation in Nafion.5,88–96 The dilemma that molecular simulations are fac-ing is that suitable methods should allow including microscopic details of chemical architecture and molecular interactions, while they should still capture the structure-related transport properties at sufficiently long time and length scales.90,97 The length and time scales of atomistic simulations, based on classical or ab initio molecular dynamics, are usually limited by excessive computational demands.

Computational modeling approaches that have been developed to under-stand the structure and transport properties of water and protons in swollen Nafion membranes89–91,98–104 include quantum mechanical,102–105 classical,91–94,97–101 and coarse-grained89,90 methods.

6.5.3.1 Atomistic Simulations

The earliest fully atomistic molecular dynamic (MD) studies of a simplified Nafion model using polyelectrolyte analogs103,106 showed the formation of a percolating structure of water-filled channels, which is consistent with the basic ideas of the cluster-network model of Hsu and Gierke.72 The first MD


simulations using a representation of the perfluorosulfonate ionomer mol-ecules by oligomers were performed by Vishnyakov and Neimark.93 In order to ensure a rapid approach to equilibrium, an annealing procedure was employed by applying a sequence of isobaric–isothermal (NPT) and canoni-cal (NVT) simulation runs. By using a coarse-grained, united atom repre-sentation for the CFx groups in both the fluorocarbon backbone and the side chains, Vishnyakov and Neimark improved the computational efficiency in their studies.

United-atom force fields have the advantage that larger systems can be simulated compared to all-atom force fields. For instance, Urata et al.107 have simulated systems of united atoms CFx containing 12,000–25,000 atoms for 1.3–2.5 ns. A possible drawback that may arise from this representation is that the results do not account accurately for the role of backbones in control-ling morphology and transport. The all-atom approach is computationally more demanding, but it is required to establish rigorous benchmarks that can validate subsequent coarse-grained simulations.

Based on a structural analysis of the microphase-separated morpholo-gies, Vishnyakov and Neimark93 developed a similar structural picture of Nafion to that obtained using earlier polyelectrolyte or oligomeric models. As a major new idea, they proposed that in the sufficiently hydrated mem-brane, the continuous passage of protons and water molecules is effected by temporary bridges between water clusters that form and break dynamically on a time scale of 100 ps. However, this conjecture is difficult to reconcile with conductivity data and with the structural models discussed in Section 6.5.2 because it supplants the permanent percolating network of water path-ways by transiently linked but otherwise disconnected water pools. Such a structural picture has important implications for water uptake as well. This peculiar model could be an artifact of the short oligomeric units that are have been used in these MD simulations.

Elliott et al.103 applied classical MD simulations to study the dynamics of small molecules in a model Nafion membrane for l 1, 3.8, and 9.7. They observed water segregation into “bound” water associated with the sul-fonate groups and more loosely attached “free” water.

Urata et al.107 used an explicit all-atom description of the fully dissociated side chains and a united-atom representation of backbones. Angular and tor-sional potential parameters and partial atomic charges were obtained from hybrid DFT and molecular orbital calculations. These authors observed that the sulfonate groups aggregated and shared water molecules at low l and that strong interactions with charged sulfonate groups suppressed the dynamics of water. At high l, frequent exchange occurred between bound and free water molecules. Moreover, the structure factors were calculated and compared with scattering data, showing that the molecular models give smaller ionic cluster sizes than those determined experimentally for PFSA membranes.

In a series of papers, Dupuis and co-workers99–101 simulated the effect of temperature and membrane hydration on membrane nanostructure and


mobility of water and hydronium ions using classical MD simulations with DREIDING108 and modified AMBER/GAFF109 force fields. In qualitative agree-ment with previous MD simulations, they showed that increasing l causes the sulfonate groups to drift apart. In their simulations, most of the water mol-ecules and hydronium ions are bound to the sulfonate groups at l < 7. This indicates that the simulation overestimates sulfonate–water interactions.

These studies showed that sulfonate groups surrounding the hydronium ion at low l sterically hinder the hydration of the hydronium ion. The inter-facial structure of sulfonate pendants in the membrane was studied by ana-lyzing structural and dynamical parameters such as density of the hydrated polymer; radial distribution functions of water, ionomers, and protons; water coordination numbers of side chains; and diffusion coefficients of water and protons. The diffusion coefficient of water agreed well with experimental data; for hydronium ions, the diffusion coefficient was found to be 6–10 times smaller than the value for bulk water.

Cui et al.98 performed similar analyses to those of Dupuis and co-workers. The side chain–side chain radial distribution functions (RDFs) reported by Cui et al. show remarkable qualitative deviation from those in Zhou et al.101 It is of note that the united atom approach used by Cui and co-workers ignored electrostatic interactions between CF2 groups of the polymeric backbone. This can lead to a poor description of the hydrated structure in the regions close to the polymeric backbones, unlike the all-atom force field used in Zhou et al.101 For the sake of limited computational resources, Cui et al. used a relatively short representation of Nafion ionomer chains consisting of three monomers as compared to the ten monomers used by Vishnyakov and Neimark91,93 or Urata et al.107 It can be expected that struc-tural correlations will strongly depend on this choice.

Jang et al.92 have used an all-atom approach in their MD simulation of phase segregation and transport in Nafion at l 16. It was shown that blocky Nafion ionomers with highly nonuniform distributions of side chains on the polymer backbones form larger phase-segregated domains compared to systems with uniform distributions of side chains on the backbones. Water diffusion coefficients at 300 and 353K were found to agree well with experi-mental values.

In a recent effort, Elliot and Paddison104 applied QM/MM calculations (using the ONIOM110 method) to understand the effects of hydration on the local struc-ture of PFSA membranes. The calculations were performed on fragments of an SSC PFSA ionomer with three side chains. Full optimization of the oligomeric fragment was carried out at the B3LYP/6-31G**:HF/3-21G** level with six to nine water molecules added. They reported a lowest energetic state with six water molecules. With more water molecules added, the energetic preference for uni-form hydration (interconnected water clusters) disappeared. The optimized structures of two oligomeric fragments at l 2.5 showed that the structure with kinked backbone was energetically preferred (~37 kJ mol–1 lower in energy) over the one in which the fluorocarbon backbone was fully extended.


Paddison and Elliott concluded that the conformation of the backbone, the side chain flexibility, and the degree of association and aggregation of the side chains under low hydration determine the formation of protonic species (Zundel and/or Eigen ions).104 These calculations for single ionomer chains do not account for ionomer aggregation. Therefore, they insufficiently repre-sent the membrane morphology and correlation effects between backbones, side chains, protons, and water.

6.5.3.2 Mesoscale Coarse-Grained Simulations

None of the architectures described before and used in atomistic molecular dynamic simulations is able to predict the structure-related properties of the membrane at long time (>10 ns) and length (>10 nm) scales. Mesoscale mod-els are needed to bridge the gap between the chemical structure of the poly-mer and the phase-segregated morphology of the self-organized membrane. The first attempt for mesoscale simulations of hydrated Nafion was based on a hybrid Monte Carlo (MC)/reference interaction site model (RISM).96 This method uses a combination of an MC routine and rotational isomeric state (RIS) theory developed originally by Flory.111 Later, Khalatur, Talitskikh, and Khokhlov96 used a highly coarse-grained representation in which each CF2 or CF3 moiety was represented by a united atom, with a uniform distribution of side groups along the backbone.

The outcome of these calculations was that the water and polar sulfonic acid groups were found to be segregated into a three-layer structure with a central water-rich region and two outer layers of side groups strongly associ-ated with water molecules. In agreement with experiments,73 Khalatur and colleagues found a linear dependency of microscopic swelling on l, attrib-uted to the swelling of the voids between the fibrils.

Coarse-grained (CG) models based on dynamic self-consistent mean field (SCMF) theory have recently been developed to study the structure of hydrated ionomers at varying l.89,90 Each side chain and backbone is con-structed of a number of CG segments (beads), which represent groups of several atoms. The interaction parameters and bead sizes were computed using the classical atomistic MD method. In the SCMF approach, the density distributions of the mesoscopic beads, r(r), evolve under the influence of a slowly varying external potential, U(r), relative to which polymer chains are equilibrated instantaneously. The main assumption of SCMF theory is that the external potential acting on the ideal system generates a density distribu-tion that matches that of the interacting system. The free energy functional consists of terms for the beads in the external potential with the addition of a Gaussian coil-stretching term and it incorporates a Flory–Huggins type of mean field mixing energy.

The bead–bead interaction parameters are generated using classical atom-istic MD or they can be calculated from Flory–Huggins parameters.89,90,111–113 In general, variation of the hydration level at a fixed temperature leads


to structural reorganization of the phase-separated morphology.89,90 Simu-lations suggested that, at low water content (l < 6), the isolated hydrophilic domains are spherical, while at higher water content (l > 8), they deform into elliptical shape. Because these levels of water content are significantly larger than the values required experimentally for achieving high bulk-like proton conductivity, it was concluded that there may be proton transport through water-depleted regions by interfacial diffusion or through a second ionic phase.

Another method applied to predict the mesoscopic structure of hydrated Nafion membranes was dissipative particle dynamics (DPD), based as well on a CG model for Nafion ionomer.112–114 In DPD simulations, the time evolu-tion for a set of interacting particles is governed by Newton’s equations.36,37 The total force acting on a particle entails contributions from a conservative force, a dissipative force, a pair-wise random force, and a binding spring force. Conservative interactions are parameterized on the basis of Flory–Huggins parameters. In agreement with SAXS measurements, Yamamoto and Hyodo114 showed in DPD simulations that the size and separation of ionic clusters increased approximately linearly with l. They also performed Lattice–Boltzmann (LB) simulations of water fluxes in the membrane based on the morphologies generated by DPD. They showed that the permeability of the porous structure, extracted from Darcy’s law, increases with water con-tent and depends strongly on the pressure gradient, fluid viscosity, and grid resolution.115

Recent DPD simulations by Vishnyakov and Neimark116 and Malek et al.117 provide the microsegregated structure of hydrated Nafion at various l. A typical structure obtained is depicted in Figure 6.5. By increasing l, the mor-phology of the membrane shows a percolation-type transition from isolated hydrophilic clusters to the three-dimensional network of randomly intercon-nected water channels. Very recently, Wu et al.118 applied extensive DPD simu-lations to study the morphologies of Nafion, SSC, and 3M PFSA membranes at various hydration levels and ionomer equivalent weights. These DPD sim-ulations suggested that 3M PFSA membranes exhibit larger water clusters compared to SSC membranes at the same water content. It was also shown that longer side chains lead to the formation of larger aggregates of sulfonate groups and consequently to larger water clusters, with cluster sizes varying from 2 to 13 nm for 5 < l < 16.118

In spite of many computational advantages, DPD and SCMF methods are not able accurately to predict physical properties that rely upon time correla-tion functions (e.g., diffusion), making them less applicable to extract struc-ture-related transport properties of phase-segregated membranes.

An alternative mesoscale approach for high-level molecular modeling of hydrated ionomer membranes is coarse-grained molecular dynamics (CGMD) simulations. One should notice an important difference between CGMD and DPD techniques. CGMD is essentially a multiscale technique (parameters are directly extracted from classical atomistic MD) and it


has a different force field handling scheme. Moreover, the angular and dihedral interactions in CGMD, which are ignored in DPD simulations, account for the conformational flexibility of ionomer molecules more appropriately.

In CGMD simulations, a model of the molecular system is defined in which spherical beads with predefined subnanoscopic length scale replace groups of atoms. Thereafter, parameters of renormalized interaction ener-gies between the distinct beads are specified. In hydrated ionomer mem-branes, polar, nonpolar, and charged beads are distinguished in order to represent water, polymer backbones, anionic side chains, and hydronium ions.119 Interactions between beads could be determined by force matching procedures from atomistic interactions120,121 or from experimental structural correlation functions.119

In Malek et al.,117 clusters of four water molecules are represented by polar beads. Clusters of three water molecules and a hydronium ion correspond to charged beads. Each of these beads has a radius 0.43 nm

FIGURE 6.5Snapshots of the final microstructure at l = 9 (RH ≈ 94%), predicted by DPD calculations. Nafion backbones are shown in black, the first side-chain beads are shown in white, and the second side chain beads, water and hydronium ions are shown in gray. (M. Malek et al. Journal of Chemical Physics 129 (2008) 204702.)


and thus a volume of 0.333 nm3. A folded ionomer chain is shown in Figure 6.6. A side chain unit in Nafion ionomer has a molecular volume of 0.306 nm3, which is comparable to the molecular volume (0.325 nm3) of a four-monomeric unit of polytetrafluoroethylene PTFE (-[-CF2-CF2-CF2-CF2-CF2-CF2-CF2-CF2-]-. Each of the four monomeric units and each side chain (represented by a charged bead) are coarse-grained as spherical beads of volume 0.333 nm3. A coarse-grained chain of 20 apolar beads, as illustrated in Figure 6.6, replaces the hydrophobic backbone. This model, considered in Malek et al.,117 is the longest backbone chain in a CG mesoscale simulation so far.

The interactions between nonbonded uncharged beads in CGMD simula-tions are modeled by the Lennard–Jones (LJ) potential:

U rr rLJ ij

ij ij( ) 4

12 6

,,

(6.2)

where the effective bead radius (a ij) is assumed as 0.43 nm for all beads at which the interbead potential is zero. The strength of interactions cij could assume five possible values, ranging from weak (1.8 kJ/mol) to strong (5 kJ/mol). Charged beads i and j interact via coulombic interactions:

U rq q

ri j

eli j( ) .

(6.3)

FIGURE 6.6Coarse-grained model of Nafion with 20-unit oligomer (length ~ 30 nm).


Interactions between chemically bonded beads in ionomer chains are mod-eled by harmonic potentials for the bond length and bond angle:

V r K r r

V r K

bond bond 0

angle angle

( ) ( )

( ) [c

12

12

2

oos( ) cos( )] ,02

(6.4)

where the force constants are Kbond 1,250 kJ mol–1 nm–2 and Kangle 25 kJ mol–1, respectively.117,119 r0 and q0 are the equilibrium bond length and angle, respectively. 119

Membrane simulations were performed with l 4, 9, and 15.117 The meso-scopic structure of the hydrated membrane is visualized in Figure 6.7, reveal-ing a sponge-like structure similar to structures obtained by other mesoscale simulations.114–116 Together with hydrophilic beads of side chains, water beads form clusters that are embedded in the hydrophobic phase of the backbones. The structural analysis indicates that the hydrophilic subphase is composed of a three-dimensional network of irregular channels. The typical channel sizes are from 1, 2, and 4 nm at l 4, 9, and 15. This corresponds roughly to linear microscopic swelling.

The site–site RDF obtained from CGMD simulations matches very well to those from the atomistic MD simulations,117,120 as shown in Figure 6.8. The RDFs between the side chain beads and the other components of the mixture shows that side chains are surrounded with water and hydrated protons. The autocorrelation functions exhibit similar dependences on bead separation at all l, even at very low relative humidity (RH), thus indicat-ing a strong clustering of side chains due to the aggregation and folding of polymer backbones.122 The degree of ordering of water near polymer–water interfaces decreases with increasing l.117

3% wt W 6% wt W 19% wt W

FIGURE 6.7Snapshots of the final microstructure in hydrated Nafion membrane at different water con-tents. Hydrophilic domains (water, hydronium, and side chains) are shown in gray, while hydrophobic domains are shown in black.


So far, CG approaches offer the most viable route to the molecular mod-eling of self-organization phenomena in hydrated ionomer membranes. Admittedly, the coarse-grained treatment implies simplifications in structural representation and in interactions, which can be systematically improved with advanced force-matching procedures; however, it allows simulating systems with sufficient size and sufficient statistical sampling. Structural correlations, thermodynamic properties, and transport parameters can be studied.

Applied to PEMs, the analysis of simulated configurations furnishes the structural picture of the self-organized, phase-segregated morphology of water channels confined by polymer aggregates. Sizes, shapes, and network properties of aqueous channels are in line with the accepted structural mod-els inferred from scattering experiments.74,75,80 Diameters of water channels vary in the range of 1–4 nm, exhibiting a roughly linear increase from low to high water content. The average separation of side chains increases as well with water content, indicating a continuous structural reorganization of polymer aggregates upon water uptake. This could involve backbones slid-ing along each other in order to adopt more stretched conformations.

The side chain separation varies in a range of 1 nm or slightly above. The network of aqueous domains exhibits a percolation threshold at a volume fraction of ~10%, which is in line with the value determined from conductiv-ity studies.24 This value is similar to the theoretical percolation threshold for bond percolation on a face-centered cubic lattice. It indicates a highly inter-connected network of water nanochannels. Notably, this percolation thresh-old is markedly smaller, and thus more realistic, than those found in atomistic simulations, which were not able to reproduce experimental values.

20

10

g (r

)

00 1

r (nm)

Atomistic-MD

CG-MD

2

FIGURE 6.8Site–site W–W radial distribution function obtained from CGMD simulation and compared with that of the atomistic MD simulation using the force-matching procedure.


The ultimate goals of molecular modeling studies for PEM materials based on fully atomistic or CG models are to develop predictive mod-els that can be used for membrane material selection and to rationalize dependence in transport properties of water and protons upon changes in the hydration level. Although experiments provide empirical insights into the structural evolution upon water uptake and the transport proper-ties, for the sake of material design we must understand how the chemi-cal architecture affects properties and performance at the device level. Any individual simulation technique described in this section falls short in making exact predictions for the morphology and effective properties of PEM materials. Coarse-grained representations of ionomer chains uti-lized in mesoscale simulation techniques provide a means to overcome length- and time-scale limitations of atomistic simulations; however, the accuracy of results of SCMF, DPD, and CG simulations hinges on appro-priate choices in defining the bead structure and the interaction param-eters between beads.

The requirements for self-consistent approaches in molecular modeling and computational materials science are (1) an appropriate structural rep-resentation of the primary polymer architecture, (2) an adequate treatment of molecular interactions between components, (3) a sufficient size (in the range of 20–50 nm) of the simulated system that allows addressing effects of nanoscale confinement and random network morphology on transport of water and protons, and (4) a sufficient statistical sampling of structural configurations or elementary transport processes for reliable determination of thermodynamic properties and transport parameters. With respect to (1), it is vital that the length of monomeric sequences of the ionomer exceed the persistence lengths of the polymer backbone, which are between 3 and 5 nm.80

Steric and electrostatic effects of charged side chains fixed at the backbone will significantly enhance the stiffness of ionomer molecules. Overly sim-plistic (often, too short) representations of the ionomer chains could lead to largely inaccurate predictions of structure and properties. Atomistic mod-els often fail in reproducing sizes and shapes of water clusters and poly-mer aggregates as well as in predicting percolation properties and swelling behavior of the hydrated membrane because the monomeric sequences they utilize are too short.

The list of competing requirements defines the need for a multiscale model-ing framework. Starting with quantum mechanical calculations at atomistic scale, one is able to develop simulation methodologies for proton transport and the resulting local electrostatic interactions to derive appropriate force fields for molecular dynamics simulations addressing larger scales. Built upon atomistic MD calculations, a coarse-grained or mesoscale descrip-tion is able to capture essential parameters in synthesis, characterization, and development of advanced membrane materials for PEFCs at the relevant time and length scales.


6.6 Water Sorption in PEMs

6.6.1 Structure of Water in PEMs: Classification Schemes

As much as the nanophase segregated morphology of Nafion has been a controversial issue in the literature over several decades, the need for under-standing the structure and distribution of water in PEMs has stimulated many efforts in experiment and theory. Major classifications of water in PEMs dis-tinguish (1) surface and bulk water,7,8,24,30 (2) nonfreezable, freezable-bound, and free water,123–125 and (3) water vapor or liquid water.126–128 Another type of water often discussed is that associated with hydrophobic regions.

