Proton Exchange Membrane Fuel Cells: Materials Properties and Performance (Green Chemistry and Chemical Engineering)

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Author: David P. Wilkinson | Jiujun Zhang | Rob Hui | Jeffrey Fergus | Xianguo Li

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Proton Exchange Membrane Fuel Cells Materials Properties and Performance

GREEN CHEMISTRY AND CHEMISTRY ENGINEERING Series Editor: Sunggyu Lee Missouri University of Science and Technology, Rolla, USA

Proton Exchange Membrane Fuel Cells: 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

Proton Exchange Membrane Fuel Cells Materials Properties and Performance

Edited by

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

Boca Raton London New York

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

CRC Press Taylor & Francis Group 6000 Broken Sound Parkway NW, Suite 300 Boca Raton, FL 33487-2742 © 2010 by Taylor and Francis Group, LLC CRC 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 paper 10 9 8 7 6 5 4 3 2 1 International Standard Book Number: 978-1-4398-0664-7 (Hardback) This book contains information obtained from authentic and highly regarded sources. Reasonable efforts have been made to publish reliable data and information, but the author and publisher cannot assume responsibility for the validity of all materials or the consequences of their use. The authors and publishers have attempted to trace the copyright holders of all material reproduced in this publication and apologize to copyright holders if permission to publish in this form has not been obtained. If any copyright material has not been acknowledged please write and let us know so we may rectify in any future reprint. Except as permitted under U.S. Copyright Law, no part of this book may be reprinted, reproduced, transmitted, or utilized in any form by any electronic, mechanical, or other means, now known or hereafter invented, including photocopying, microfilming, and recording, or in any information storage or retrieval system, without written permission from the publishers. For permission to photocopy or use material electronically from this work, please access www.copyright. com (http://www.copyright.com/) or contact the Copyright Clearance Center, Inc. (CCC), 222 Rosewood Drive, Danvers, MA 01923, 978-750-8400. CCC is a not-for-profit organization that provides licenses and registration for a variety of users. For organizations that have been granted a photocopy license by the CCC, a separate system of payment has been arranged. Trademark Notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation without intent to infringe. 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 2010 621.31’2429--dc22 Visit the Taylor & Francis Web site at http://www.taylorandfrancis.com and the CRC Press Web site at http://www.crcpress.com

2009036642

Contents Series Preface ........................................................................................................ vii Preface ......................................................................................................................ix Editors ......................................................................................................................xi Contributors ......................................................................................................... xiii 1.

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

2.

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

3.

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

4.

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

5.

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

6.

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

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

v

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 philosophy, 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 hazardous 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 engineering 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, thermodynamics, 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 environmental 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 innovative 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.

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Some successful examples and key technological developments include the use of supercritical carbon dioxide as green solvent in separation technologies, application of supercritical water oxidation for destruction of harmful 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 efficient oxidation, development of hydrogen proton exchange membrane (PEM) fuel cells for a variety of power generation needs, advanced biofuel production, 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 growing 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 century 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 sustainable 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

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 positive global change, increase fuel efficiency, and decrease dependency on traditional 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 security 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 polymer 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 applicability 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 consumer 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 understanding, 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 technical aspects needs to be covered. This book will provide updated, detailed background material on key developments in the PEMFC area. In particular, ix

x

Preface

the book reviews the progress and current aspects of materials and performance 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, diffusion 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 properties, 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

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 presently 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 technology 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 electrochemical 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 electrochemistry 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 electrochemistry, 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 electrochemical 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. xi

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 characterization for more than 20 years. Dr. Hui has worked on various projects, including chemical sensors, solid oxide fuel cells, magnetic materials, gas separation membranes, nanostructured 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, protective ceramic coatings on metallic substrates, ceramic electrode materials for batteries, and ceramic proton conductors. Dr. Hui is also an active member 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, confidential 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 postdoctoral 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 sensors and fuel cells) and the chemical stability of materials (e.g., high-temperature 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.