The distinction of surface and bulk water in Eikerling, Kornyshev, and Stimming24 and Eikerling et al.129 is related to the strength of hydrogen-bond interactions of water with the polar surface groups that are fixed at the poly-meric backbones, as indicated in Figure 6.3 (middle). Surface water strongly interacts with these surface groups and it forms a highly oriented, strongly hydrogen-bonded network at polymer–water interfaces in the water-filled channels.105,130,131 Bulk-like water is mainly identified by the liquid water-like dynamics of protons and water molecules, as discussed in Section 6.5.2.33

Distinguishing surface and bulk water has proven useful in explaining the effect of the water content on the microscopic mobility of protons, indicated by the dramatic increase of the activation free energy of proton transport from ~0.12 eV at l > 4 to ~0.36 eV at l < 2.132,133 Moreover, a statistical model of membrane water uptake and proton conductivity7,24,129 suggests that conduct-ing elements with strongly restricted, surface-like mobility of protons control membrane conductivity at low l, while proton current at large l will be car-ried by percolating clusters of nanoscale elements that exhibit high, bulk-like proton mobility. Variants of pore network models based on Gierke’s structural model and on cylindrical pore-type models were developed to account for the transition from surface- to bulk-like conductivity.129 Relations of conduc-tivity versus water content, calculated with the random network type models and cylindrical pore type model, were found to agree well with experimental data for Nafion and Dow membranes. It is evident that the same concept of surface-bulk distinction could be straightforwardly adapted to the structural models of Gebel and Schmidt-Rohr.74,75,80

The categorization into nonfreezable, freezable-bound, and free water is based on observations of the freezing behavior of water by differential scanning calorimetry (DSC) and NMR.134,135 DSC has been used to deter-mine the amounts of the different types of water. Nonfreezable water is tightly bound to sulfonic acid head groups; it plasticizes the polymer and lowers its glass transition temperature, Tg. Freezable water is loosely bound to the polymer, exhibiting a freezing point suppression by up to ~20 C. Notably, a freezing point suppression has also been observed by Cappadonia et al.132 and Cappadonia, Erning, and Stimming133 in Arrhenius


plots of conductivity data. This phenomenon was explained as an effect of the nanoscale confinement that leads to an increased activity of water in ultrasmall pores. The free water possesses the same melting point as bulk water and it is anticipated to facilitate high, bulk-like mobility of protons and water. There seems to be some correlation between surface water and nonfreezable/freezable-bound water, but the assignment is not unique. Moreover, the distinction of freezable-bound and free water is somewhat vague.

The third type of classification (mentioned previously) into vapor and liq-uid water inside PEMs is unphysical and misleading. It is guided by empiri-cal efforts in understanding the role of externally controlled conditions on vapor sorption isotherms. In employing this distinction, the state of water in the adjacent phase outside the membrane, which determines the boundary conditions and mechanism of water sorption, is confused with the state of water in the membrane. Further contributing to this misconception is the fre-quently cited Schroeder’s paradox,136,137 which indeed refers to a rather logi-cal difference in membrane water uptake under equilibration with varying external conditions. It is unrealistic to assume the existence of water vapor in the membrane, and referring to Schroeder’s observation as a paradox is unjustified.137 A consistent physical model of the water sorption equilibrium of a PEM should dispense with both of these issues.

6.6.2 Phenomenology of Water Sorption

The experimental basis of sorption studies includes structural data (SANS, SAXS, USAXS),74,75,77,80 isopiestic vapor sorption isotherms,138–141 and capil-lary isotherms, measured by the method of standard porosimetry.7,8,142–144 Thermodynamic models for water uptake by vapor-equilibrated PEMs have been suggested by various groups.127,128,145–149 The models account for interfacial energies, elastic energies, and entropic contributions. They usu-ally treat rate constants of interfacial water exchange and of bulk transport of water by diffusion and hydraulic permeation as empirical functions of temperature.

The shortcoming of the majority of water sorption models is that they employ a single equilibrium condition expressed through the activities or chemical potentials of water in the membrane and in the adjacent vapor or liquid phases. As noted later, this treatment is insufficient to describe the physical state of a water-equilibrated membrane as a function of external conditions. The frequent but inept citation of Schroeder’s paradox is a conse-quence of this shortcoming. Moreover, explanations of water sorption data often invoke the existence of water vapor in the interior of the membrane, which is then further justified by postulating the existence of hydrophobic pores inside PEMs, with contact angles, q, slightly exceeding 90 .146

These assumptions are, however, largely uncorroborated. There is evi-dence neither for a significant hydrophobic gas porosity nor for the existence


of water vapor inside PEMs. “Gas-tightness” of the pertinent membranes and the collapse of the pore structure upon complete dehydration are clear arguments against these hypotheses. Furthermore, data that are invoked to corroborate hydrophobicity in Nafion channels refer to measurements of wetting angles at external membrane surfaces, which are expectedly pre-dominantly hydrophobic.150 These data are of little validity for microscopic polymer–water interfaces inside PEMs.

For the membrane interior, independent sources of information exclude the existence of vapor and of hydrophobic pores:

The structural model of Gebel74,75,77 and its recent overhaul by Schmidt-Rohr80 show no indication of hydrophobic pores; these models cor-respond to hydrated cylindrical fibrils or water-containing inverted cylinders with rather uniform distribution of charged surface groups at polymer–water interfaces.

Gibbs free energies of water sorption, ΔGs(l), can be extracted from isopiestic vapor sorption isotherms138–141; this analysis shows that ΔGs(l) < ΔGw, where ΔGw –44.7 kJ mol–1 is the Gibbs free energy for vapor sorption at a free water surface at ambient conditions. Water absorbed by the membrane is therefore more strongly bound than water at a free bulk water surface; this affirms the hydrophilic nature of water sorption in PEMs.

DFT calculations of water binding to a dense interfacial array of proto-genic surface groups, which represent acid-terminated side chains in PEMs, have been performed105; in the relevant range of surface group densities, water molecules are strongly bound to the interfa-cial array. Only in the case of high packing density of surface groups can the minimally hydrated interfacial array exhibit hydrophobicity; this transition occurs for separations of side chains < 7 Å, which is below the value at which the known PFSA PEMs dissolve in water.

In current PEMs, hydrophilicity of internal polymer–water interfaces is thus warranted.

6.6.3 Thermodynamic Model of Water Sorption

The subsequently presented model of water sorption in PEMs reconciles vapor sorption and porosity data. At sufficiently large water contents exceeding the amount of surface water, l > ls, equilibrium water uptake is controlled by capillary forces. Deviations from capillary equilibrium arising at l < ls can be investigated by explicit ab initio calculations of water at dense interfacial arrays of protogenic surface groups.105 In the presented model, the problem of Schroeder’s paradox does not arise and there is no need to invoke vapor in pores or hydrophobicity of internal channels. Here, we will present a general outline


of this model and explore its ramifications for membrane transport properties and water balance. Details will be discussed in another publication.

External conditions of the membrane-water system are given by tempera-tures (Ta,Tc), vapor pressures (Pa

v,Pcv), and total gas pressures (Pa

l, Pcl) in the

compartments flanking the membrane, corresponding to anode (index a) and cathode (index c) compartments in an operational fuel cell. Equilibrium in this system requires three independent conditions151:

1. We will adopt the assumption of thermal equilibrium under consid-ered conditions of membrane operation; this implies uniform tem-perature and zero heat flux in the system.

2. Chemical equilibrium corresponds to zero water flux and uniform chemical potential of water in the membrane interior and in the external vapor phase,

wPEM

wext v v v v s,( ) ( ) ln / ,a RT a a P P (6.5)

where av is the activity of external vapor, P v is the vapor pressure, and Ps is the saturated vapor pressure of a free bulk water surface.

3. Mechanical equilibrium corresponds to the balance of pressures in the system, involving total gas pressure, P g; liquid pressure in the water phase inside the membrane, Pl; capillary pressure, P c; and elastic pressure exerted by the polymer matrix, Pel.

The Young–Laplace equation gives the equilibrium pressure difference (mechanical equilibrium) at the menisci between liquid water in membrane pores and vapor in the adjacent phase:

P P Pc g l . (6.6)

In this description, the local state of water in the membrane is thus deter-mined by two independent variables, mw

PEM and Pl.In the pertinent literature, conditions 2 and 3 are often fused into a

single condition by defining a generalized chemical potential of water. However, this makes it impossible accurately to predict the response of the membrane state of hydration to changing external conditions. Under nonequilibrium conditions, applying a difference in vapor pressures, ΔPv Pc

v – Pav 0, or in total gas pressures, ΔPg Pc

g – Pag 0, between cathode and

anode compartments will generate a finite water flux, jw 0, through the membrane. Such perturbations of equilibrium are used in measurements of water transport through the membrane (by diffusion and hydraulic permeation) and of kinetic parameters of vaporization and condensation at its surfaces.152,153


Under fuel cell operation, a finite proton current density, jp 0, and the associated electro-osmotic drag effect will further affect the distribution and fluxes of water in the PEM. After relaxation to steady-state operation, mechanical equilibrium prevails locally to fix the water distribution, while chemical equilibrium is rescinded by the finite flux of water across the mem-brane surfaces. External conditions defined by temperature, vapor pressures, total gas pressures, and proton current density are sufficient to determine the stationary distribution and the flux of water.

To begin, it is essential to rationalize the equilibration of water within the membrane at ΔP v 0, ΔP g 0, jw 0, and jp 0. The suggested scenario of membrane swelling is based on the interplay of capillary forces and poly-mer elasticity. In order to justify a scenario based on capillary condensation, isopiestic vapor sorption isotherms for Nafion138 in Figure 6.9(a) are com-pared with data on pore size distributions in Figure 6.9(b) obtained by stan-dard porosimetry.142–144 In Figure 6.9(a), a simple fit function,

3 0 11 00 2 4

. . ,.

PP

PP

v

s

v

s

(6.7)

provides very good agreement with experimental sorption data.138 The first term in Equation (6.7) could be assigned to strongly bound water near the charged polymer surface (or freezable-bound water), which exhibits only a weak dependence on external vapor pressure. Equation (6.7) implies that the amount of surface water corresponds to l s 3. The second term could be identified with bulk-like or free water.

Figure 6.9(b) reproduces porosity data142 with a log-normal pore size distri-bution, similar to the function suggested in Eikerling et al.7:

max mdrr r

s

r

explog( / )

log

2

0

cc

, with d mrr r

sexp

log( / )log

2

0,

(6.8)

where rm 0.75 nm, s 0.15, and lmax 14. We can compute the Gibbs free energy of water sorption139 using the relationship

G G RT

PP

s wv

sln ,

(6.9)

For water uptake by capillary condensation, the Gibbs free energy of water in pores is given by

G G P Vc w cw , (6.10)


with the capillary pressure

P

rc

c

2 cos,

(6.11)

where s and q are the surface tension and the contact angle of water in pores, respectively, and Vw is the molar volume of water. Inversion of the

Pv/P s

FIGURE 6.9Water uptake of Nafion 117. (a) Isopiestic water sorption data (extracted from T. E. Springer et al. Journal of the Electrochemical Society 138 (1991) 2334–2342) fitted by Equation (6.7); (b) capil-lary isotherms (extracted from J. Divisek et al., Journal of the Electrochemical Society 145 (1998) 2677–2683) fitted by Equation (6.8).


experimental relations f PP

( )v

s and l g(rc) gives two expressions of the Gibbs free energy of water sorption as functions of l:

G G RT f G G

Vs w c w w( ) ln ( ) ( )

cos1

2and

gg 1( ).

(6.12)

In Figure 6.10 we compare these two expressions obtained from the inde-pendent sets of experimental data. At l/lmax > 0.2, the two graphs are indistinguishable—ΔGs(l) ΔGc(l)—exhibiting a modest increase with decreasing l due to the confinement of water in hydrophilic pores. This supports the conjecture that, in this range, capillary condensation is indeed the relevant mechanism of water uptake. The agreement fails for low water content, l/lmax < 0.2, where the steeply increasing strength of water binding, seen in |ΔGs(l)|, is caused by predominant interfacial effects unaccounted for in |ΔGc(l)|. The large energies of water binding at low l, observed in Figure 6.10, are consistent with values found in ab initio quantum mechani-cal calculations of water molecules at hydrated arrays of charged surface groups by Roudgar, Narasimachary, and Eikerling.105

Upon capillary condensation of water in PEMs, the relative humidity, Pv/Ps, determines capillary pressure, Pc, and capillary radius, rc, via the Kelvin–Laplace equation:

G G P V

Vr

RTPP

c w cw

w

c

v

s

2 cosln .

(6.13)

G G

FIGURE 6.10Gibbs free energy of water sorption by Nafion 117. Comparison of energies obtained from sorption isotherms (solid line), corresponding to Figure 6.10(a), and from capillary isotherms (dashed line), corresponding to Figure 6.10(b).


For given external vapor pressure and total gas pressure, the liquid pres-sure (or swelling pressure) in the membrane is obtained from Equations (6.6) and (6.13):

P P

RTV

PP

l g

w

v

sln .

(6.14)

This internal fluid pressure in aqueous domains in the membrane interior is balanced by the elastic pressure exerted by the polymer matrix:

P Pl el . (6.15)

Effects of membrane elasticity on swelling due to water uptake were incorporated in several models of water sorption.127,128,149 Choi, Jalani, and Datta discussed distinct approaches to establish relations between P el and the degree of swelling upon water uptake for microphase-segregated ion-exchange resins,127 including the statistical mechanical theory of Flory and Rehner154 and the more recent model of Freger.155 Flory–Rehner theory pro-vides the following relation for the elastic (or swelling) pressure:

P G

celp p1 3

2,

(6.16)

where G is the shear modulus of the polymer matrix156 and fp is the volume fraction of polymer. The volume fraction of excess water, which corresponds to lb l – l s, is

w p

b

b

1 .

Here it is assumed that only excess water causes swelling. The parame-ter V Vp w/ is the ratio of partial molar volumes of ionomer molecules and water and v is the number of polar head groups (SO3

–) per ionomer molecule.

If we define the molar volume of ionomer based on one monomer unit, we have v 1 and V Mp

1L molp p/ .0 55 , with the molar mass of a mono-mer unit Mp 1,100 g mol–1 (i.e., the equivalent weight) and the density of the dry polymer rp 2.0 g cm–3. We thus obtain r 30.6, where we have used the value Vw

1L mol0 018. of bulk water. The constant c in Equation (6.16) varies between 0 (so-called phantom limit, for which the internal energy is independent of the volume157) and 1 (affine limit). In either case, this theory predicts a finite swelling pressure at zero swelling (i.e., for fp 1). Moreover, Pel remains nearly constant at moderate swelling. These characteristics of the Flory–Rehner theory are in striking contradiction to experimental findings for swelling of ionomer membranes as discussed by Freger.155


Freger’s model treats swelling as a nonaffine inflation of the hydrophobic polymer matrix by small aggregates of water molecules (droplets) separated by polymer films. This model results in a relation,

P Gel

p p23

1 3 7 3

(6.17)

with

G

N k TV

B

0

,

where V0 is the volume of the dry resin and N could be interpreted as the effective number of polymer chains in all films. Equation (6.17) exhibits decreasing swelling pressure upon dehydration and zero swelling pressure under dry conditions, in correspondence to experiment.

The relation between P el and fp is sensitive to the structural model of the membrane and the mechanism of swelling. Equation (6.17) may be appropri-ate for the swelling of spherical domains as suggested in the morphological models of Eisenberg and Gierke.46,71–73 It can be expected that the pertinent structural models of Gebel74,75,77 and Schmidt-Rohr,80 with elongated cylindri-cal domains of polymer fibrils or water channels, will lead to modified rela-tions of P el versus fp.

Nevertheless, Equation (6.17) can serve as a qualitative tool for explaining membrane swelling upon water uptake. Under equilibrium conditions, the elastic pressure in Equation (6.15) increases with increasing vapor pressure due to Equation (6.14). Overall, by invoking Equations (6.14)–(6.17), we obtain a relation between external conditions (P g and P v) and l:

23

1 3 7 3

G Pb b

gg

w

v

s

RTV

PP

ln .

(6.18)

This relation can adequately reproduce the shape of experimental vapor sorp-tion isotherms in the regime corresponding to capillary condensation. The approach to the limit of lb 0 (only surface water remaining) is given by

b

g

w

v

s

34 G

PRTV

PP

ln .

(6.19)

The causal chain of capillary condensation of excess water and swelling in PEMs is thus as follows:

Temperature and vapor pressure of the adjacent gas phase determine the capillary radius, rc, up to which pores are swollen via Equation (6.13).


The corresponding capillary pressure, P c, Equation (6.11), and the external gas pressure, P g, determine the liquid pressure, Pl, via Equation (6.14).

This liquid pressure is balanced by the elastic pressure P el.The self-organized morphology of the membrane determines the rela-

tion between P el and the volume fraction of polymer fp given by Equation (6.16) or (6.17).

fp can be related to the water uptake by

b

p

p

1,

where

V Vp w

and v is the number of polar head groups (SO3–) per considered unit

of the ionomer molecule.

A problem with Equation (6.18) is the prefactor of the logarithmic term. Using the molar volume of liquid water, this factor is

RT Vw atm.1 4 103.

This implies that vapor equilibration of PEMs corresponds to large nega-tive liquid pressures inside the membrane or that lb would increase from zero to the saturation value for Pv very close to Ps. Moreover, the effect of the total gas pressure on water uptake should be insignificant at normal values of Pg 1 atm. Heuristic solutions out of this dilemma would be to recalibrate the value of Vw

or to normalize Pv to a reference value of a porous standard with known relation between lb and water activity. The latter option was employed by Freger.155

In conclusion of this part, it can be stated that the task to reconcile struc-tural membrane models with models of swelling upon water uptake is not yet accomplished. Vital refinements in theory and corresponding experimental studies are needed to include the pertinent structural model, to account for the coupling to external conditions, and to validate the applicability of the Kelvin–Laplace equation in nanopores.

What happens upon equilibration with liquid water instead of water vapor? According to Equation (6.13), the capillary radius would go to infinity for Pv/Ps 1. Thus, in terms of external conditions, swelling would be thermodynamically unlimited, corresponding to the formation of an infinitely dilute aqueous solution of ionomer. However, the self-organized polymer is an effectively cross-linked elastic medium. Under liquid-equilibrated conditions, swelling is not controlled by external vapor


pressure because capillary condensation does not occur at the membrane surfaces, but rather by the elastic pressure of the electrostatically cross-linked polymer. The equilibrium condition reduces to a purely mechani-cal balance of pressures:

23

1 3 7 3

G Pb b

gg ,

(6.20)

where P c 0 and Pl P g; that is, the liquid pressure is equal to the externally controlled gas pressure (of a flat liquid–gas interface).

Equation (6.20) determines the maximum degree of swelling and the maxi-mum pore radius of a liquid-equilibrated membrane. This relation suggests that the external gas pressure over the bulk water phase, which is in direct contact with the membrane, controls membrane swelling. The observa-tion of different water uptake by vapor-equilibrated and by liquid water-equilibrated PEMs, denoted as Schroeder’s paradox, is thus not paradoxical because an obvious disparity in the external conditions that control water uptake and swelling lies at its root cause.

Under steady-state operation with a constant water flux through the membrane, mechanical equilibrium of water will prevail locally at external membrane faces and inside the membrane that involves the balance of local liquid, gas, capillary, and elastic pressures. This condition corresponds to a stationary distribution of water in the membrane. However, the condition of chemical equilibrium, Equation (6.5), will be violated due to the chemical flux of species.

Continuity of the water flux through the membrane and across the exter-nal membrane interfaces determines gradients in water activity or concen-tration; these depend on rates of water transport through the membrane by diffusion, hydraulic permeation, and electro-osmotic drag, as well as on the rates of interfacial kinetic processes (i.e., vaporization and condensation). This applies to membrane operation in a working fuel cell as well as to ex situ membrane measurements with controlled water fluxes that are conducted in order to study transport properties of membranes.

The flux boundary condition accounting for vaporization and condensa-tion kinetics at membrane–vapor interfaces is

j

FRT

k P Pw vv,eq

intv

2( ) ,

(6.21)

where kv is the vaporization rate constant, Pv is the actual vapor pressure in the adjacent gas phase, and pv,eq is the equilibrium vapor pressure that cor-responds to the membrane water content, lint, at the interface. The relation


Pv,eq(lint) is the vapor sorption isotherm, as depicted in Figure 6.9(a), applied at the membrane–gas interface. Equally, we could write

j

FRT

k PG

RTPP

s

w vs int

v

s2exp

( )

(6.22)

where the only empirical input is the relation ΔGs(l), given in Equation (6.12) and plotted in Figure 6.10.

This condition has been recently used in a vaporization-exchange model for water sorption and flux in phase-separated ionomer membranes. The model allows determining interfacial water exchange rates and water perme-abilities from measurements involving membranes in contact with flowing gases.153 It affords a definition of an effective resistance to water flux through the membrane that is proportional to

m

kP

RTD cLv

v

s

effmax

PEM ,

(6.23)

with v 2 if the PEM is in contact with vapor on both sides and v 1 if it is in contact with vapor on one side and liquid water on the other side. The maximum water concentration, cmax, corresponding to the state of complete hydration, depends on ambient temperature and pressure. Trends in steady-state flux data that this model predicts agree very well with experiments.152,153 The vaporization exchange rate, kv, and the effective permeability, Deffcmax, can be independently determined.

A characteristic thickness can be defined on the basis of Equation (6.23):

L

RTD ck PPEM

c veff

max

vs

(6.24)

This parameter helps distinguishing the relative importance of interfacial kinetics and bulk transport. For LPEM < LPEM

c, water transport through the PEM is dominated by interfacial water exchange, whereas for LPEM > LPEM

c, bulk permeation of water prevails. The data obtained in Monroe et al.153 yield LPEM

c 100–300 mm. This indicates that the interfacial vaporization resistance exceeds the resistance due to bulk transport in the membrane when the membrane thickness is LPEM < 100 mm.

For typical catalyst layers impregnated with ionomer, sizes of hydrated ionomer domains that form during self-organization are of the order of 10 nm. The random distribution and tortuosity of ionomer domains and pores in catalyst layers require more complex approaches to account properly for bulk water transport and interfacial vaporization exchange. A useful approach for studying vaporization exchange in catalyst layers could be to exploit the analogy to electrical random resistor networks of


composite electrodes that consist of ohmic resistances and charge transfer resistances.158

6.7 Proton Transport from the Bottom Up

6.7.1 Proton Transport Phenomena in Membranes

In this section, we describe the role of the specific membrane environment on proton transport. As we have already seen in previous sections, it is insufficient to consider the membrane as an inert container for water path-ways. The membrane conductivity depends on the distribution of water and the coupled dynamics of water molecules and protons at multiple scales. In order to rationalize structural effects on proton conductivity, one needs to take into account explicit polymer–water interactions at molecular scale and phenomena at polymer–water interfaces and in water-filled pores at meso-scopic scale, as well as the statistical geometry and percolation effects of the phase-segregated random domains of polymer and water at the macro-scopic scale.

A question of utmost interest is whether high proton mobility in aqueous-based PEMs is possible under conditions of minimal hydration and at ele-vated temperature. Obviously, the answer could have a tremendous impact on promising design strategies in membrane research.37,40 This calls attention to interfacial mechanisms of proton transport (PT).

A look aside to the plethora of experimental studies on lateral proton transport at biomembranes and Langmuir monolayers provides encourag-ing insights for this endeavor. These experiments suggest that lateral pro-ton mobility at interfaces could be rather high—as high as half of the value of proton mobility in bulk water, provided that the packing density of pro-togenic surface groups (SGs) at the interface exceeds a critical value.172,176,177 Later, we will discuss recent studies that focus on the role of acid-functional-ized SGs densely packed at polymer–water interfaces on proton conduction mechanisms.

Proton conductivities of ~0.1 S cm–1 at high excess water contents in cur-rent PEMs stem from the concerted effect of a high concentration of free protons, high liquid-like proton mobility, and a well-connected cluster network of hydrated pathways.94,129,132,133,159,160 Correspondingly, the detri-mental effects of membrane dehydration are multifold. It triggers morpho-logical transitions that have been studied recently in experiment51,74–76 and theory.24,129,161,162 At water contents below the percolation threshold, the well-hydrated pathways cease to span the complete sample, and poorly hydrated channels control the overall transport.24,162 Moreover, the structure of water and the molecular mechanisms of proton transport change at low water contents.


The value of the activation energy of proton transport in well-humidified PEMs, ~0.12 eV,132,133 suggests that the widely studied relay-type mechanism of prototropic mobility in aqueous media prevails; this is also often referred to as structural diffusion or the Grotthuss mechanism.24,163–167 In his semi-nal publication in 1806—long before the discovery of the proton by Ernest Rutherford, which can be dated around 1918—the ingenious physical chem-ist Freiherr Theodore von Grotthuss introduced the rudimentary concept of structural diffusion of positively charged moieties (protons) in acidic solu-tions.168 Ab initio molecular dynamics calculations by Tuckerman et al. 200 years later have established vital molecular details of the structural diffusion of protons in bulk water.165–167

The mechanism is understood nowadays in the following way: Once an excess proton is in water, this proton or any other neighboring pro-ton of the hydrated proton complex can act as a positive-charge carrier. Protons can exchange between localized states in which excess protons reside on hydronium ions (H3O ). H3O has hydrogen bonds to three neighboring water molecules, forming a so-called Eigen ion, H9O4 . The transferring proton passes through an intermediate state, in which it is delocalized between two water molecules, forming a Zundel ion, H5O2 . Transformations H9O4 H5O2 are triggered by hydrogen bond fluc-tuations in the second hydration shell of the central H3O . Subsequent destruction of the metastable H5O2 complex leads to the formation of new H3O or H9O4 moieties (i.e., H5O2 H9O4 ), completing the net transfer. Overall, sequences of breaking and making of hydrogen bonds, local reorientations of water molecules, and barrierless proton transfer in H5O2 establish the mechanism of the anomalously high prototropic mobility in bulk water.

In PEMs, conditions for the genuine bulk-water-like PT are hardly ever encountered.105,169 Similar to proton transport in biophysical systems, rates of PT in PEMs are strongly affected by confinement of water in nanochan-nels, electrostatic effects at interfaces, and desolvation phenomena.129,160,170–174 In the case of cellular membranes and lipid monolayers, systematic experi-mental studies have revealed strong dependence of lateral proton migration on the packing density of proton-binding surface groups tethered at their tails to the interface, as well as by the length, the chemical structure, and the flexibility of these groups.172,175–177

In a similar way, in PEMs, the interfaces between charged polymeric side chains and water account for differences in membrane morphology, stability, state of water, and proton-conductive abilities. Evaporation of weakly bound liquid-like water at temperatures exceeding 90 C extinguishes the most favorable mechanism of proton transport through bulk water. Inevitably, aqueous-based PEMs for operation at higher temperatures (120–200 C) have to attain high rates of proton transport with a minimal amount of tightly bound surface water.


6.7.2 Pore-Scale Models of Proton Conduction

Extensive water loss triggers the observed increase in the activation energy of PT from 0.12 eV at high levels of hydration to >0.35 eV at lowest water uptakes of PEMs.132,133 As of today, it is still unclear whether this transition is due to a change in the molecular mechanism of proton mobility, a morpho-logical transition, or both.161,162 The small but finite residual value of proton conductivity at minimal levels of hydration suggests that sample-spanning pathways of proton transport persist even in the almost dry membrane. The microscopic mechanism of proton transport changes because narrow channels in minimally hydrated PEMs could not perpetuate the high bulk-like proton mobility. In this regime, interactions and correlation effects at polymer–water interfaces become vital.

The complications for the theoretical description of proton transport in the interfacial region between polymer and water are caused by the flexibility of the side chains, their random distributions at polymeric aggregates, and their partial penetration into the bulk of water-filled pores. The importance of an appropriate flexibility of hydrated side chains has been explored recently in extensive molecular modeling studies.104,178 Continuum dielectric approaches and molecular dynamics simulations have been utilized to explore the effects of static interfacial charge distributions on proton mobility in single-pore environments of PEMs.30,160,173,179 Molecular level simulations were employed in order to study side chain correlations and examine direct proton exchange between water of hydration and surface groups.180,181

The empirical valence bond (EVB) approach introduced by Warshel and co-workers182–184 is an effective way to incorporate environmental effects on breaking and making of chemical bonds in solution. It is based on parame-terizations of empirical interactions between reactant states, product states, and, where appropriate, a number of intermediate states. The interaction parameters, corresponding to off-diagonal matrix elements of the classical Hamiltonian, are calibrated by ab initio potential energy surfaces in solu-tion and relevant experimental data. This procedure significantly reduces the computational expenses of molecular level calculations in comparison to direct ab initio calculations. The EVB approach thus provides a power-ful avenue for studying chemical reactions and proton transfer events in complex media, with a multitude of applications in catalysis, biochemistry, and PEMs.

Petersen et al.,169 Petersen and Voth,185 Spohr,88 Spohr et al.,94 and Walbran and Kornyshev186 developed EVB-based models to study the effect of con-finement in nanometer-sized pores and the role of acid-functionalized poly-mer walls on solvation and transport of protons in PEMs. The calculations by the Voth group revealed an inhibiting effect of sulfonate ions on pro-ton motions. The EVB model by Kornyshev, Spohr, and Walbran88,94,186 was specifically designed to study effects on proton mobility due to charge delo-calization within SO3

– groups, side chain packing density, and fluctuations


of side chains and head groups. It was found that proton mobility increases with increasing delocalization of the negative countercharge on SO3

–. The motion of sulfonate groups increases the mobility of protons. Conformational motions of the side chains facilitate proton motion as well. EVB-based studies were able to explain the increase in proton conductivity with increasing water content in PEMs qualitatively.

Continuum dielectric approaches have been used to study proton con-ductivity in model pores of PEMs with well-defined geometries.30,129,160,173 In Eikerling and Kornyshev,173 polymer side chains and anionic countercharges were represented by a regular array of immobile point charges. Poisson–Boltzmann theory was used to calculate the distribution of proton density, r (z), and charge transfer theory was applied to determine electrostatic con-tributions to activation energies of proton transport in slab-like model pores. In the absence of correlation effects in proton transport, the conductance of a slab-like model pore is given by the product of proton mobility (m (z)) and proton density (r (z)) integrated over the pore thickness dimension,

LL

z z zz

z

x

pd

0

0

( ) ( ),

(6.25)

where L is the length of the pore and z0 denotes the positions of the inter-facial layers.173

The main component of the model is the distinction between surface and bulk contributions to pore conductance. It was found that the region for surface conduction is confined to the thickness of about one monolayer of water near the interfaces. The bulk contribution is mainly affected by the density of protons, r (z), which increases from the pore center toward the interfaces.

On the other hand, surface mobility of protons in the vicinity of the SO3–

groups is suppressed due to large coulomb barriers. A higher density of SO3

– groups diminishes coulomb barriers and thus facilitates proton motion near the surface. With increasing water content in the pore, the trade-off between proton concentration and mobility shifts in favor of the bulk con-tribution. A refinement of this single-pore model considered in Commer et al.160 incorporated charge delocalization and thermal fluctuations of SO3

– groups. These effects significantly reduce the interfacial coulomb barriers, thereby facilitating proton motion near the surface. The refined calculations suggest that the transition from surface to bulk conductivity occurs at rather low l.

EVB-based MD simulations, as well as continuum dielectric approaches, involve empirical correlations between the structure of acid-functionalized interfaces in PEMs and proton distributions and mobilities in aqueous domains. The results remain inconclusive with respect to the role of packing


density, conformational fluctuations, and charge delocalization of side chains and SO3

– groups on molecular mechanisms and rates of proton conduction. Most importantly, they do not describe proton conduction in PEMs under conditions of low hydration with l < 3, where interfacial effects prevail.

6.7.3 Proton Mobility near the Polymer–Water Interface

Overall, the effects of confinement and low hydration still represent great challenges for theory and molecular modeling. The approaches described so far provide only an incomplete understanding of fundamental interac-tions of polymer groups, ionized side chains, water, and protons. It is not known how length, density, chemical structure, and the random distribution of charged side chains determine water binding and molecular mechanisms of PT in hydrated channels or pores of nanoscale dimensions.

On the other hand, the merits of such insights are obvious. It would become possible to evaluate the relative importance of surface and bulk mechanisms of PT. The transition from high to low proton mobility upon dehydration could be related to molecular parameters that are variable in chemical synthesis. It could become feasible to determine conditions for which high rates of inter-facial PT could be attained with a minimal amount of tightly bound water. As an outcome of great practical value, this understanding could direct the design of membranes that operate well at minimal hydration and T > 100 C.

Molecular modeling of PT at dense interfacial arrays of protogenic surface groups in PEMs needs ab initio quantum mechanical calculations. In spite of the dramatic increase in computational capabilities, it is still “but a dream” to perform full ab initio calculations of proton and water transport within real-istic pores or even porous networks of PEMs. This venture faces two major obstacles: structural complexity and the rarity of proton transfer events. The former defines a need for simplified model systems. The latter enforces the use of advanced computational techniques that permit an efficient sampling of rare events.187–191

Molecular level simulations in PEM modeling based on density functional theory were employed by Elliott and Paddison,104 Paddison and Elliott,178 and Paddison180 in order to study side chain correlations and examine direct pro-ton exchange between water of hydration and surface groups. The detailed calculations by Paddison and Elliott of hydrated polymeric fragments, includ-ing several side chains attached to a single polymer backbone, are insight-ful in view of fundamental ionomer–water interactions. These approaches ignore, however, correlation effects that arise in two-dimensional interfacial conformations, as encountered in self-organized membrane architectures. As we will discuss later, such effects dramatically influence hydrogen-bond formation, acid dissociation, and flexibility of surface groups at hydrated interfaces.

A trifluoromethane sulfonic acid monohydrate (TAM) solid was studied by Eikerling et al.181 The regular structure of the crystal192 provides a proper


basis for performing ab initio molecular dynamics calculations. The Vienna Ab Initio Simulation Package (VASP) based on density functional theory was used to study the dynamics in the system.193–196 Overall, an MD trajectory of >200 ps was simulated. This trajectory is still too short for direct observations of proton transfer events, which occur on time scales > 1 ns. Intermittent introduction of a proton–hole defect triggered the transition from the native crystal structure with localized excess proton states to an activated state with two delocalized protons, as indicated in Figure 6.11.

One of these protons resides within a Zundel ion, H5O2 , whereas the other is accommodated between two SO3

– groups, which approach each other at hydrogen-bond distance. The formation of this sulfonate O · · · H · · · O complex requires a considerable rearrangement of the crystal structure. The two pro-ton complexes, which are formed almost simultaneously, stabilize the inter-mediate state. The energy of formation for the defect state is approximately 0.3 eV. These calculations suggest that an appropriate flexibility of anionic side chains could be vital for high proton mobility in PEMs under condi-tions of minimal hydration and high density of fixed anions. Furthermore, a drift of the Zundel ion was observed, which alludes to its possible role as a relay group for proton shuttling between hydronium ions and/or sulfonate anions.

For the model system considered in Eikerling et al.,181 the chemical com-position and water content are fixed. Only minimal hydration could be con-sidered. A more recently begun work aims explicitly at the understanding of structural correlations and dynamics at acid-functionalized interfaces between polymer and water in PEMs.105 It directly addresses the question of

S1 S2O3

O4

O2

O1

(b) Intermediate State(a) Native Crystal

H1

H2

2.5 Å

4.5 Å

2.4 Å

FIGURE 6.11Ab initio molecular dynamics simulation of a triflic acid monohydrate crystal. The intermediate state (right) with two delocalized protons is ~0.3 eV higher in energy than the ordered confor-mation of the native crystal (left).


how to increase proton conductivity in PEMs under conditions of minimal hydration.

The model system introduced in Roudgar, Narasimachary, and Eikerling105 emerges from the self-organized morphology of the membrane at the meso-scopic scale that is shown in Figure 6.12(a). The random array of hydrated and ionized side chains is tethered to the surface of aggregated hydropho-bic polymer backbones. Relevant structural properties that define this array include the shape, the thickness, and the persistence length of aggregates as well as the density and the effective lengths of side chains on their surface. In order to obtain a computationally feasible model for ab initio calculations, it is assumed that, to first approximation, the highly correlated interfacial dynamics of side chains, protons, and water decouples from the dynamics of the polymeric aggregates. The supporting aggregate layer is assumed to form an inert basal plane, anchoring the side chains or SGs. Terminal carbon atoms of the SGs are fixed at the positions of a regular hexagonal grid on this basal plane, as depicted in Figure 6.12(b).

In spite of the highly simplified structure, this model retains essential characteristics for studying stable structural conformations and the con-certed dynamics of polymer side chains, water, and protons at polymeric aggregates in PEMs. The approach implies that the effect of polymer dynam-ics on processes inside pores is primarily due to variations in chemical archi-tecture, packing density, and vibrational flexibility of SGs.

Calculations in Roudgar et al.105 focused on the shortest SGs (i.e., CF3SO3H) under conditions of minimal hydration (i.e., with one H2O per SG). The main parameter considered is the nearest neighbor distance of the terminal, fixed C atoms. It was varied from 5 Å dCC 12 Å, which encompasses the range of side chain separations found in prototypical ionomer membranes. The VASP based on DFT was used.193–196 Figure 6.13 displays the formation

(a) (b)

Fixed Carbon

Side chain

“length”dCC

FIGURE 6.12(a) Self-organized morphology of the membrane at the mesoscopic scale. (b) The resulting primitive model of the polymer–water interface consists of a regular hexagonal array of surface groups with fixed endpoints. (Reprinted from S. P. Narasimachary et al. Electrochimica Acta 53 (2008) 6920–6927. Copyright 2008, with permission from Elsevier.)


energy (Efuc) per unit cell as a function of dCC for arrays of CF3CF2SO3H and

CF3OCF2CF2SO3H.130 The formation energy is defined as the energy gained upon bringing three SGs and three water molecules together from infinite separation:

E E d Efuc

total CC total( ) , (6.26)

where Etotal(dCC) and Etotal are the total energy of the system at separation dCC

and at infinite separation of SGs, respectively.As can be seen in Figure 6.13, at high density (dCC < 6.7 Å), ionized SGs and

hydronium ions (H3O ) form an ordered “upright” conformation with full dis-sociation of all acid groups. At dCC > 6.7 Å, cluster-like “tilted” conformations were found. The conformational transition at dCC 6.7 Å that occurs upon decreasing dCC is accompanied by a sharp transition from strong (>0.6 eV) to weak (<0.1 eV) binding of additional H2O. These results suggest that the highly charged, minimally hydrated interface becomes hydrophobic at

0.0

–0.5

–1.0

–1.5

–2.0

–2.5

–3.0

5.0 6.0 7.0 8.0 9.0 10.0

C-C Distance (Å)

Point c

Point c

Partially

Dissociated

Partially Dissociated

CF3CF2SO3H

CF3OCF2CF2SO3H

Point e

Point e

Point d

Point d

Point b

Point b

Point a

Point a

11.0 12.0 13.0 14.0 15.0 16.0

Fo

rmat

ion

En

erg

y (e

V)

0.0

–0.5

–1.0

–1.5

–2.0

–2.5

–3.05.0 6.0 7.0 8.0 9.0 10.0

C-C Distance (Å)

11.0 12.0 13.0 14.0 15.0 16.0

Fo

rmat

ion

En

erg

y (e

V)

FIGURE 6.13The formation energy (Ef

uc) per unit cell as a function of dCC for arrays of CF3CF2SO3H and CF3OCF2CF2SO3H. (Reprinted from S. P. Narasimachary et al. Electrochimica Acta 53 (2008) 6920–6927. Copyright 2008, with permission from Elsevier.)


dCC < 6.7 Å. This intriguing effect is due to strong interfacial correlations and the trigonal symmetry of H3O and SO3

– head groups.Notably, dCC 7 Å has been identified in experiments as a critical value

for the occurrence of long-range proton conduction at lipid and stearic acid monolayers.172,176,177 According to the results in Roudgar et al.,105 this value represents a favorable trade-off between long-range correlations and flex-ibility of SGs. Arrays with longer SGs, which resemble the side chains in real PEMs more closely, exhibit similar interfacial conformations and structural transitions.130 For the longest SGs considered so far, which correspond to the side chains in Dow membranes, two-dimensional correlations and partial dissociation were found to persist up to dCC 15 Å.

The plots of Efuc in Figure 6.13 for the longer SGs look similar to the plots

for the shortest SGs—CF3SO3H and CH3SO3H—presented in Roudgar et al.105 This corroborates that basic interfacial conformations and correlations are independent of the chemical architecture of polymeric side chains. It implies that the main structural effects at hydrated interfacial arrays in PEMs are due to the structure of the acid head group and the packing density of SGs.