Contributors Mauricio Blanco Chemical and Biological Engineering Department and Clean Energy Research Center (CERC) University of British Columbia Vancouver, British Columbia, Canada Institute for Fuel Cell Innovation National Research Council of Canada Vancouver, British Columbia, Canada Tim Cheng Automotive Fuel Cell Cooperation (AFCC) Vancouver, British Columbia, Canada Michael H. Eikerling Department of Chemistry Simon Fraser University Burnaby, British Columbia, Canada Institute for Fuel Cell Innovation National Research Council of Canada Vancouver, British Columbia, Canada Steven Holdcroft Department of Chemistry Simon Fraser University Vancouver, British Columbia, Canada

Kourosh Malek Institute for Fuel Cell Innovation National Research Council of Canada Vancouver, British Columbia, Canada Timothy J. Peckham Department of Chemistry Simon Fraser University Vancouver, British Columbia, Canada Chaojie Song Institute for Fuel Cell Innovation National Research Council of Canada Vancouver, British Columbia, Canada David Thompsett Johnson Matthey Technology Center Reading, England David P. Wilkinson Chemical and Biological Engineering Department and Clean Energy Research Center (CERC) University of British Columbia Vancouver, British Columbia, Canada Institute for Fuel Cell Innovation National Research Council of Canada Vancouver, British Columbia, Canada

xiii

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Zhong Xie Institute for Fuel Cell Innovation National Research Council of Canada Vancouver, British Columbia, Canada Yunsong Yang Ballard Power Systems Vancouver, British Columbia, Canada

Contributors

Jiujun Zhang Institute for Fuel Cell Innovation National Research Council of Canada Vancouver, British Columbia, Canada

1 Recent Developments in Electrocatalyst Activity and Stability for Proton Exchange Membrane Fuel Cells David Thompsett CONTENTS 1.1 Introduction ....................................................................................................2 1.2 The State of the Art ........................................................................................2 1.3 Application Targets........................................................................................4 1.4 Electrocatalyst Discovery .............................................................................5 1.4.1 High-Throughput Screening ............................................................ 5 1.4.2 Computation Studies .........................................................................7 1.5 Electrocatalyst Preparation...........................................................................9 1.5.1 Conventional Routes .........................................................................9 1.5.2 Colloidal Routes ............................................................................... 10 1.5.3 Molecular Precursor Routes ........................................................... 11 1.5.4 Vapor Phase Routes ......................................................................... 12 1.5.5 Surface Modification Routes .......................................................... 12 1.6 Electrocatalyst Testing ................................................................................ 13 1.7 Enhanced Activity Cathode Catalysts ...................................................... 14 1.7.1 Pt Alloy Catalysts............................................................................. 14 1.7.1.1 Model Surface Studies ...................................................... 15 1.7.1.2 Pt Alloy Nanoparticles and Particle Size Effects .......... 16 1.7.2 Pt Core-Shell Catalysts .................................................................... 20 1.7.2.1 Use of Alternative Promoters to Pt ................................. 24 1.7.3 Non-Pt Catalysts .............................................................................. 24 1.7.4 Pd-Based Catalysts ........................................................................... 25 1.7.4.1 Fe- and Co-Based Materials............................................. 25 1.7.4.2 MeOH-Tolerant Oxygen Reduction Catalysts ............... 27 1.8 Cathode Catalyst Stability .......................................................................... 29 1.8.1 Pt Electrochemical Area Loss ........................................................ 29 1.8.2 Stabilization of Pt Catalysts toward Potential Cycling............... 31 1.8.3 Effect of High-Cathode Voltages on Catalyst Stability............... 32 1.8.4 Stabilization of Pt Catalysts toward High-Voltage Excursions .........................................................................................34 1