Ab initio Car–Parinello molecular dynamics calculations197,198 were used to study proton transfer at minimally hydrated interfaces with dense packing of proton-binding surface groups.131 It was found that concerted tilting-rotation modes of the group that acts as acceptor or donor site for the transferring proton facilitate elementary interfacial proton exchange (depicted in Figure 6.14). The analysis of frequency spectra suggests that local fluctuating modes couple only weakly to the relaxation of the remain-ing interfacial system. The analysis of potential energy surfaces provides the lowest energy path and the activation energy of the elementary proton exchange. These properties are currently being explored for varying sur-face group densities

The goal of such calculations is to discern the pathways of long-range PT along the minimally hydrated interface, for which the ionized SGs could act as relay groups. This mechanism will probably involve several steps of molecular rearrangements of the type in Figure 6.14 and eventually it will require the creation of proton–hole defects. At larger dCC and higher degrees of hydration, surface groups will play a less active role in facilitat-ing long-range PT. Nevertheless, mechanisms of proton exchange between the interfacial network and adjacent water layers will be of great impor-tance for the interplay of bulk and surface conduction of protons as well as for the coupling between proton and water mobilities in well-hydrated membranes.

The overall objective of these studies is to unravel mechanisms of interfacial PT. This requires identification of collective coordinates (or reaction coordi-nates) and transition pathways of transferring protons. Differences in activa-tion energies and rates of corresponding mechanism due to distinct polymer constituents, acid head groups, side chain lengths, side chain densities, and levels of hydration have to be examined. Comparison with experimental


data (e.g., NMR, IR, pulse experiments, and SECM) for molecular mobilities in PEMs, at biomembranes, and at lipid bilayers and for conductivities (e.g., Arrhenius plots) will help rationalize the effects of interfacial structure and segmental motions of side chains on PT. In turn, calculations could stimulate new experiments on well-defined, two-dimensional structures. These stud-ies could furnish constitutive relations between polymer architectures and the dynamics at the interface.

6.7.4 Network Model of Membrane Conductivity

The effective conductivity of the membrane depends on its random hetero-geneous morphology—namely, the size distribution and connectivity of the proton-bearing aqueous pathways. On the basis of the cluster network model,72,73 a random network model of microporous PEMs was developed in Eikerling et al.24 It included effects of varying connectivity of the pore net-work and of swelling of pores upon water uptake. The model was applied to exploring the dependence of membrane conductivity on water content and

d

FIGURE 6.14Elementary interfacial proton transfer and configuration energy along the reaction path, iden-tified in ab initio calculations.


temperature. It could account for distinct swelling properties and conduc-tivities of different PEMs

In general, pores swell nonuniformly. As a simplification, the random network was assumed to consist of two types of pores. In this two-state model, nonswollen or “dry” pores (referred to later as “red” pores) per-mit only a small residual conductance due to tightly bound surface water, which solvates the charged surface groups. Swollen or “wet” pores (referred to later as “blue” pores) contain extra water in the bulk, allowing them to promote the high bulk-like conductance. Water uptake by the membrane corresponds to the swelling of wet pores and to the increase of their rela-tive fraction.

In this model, proton transport in the membrane is mapped on a per-colation problem, wherein randomly distributed sites represent pores of variable sizes and thus variable conductance. The distinction of pores of different color (red or blue) corresponds to interfacial or by bulk-like pro-ton transport. Water uptake by wet pores controls the transition between these mechanisms. The chemical structure of the membrane is factored in at the subordinate structural levels, as discussed in the previous subsections.

The relative number fraction of wet pores as a function of the water content l is given by the law of swelling:

x

NN

( )number of blue porestotal number

b

oof pores,

(6.27)

in which both N and Nb are functions of l. It determines bond probabilities,

p x p x x p xbb br rr( ) ( ) , ( ) ( )( ( )), ( (2 2 1 1 )) ,2 (6.28)

between two wet pores, a wet and a dry pore, and two dry pores, respectively.At the time the model was developed (about 10 years ago), the phenomeno-

logical law of pore swelling that it employed was based on Gierke’s experi-mental data for the structural reorganization of the membrane upon water uptake.46,72,73 The following empirical relations for the number of SO3

– groups in an average pore, n(l), and for the average volume of water-filled pores, v(l), were used:

n n v v( ) ( ) ( ) ( ) ,0 031 1and (6.29)

where n0 is the number of SO3– groups in the average pore of a dry mem-

brane, v0 is the average pore volume in the dry membrane, and a and b are fitting parameters. Large values of a and b reflect pronounced membrane reorganization upon water uptake.


Invoking the conservation of the total number of dissociated SO3– groups

in the membrane and the proportionality between total water content and the volume increase of blue (or wet) pores, the swelling law can be derived:

x( )

( ).

1 3 2

(6.30)

The parameters a, b, and g can be adjusted in order to reproduce the extent of swelling and of reorganization of the polymer matrix upon water uptake. This law accounts for the possibility of merging of smaller pores into larger pores upon swelling. It could represent distinct elasticity of the polymeric membrane matrix that leads to distinct water distributions. In a soft polymer matrix, water would be distributed rather heterogeneously, with individual pores swelling to large equilibrium radii and thus taking up a lot of water. In a more elastic polymer matrix, pores swell more homogeneously with smaller equilibrium radii.

A rigid microporous morphology, which does not reorganize upon water uptake, corresponds to a simple linear relation x(w) g w. In this limiting case, the model resembles the archetypal problem of percolation in bicontin-uous random media.199 Due to deviations of swelling from this law, universal percolation exponents in relations between conductivity and water content are not warranted.

Pore-size-dependent conductances are assigned to individual pores and channels. Three possible types of bonds between pores exist. The corre-sponding bond conductances—s bb(l), s br (l), and srr(l)—can be established straightforwardly. The model was extended toward calculation of the com-plex impedance of the membrane by assigning capacitances in parallel to conductances to individual pores. The probability distribution of bonds to have conductivity s bb, s br, or srr is

f p p p( ) ( ) ( ) ( ) ( ) (b bb b bb br b br rr )) ( ).b rr (6.31)

The simplest method of solution of the Kirchhoff equations that correspond to the random network of conductance elements in three dimensions is in the single-bond effective medium approximation (SB-EMA), wherein a sin-gle effective bond between two pores is considered in an effective medium of surrounding bonds. The conductivity, s b, of the effective bond is obtained from the self-consistent solution of the equation:

d b

b

fd

( )( )

,1

0

(6.32)

which corresponds to an averaging of voltage fluctuations across a single effective bond, where d z/2 and z is the number of channels leading to a


single pore. The effective bond conductivity is obtained from

p

qpbb

b bb

bb bbr

b br

b

( )( )

( )( )

( )

rr brr

b rr

rr b( )( )

( )( )

,q

pq

0

(6.33)

where q was introduced in lieu of (d – 1) to represent the connectivity of the pore network.

When the conductivities of “dry” pores and channels vanish, the true per-colation behavior is obtained:

b bb( ) [( ) ( ) ] ( ),

11 12

qq x

(6.34)

with percolation threshold

x

qc1

1

(6.35)

A value q 24 reproduces well the quasi-percolation behavior of Nafion with equivalent weight of 1,100g mol-1.

This random network theory could explain differences in s b(l) relations for various sulfonated ionomer membranes.24 It could rationalize effects of membrane elasticity and swelling behavior on performance under vary-ing degrees of hydration. The EMA solution of the random network model, if implemented correctly, reproduces the percolation behavior observed in Nafion-type membranes and Nafion-composite membranes.24,200 High elas-ticity of the membrane matrix and high connectivity of pores give small values of the percolation threshold and thus favorable relations, s b(l). In a soft polymer matrix, on the other hand, the fraction of water-filled pores, x(l), increases slowly at low to intermediate water contents, indicating that pores swell rather heterogeneously. In this respect, Nafion seems to offer a favorable elasticity of the polymer matrix.

In Eikerling et al.,129 different model variants of pore space evolution (ran-dom network, serial, and parallel pore models) were compared to each other. A morphology of equally swelling parallel cylindrical pores gives the most beneficial relations, s b(l), with the steepest increase of proton conductiv-ity at small water contents. Results obtained for such a morphology are in good agreement with conductivity data for Dow membranes, which possess shorter pendant side chains than Nafion. In order to adapt the model of pore space evolution to different structural membrane models, the major task is to find the law of swelling, x(l).


6.7.5 Electro-Osmotic Drag

Proton flow in water-filled nanopores of PEMs induces water transport through electro-osmotic coupling. The phenomenological coupling coeffi-cient—the so-called electro-osmotic drag coefficient, nd—is a function of l. It incorporates contributions of molecular diffusion of protonated water clus-ters (i.e., H3O ) and of the hydrodynamic coupling in nanometer-sized water channels.201,202 Typical values of nd (obtained, for example, by volume flux measurements, radiotracer methods, streaming potential measurements, or electrophoretic NMR) are found in the range of nd ~ 1–3.33,203

Electro-osmotic drag phenomena are closely related to the distribution and mobility of protons in pores. The molecular contribution can be obtained by direct molecular dynamics simulations of protons and water in single iono-mer pores, as reviewed in Section 6.7.2. The hydrodynamic contribution to nd can be studied, at least qualitatively, using continuum approaches. Solution of the Poisson–Boltzmann (PB) equation,

( ( )) exp

( ),0 0r

F rRT

(6.36)

provides the proton distribution,

( ) exp

( ),r

F rRT0

(6.37)

and the electric potential, (r), in pores. Here we evaluate cylindrical pore geometry with radius rp. The normalized PB equation in cylindrical pores is

1 2

x xx

xx

r xdd

dd p

( )( ) exp( ( ))

(6.38)

with

x r r F RT/ , ( ) ( ) ,p

and

FRT

0

0

1 2

The radial component of the electric field is

E

RTFr xr

p

dd

,

(6.39)


with nd(x 0) 0 in the pore center and Er(x 1) /e e0 at the pore wall, where is the surface charge density.

With the substitution u ln x and g u2 , Equation (6.38) can be trans-formed into the one-dimensional PB equation

dd p

2

22g

ur g( ) exp( ),

(6.40)whose solution is known.

Pore size and dielectric constant e of water in pores exhibit a strong effect on proton distributions, as studied in Eikerling.204 Model variants that take into account the effect of strongly reduced e near pore walls205 and the phe-nomenon of dielectric saturation206,207 lead to nonmonotonous profiles in pro-ton concentration with a maximum in the vicinity of the pore wall.

The hydrodynamic equation of motion (Navier–Stokes equation) for the sta-tionary axial velocity, vz(r), of an incompressible fluid in a cylindrical pore under the influence of a pressure gradient, dPl/dz, and an axial electric field, Ez, is

1 1r r

rvr

Pz

Er

dd

dd

dd

zl

z ( ),

(6.41)

where m is the dynamic viscosity. For vanishing pressure gradient, water transport is entirely driven by electro-osmotic drag, with a velocity deter-mined by

1 10r r

rvr

Er r

rr

dd

dd

dd

dd

z z .

(6.42)

With the boundary conditions

v r

vr

rz p

zdd

( ) 0 00

and

the velocity profile is

v r

Er

rrp

pz

z( ) ( )( )( )

.0 1

(6.43)

The volumetric water flux due to electro-osmosis is given by

V rdr v r

rr I E

r

peo z p g z

p

20

0 2( )( )

,

(6.44)


where

Ir

rdrrrp

r

gp

p

21

20

( )( )

is a geometry factor of order 1.Using Ohm’s law of proton transport in a pore, jpore eporeEz, and the hydro-

dynamic flux density,

j

Fc rI Ep

hydrw

g z0 ( )

,

where cw is the concentration of water, which is assumed constant, we find

n

j

j

Fc r Ihydr

hydr

pore

w p g

pore

0 ( ).

(6.45)

This hydrodynamic contribution to nd is determined by the dielectric con-stant (e) and the viscosity of water (m), the surface charge density of the pore ( ), the pore radius (rp), and the proton conductivity of the pore (spore). The hydrodynamic electro-osmotic coefficient for a typical pore with rp 1 nm is found in the range of [i.e., nhydr 1–10].

The total electro-osmotic coefficient nd nhydr nmol includes a contribution of hydrodynamic coupling (nhydr) and a molecular contribution (nmol) related to the diffusion of mobile protonated complexes—namely, H3O . The rela-tive importance, nhydr and nmol, depends on the prevailing mode of proton transport in pores. If structural diffusion of protons prevails (see Section 6.7.1), nmol is expected to be small and nd nhydr. If, on the other hand, proton mobility is mainly due to the diffusion of protonated water clusters via the so-called “vehicle mechanism,”208 a significant molecular contribution to nd can be expected. The value of nd is thus closely tied to the relative contribu-tions to proton mobility of structural diffusion and vehicle mechanism.33

Structural diffusion is favored by conditions that enhance the stiffness of the hydrogen-bonded network between water molecules: low tempera-tures and low acid concentration.209 The decrease in water content leads to an effective increase in the concentration of acid protons, which in turn sup-presses the contribution of structural diffusion, as found in aqueous acidic solutions.209 This agrees with the finding of an enhanced contribution of vehicular transport in PEMs at low hydration.33 Such an observation is also supported by recent studies of molecular mechanisms of proton transport in PEMs at minimal hydration.131

Similar to proton conductivity, the effective membrane parameter, nd, is determined by the membrane molecular architecture, in particular by the


equivalent concentration of acid protons in water-filled channels. The struc-tural effects translate into characteristic dependence of nd on l. In general, it is observed that nd increases with increasing l. Prevalence of the vehicle mechanism of proton transport for low l suggests nd nmol 1 in this limit. nd is reduced in membranes with narrow channels and strong polymer–solvent interactions (e.g., in S-PEK as compared to Nafion). These trends can be explained with the decrease of structural diffusion, which reduces nhydr, as discussed previously. Kreuer et al. recently published a review of experimen-tal data on nd.33

6.8 Membrane in Fuel Cell Modeling

Under ideal operation of PEFCs, the membrane would retain a uniformly saturated level of hydration, providing the highest proton conductivity, sp

s. The PEM would therefore perform like a linear ohmic resistance, with irre-versible voltage losses:

PEM

PEM

ps

Lj0 ,

where j0 is the fuel cell current density.In reality, this behavior is only observed in the limit of small j0. At cur-

rents j0 ~ 1 A cm–2 that are relevant for fuel cell operation, the electro-osmotic coupling between proton and water fluxes causes nonuniform water distri-butions in PEMs, which lead to nonlinear effects in hPEM. These deviations result in a critical current density, jpc, at which the increase in hPEM causes the cell voltage to decrease dramatically. It is thus crucial to develop membrane models that can predict jpc on the basis of experimental data on structure and transport properties.

To be able to rationalize the role of the PEM in regulating the water fluxes through the cell, it is vital to understand spatial distributions of water and water fluxes. Eikerling et al.7 and Weber and Newman34 have recently reviewed modeling approaches that focus on membrane water management.

The physical mechanism of membrane water balance and the formal struc-ture of modeling approaches are straightforward. Under stationary opera-tion, the inevitable electro-osmotic flux has to be compensated by a back flux of water from cathode to anode, driven by gradients in concentration, activ-ity, or liquid pressure of water. The water distribution in PEMs that is gener-ated in response to these driving forces decreases from cathode to anode. With increasing j0, the water distribution becomes more nonuniform. At jpc, the water content near the anode falls below the percolation threshold of pro-ton conduction, l < l c. This leaves only a small conductivity due to surface transport of water. As a consequence, hPEM increases dramatically; this can lead to failure of the complete cell.


The random phase-segregated morphology of PEMs has been integrated to varying degrees of complexity into models of membrane operation. The genuine structural picture of the PEM as a self-organized, phase-segregated polymer is too complex for the modeling of PEM operation in PEFCs. Performance model-ing on the basis of this structure would warrant predictive multiscale modeling capabilities, which are beyond currently available computational resources.

The simplest practicable approach considers the membrane as a continu-ous, nonporous phase in which water of hydration is dissolved.35,210,211 In such a scenario, which is based on concentrated solution theory, the sole thermo-dynamic variable for specifying the local state of the membrane is the water activity; the relevant mechanism of water back-transport is diffusion in an activity gradient. However, pure diffusion models provide an incomplete description of the membrane response to changing external operation con-ditions, as explained in Section 6.6.2. They cannot predict the net water flux across a saturated membrane that results from applying a difference in total gas pressures between cathodic and anodic gas compartments.

Structural models of membrane operation, on the other hand, treat the mem-brane as a heterogeneous porous medium with percolating water networks for proton and water transport. This structural concept warrants hydraulic perme-ation (D’Arcy flow212) as the prevailing mechanism of water transport, which supersedes diffusion except at rather small water contents. Analogous to water flux in porous rocks, the capillary pressure controls the water saturation in the membrane, and gradients in capillary pressure are responsible for the hydraulic water flux from cathode to anode. Thus, this approach integrates the previously discussed understanding of membrane structure, water sorption, membrane swelling, and proton transport mechanisms.

An external gas pressure gradient applied between anode and cathode sides of the fuel cell may be superimposed on the internal gradient in liquid pressure. This provides a means to control the water distribution in PEMs under fuel cell operation. This picture forms the basis for the hydraulic per-meation model of membrane operation that has been proposed by Eikerling et al.7,8 This basic structural approach can be rationalized on the basis of the cluster network model.46,72,73 It can also be adapted to include the pertinent structural pictures of Gebel et al.74,75,77,78 and Schmidt-Rohr et al.80

Notwithstanding any particular structural model, water transport in PEMs, in general, should be considered a superposition of diffusion in gradients of activ-ity or concentration and hydraulic permeation in gradients of liquid or capillary pressure. Hydraulic permeation is the predominant mechanism under condi-tions for which water uptake is controlled by capillary condensation, whereas diffusion contributes significantly if water strongly interacts with the polymeric host. The molar flux of liquid water in the membrane,

vNl, is thus given by

vv

N nj

FD c

kPd

p m pm

l( ) ( )( )

,

(6.46)


wherec, Pl, and m are the concentration (relative to a unit volume of the membrane),

the pressure, and the viscosity of liquid water in pores, respectivelyDm(l) is the membrane diffusivitykp

m(l) the hydraulic permeabilitynd(l) the electro-osmotic drag coefficient.

It is indicated that these transport parameters are functions of l. The hydraulic permeability (D’Arcy coefficient), kp

m(l), exhibits strong depen-dence on l because larger water contents result in an increased number of pores used for water transport and better connectivity in the porous net-work, as well as in larger mean radii of these pores. A modification of the Hagen–Poiseuille–Kozeny equation was considered by Eikerling et al.7,8 to account for these structural effects:

k

rpm c c

c( )( ) ( )

( ),8

(6.47)

where x(<1) is the inverse tortuosity factor.For an isotropic tortuosity in three dimensions, x 1/3. is the Heaviside

step-function that accounts for the existence of a percolation threshold, l c, in the water-permeating network. l c can be obtained from the specific law of swelling established for the membrane under consideration, as discussed in Section 6.7.4. The mean square radius of pores that contribute to the water flow in a local volume element with water content l is

( )( )

.rr

rr r

r

cd

dd

c1 2

0 (6.48)

Due to mass conservation of water in the membrane, we have

vNl 0. (6.49)

Proton current is determined by Ohm’s law and by the continuity of pro-ton flux,

v vj jp p p( ) ,and 0

(6.50)

with the membrane conductivity sp(l). Proton current in the membrane is transported predominantly in the through-plane direction. It is thus expedi-ent to consider a one-dimensional problem with scalar variables j and Nl.

The model demands knowledge of the functions Dm(l), kpm(l), nd (l), and

sp(l) and their dependence on membrane morphology. It is impossible to


calculate these properties from first principles; they must be obtained from experiments. Moreover, boundary conditions on vapor pressures, gas pres-sures, water, and proton fluxes must be fixed.

The hydraulic permeation model in Eikerling et al.7,8 helped rationalize main dependence of the critical current density on membrane parameters. A sharply peaked -function-like pore size distribution,

dd max

( )( ),

rr

r r1

which is completely determined by the maximum water uptake, l max, and by the first moment of the pore size distribution (i.e., the average pore size), r1, was used in Eikerling et al.7 to derive an explicit expression for jpc:

j

nj Jmpc

dw

c

max

11 ,

(6.51)

where nd is assumed to be constant and jw is the net water flux through the membrane.

The important membrane parameter, Jm, given by

JF c r

Lmw max

PEM41 , (6.52)

where s is the surface tension of water in ionomer channels, is mainly determined by r1 and the membrane thickness, LPEM. Equations (6.51) and (6.52) suggest that membranes with higher water uptake, larger pore radii, reduced thickness, and suppressed electro-osmotic drag are less prone to dehydration.

The theoretical analysis of the hydraulic permeation model, moreover, pro-vided an expression for the current density, jps, at which membrane dehydra-tion commences:

j

jn

Fc k

nP

Lpsw

d

w pm

max

d

g

PEM

( ).

(6.53)

Below jps, the membrane performs under uniform saturation conditions. like a linear ohmic resistance. According to Equation (6.53), two modes of water management can be applied to compensate for electro-osmotic drag and keep the membrane in a well-hydrated state. Sufficient replenishment of water in the membrane can be accomplished by (1) providing a steady exter-nal water supply jw ndjp at the anode side, or (2) applying an external gas


pressure gradient that enforces a sufficiently high water internal back flux in the membrane from cathode to anode,

P n

LFc k

jgd

PEM

w pm

maxp( )

( ),

where 0 if product water is removed via the cathode and ½ if it is removed via the anode.

The diffusion model and the hydraulic permeation model differ decisively in their predictions of water content profiles and critical current densities. The origin of this discrepancy is the difference in the functions Dm(l) and kp

m(l). This point was illustrated in Eikerling et al.,7 where both flux terms occurring in Equation (6.46) were converted into flux terms with gradients in water content (i.e., l) as the driving force and effective transport coefficients for diffusion, Ddiff

eff ( ) , and hydraulic permeation, Dhydr

eff ( ) ,

vv

N nj

FD Dd

pdiffeff

hydreff( ) ( ) ( ) .

(6.54)

Direct comparison of Ddiffeff ( ) and Dhydr

eff ( ) showed that hydraulic permeation dominates at high l, whereas diffusion prevails at low l.

The hydraulic permeation model predicts highly nonlinear water content profiles, with strong dehydration arising only in the interfacial regions close to the anode. Severe dehydration occurs only at current densities closely approaching jpc. The hydraulic permeation model is con-sistent with experimental data on water content profiles and differential membrane resistance,213,214 as corroborated in Eikerling et al.7,8 The bare diffusion models exhibit marked discrepancies in comparison with these data.

Recently, it was shown that the hydraulic permeation model could explain the response of the membrane performance to variations in external gas pressures in operating fuel cells.215 Figure 6.15 shows data for the PEM resis-tance in an operational PEFC,

R

R L zz

L

PEM

s

p max

PEM p

dPEM( )

( ( )),

1

0

(6.55)

normalized to the resistance of the uniformly saturated membrane (at open circuit conditions), Rs LPEM/sp(l max), for various applied gas pressures in anode and cathode compartments (from Renganathan et al.215).


The experimental data (dots) are reproduced very well within the frame-work of the hydraulic permeation model (solid lines). For the basic case with zero gas pressure gradient between cathode and anode sides, ΔP g 0, the model (solid line) predicts uniform water distribution and constant mem-brane resistance at jp < 1 A cm–2 and a steep increase in R/Rs beyond this point. These trends are in excellent agreement with experimental data (open circles) for Nafion 112 in Figure 6.15. A finite positive gas pressure gradi-ent, ΔP g Pc

g – Pag > 0, improves the internal humidification of the mem-

brane, leading to more uniform water distribution and significantly reduced dependence of membrane resistance on l. The latter trends are consistent with the predictions of the hydraulic permeation model.

The previous discussion suggests that hydraulic permeation should be the dominant mode of water transport in Nafion at sufficiently large l, whereas a diffusive contribution to water transport will dominate at low l. This change in the prevailing mechanism of water transport with l could explain the peculiar transition in transverse water concentration profiles through operating PEMs observed in recent neutron-scattering experiments.216 Water management models that could account for this interplay of diffusion and hydraulic perme-ation have been discussed in Eikerling et al.7,8 and Weber and Newman.34

So far, water management models have assumed a controlled net water flux, jw, through the PEM. The basic case in Eikerling et al.7,8 considered jw 0. This approach is incomplete because it does not allow coupling of water fluxes in the membrane to water fluxes in other components and to externally

1.16

1.14

1.12

1.1

1.08

ΔPg = Pgc – Pg

a = 0

ΔPg = Pgc – Pg

a

Increasing

R/R

*

1.06

1.04

1.02

10 0.2 0.4 0.6 0.8 1

Current Density, i (A/cm2)

1.2 1.4 1.6 1.8 2

FIGURE 6.15PEM resistance in operational PEFC as a function of the fuel cell current density, compar-ing experimental data (dots) and calculated results from a performance model based on the hydraulic permeation model for various applied gas pressure differences between anode and cathode compartments. (Reprinted from S. Renganathan et al. Journal of Power Sources 160 (2006) 386–397. Copyright 2006, with permission from Elsevier.)


provided conditions of gases in electrodes. Continuity of water fluxes at the PEM interfaces with adjacent media requires a condition of interfacial water exchange, as suggested in Equation (6.22) on the anode side:

j

FRT

k PPP

GRT

s

w v a as a

v

as

a

2exp

( ).

(6.56)

and on the cathode side,

j

FRT

k PGRT

PP

s

w v c cs c c

v

cs2

exp( )

.

(6.57)

The factors xa and xc account for the heterogeneity of the interface. The interfacial flux conditions, Equations (6.56) and (6.57), can be straightforwardly applied at plain interfaces of the PEM with adjacent homogeneous phases of water (either vapor or liquid). However, in PEFCs with ionomer-impregnated catalyst layers, the ionomer interfaces with vapor and liquid water are randomly dispersed inside the porous composite media. This leads to a highly distributed heteroge-neous interface. An attempt to incorporate vaporization exchange into models of catalyst layer operation has been made217 and will be described in Section 6.9.4.

6.9 Ionomer in Catalyst Layers: Structure Formation and Performance

This section provides a comprehensive overview of recent efforts in physical theory, molecular modeling, and performance modeling of CLs in PEFCs. Our major focus will be on state-of-the-art CLs that contain Pt nanoparticle electrocatalysts, a porous carbonaceous substrate, and an embedded network of interconnected ionomer domains as the main constituents. The section starts with a general discussion of structure and processes in catalyst layers and how they transpire in the evaluation of performance. Thereafter, aspects related to self-organization phenomena in catalyst layer inks during fabrica-tion will be discussed. These phenomena determine the effective properties for transport and electrocatalytic activity. Finally, physical models of catalyst layer operation will be reviewed that relate structure, processes, and operat-ing conditions to performance.

From the insights presented, conclusions can be drawn about reserves for improvements in catalyst effectiveness, voltage efficiency, water handling capabilities, stability through optimized operating conditions, and advanced structural design. Due to recent comprehensive coverage,7,25,26 these topics will be revisited here with appropriate brevity.


6.9.1 Challenges for Design and Operation of Catalyst Layers

The foremost objective is to obtain the highest reactivity or transfer current density of desired electrochemical reactions with a minimum amount of Pt-based catalyst (DoE target for 2010: 0.29 g Pt per kilowatt). This demands a huge electrocatalytically active surface area; small kinetic barriers to the transport of protons, electrons, and reactant gases; and proper handling of product water and waste heat.

Because no single homogeneous phase could fulfill these conflicting needs simultaneously, CLs require composite morphologies that consist of sev-eral interpenetrating phases. A minimum of two distinct phases is needed, including a solid phase of nanoparticle catalyst (Pt) and electronically con-ducting substrate (carbon) and a liquid water phase in the void spaces of the substrate for diffusion and permeation of protons, water, and reactant molecules.

Impregnating these layers with PFSA ionomer for enhanced proton con-duction or hydrophobizing agents like Teflon for sufficient gas porosity is optional. However, ionomer impregnation is indispensable in CLs with thicknesses of LCL > 1 mm. Ultrathin CLs with LCL 100–200 nm, on the other hand, can operate well without these additional components, based on suf-ficiently high rates of transport of dissolved reactant molecules and protons in liquid water, which could ensure uniform reaction rate distributions over the entire thickness of the layer.

The typical thickness of currently used CLs is LCL 10 mm. They therefore require impregnation with an ionomer, usually Nafion. Self-organization of ionomer and carbon/Pt in colloidal ink solution leads to the formation of agglomerated morphologies. Agglomerates (with radii Ra 30–100 nm) con-sist of primary particles of carbon (with sizes in the range of 5–10 nm) onto which Pt nanoparticles are deposited. Resulting structures possess bimodal, bifunctional pore size distributions (PSDs), with primary pores (with pore radii rm 1–10 nm) inside agglomerates between the primary Pt/C particles and secondary pores (rM 10–50 nm) between agglomerates. Secondary pores and ionomer domains compete to occupy the voids between agglom-erates. The composition is specified in terms of volume fractions of the solid carbon/Pt phase, XPtC, the ionomer phase, Xel, and the remaining pore space, Xp 1 – XPtC – Xel. The total porosities lie in the range of XP 30–60%. The vol-ume fractions of primary and secondary pores are Xm and XM, respectively.

The primary optimization target of CLs is the effectiveness factor of Pt uti-lization, CL. It includes a factor, stat, that accounts for statistical limitations of catalyst utilization that arise on a hierarchy of scales, as specified in the following equation. stat determines the exchange current density217:

j m j g0 3 02 10 [ ] , (*

Pt stat statS/V

awith SSf X X

XrPtC, el

PtC

)( )

.

(6.58)


The structure-related statistical factors include the surface-to-volume atom ratio of Pt nanoparticles, eS/V, the effectiveness factor of catalyst utilization of mesoscopic agglomerates, a, and percolation and wetting effects at the macroscopic level, represented by the functions f(XPtC,Xel) and g(Sr), where Sr is the liquid saturation.

In addition to these structural effects, the effectiveness factor, CL, accounts for nonuniform reaction rate distributions due to mass transport limitations at finite operating current densities of PEFCs. In simple one-dimensional electrode theory, the interplay of j0 with transport parameters of reactants, protons, and electrons determines the “reaction penetration depth,” CL.7 The criterion for uniform reaction rate distributions is that the reaction pen-etration depth is comparable to or larger than the thickness of the CL (i.e., dCL > LCL).

A one-dimensional model for the physical evaluation of effects of structure and distributed processes on stationary catalyst layer performance requires a minimum of two phenomenological parameters: the exchange current density, j0, and the reaction penetration depth, dCL. Each is a complicated but unique function of structure and operating conditions. Dealing with the CCL only and neglecting further complicating traits—for instance, issues related to the liquid water balance—these two parameters uniquely deter-mine the irreversible voltage loss incurred by the CCL, h0. Therefore, rela-tions to voltage efficiency, energy density, power density, and effectiveness of catalyst utilization at given current density and catalyst loading can be easily established.

For illustration purposes, we consider here a simple scenario of this interplay. We evaluate the effectiveness factor at a fixed cell voltage and thus at a fixed h0. We can express the corresponding current density as a two-variable function, j0 f(j0, dCL), where the reaction penetration depth, dCL, depends on h0. This function can be used to determine the effective-ness factor, CL. In the case of severely limited oxygen diffusion, the follow-ing relations for local oxygen partial pressure and current density can be obtained:

,CL

CL CL

p zL z

L

j z IL

( ) exp

( )

1

CCL

CL

CL

CL CL

1 1expL z

L

,CL

CLLjI b

00

2exp

(6.59)


wherez 0 is the PEM–CCL boundaryz LCL is the CCL–GDL boundaryp P PO O2 2

/ is the nondimensional oxygen partial pressure

PO2 is the O2 partial pressure at the CCL–GDL interface, b RT/a cntF, the

Tafel-parameter.

The combined parameter,

I

FP D

RTL

4 Oo

CL

2

is a characteristic current density of diffusive flux through the layer (with a unit of amperes per square meter). For a fixed h0, the overall effectiveness of Pt utilization can be defined by

CL stat

Pt 0id

jm j

jj

0

3 00

2 10 [ ],

*

(6.60)

where stat accounts for statistical factors as discussed before, and d is the ratio of the actual current density, including transport limitations, relative to the ideal current density, j0

id, that would be obtained if reaction rates were distributed in an ideal, uniform way without any transport losses. Using Equations (6.59) and (6.60), it can be demonstrated that

CL stat

CL

CL

CL

CLLL

1 exp .

(6.61)

In general, Equations (6.59) and (6.61) highlight the importance of adjusting thickness and effective properties of transport and reaction in CLs in such a way that dCL LCL. If we replace dCL by h0, using Equation (6.59), we obtain an explicit expression for CL as a function of the catalyst layer voltage loss:

CL stat

jI b

jI

00

00

21

2exp exp exp

bb.

(6.62)

Equations (6.59)–(6.61) represent a highly simplified scheme for evaluating various catalyst layer designs. Refinements of this crude framework for evaluat-ing catalyst layer performance should address all transport limitations, account for water accumulation, and include two- and three-dimensional effects.

6.9.2 Multiscale Modeling Scheme of Catalyst Layers

In general, we expect valuable insights for the advanced design of catalyst layers from understanding the microstructure of interconnected phases of


agglomerates, ionomers, and pore spaces. The random morphology of CLs has a marked impact on transport and electrochemical properties of the resulting composite medium. In principle, only catalyst particles at Pt–liquid water interfaces are electrochemically active. Electrochemical reactions thus occur on the walls of wetted pores inside and between agglomerates as well as at interfaces between Pt and wetted ionomer. Relative contributions of these distinct types of interfaces—Pt–water in carbon pores and Pt–water in ionomer pores—will depend on the corresponding interfacial areas and the percolating pathways for protons and reactants that lead to them.

Microstructures of CLs vary depending on applicable solvent, particle sizes of primary carbon powders, ionomer cluster size, temperature, wetting properties of carbon materials, and composition of the CL ink. These factors determine the complex interactions between Pt/carbon particles, ionomer molecules, and solvent molecules, which control the catalyst layer formation process. The choice of a dispersion medium determines whether the iono-mer is to be found in solubilized, colloidal, or precipitated forms. This influ-ences the microstructure and the pore size distribution of the CL.217 It is vital to understand the conditions under which the ionomer is able to penetrate into primary pores inside agglomerates. Another challenge is to character-ize the structure of the ionomer phase in the secondary void spaces between agglomerates and obtain the effective proton conductivity of the layer.

The modeling of structure and operation of CLs is a multiscale problem. Generally, physical modeling of CL operation takes place in two steps: (1) relating structure to physical properties of the layer (assumed as an effec-tive medium), and (2) relating effective properties to performance. The main structural effects in CLs occur at well-separated scales: at catalyst nanopar-ticles (rPt 2 nm), at agglomerates of carbon/Pt (Ra ~ 100 nm), and at the macroscopic device level (LCL ~ 10 mm), at which CLs can be considered as effective homogeneous media. Separate approaches in theory and modeling can be developed at these different scales.

At the smallest length scale, the specific exchange current density depends on the size, surface morphology, and surface electronic structure of the Pt nanoparticles as well as on the properties of the substrate.218–222 A refined understanding of the relations between particle size and electrocata-lytic activity is critical in view of the design of highly performing catalyst systems.223–225 Obviously, a reduction in particle size improves the surface-to-volume ratio, eS/V, of catalyst nanoparticles. Yet, the relation between particle size and activity is highly nontrivial because the size of the particles also affects electronic and geometric properties at their surface.226

Computational efforts using DFT calculations as well as kinetic model-ing of reactivities based on Monte Carlo simulations or mean field meth-ods have been employed to study elementary processes on Pt surfaces.227,228 Unraveling systematic trends in structure versus reactivity relations remains a formidable challenge due to the complex nature of structural effects in electrocatalysis.


The d-band model of Hammer and Norskov229–231 has been successful in relating trends in chemisorption energies for various adsorbates on transi-tion metal surfaces to the position of the d-band center—the first moment of the density of states from the Fermi level. Exploring this model, system-atic ab initio calculations based on DFT have been performed on a series of polycrystalline alloy films of the type Pt-M (with M Ni, Co, Fe, and Ti) and results have been compared with corresponding experiments. Predicted correlations among the position of the d-band center, oxygen chemisorp-tion energies, and electrode activities for the oxygen reduction reaction have been confirmed.232 This success of the d-band model has catalyzed efforts in devising DFT-based, high-throughput combinatorial screening schemes for identifying highly active electrocatalyst materials.233

So far, the success of DFT-based modeling schemes in electrocatalysis has been limited to studying elementary surface processes on catalyst systems with well-defined periodic slab geometries that mimic single-crystalline surface structures. These insights are not straightforwardly applicable to the design of supported nanoparticle electrocatalysts, however. The latter systems represent special scientific challenges due to effects of quantum confinement, irregular surface structures with a large portion of low-coordination atoms, and the widely unexplored role of the substrate.23,234–236

At the mesoscopic scale, interactions between molecular components control the self-organization phenomena between molecular compo-nents that lead to random phase segregation during fabrication of CLs.22 Mesoscale simulations allow evaluating key factors during fabrication of CLs. These simulations rationalize structural factors such as pore sizes, internal porosity, and wetting properties of internal/external surfaces of agglomerates. Moreover, dispersion media with distinct dielectric prop-erties can be evaluated in view of capabilities for controlling sizes of carbon/Pt agglomerates, ionomer domains, and the resulting pore net-work topology.

At macroscopic level, the overall relations between structure and perfor-mance are strongly affected by the formation of liquid water. Solution of such a model that accounts for these effects provides full relations among struc-ture, properties, and performance, which in turn allow predicting architec-tures of materials and operating conditions that optimize fuel cell operation. For stationary operation at the macroscopic device level, one can establish material balance equations on the basis of fundamental conservation laws. The general ingredients of a so-called “macrohomogeneous model” of cata-lyst layer operation include:

source terms for electrochemical current conversion, employing Butler–Volmer equation or first principles of transition state theory, and for the transformation of water at interfaces (vaporization, condensa-tion); and


terms that account for the transport of species—that is, the migration of electrons and protons in conduction media, the diffusion of dis-solved oxygen and protons in water-filled pores, the diffusion of oxygen and vapor in gas-filled pores, and the permeation of liquid water in water-filled pores.

6.9.3 Mesoscale Simulations of Self-Organization in Catalyst Layers

Mesoscale calculations, discussed for the membrane in Section 6.5.3, provide insights into segregation behavior, structural correlations, and dynamical behavior of different phases in CLs. They contribute to furnishing relations among structure, transport properties, and reactivity. Compared to hydrated ionomer membranes (Section 6.5), structural complexity is more pronounced in CLs.

Coarse-grained molecular dynamics simulations in the presence of solvent provide insights into the effect of dispersion medium on microstructural properties of the catalyst layer.22 To explore the interaction of Nafion and sol-vent in the catalyst ink mixture, simulations were performed in the presence of carbon/Pt particles, water, implicit polar solvent (with different dielectric constant e), and ionomer. Malek et al. developed the computational approach based on CGMD simulations in two steps.22 In the first step, groups of atoms of the distinct components were replaced by spherical beads with predefined subnanoscopic length scale. In the second step, parameters of renormalized interaction energies between the distinct beads were specified.

Figure 6.16 shows a snapshot of the carbon–Nafion–water–solvent (CNWS) blend. The final microstructure was analyzed in terms of density map profiles, RDFs, pore size distributions, and pore shapes. The interaction parameters of the carbon particles were selected to mimic the properties of VULCAN-type C/Pt particles.

Structural analysis based on site–site RDFs is shown in Figure 6.17. There is a strong correlation between carbon particles. As expected, hydrated pro-tons (H) and water (W) behave similarly. The correlation between hydro-philic species (H and W) and ionomer (N) is significantly stronger than that between those species and carbon (C). The autocorrelation functions, gSS and gHH, exhibit a similar structure as gWW. This indicates a strong cluster-ing of side chains and hydronium ions due to the aggregation and folding of polymer backbones. However, the primary S-S and H-H peaks are sup-pressed compared to the primary peak in gWW due to electrostatic repulsion between these charged beads. gBB and gCC exhibit upturns toward larger r, superimposed on the primary bead–bead correlations, which correspond to the characteristic dimensions of carbon particles (~5 nm diameter) and back-bone clusters (~2–3 nm).

More detailed analysis of RDFs revealed a strong correlation of carbon par-ticles and polymer backbones (gCB). This suggests that polymer backbones


are attached to the surfaces of carbon particles, while side chains strive to maximize their separation from the surface of carbon agglomerates.