2

Proton Exchange Membrane Fuel Cells

1.8.5 Alternative Supports ....................................................................... 35 1.8.5.1 Oxides ................................................................................. 35 1.8.5.2 Carbides and Nitrides ...................................................... 36 1.8.5.3 Nonconductive Whiskers................................................. 36 1.9 Carbon Support Materials .......................................................................... 37 1.9.1 Conventional Carbon Blacks .......................................................... 37 1.9.1.1 Modification of Carbon Blacks ........................................ 38 1.9.2 Synthetic Carbon Materials ............................................................ 38 1.9.2.1 Carbon Nanotubes and Nanofibers ............................... 38 1.9.2.2 Synthetic Mesoporous Carbons ...................................... 39 1.10 Reformate-Tolerant Anode Catalysts ........................................................ 41 1.10.1 Mechanistic Studies.........................................................................42 1.10.2 Improved Reformate-Tolerant Catalysts .......................................43 1.10.2.1 PtRu Variants .....................................................................43 1.10.2.2 PtMo Catalysts ...................................................................44 1.11 Reformate-Tolerant Catalyst Stability ....................................................... 45 1.11.1 Ru and Mo Stability ......................................................................... 45 1.11.2 Cell Reversal Tolerance ................................................................... 46 1.12 MeOH Oxidation Catalysts ........................................................................ 47 1.12.1 Mechanistic Advancements ........................................................... 47 1.12.2 PtRu Variants .................................................................................... 48 1.12.3 Alternative Pt Catalysts................................................................... 50 1.13 MeOH Oxidation Catalyst Stability .......................................................... 52 References............................................................................................................... 53

1.1 Introduction The aim of this chapter is to review recent developments in electrocatalyst technology 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 nanoparticles (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

Recent Developments in Electrocatalyst Activity and Stability

3

A

2 nm

111 0.10

Step A Step Pt(s)4 5(111) × (100) Island 3

11–1

B Step Terrace

2 B

111/111 Edge 1 A

B

111/100 Edge Pt(s)-4(100) × (111)

4 Step Vacancy A Step 3 B Step 2 Kink [110] 1 C

FIGURE 1.1 Restored 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.)

4


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 loading result in activity losses consistent with kinetic losses due to the oxygen reduction reaction. For reformate applications, cathodes are similar to H 2 PEMFCs; however, anode catalysts are typically high-loading PtRu on high surface blacks used at higher electrode loadings than H 2 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, unsupported 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 assembly (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 development (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 performance 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 component 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

5


TABLE 1.1 Stack Targets Property a

MEA PGM content MEA PGM content MEA catalyst cost Durability with cycling (≤80pC) Cathode ECA loss Cathode support loss Cathode mass activity Cathode specific activity Cathode non-Pt volume activity a

Units

2005 Status

2010 Target

2015 Target

Grams/kilowatt Milligrams PGM/cm2 geometric area Dollars/kilowatt Hours

1.1 0.8

0.3 0.3

0.2 0.2

9 >2,000

5 5,000

3 5,000

90 >>30

150 mV), even though this was a steadystate 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 alternative 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 catalysts 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


23

Pt20Cu60Co20 Pt20Cu20Co60

0.94

Pt20Cu40Co40 Pt25Co75

0.92

Pt 45 wt% Pt 30 wt%

E/V

0.90

0.88

0.86

0.84 0.1

1 –1 Im/A mgPt

(a) 0.94

Pt20Cu60Co20 Pt20Cu20Co60 Pt20Cu40Co40 Pt25Co75

0.92

Pt 30 wt% 0.90 E/V

Pt 45 wt%

0.88

0.86

0.84 100

1000 Is/μA cm–2 Pt (b)

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

(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 systems can be developed.

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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 considered. Although use of promoters in controlling activity and selectivity is common 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 identify 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 constraints.71 Assuming that the catalyst layer thickness could be increased to ~100 μm from the currently used 10 μm, GM has calculated that the minimum 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 generous, given the recent trend to reduce catalyst layer thicknesses to optimize high-current performances. The DoE has developed a series of volume activity 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 oxygen 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 catalysts. The most promising systems will now be reviewed. However, very little work has been reported on the stability of these systems and virtually


25

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 differences 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 substrates 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 significantly higher activity than Pt and Pt/Pd core-shell materials.77 If these materials 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 reduction activity in acidic media. In general, these materials were based on molecular macrocyclic compounds (e.g., porphyrins, phthalocyanines, tetraazaannulenes) 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 mixing 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

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0.15

1.0

0.12

Cell Voltage (V)

0.8 Co-PPY-C

0.6

0.09

Co/C 0.4

0.06

0.2

0.03

Power Density (W cm–2)

H2-O2

0.00

0.0 0.0

0.2 0.4 Current Density (Acm–2)

0.6

FIGURE 1.13 Oxygen reduction polarization and power density curves for Co/C and Co-polypyrrole/C cathodes in MEAs at 80pC, 210 kPa H2/O2, 100% RH, 0.06 mg (Co) cm–2. (Reprinted by permission from Macmillan Publishers Ltd: Nature, 443, 63 (2006). Copyright 2006.)