Overall, the correlation functions discussed in detail in Malek et al.22 pro-vide valuable structural information at the nanometer scale that allows refin-ing the picture of the phase-segregated catalyst layer morphology. Ionomer

8

7

6

W-W

W-W

S-S

S-S

H-H

H-H

5

4

3

g (

r)

2

1

00 1 2

r (nm)

3 4

10

8

6

B-B

B-B

C-C

C-C4

g (

r)

2

00 1 2

r (nm)

3 4

2.5

2

1.5

C-B

C-B

C-S

C-S

1g (

r)

0.5

00 1 2

r (nm)

3 4

FIGURE 6.17Site–site radial distribution functions for the CNWS system (C: carbon; P: polymer backbones; W: water; H: cluster containing hydronium). (Reprinted from K. Malek et al. Journal of Physical Chemistry C 111 (2007) 13627. Copyright 2007, with permission from ACS.)

FIGURE 6.16Equilibrium structure of a catalyst blend composed of carbon (black), Nafion (dark gray), water (light gray), and implicit solvent. (Reproduced from K. Malek et al. Journal of Physical Chemistry C 111 (2007) 13627. Copyright 2007, with permission from ACS.)


backbones form clusters or fibers that are attached to the carbon agglomer-ates. Water and hydronium ions tend to maximize their separation from the carbon while trying to stay in the vicinity of the side chains. A key finding of these simulations is that no evidence of ionomer penetration into primary pores inside carbon agglomerates has been found.

Malek et al. examined the effect of the solvent dielectric constant on struc-tural correlations.22 Overall, the magnitude of the RDF peaks decreases from low to high polar solvent. Low increases correlations between carbon particles and hydrophobic polymer backbones, enhancing the tendency to phase-segregate into hydrophobic and hydrophilic domains. In the presence of an apolar solvent, structural correlations of hydrophilic and hydrophobic domains extend to larger separation distances compared to polar solvents. The dielectric properties of the solvent strongly affect the size and connec-tivity within ionomer and carbon clusters. The average agglomerate size decreases markedly, from 33 nm for e 2 to 15 nm for e 80. Concomitantly, ionomer domain sizes decrease from 11 to 10 nm. The effect of solvent on domain sizes is thus more pronounced for carbon agglomerates than for ion-omer aggregates. Figure 6.18 depicts size distributions of ionomer domains in CLs obtained with different e.

These coarse-grained MD calculations helped consolidate the main fea-tures of microstructure formation in CLs of PEFCs. They showed that the final microstructure depends on carbon particle choices and ionomer–carbon

Dielectric = 2Dielectric = 20Dielectric = 80

8000

6000

4000a.u

.

2000

00 5 10

Ionomer Cluster Size (nm)

15

FIGURE 6.18Effect of polarity of the solvent on calculated domain size distributions of ionomer. (Reprinted from K. Malek et al. Journal of Physical Chemistry C 111 (2007) 13627. Copyright 2007, with per-mission from ACS.)


interactions. For carbon meterials with hydrophilic surface, ionomer side chains are buried inside hydrophilic domains with a weak contact to carbon domains and the ionomer backbones are attached to the surface of carbon agglomerates. The correlation between hydrophilic species and ionomer is significantly stronger than that between those species and carbon particles. A densely packed layer of ionomer of ~10 nm thickness formed around Pt/C agglomerates was evident in the presence of polar solvents.

The evolving structural characteristics of CLs are particularly important for further analysis of transport of protons, electrons, reactant molecules (O2), and water as well as for the distribution of electrocatalytic activity at Pt–water interfaces. In principle, the mesoscale simulations allow relating these properties to the choices of solvent, ionomer, carbon particles (sizes and wettability), catalyst loading, and hydration level. Explicit experimental data with which these results could be compared are still lacking. Versatile experimental techniques have to be employed to study particle–particle interactions, structural characteristics of phases and interfaces, and phase correlations of carbon, ionomer, and water in pores.

Recently, 500 MHz 19fluorine NMR was used to study adsorption of Nafion ionomer on PEFC catalysts and the supporting carbons in aqueous solution.237 It was observed that Nafion adsorbs strongly on carbon as well as on Pt and PtRu. The adsorption was classified into primary and secondary adsorption. At low concentration of Nafion ionomer, the adsorption was found to follow a Langmuir isotherm (primary adsorption). Although there was uncertainty in the types of adsorption isotherms at high concentration of Nafion iono-mer, the secondary adsorption isotherms were fitted to a Langmuir isotherm as well.

6.9.4 Main Results of Macrohomogeneous Catalyst Layer Models

In the past, studies using macrohomogeneous models of CL operation have explored the effects of thickness and composition on performance and cata-lyst utilization. The specific effects due to the complex coupling of porous morphology, liquid water formation, oxygen transport, and reaction rate distributions will be discussed in a separate section later. In this section, results will be presented assuming conditions under which the effects of liquid water accumulation on performance can be neglected. This assump-tion is valid for current densities below a critical value, which is roughly ~1 A cm–2.

The relevant solutions of the macrohomogeneous model for the case of negligible agglomerate effects have been discussed in detail.7,238–240 Analytical relations for reaction rate distributions and relations between fuel cell current density and overvoltage losses in the CCL were obtained for limiting cases of fast oxygen diffusion and fast proton transport. Eikerling and Kornyshev241 included double layer charging into the model.


With this extension, the complex impedance response of the CCL could be calculated. The model of impedance amplifies diagnostic capabilities—for example, providing the proton conductance of the CCL from the linear branch of impedance spectra (in Cole–Cole representation) in the high-frequency limit.

Three regimes of current density exist in electrode polarization curves:

1. a kinetic regime at small current densities, j0 << I, with simple Tafel dependence, h0 bln(j0/j0), where b RT/acF;

2. an intermediate regime for j0 ~ I with prevailing double Tafel-slope characteristic,

0

024

bjI

IFD p

RTLln , ,with

oL

CL (6.63)

where pL p(z LCL) is the oxygen partial pressure at the CCL–GDL interface; and

3. an oxygen depletion regime for j0 > I, in which all oxygen is con-sumed in a sublayer of thickness dCL, defined in Equation (6.64).

For a given composition, the thickness and the target current density of fuel cell operation should be adjusted in order to operate the catalyst layer in the intermediate regime because this represents the best compromise between transport losses and kinetic losses. Although reaction rate distribu-tions exhibit a pronounced nonuniformity in this regime, the layer uses all parts. There are thus no inactive parts. As long as the CCL is operated in the intermediate regime, overvoltage losses are almost independent of the thick-ness. In Eikerling et al.7 and Eikerling, Kornyshev, and Ioselevich,240 these findings were displayed in the form of a phase diagram.

The existence of a maximum thickness beyond which the performance deteriorates is due to the concerted impact of oxygen and proton transport limitations. Considered separately, each of these limitations would only serve to define a minimum thickness below which performance worsens due to an insufficient electroactive surface. The thickness of the effective layer, in which current density is predominantly generated, is given by the reaction penetration depth:

CL CL

Ij

L0

.

(6.64)

In the oxygen depletion regime, j0 >> I, only a thin sublayer with thickness dCL << LCL, adjacent to the GDL, is active. The remaining sublayer with thickness (L – dCL) LCL, adjacent to the PEM side, is not used for reactions,


due to the starvation in oxygen. This highly nonuniform reaction rate dis-tribution implies underutilization of the catalyst. The inactive part incurs ohmic losses of proton transport in the embedded ionomer phase, which could become rather detrimental if the proton conductivity is poor. In sum-mary, dCL LCL should be warranted. This corresponds to operation of the CCL in the intermediate regime, which ensures a good compromise of high performance (i.e., minimal overpotential losses) and effective catalyst utilization.

The macrohomogeneous model was exploited in optimization studies of the catalyst layer composition. The theory of composition-dependent per-formance reproduces experimental findings very well.242–246 The value of the mass fraction of ionomer Yel that gives the highest voltage efficiency for a CCL with uniform composition depends on the current density range. At intermediate current densities, 0.5 A cm–2 j0 1.2 A cm–2, the best perfor-mance is obtained with Yel 35 wt%. The effect of the Nafion weight fraction on performance predicted by the model is consistent with the experimental trends observed by Passalacqua et al.246

The model was also used to explore novel design ideas. It was predicted and subsequently confirmed in experiment that functionally graded layers result in improved performance compared to standard CCLs with uniform composition.243,244 In this design, the catalyst layer is fabricated as a sublayer structure with gradually varying composition. Compared with the layer of uniform composition with Yel 35 wt%, a three-sublayer structure with 30 wt% Nafion content in the sublayer toward the GDL side, 35 wt% Nafion content in the middle sublayer, and 40 wt% Nafion content toward the mem-brane side improves E( j0) by about 5%. Another advantage of this design is the reduced ohmic resistance at the PEM–CCL interface due to the improved contact area between electrolyte phases at higher Nafion loading on the CCL side. In the interface between GDL and CCL, lower Nafion loading will decrease the probability of blockage of pores by Nafion and thereby facilitate water removal via the GDL.

6.9.5 Water Management in Catalyst Layers

Liquid water arrives in the CCL via transport through the PEM or it is gen-erated in the electrochemical reaction. Invariably, PEFCs require a medium that is highly effective in transforming liquid water into water vapor; other-wise, liquid water will clog pores and channels in gas diffusion layers and flow fields that are needed for the gaseous supply of reactants.

The net rate of vaporization of a porous electrode with thickness LCL is given by

Q z

eL

S q T q zCL

lve

lvr r

s( ) ( ) ( ) ( ) ,0

(6.65)


with local vapor pressure, q(z), and the saturated vapor pressure in pores with capillary radius, rc:

q T q T

VRTr

qrs s, m

cs,( ) ( )exp

cos( ),

2(( ) exp ,T q

Gk T

0a

B (6.66)

where qs,∞(T) is the saturation pressure for a planar vapor–liquid inter-face s is the surface tension q is the wetting angle of catalyst layer pores e0 is the elementary charge ke is an intrinsic rate constant of evapora-tion247 and lv is the ratio of the distributed liquid–vapor interfacial area to the apparent electrode surface area. For a porous medium with total porosity, Xp,rc

, at a given rc, a simple estimate of this factor gives

lv

p,r cc~ ,X L rCL 2

assuming cylindrical pores with radius rc and lengths lc ~ 4rc. An activation energy of evaporation, Ga 0.44 eV, and a preexponential factor, q0 1.18 106 atm, reproduce saturation pressures of water in the range of 0–100 C. Assuming a typical catalyst layer thickness of 10 mm, Eikerling217 demon-strated that evaporation rates in CCLs are high enough to keep up with rates of water production. The comparison depends strongly, however, on T and on the porous structure. Higher temperatures and smaller pore sizes increase evaporation rates.

Obviously, the CCL not only determines the rate of current conversion and the major portion of irreversible voltage losses in a PEFC, but also plays a key role for the water balance of the whole cell. Indeed, due to a benign porous structure with a large portion of pores in the nanometer range, the CCL emerges as favorite water exchanger for PEFCs. Once liquid water arrives in gas diffusion layers or flow fields, PEFCs are unable to handle it.

The challenge for modeling the water balance in CCL is to link the com-posite, porous morphology properly with liquid water accumulation, trans-port phenomena, electrochemical kinetics, and performance.217,248 At the materials level, this task requires relations between composition, porous structure, liquid water accumulation, and effective properties. Relevant properties include proton conductivity, gas diffusivities, liquid permeability, electrochemical source term, and vaporization source term. Discussions of functional relationships between effective properties and structure can be found in the literature.217,239,240,248 Because the liquid water saturation, Sr(z), is a spatially varying function at j0 > 0, these effective properties also vary spa-tially in an operating cell, warranting a self-consistent solution for effective properties and performance.

It is assumed that capillary forces at the liquid–gas interfaces in pores equilibrate the local water content in the catalyst layer. Pore-filling under stationary conditions is therefore expressed through the Young–Laplace


equation, which relates capillary pressure, pc, and capillary radius, rc, to local gas pressure, pg, and liquid pressure, pl:

p

rp pc

cg l2 cos( )

.

(6.67)

Under normal operating conditions, the problem on the cathode side is to deal with excessive amounts of liquid water due to a net water flux through the PEM and the production of water on this side. Under these conditions, it can be assumed that the ionomer phase in the CCL is fully hydrated. Moreover, an assumption was made that small primary pores inside agglomerates are hydrophilic, ensuring that these pores are filled completely with water. The liquid water front therefore advances in secondary pores between agglom-erates. The wettability of these pores is a vital property that controls the water formation in CCLs. Assuming that these pores are partially hydro-philic (typical wetting angles of carbonaceous materials are in the range of q ~ 80–90 ), the liquid saturation is

SX

rX r

r

r1

0p

pdd

d

c

'( )

.

(6.68)

Equations (6.67) and (6.68) establish relations among local water content, operating conditions, wetting properties, and pore space morphology in the CCL.

Finally, continuity equations that account for mass conservation laws,

( , )( , ) ( , )

v v v vr tt

j r t Q r t

(6.69)

—including fluxes of protons, gases, and water as well as transformations between these species—have been established to relate the CCL structure and effective properties to performance.

Eikerling217 has demonstrated capabilities of this approach. A simple rep-resentation of the pore space by a bimodal -distribution reveals the role of the CCL as a “watershed” in PEFCs. For this case, a full analytical solution could be found. At the same time, it still captures essential physical processes and major structural features such as typical pore sizes (rm, rM), and distinct contributions to porosity from primary and secondary pores (Xm, XM).

In terms of liquid water saturation and water management in the CCL, the bimodal -distribution leads to a three-state model. Effective properties are constant in each of these states. In the dry state, the porous structure is water-free (Sr 0). Gaseous transport is optimal. Electrochemical reaction and evaporation rates are poor, however, because g 0 and xlv 0. In the optimal wetting state (Sr Xm/Xp), primary pores are completely water filled while secondary pores are water free. Catalyst utilization and exchange


current density are high (g 0.8–0.9); the surface area for evaporation, xlv, is large as well. Moreover, diffusion coefficients will still be high because secondary pores remain water-free. In the fully flooded state (Sr 1), all pores are filled with water and the affected parts of the CCL are deactivated due to impeded gas transport at macroscopic scale.

An expression for the current density, below which complete liquid-to-vapor conversion is possible, was obtained. This characteristic parameter is related to saturated vapor pressure and vapor diffusion. Moreover, the CCL fulfills an important function in regulating hydraulic fluxes toward PEM and GDL sides.

The results also suggest that the CCL is a critical fuel cell component in view of excessive flooding. Critical liquid water formation arises first in the interior of the layer, close to the CCL–GDL interface and not at the PEM–CCL boundary. The model reveals sensitive dependence of CCL operation on porous structure, thickness, wetting angle, total cathodic gas pressure, and net liquid water flux from the membrane. With rather favorable param-eters (10 mm thickness, 5 atm cathodic gas pressure, 89 wetting angle), the critical current density of CCL flooding is found in the range of 2–3 A cm–2. For increased thickness, smaller gas pressure, or slightly reduced wetting angles of secondary pores, CCLs could be flooded at current densities well below 1 A cm–2. The contact angle is an important parameter in this con-text, highly sensitive in view of water handling, but difficult to control in fabrication.249,250

Liu and Eikerling248 presented solutions of the physical model of CCL operation for general continuous pore size distributions. With this decisive extension, the full coupling of composite porous morphology, liquid water accumulation, transport of reactants and products, and electrochemical conversion in the oxygen reduction reaction could be explored. Continuous PSDs allow relating global performance effects (limiting currents, bistabil-ity) to local distributions of water, concentrations of reactants, and reaction rates in the layer.

It was found that a CCL alone cannot give rise to limiting current behavior in voltage–current response curves. The explanation is simple and intuitive: In the fully saturated state, the CCL retains a residual oxygen diffusivity through liquid water-filled pores, Dfl

o 2.0 10–6 cm2 s–1; the main effect of flooding in the CCL will be a reduction of the reaction penetration depth by about a factor of 100 (i.e., from ~10 mm in the nonflooded CCL to ~100 nm in the flooded CCL); the voltage losses incurred by the CCL will increase accordingly so that the smaller fraction of active Pt atoms in the active part can generate the fixed total current. The increase in voltage losses depends roughly logarithmically on the factor by which the reaction penetration depth is reduced. This dependence is a consequence of the Butler–Volmer equation.

Upon increase of the current density generated by the fuel cell, a tran-sition between two principal states of operation occurs, as illustrated in


Figure 6.19. The ideally wetted state at low current densities exhibits levels of liquid water saturation well below the critical value for pore blocking, corresponding to relatively uniform distributions of reactants and reaction rates. In the fully saturated state, liquid water saturation exceeds the critical value in parts of the layer. These parts could sustain only low residual gas diffusivity. Corresponding reactant and reaction rate distributions will be highly nonuniform, rendering the main part of the CCL inactive. The tran-sition between the two states of operation can occur monotonously or can involve bistability as a signature of nonlinear coupling of liquid water accu-mulation, gaseous diffusion, and electrochemical conversion rate. Bistability means that two steady-state solutions of the continuity equations coexist in the transition region.

The current density of the transition from ideally wetted state to transition region or fully saturated state is a key parameter for optimization of CCLs in view of their water-handling capabilities. A larger value of this critical cur-rent density allows extracting higher voltage efficiencies and power densities from PEFCs. Critical current densities depend on structural parameters and operating conditions.

Stability diagrams have been introduced for assessing effects of param-eters on performance. The stability diagram in Figure 6.20 displays the effect of the wetting angle, distinguishing among ideally wetted state, bistability region, and fully saturated state. Simply stated, the task of water manage-ment in CCLs is to push back capillary equilibrium to small enough pores so that liquid water formation cannot block gaseous transport in secondary

PrimaryPores Filled

Transition

Region

Fully

Flooded

1

0.9

0.8

0.7

E/V

0.6

0.5

j0/A cm–2

0.4 0.60.20 0.8 1

FIGURE 6.19Plot of fuel cell voltage versus current density showing the transition between two principal states of operation corresponding to ideally wetted conditions with primary pores filled and secondary pores available for gas diffusion and fully flooded conditions. In the depicted case, the transition involves a bistability.


pores. Beneficial conditions in view of this objective are high total porosity, large volume fraction of secondary pores (XM > 0.3), wetting angle that closely approaches 90 , high total gas pressure, and high temperature of operation.

6.10 Concluding Remarks

Numerous demonstrations in recent years have shown that the level of per-formance of present-day polymer electrolyte fuel cells can compete with current energy conversion technologies in power densities and energy effi-ciencies. However, for large-scale commercialization in automobile and por-table applications, the merit function of fuel cell systems—namely, the ratio of power density to cost—must be improved by a factor of 10 or more. Clever engineering and empirical optimization of cells and stacks alone cannot achieve such ambitious performance and cost targets.

Rather, the success of fuel cell technology hinges on major break-throughs, not incremental improvements, in design and implementation of advanced materials that are specifically optimized to meet targets in performance, operating conditions, lifetime, and cost. In particular, the required improvements would be impossible without advances in the

FIGURE 6.20The stability diagram displays the effect of the wetting angle on CCL operation, distinguish-ing among ideally wetted state, bistability region, and fully saturated state. (Reprinted from J. Liu and M. Eikerling. Electrochimica Acta 53 (2008) 4435–4446. Copyright 2008, with permission from Elsevier.)


concepts of proton-conducting membranes and catalyst layers. Due to recent detailed accounts of the status of catalyst layer research, the pres-ent chapter has mainly highlighted the challenges and the progress in understanding hydrated ionomer membranes.

Physical models of fuel cell operation contribute to the development of diagnostic methods, the rational design of advanced materials, and the sys-tematic optimization of performance. The grand challenge is to understand relations of primary chemical structure of materials, composition of hetero-geneous media, effective material properties, and performance. For polymer electrolyte membranes, the primary chemical structure refers to ionomer molecules, and the composition-dependent phenomena are mainly deter-mined by the uptake and distribution of water.

For several reasons, our account of the state of membrane research focused on Nafion-type materials. This membrane type still represents the bench-mark in fuel cell science and technology. Theory and modeling rely on a consistent set of experimental data to be able to develop and corroborate the multiscale relations between micromorphology, transport, and operation, encompassing scales from angstroms to meters. Although many experimen-tal studies exist for alternative membranes, Nafion provides the most com-plete experimental characterization of structure and phenomenology. As a matter of course, it is expected that the bulk of insights from physical theory and modeling presented in this chapter is of a general value, easily adaptable for alternative membrane materials.

Knowledge of the supramolecular morphology of Nafion is the basis for understanding the principles of membrane operation. Self-organization of the hydrophobic backbones and hydrophilic ionic side chain branches leads to the formation of elongated structures, visible in small-angle scattering data. A recently proposed model by Schmidt-Rohr and Chen of water chan-nels inside cylindrical inverted micelles, confined by polymer walls, rebuts Gierke’s spherical cluster model and refines the picture of elongated polymer bundles of Gebel and co-workers.

Successful strategies in modeling should integrate knowledge of the pri-mary chemical architecture of the ionomer and allow predicting the mor-phology, transport properties, and operation of the self-organized medium. This warrants a well-devised hierarchy of methods, which fall into two major categories:

fully self-consistent molecular-level simulations of the membrane architecture; and

models that predefine certain critical structural elements of the membrane (e.g., arrays of hydrated surface groups, polymer fibrils, single pore environments, or a random network of pores) and study the influence of well-defined structural parameters (e.g., length and grafting density of polymeric side chains or pore sizes) on properties and performance.


As progress is made on both fronts, the different approaches can inform each other and generate a coherent picture of membrane structure and operation.

Self-consistent approaches in molecular modeling have to strike a balance of appropriate representation of the primary polymer chemistry, adequate treatment of molecular interactions, sufficient system size, and sufficient sta-tistical sampling of structural configurations or elementary transport pro-cesses. They should account for nanoscale confinement and random network morphology and they should allow calculating thermodynamic properties and transport parameters.

As discussed in Section 6.5.3, coarse-grained molecular modeling approaches offer the most viable route to the molecular modeling of hydrated ionomer membranes. The coarse-grained treatment implies simplification in interactions, which can be systematically improved with advanced force-matching procedures, but allows simulations of systems with sufficient size and sufficient statistical sampling. Structural correlations, thermodynamic properties, and transport parameters can be studied.

The simulations reviewed in Section 6.5.3.2 furnish the picture of the self-organized, phase-segregated morphology. Structural features of domains of polymer and water and of the interfaces between them are in good agree-ment with the accepted structural models that were inferred from scattering experiments. Moreover, the obtained percolation threshold is in line with experimental conductivity data. The low percolation threshold indicates a highly interconnected network of water nanochannels. Notably, this perco-lation threshold is markedly smaller than those found in other simulations that employed significantly shorter representations of the ionomer. This find-ing emphasizes the importance of an appropriate length of the monomeric sequence in simulations.

Models composed of a two-dimensional array of polymer with predefined distribution of side chains or surface groups can mimic structure and trans-port properties at acid-functionalized polymer–water interfaces, as discussed in Section 6.7.3. They provide insights into the structure of surface water, cor-relation effects between surface groups, and fundamental transport mecha-nisms at the interface. It was found that the model of the minimally hydrated interface exhibits transition from hydrophilicity to hydrophobicity. This transition point is above a critical density of surface groups and corresponds to surface group separations of 7 Å. At these high densities of surface groups, rates of interfacial proton transport could be rather high.

The interesting transitions in interfacial conformation, water binding, and proton mobility occur at high surface group densities. The exploitation of these findings for the design of advanced polymers warrants systematic experimental studies at hydrated two-dimensional arrays with controlled surface group densities. Positive results of such studies could guide system-atic efforts in the synthesis of hierarchically structured polymer assemblies (e.g., utilizing block-copolymer architectures). The DFT-based calculations for the model system of a two-dimensional array could thus facilitate the


bottom-up design of advanced polymer materials that are optimized for operation under minimal hydration—that is, low relative humidity and/or elevated temperature (>100 C).

The distinction of surface and bulk water has served as a concurrent concept for this chapter. In sorption studies, surface water exhibits large Gibbs free energies of water binding, while Gibbs free energies of additional water mol-ecules correspond to capillary condensation of bulk-like water in pores. Based on this mechanism of water uptake, the thermodynamic model of water sorp-tion and swelling in PEMs was presented in Section 6.6.3. Vapor pressure and temperature of the adjacent gas phase determine the capillary radius, rc, up to which pores are swollen. The corresponding capillary pressure, P c, and the external gas pressure, P g, determine the liquid pressure, Pl, in swollen pores. This liquid pressure is balanced by the elastic pressure P el. The self-organized morphology of the membrane determines the relation between P el and the volume fraction of polymer fp. Overall, this causal chain thus describes how external conditions specified by temperature, relative humidity, and total gas pressure determine the equilibrium water uptake of the membrane. The model should be able to account for liquid and vapor equilibrated conditions.

The distribution of water in the membrane determines the transport prop-erties, which can be included in the model of membrane operation in PEFCs. As discussed in Section 6.8, hydraulic permeation transpires as the domi-nant mode of water transport at sufficiently large l, in analogy to a porous medium; a diffusive contribution to water transport will dominate at low l. Water management models that incorporate diffusion and hydraulic perme-ation are consistent with the physics of water sorption in the membrane and they can explain observations of membrane dehydration and nonlinearities in the differential membrane resistance under operation; the response of the membrane to changing operating conditions can be predicted and the role of thickness and porous morphology can be evaluated. Extensions of the exist-ing water management models of PEMs should include the coupling of water fluxes in the membrane and the adjacent porous electrodes.

In spite of significant progress in understanding of structure formation, water sorption, transport phenomena, and operation, the theory and model-ing of proton-conducting ionomer membranes for fuel cells continues to be a vibrant field of research. In order to furnish rational strategies in design and implementation of new membranes, major progress is needed in

modeling self-organization phenomena in hydrated ionomer systems;developing models of water sorption that link polymer composition,

morphology, and elastic properties with external conditions (RH, pressure, temperature);

rationalizing effects of polymer morphology on transport mechanisms with particular emphasis on the feasibility of proton transport at low relative humidity and high temperature;


understanding the coupling of water distribution and fluxes in the mem-brane with processes and conditions in adjacent porous media; and

understanding the factors that determine the stability of membranes.

Major performance targets for catalyst layers involve minimization of par-asitic voltage losses, balancing of water distribution and fluxes, and optimi-zation of the effectiveness factors of catalyst utilization. Due to the random composition, complex spatial distributions of electrode potential, reaction rates, and concentrations of reactants and water evolve under PEFC opera-tion. In spite of the structural complexity of catalyst layers, existing tools in molecular modeling and physical theory have contributed to the fundamen-tal understanding of structure formation and operation. Sources of voltage losses due to underutilization of the catalyst, impaired mass transport, and insufficient water management have been identified. Routes for the minimi-zation of these performance losses by optimized thickness, composition, and porous structure have been explored with some success.

Because the theory inevitably has to invoke quite a number of simplifying assumptions, often of an uncontrollable nature, offering a pure theoretically driven optimization would be irresponsible. The best strategy to approach the optimization of catalyst layers would be a concerted experimental–theoretical effort. Ex situ diagnostics is needed to characterize structural details and explore their relations to effective properties. The availability of such experimental data defines the level of detail of structure–property relationships that the theory could employ. In situ experimental studies, exploring the performance and com-paring it with the theoretical predictions, provide the essential benchmark for the modeling-based optimization of fuel cell efficiencies and power densities. Corroborated by these systematic experimental procedures, the theory could then be applied to

identify salient features of good or bad catalyst layer performance;detect causes of catalyst layer failure; andidentify requirements for advanced design of catalyst nanoparticles,

porous substrates (porosity, wetting properties), and composite architectures of catalyst layers (composition, thickness).

Symbols

av: activity of external vaporcmax: maximum water concentration in PEM (mol L)Dt: local diffusivity in PEM (QENS) (cm2 s–1)Dlr: long-range diffusivity in PEM (QENS) (cm2 s–1)Ds: self-diffusion coefficient in PEM (PFG-NMR) (cm2 s–1)


dCC: side chain separation in array of surface groups (Å)E: fuel cell voltage (V)Eeq: equilibrium electromotive force (V)ΔG: reaction Gibbs energy (kJ mol–1)ΔGs: Gibbs free energy of water sorption (kJ mol–1)ΔH: reaction enthalpy (kJ mol–1)I: characteristic current density of diffusive oxygen flux through CCL

(A cm–2)j0: exchange current density of CCL (A cm–2)jp: local proton current density (A cm–2)jw: water flux (A cm–2)Kangle: angle force constant (kJ mol–1 nm–2)Kbond: bond force constant (kJ mol–1)kP

m: hydraulic permeability (cm2)kv: vaporization exchange rate (cm s–1)LCL: CL thickness (mm)LPEM: thickness of PEM (mm)L: length of the pore (nm)m: effective resistance to water fluxMp: molar mass of polymer (g mol–1)N : effective number of polymer chains in resinrNl : molar flux of liquid water in the membranen0: number of SO3

– groups in the dry membranend: electro-osmotic drag coefficient in PEMnt: number of electrons transferred in the overall fuel cell reactionP: power density (kW cm–2)P c: capillary pressure (PEM modeling; atm)P el: elastic pressure in PEM (atm)P g: total gas pressure (PEM modeling; atm)Pl: liquid pressure (PEM modeling; atm)Ps: saturated vapor pressure (PEM modeling; atm)Pv: vapor pressure (PEM modeling; atm)pc: capillary pressure (CCL modeling; atm)pg: total gas pressure (CCL modeling; atm)pl: liquid pressure (CCL modeling; atm)q: coulombic charge (C)qs: saturated vapour pressure (CCL; atm)q: vapor pressure (CCL modeling; atm)Ra: agglomerate radius (nm)RPEM: membrane resistance (Ω cm2)r0: equilibrium bond length (nm)r c: capillary radius (CCL modeling) (nm)rc: capillary radius (PEM modeling) (nm)rij: effective bead radius (nm)


r rM: radius of primary and secondary pores in CCL (nm)ΔS: reaction entropy (J mol–1 K–1)T: temperature (K)Sr: liquid saturationU(r): external potential (kJ mol–1 K–1)Uel: coulombic potential (kJ mol–1)V0: volume of the dry resin (L)Vangle: harmonic angle potential (kJ mol–1)Vbond: harmonic bond potential (kJ mol–1)Vp : molar volume of polymer (L mol–1)Vw : molar volume of polymer (L mol–1)Xi: volume fraction of component i

Greek Symbols

aij: effective bead radius (nm)CL: effectiveness factor of CL

dCL: reaction penetration depth (mm)efc: fuel cell efficiencyeth: ideal theoretical efficiencyeV: voltage efficiencyhCCL: voltage losses incurred in the CCL (V)hother: parasitic voltage losses incurred in components other than PEM and

CCL (V)q: contact angle ( )l: water content in PEM, number of moles of water molecules per moles of

acid head groupslb: water content corresponding to bulk-like waterls: percolation threshold of water contentls: water content corresponding to surface watermw

PEM: chemical potential of water (kJ mol–1)n0: average pore volume in the dry membraner(r): density distribution (nm–3)rp: density of polymer (g cm–3)s: surface tension (N m–1)s bb: percolation bond conductance (S)sp: proton conductivity of PEM (S cm–1)fp: volume fraction of polymerfw: volume fraction of water

ij0: depth of the LJ potential well (kJ mol–1 K–1)


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89–106. 240. M. Eikerling, A. A. Kornyshev, and A. S. Ioselevich. Fuel Cells 4 (2004) 131–140. 241. M. Eikerling and A. A. Kornyshev. Journal of Electroanalytical Chemistry 475 (1999)

107–123. 242. M. Uchida, Y. Aoyama, E. Eda, and A. Ohta. New preparation method for

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243. Q. Wang, M. Eikerling, D. Song, S. Liu, T. Navessin, Z. Xie, and S. Holdcroft. Journal of the Electrochemical Society 151 (2004) A950–A957.

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246. E. Passalacqua, F. Lufrano, G. Squadrito, A. Patti, and L. Giorgi. Electrochimica Acta 46 (2001) 799–805.

247. K. J. Beverley, J. H. Clint, and P. D. I. Fletcher. Physical Chemistry Chemical Physics 1 (1999) 149.

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B 105 (2001) 7998. 250. M. L. Studebaker and C. W. Snow. Journal of Physical Chemistry 59 (1955) 973.

435

Index

2050-A, 1992050-HF, 1992050-L, 199

A

Acetylene black, 247Aciplex®, 138, 353AFM. See Atomic force microscopy

(AFM)Aluminum, 335Aluminum-magnesium alloy, 335Anode catalyst(s), 4

loadings, 4reformate-tolerant, 41–45

improved, 43–45mechanistic studies, 42–43

Atom-transfer radical polymerization (ATRP), 154

Atomic force microscopy (AFM), 117–118, 126, 145, 153

Atomistic simulations, 359–362ATRP. See Atom-transfer radical

polymerization (ATRP)Avcarb 1071HCB, 215

B

Bac2 Conductive Composites, 322BAM®, 112

Nafion vs., 114oxygen reduction and transport

characteristics, 120proton mobility, 113SPEEKK vs., 114

Bipolar plate(s)based composite, 316basic structure, 309–311comparative costs, 332competition between different

candidate materials and processing, 336

Fe-27Cr-6V alloy, 334

functions and performance requirements, 311–314

graphite, 315 (See also Graphite)graphite vs. SS316 for, 315from lab-scale to large quantity

production, 336–337materials

with better performance and durability, 334–335

comparison of major properties and cost of, 338

composite, 316–325carbon/carbon, 317–319filler, 316matrix, 316thermoplastic-based, 321–325thermoset-based, 319–321

metal, 325–333key fabrication processing of,

328–333materials used for, 326–327welding, 332–333

progress and challenges in development of, 315–333

traditionally used, 314–315new activities and development

trends, 333–337new manufacturing processes,

335–336performance targets of, 312requirements, 313schematic indication in simplified

PEM fuel cells, 310SS316 vs. graphite for, 315stack cost, 315transportation applications, 311, 312,

313Borides, 328, 334Boronization, 330, 331BP2000, 34BPSH. See Disulfonated biphenol

poly(arylene ether sulfone) (BPSH)

436 Index

C

Cabot Fuel Cells, 12Capillary flow porometry, 259Capillary pressure, 259Capillary rise method, 69Carbides, 36Carbon(s)

PTFE and, 223synthetic mesoporous, 39, 41

Carbon black(s), 37conventional, 37–38modification of, 38

Carbon cloth, 207–208fabrication, 207–209, 210SEM of, 208–209

Carbon corrosion, 30measurements, 34mechanism, 33–34Pt promoted, 36

Carbon fiber paper, 196–207with aerogels, 206–207large-scale production, 204–205manufacturing, 204–206SEM picture, 202–203

Carbon nanofibers, 38Carbon nanotube (CNT), 38

catalyst layer and, 77Carbon support materials, 32, 37–41

synthetic, 38–41mesoporous, 39, 41

Catalyst-coated gas diffusion electrode (CCGDL), 70–75

CCM vs., 77gradient, 71–75

catalyst, 71–73dual, 73–75Nafion, 73

uniform, 70Catalyst-coated membrane (CCM),

76–77CCGDL vs., 77conventional, 76–77nanostructural thin-film electrode,

77Catalyst layer (CL), 62

with additives, 80–81CNT-based, 77columnar oxide supported, 77–78

components and their corresponding function, 64–70

dual-bound composite, 75–76fabrication, 81–91

first generation, 81–82Pt black catalyst, 81–82PTFE-bound, 82

gas porosity, 404ionic conductivity, 70ionomer in, 403–419

challenges for design and operation, 404–406

macrohomogenous, 412–414mesoscale simulations of self-

organization in, 409–412multiscale modeling scheme,

406–409water management in, 414–419

location of, 62Nafion loading, 68nanowire-based three-dimensional

hierarchical core/shell, 79with novel ionomers, 81novel structural, 77–81optimization, 91–96

composition, 92–95experimental studies on, 93–95microstructure, 95–96modeling and simulation in, 92–93

porosity, 69properties, 66–70prospects and conclusions, 96–97Pt utilization and, 66–68PTFE-bound, 64self-supported, 80thin-film, 64, 70

dual-bound composite, 75–76fabrication, 83–91

DLR process in, 89–90dual ion-beam assisted

deposition in, 87–88electrodeposition in, 88electrophoretic deposition in, 91electrospray technique in,

90–91ink-based, 83–86inkjet printing in, 85–86pulsed laser deposition, 89

Index 437

RSDT in, 88–89screen printing in, 84–85in situ, 86–89sol-gel Pt application in, 91spray coating in, 85sputter deposition in, 86–87

nanostructured, 77, 78preparation of, 65types of, 70–76

CCGDL, 70–75 (See also Catalyst-coated gas diffusion electrode (CCGDL))

CCM, 76–77 (See also Catalyst-coated membrane (CCM))

types of, 70–81wetting property, 68–69

Catalyst performance targets, 4Cathode catalyst(s)

enhanced activity, 14–29Pt alloy, 14–20

loadings, 4stability, 29–37

alternative supports for, 35–37high voltage and, 32–34potential cycling and, 31, 32Pt electrochemical area loss and,

29–30toward high-voltage excursions,

34–35Cathode loadings, 4CCGDL. See Catalyst-coated gas

diffusion electrode (CCGDL)

CCM. See Catalyst-coated membrane (CCM)

Cell reversal tolerance, stability and, 46–47

Chemical vapor deposition, 223Chemical vapor-infiltrated graphite

carbon, 317Chromium, 330CL. See Catalyst layer (CL)Cladding process, 331CNT. See Carbon nanotube (CNT)CO2 poisoning, 42Coarse-grained molecular dynamics

(CGMD) simulations, 363–367, 409

Coating compounds, 328Cold-start tests, 279Contact angle of moving droplet,

254Contact angle tests, 254, 281Copolymerization, 144Copper diffusion layers, 220–221Corrosion studies, 279–280Cross-link density, 163Cross-linkers, 150, 156, 163Cyclic voltammogram, 67

D

Darcy’s coefficient, 399Darcy’s law, 261Density functional theory (DFT), 351,

408, 421calculations of water binding, 371hybrid, 360VASP and, 387

Desulfonation, 136Diffusion layer (DL), 192–288

capillary pressure, 259cold-start tests, 279contact angle of moving droplet, 254contact angle tests, 281corrosion studies, 279costs, 194direct visualization in DMFC, 267electricity conductivity, 273–274

in-plane, 273through-plane, 273–274

engineered, 215–221flow field interaction, 282–286future direction, 286–288gas transport properties, 260–267general transport properties,

255–260hydrophobic treatment, 227–233

(See also Hydrophobic treatments)

hydrophobicity and hydrophilicitiy, 251–255

in-plane permeability, 261–264internal contact angle, 254–255liquid transport properties, 267–272

permeability, 270–271mechanical properties, 276–278

438 Index

performance, 224–227in PEMFCs, 224–226

polarization curves showing effect of, 225

pore size distribution, 256–259porosity, 255–256pressure drop measurements,

260pressure drop tests, 282–284properties and measurements,

248–286sessile drop for, 251silicon-based, 221, 223thermal conductivity, 274–276thickness, 249–251through-plane permeability,

264–266transport, 255–272treatment and coatings, 227–248

hydrophilic, 233–234hydrophobic, 227–233microporous layers, 234–248

types, 196–227carbon cloth

commercially available, 210fabrication, 207–208SEM of, 208

carbon cloth fabrication, 207–209, 210

SEM of, 208, 209carbon fiber paper

commercially available, 198–201

fabrication, 196–207SEM of, 202–203

engineered, 215–221metal, 209, 211–221

foams, 215meshes, 211micromachined, 214SEM of, 212sintered, 213–214

silicon-based materials, 221–223visualization techniques, 284–286water balance analysis, 271–272Wilhelmy method for, 252–254

Diffusion medium. See Diffusion layer (DL)

Direct ethanol fuel cell, 211

Direct membrane fuel cells (DMFCs), 4, 28, 39, 52, 120

cell configurations, 29liquid-feed polarization curve, 50methanol crossover, 122performance degradation, 52PtRu particles for, 39, 52reviews of technology, 47

Direct methanol fuel cell, 211, 215cathode loadings, 4diffusion layers (See Diffusion layer

(DL))performance comparison of,

226–227hydrophilic treatment, 234hydrophobic treatment, 232–233microporous layers, 246–248visualization of gas bubbles in, 267

Dissipative particle dynamics (DPD), 363Disulfonated biphenol poly(arylene

ether sulfone) (BPSH), 120, 144–145, 153

mechanical properties, 130MeOh diffusion coefficients, 126Nafion 117 vs., 121

Divinylbenzene (DVB), 131, 156, 158, 159DLR process, in thin-film CL

fabrication, 89–90DMFC. See Direct membrane fuel cell

(DMFC)DMFCs. See Direct methanol fuel cells

(DMFCs)DoE Hydrogen Program, 318Double-walled nanotubes (DWNTs), 39DPD. See Dissipative particle dynamics

(DPD)Dual ion-beam assisted deposition, in

thin-film CL fabrication, 87–88

DVB. See Divinylbenzene (DVB)DWNTs. See Double-walled nanotubes

(DWNTs)

E

Ebonex, 35Electricity conductivity, 273–274

in-plane, 273through-plane, 273–274

Index 439

Electro-osmotic drag, 122–123, 394–397

Electrocatalyst(s), 5–52anode loadings, 4carbon support materials, 37–41cathode loadings, 4computation studies, 7–9discovery, 5–9electrochemical screening of, 8preparation, 9–13

colloidal routes for, 10–11conventional routes for, 9–10molecular precursor routes for,

11–12vapor phase routes, 12–13

scanning electron micrographs, 6testing, 13–14

Electrochemical impedance spectroscopy, 355

Electrodeposition, in thin-film CL fabrication, 88

Electron micrograph, restored high-resolution transmission, 4

Electron microscopy, 355Electrophoretic deposition, in thin-film

CL fabrication, 91Electrospray technique, in thin-film CL

fabrication, 90–91Enhancement layer, 214EP40, 198EP40T, 198ETFE-g-PSSA. See Ethylene-tetra-

fluoroethylenecopolymer (ETFE-g-PSSA)

Ethylene-tetrafluoroethylenecopolymer (ETFE-g-PSSA), 125

BAM vs., 114methanol uptake, 125Nafion vs., 114, 125oxygen permeability, 121oxygen reduction and transport

characteristics, 120proton conductivity, 111, 112

comparative studies of, 113, 114, 120, 121

water content and, 112radiation grafted, 157SPEEKK vs., 114

Ex situ visualization techniques, 270

F

Fe-and Co-based catalysts, 25–27Fe-Cr-V alloy, 334FEA. See Finite element analysis (FEA)Fenton test, 134FEP. See Fluoroethylenepropylene (FEP)Finite element analysis (FEA), 307Flemion®, 353Flow field interaction, 282Flow field plate. See Bipolar plate(s)Fluoride evolution rate (FER), 134–135Fluorinated ethylene-propylene

copolymer, 130–131Fluoroethylenepropylene (FEP), 227Forchheimer equation, 261Fortron PPS, 322Fourier transform infrared

spectroscopy, 355Furnace blacks, 37

G

Gas diffusion layer, 62, 306, 309. See also Diffusion layer (DL)

Gas-permeable elastomer, 223Gas transport properties, 260–267

in-plane permeability, 261–264relative permeability, 266through-plane permeability, 264–266

GD05505G, 200GD07508G, 200GD12012G, 200GD65055G, 200GD05505T, 200GD07508T, 200GD12012T, 200GD65055T, 200GDL10BA, 200GDL24BA, 201GDL25BA, 201GDL34BA, 201GDL10BB, 201GDL10BC, 201GDL24BC, 201GDL25BC, 201GDL34BC, 201GDS1120, 199GDS2120, 199

440 Index

GDS3215, 198GDS22100, 199Gebel’s calculations, 115GLAD. See Glancing angle deposition

(GLAD)Glancing angle deposition (GLAD), 78GRAFCELL, 320, 321Grafcell flexible graphite diffusion

layers, 216–220Graphite, 323, 328

alternative materials to, 315–316bulk electrical conductivity, 315chemical structure, 314cost, 315, 319, 330fillers, 316, 317

functional requirement of plates and, 320

ratio between matrix and, 320high-density, 314layered metals and thermal

expanded, 316metal vs., 333, 335natural, 314, 319

artificial vs., 320nonporous, 314normal carbon vs., 314physical and chemical properties,

314reinforced, 324shortcomings, 314SS316L vs., 315, 326, 327, 333synthesized, 319synthetic, 314thermoplastic-based composite,

316thermoset-based composite, 317, 319,

321Grotthuss mechanism, 109

H

H1-Pt catalysts, 11Hafnium, 326Hagen–Poiseuille–Kozeny equation,

399Hexafluoropropylene, 228High-throughput screening, 5–7High-velocity electromagnetic forming

process (HVEF), 335–336

HVEF. See High-velocity electromagnetic forming process (HVEF)

Hydrophilic treatment(s), 233–234measurement of, 68–69, 251–255for PEMFC, 233–234

Hydrophobic treatment(s), 227–233for DMFCs, 232––233effect, 229–232fabrication processes and

procedures, 227–229gas permeability and, 229measurement of, 251–255thermal conductivity and, 231water flooding and, 230

Hyflon® Ion E83, 353

I

Imidazole, proton conduction and, 169

In-plane conductivity, 273In-plane permeability, 261–264Infrared spectroscopy (IR), 355Inkjet printing, in thin-film CL

fabrication, 85–86Internal contact angle, 254–255Ionomer(s), 165–166

in catalyst layer, 403–419challenges for design and

operation, 404–406macrohomogenous, 412–414mesoscale simulations of self-

organization, 409–412multiscale modeling scheme,

406–409water management in, 414–419

membranesmolecular modeling of self-

organization in, 359–368perfluorinated sulfone, 96perfluorosulfonic acid, 353status and directions in research,

353–355structural organization and

dynamic properties of, 352–368

structure and dynamics in, 355–359

Index 441

K

Ketjen, 32, 34Key fabrication processing, 328–333

coating materials, 328–330coating process, 330–332forming process, 328

L

Lattice–Boltzmann simulations, 363Lennard–Jones potential, 365Liquid crystal polymer (LCP), 322Liquid transport properties, 267–272

permeability, 270–271Long term durability tests, 169LT1200-N, 199LT1300-N, 199

M

MEA. See Membrane electrode assembly (MEA)

Mechanical properties, 276–278Membrane electrode assembly (MEA),

4, 62, 63constituents, 192, 193cyclic voltammogram for cathode

of, 67description, 308targets, 4, 5

durability of, 6performance of, 5

MeOH oxidation catalyst(s), 43, 47–52mechanistic advancements, 47–48

MeOH-tolerant oxygen reduction catalysts, 27–29

Mercury porosimetry, 256Mesoscale simulations

coarse-grained, 362–368DPD in, 363Lattice–Boltzmann, 363SCMF theory in, 362

of self-organization, 409–412Metal(s)

for bipolar plates, 325–333for DL, 209, 211–221

foams, 215meshes, 211micromachined, 214

SEM of, 212sintered, 213–214

graphite vs., 333, 335welding, 332–333

Methanol crossover, 122–126diffusion mode vs. electro-osmotic

drag, 122–123electro-osmotic drag vs.diffusion

mode, 122–123sulfonation and, 123–124

Method of standard porosimetry, 257–258

Microelectromechanical system fabrication techniques, 221

Micromachined meshes, 214Microporous layers, 234–248

in DLFC, 246–248effect on fuel cell performance,

237–239fabrication processes, 236–237hydrophobic content, 240, 242multilayered, 244–246parameters affecting, 239–244pore formers in, 244thickness and carbon loading,

239–240types of carbon particles, 242–244

Molybdenum, 330Monte Carlo simulation, 407Multiple-walled nanotubes (MWNTs),

39MWNTs. See Multiple-walled

nanotubes (MWNTs)

N

Nafion®, 64, 65, 138coarse-grained model, 365degradation mechanism, 134MeOH diffusion flow, 123, 126microphase separation, 117microstructure, 115, 116

in hydrated membranes at different water contents, 366

morphology, 118N117, 117

BPSH vs., 121MeOH permeability, 123S-SEBS vs., 124

442 Index

oxygen reduction and transport characteristics, 120

phase separation, 114, 117proton conduction, 109

comparative studies of, 119proton mobility, 113PVDF and, 161QENS data for hydrated, 357source of protons, 111SPEEKK vs., 113, 114, 115, 117water absorption, 123

Navier–Stokes equation, 395Nb-doped TiO2, 36NEDO. See New Energy and Industrial

Technology Development Organization (NEDO)

Nernst–Einstein relationship, 110Neutron imaging in fuel cells,

268–269New Energy and Industrial Technology

Development Organization (NEDO), 4

Niobium, 326Nitridation, 330Nitrides, 36, 328, 334NMR. See Nuclear magnetic resonance

(NMR)NMR relaxometry, 357Non-Pt catalysts, 24–29Nonconductive whiskers, 36–37Nuclear magnetic resonance (NMR),

355

O

ONIOM method, 361Optical contact measurement system,

68Oxides, 35–36, 328Oxygen permeability, 119–122Oxygen reduction reaction, 238, 242

P

P50, 198P75, 198PAN. See Polyacrylonitrile (PAN)PBI. See Poly(benzimidazole) (PBI)Pd-based catalysts, 25

PEM. See Proton exchange membrane (PEM)

PEMFC. See Polymer electrolyte membrane fuel cell (PEMFC); Proton exchange membrane fuel cell (PEMFC)

Perfluorinated PEM, 140, 141Perfluorinated sulfone ionomers

(PFSIs), 96Perfluoroalkoxy polymer, 130Perfluorosulfonic acid ionomer

membranes, 353, 360, 363, 371

effect of hydration on local structure, 361

enhancing proton conduction with, 404

transport properties, 359Perforated pores, 217–219Perylene red, 37PFSA. See Perfluorosulfonic acid

ionomer membranesPFSIs. See Perfluorinated sulfone

ionomers (PFSIs)Phase separation, 114Platinum

alternative promoters to, 24electrochemical area loss and,

electrocatalyst stability and, 29–30

Platinum alloy catalysts, 14–20, 404. See also specific catalysts, e.g., PtRu catalyst

model surface studies, 15–17nanoparticles and particle size

effect, 17–20Platinum catalyst(s)

alternative, 50–52stabilization, 31, 34–35

layer for, 65Platinum core-shell catalysts, 20–24Polphosphazenes, 149Poly(4-vinylpyridine) (P4VP), 163Polyacrylonitrile (PAN), 196, 197, 207

carbon fiber manufacturing from, 204

Poly(aryl ether ketones), 143Polyarylenes, 142–149

Index 443

Poly(benzimidazole) (PBI), 143, 144, 163, 165, 169, 355

Poly(dimethylsiloxane), 223Poly(ether sulfone) (PES), 161, 163,

164Polyethylene substrates, 125Poly(ethyleneimine) (PEI), 163, 164Polyimides, 143, 144Polymer electrolyte fuel cell (PEFC)

catalyst layers, 348challenges for materials and

operation, 346–347energy conversion, 344–346hierarchy y of scales, 351–352membrane, 348

atomistic simulations of, 359–362conductivity, network model of,

390–393in fuel cell modeling, 397–403mesoscale coarse-grained

simulation, 362–368molecular modeling of self-

organization of, 359–368structural evolution of, 354structural organization and

dynamic properties of ionomer, 352–368

structure and dynamics in, 355–359

physical theory and molecular modeling of materials, 347–349

role of water in, 349–352, 369–381seven-layer structure and basic

processes in, 345water sorption in, 369–381

phenomenology of, 370–371thermodynamic model of,

371–381Polymer electrolyte membrane. See

Polymer electrolyte fuel cell (PEFC), membrane

Polymer electrolyte membrane fuel cell (PEMFC), 192

cost component distribution, 195POLYMET®, 224Polyol method, 11Polyphenylene oxide (PPO), 143, 162Polyphenylene sulfide (PPS), 322, 324

Polypropylene (PP), 321Polysiloxane, sulfonated, 150Polytetrafluoroethylene (PTFE), 192,

233carbon and, 223, 236catalyst, 64fabrication process and procedure,

227–228effects of, 229–231

thermal conductivity and, 231Poly(vinyl imidazole), 169Poly(vinyl phosphanate)-b-polystyrene,

162Poly(vinyl triazole), 169Poly(vinylidene difluoride-co-

hexafluoropropylene), 154Poly(vinylidene fluoride) (PVDF), 125,

131, 322as base substrates, 156dehydrofluorination, 161Nafion and, 161radiation grafted, 157SEBS and, 162sulfonated, 153

Pore size distribution, 256–259Porosity, 255–256Positron annihilation spectroscopy,

355PP. See Polypropylene (PP)PPO. See Polyphenylene oxide (PPO)PPS. See Polyphenylene sulfide (PPS)Pressure drop tests, 260, 282–284Proton conduction, 108–119

acid content and, 111AFM in study of, 118bulk water for, 354connectivity of aqueous domains, 110distance between acid groups, 110enhancing, 404imidazole and, 169measurement, 118morphology and, 114Nernst–Einstein relationship and, 110polymer microstructure and, 114pore-scale models of, 383–385of SPEKK/PEI blends, 164sulfonation and, 123–124triazole and, 169water and, 109, 112–113, 354

444 Index

Proton exchange membrane (PEM), 62, -trifluorostyrene-based, 139

block copolymer, 151–155chemical stability, 131–136composite, 165–166cross-linked,sulfonic-

acid-substituted, polyphosphazene-based, 150

development of new, 354fluorinated block copolymer, 154fluorosulfonic acid-based, 139fuel cell (PEMFC)

air-breathing, 231cost component distribution, 195diffusion layers (See Diffusion

layer (DL))performance comparison of,

224–226direct hydrogen, 195factors affecting power and

density, 120hydrophobic treatment for,

227–232polarization curves with

different carbon loadings, 241

water flooding, 230–231future directions, 170Gierke model, 355graft copolymer, 155–159for high-temperature operation

and alternative proton conductors, 166–170

hydrocarbon block copolymer, 152

ionically cross-linked acid-base blend, 163

ionomer-filled porous substrates and reinforced, 165–166

materials, 137–169mechanical properties, 129–131methanol crossover, 122–126network model of conductivity,

390–393oxygen permeability, 119–122perfluorinated, 140, 141polyarylene-based, 143polymer blends, 161–164

properties and structure-property relationships, 108–136

proton conduction, 108–119 (See also Proton conduction)

QENS studies, 357, 358radiation-grafted, 156random network model, 355simple water channel models, 356sites for radical attack in, 134stability, 356statistical copolymers, 137–150structure of water in, 369–370synthesis, reviews of, 355water sorption in, 369–381

phenomenology of, 370–371thermodynamic model of,

371–381water transport, 127–129

Proton mobility, near polymer-water interface, 385–390

Proton transport, 381–397near polymer-water interface,

385–390Ohm’s law of, 396pore-scale models of, 383–385

P50T, 198P75T, 198PTFE. See Polytetrafluoroethylene

(PTFE)PtMo catalyst(s), 44–45

durability, 46PtRu catalyst(s), 38, 47, 52

durability, 46layer for, 46preparation, 12variants, 43–44, 48–50

Pulsed field gradient-NMR experiments, 357, 358

Pulsed laser deposition, 4, 89PVDF. See Poly(vinylidene fluoride)

(PVDF)P4VP. See Poly(4-vinylpyridine) (P4VP)PWB-3, 224PyroCell, 314

Q

Quasi-elastic neutron scattering (QENS), 355

Index 445

R

Radial flow permeability testing apparatus, 263

Raman spectroscopy, 355Reactive spray depositon technology

(RSDT), 88–89Reference interaction site model

(RISM), 362Reformate-tolerant catalyst(s), 41–47

anode, 41–45improved, 43–45mechanistic studies, 42–43

stability, 45–47RISM. See Reference interaction site

model (RISM)Rotating disk electrode, 7RSDT. See Reactive spray depositon

technology (RSDT)Ru/C catalysts, 28

S

SANS. See Small-angle neutron scattering (SANS)

SAXS. See Small-angle x-ray scattering (SAXS)

Scanning electrochemical microscopy (SECM), 355

Scanning probe microscopy, 355SCMF. See Self-consistent mean field

theoryScreen printing, in thin-film CL

fabrication, 84–85SDAPP. See Sulfonated Diels-Alder

poly(phenylene)SECM. See Scanning electrochemical

microscopy (SECM)Self-consistent mean field theory,

362SENS. See Small-angle neutron

scattering (SENS)Separation plate. See Bipolar plate(s)Sessile drop method, 68, 251Shutdown-start-up tests, 169SIGRAFLEX, 324Silicon diffusion layers, 220–221Single-walled nanotubes (SWNTs), 39

Sintered metals, 213–214Small-angle neutron scattering (SANS),

115, 140, 355Small-angle x-ray scattering (SAXS),

115, 140, 355Sol-gel Pt application, in thin-film CL

fabrication, 91SPAEKs. See Sulfonated poly(arylene

ether)sSPAES. See Sulfonated poly(aryl ether

sulfone)sSPEEKK. See Sulfonated poly(ether

ether ketone ketone) (SPEEKK)

SPES. See Sulfonated poly(ether sulfone)

Spray coating, in thin-film CL fabrication, 85

Spray pyrolysis, 12sPSO2, 145Sputter deposition, in thin-film CL

fabrication, 86–87SS316, 315, 326, 327SS430 core, 331S-SEBS. See Sulfonated polystyrene-b-

poly(ethylene-r-butylene)-b-styrene (S-SEBS)

S-SIBS. See Sulfonated polystyrene-b-(isobutylene)-b-sulfonated polystyrene (S-SIBS)

SS316L, 333Stack cost, 315Stack targets, 5Starbons, 41Sulfonated Diels–Alder

poly(phenylene), 147Sulfonated poly(aryl ether ketones),

144Sulfonated poly(aryl ether sulfone)s,

143–144, 162Sulfonated poly(arylene ether)s, 142,

143Sulfonated poly(ether ether ketone

ketone) (SPEEKK), 112BAM vs., 114ETFE-g-PSSA, 114microphase separation, 117Nafion vs., 113, 114, 115, 117SAXS analyses, 115

446 Index

Sulfonated poly(ether sulfone), 161Sulfonated polysiloxane, 150Sulfonated polystyrene-b-(isobutylene)-

b-sulfonated polystyrene (S-SIBS), 124

casting solvent, 153Sulfonated polystyrene-b-

poly(ethylene-r-butylene)-b-styrene (S-SEBS), 120

Nafion 117 vs., 124ratios of MeOH to water uptakes

for, 125ratios of proton conductivity to

MeOH permeability, 124SuPAES. See Sulfonated poly(aryl ether

sulfone)sSuperior MicroPowder, 12Superprotonic conductor, 355SWNTs. See Single-walled nanotubes

(SWNTs)

T

Tantalum, 326TEM. See Transmission electron

microscopy (TEM)TG-090, 224TGP-H030, 199TGP-H060, 199TGP-H090, 199, 228TGP-H120, 199Thermal conductivity, 231, 274–276Through-plane permeability, 264–266TiO2, 35Ti4O7, 35Titanium, 326Transmission electron microscopy

(TEM), 115Transparent fuel cells, 267–268Transportation fuel cells, 312Triallyl cyanuarate, 156Triazole, as proton conductor, 169Tungsten carbide, 36Tungsten nitride, 36

U

U. S. Department of Energy (DoE), 4Ultrasmall-angle x-ray scattering

(USAXS), 356USAXS. See Ultrasmall-angle x-ray

scattering (USAXS)

V

Vanadium, 330, 334VASP. See Vienna Ab Initio Simulation

Package (VASP)Vectra LCP, 322Vienna Ab Initio Simulation Package

(VASP), 386Visualization techniques, 284–286Vulcan XC-72R, 247

W

Water balance analysis, 271–272Water channel models, 356Water flooding, 230–231Water management layer, 243Water transfer region, 245Water transport, 127–129WAXS. See Wide-angle x-ray scattering

(WAXS)Wide-angle x-ray scattering (WAXS),

355Wilhelmy method, 68, 252–254Wilhelmy plate gravimetric technique, 69Williamson ether synthesis, 153

X

X-ray radiography, 269–270XC72R, 32, 34

Y

YLP-100 single-mode ytterbium fiber laser, 332

Z

Zirconium, 326

crc press - proton exchange membrane fuel cells - materials properties and performance - 2010

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