Gasteiger et al. reviewed the best performing Fe-based catalysts in the literature up to 200471. Even the best of these catalysts (Fe on pyrolyzed perylenetetracarboxylic 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 pretreating 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 macrocycle 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 nondisclosed 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 containing graphite and well-defined Fe3C and b-Co species.82 It was concluded that


27

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 durability (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 optimization 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 magnitude 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; however, given Pt’s ability to reduce oxygen and oxidize methanol, this lack of selectivity makes them sensitive to methanol crossover from anode to cathode 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 unaffected by the presence of MeOH. This has led to a large body of work investigating 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 reduction 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 selenides. 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

28


1.0 Ru1.92Mo0.08SeO4 Ru/Vulcan Pt/Vulcan

E (V/RHE)

0.9

0.8

0.7

60°C, 0.5 M H2SO4 0.6 0.01

0.01 im (A/mgnoble metal)

1

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

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 significant 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 (Pt xSy/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 closest 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


29

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 nearterm 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 (cathode) catalyst to these. In particular, for automotive systems as well as steadystate 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 0.8 V). Evidence of gross soluble Pt migration is found by appearance of Pt particles within the membrane, formed by reduction of soluble Pt species by H2 gas crossover from the anode. r Pt particle coalescence is due to migration. This mechanism is supported by observations that, upon cycling, Pt particle size distributions 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. r 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 agglomeration and electrical isolation, leading to loss in activity. It has been well established that Pt can dissolve under oxidizing conditions, 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 soluble 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.


31

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 elsewhere 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 particles. 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 similar 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.

32


1200

90.00 Pt - eca

80.00

Pt - specific activity PtCo - specific activity

Electrochemical Area/m2g–1 Pt

70.00

1000

800

60.00 50.00

600 40.00 400

30.00 20.00

200

Specific Activity/??A cm–2 Pt @ 900 mV

PtCo - eca

10.00 0.00

1

10

100 1000 Cycle Number

10000

0 100000

FIGURE 1.15 Change in electrochemical area and oxygen reduction specific activity for Pt and PtCo cathodes 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, 80pC. (S. C. Ball et al., Electrochemical Society Transactions, 11, 1267 (2007). Reproduced by permission of The Electrochemical Society.)

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 sufficiently 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 performance 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


33

= ≈50% Pt/Cstandard = ≈30% Pt-alloy/Ccorr.-resist.

Carbon-loss (wt%)

25

20

15 1.2V 10

5

0.9V

0 1

10

100

1000

Time (h)

FIGURE 1.16 Carbon weight loss as a function of time for two different catalysts in MEAs at both 0.9 and 1.2 V, 80pC, 100% RH, 120 kPaabs. (M. F. Mathias et al., Electrochemical Society Interface, 14, 24 (2005). 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 n 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 n Cs e–; 2. hydrolysis, Cs ½H2O n CsO H ; and 3. gasification to carbon dioxide, 2CsO H2O n 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 corroded 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

34


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 investigated. 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 steadystate 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 excursions 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 standard carbon blacks (e.g., XC72R, Ketjen, BP2000) that are subjected to hightemperature graphitization conditions (1,800–2,800°C). This has the effect of reducing surface area by removal of micropores ($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 investigated 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

40


A

Frequency (%)

40 Pt/C Mean = 2.78 nm

30 20 10 0

1

20 nm

4 2 3 Particle Size (nm)

5

B

Frequency (%)

40

PUSWNT Mean = 2.72 nm

30 20 10 0 1

20 nm


C

Frequency (%)

40 Pt/MWNT Mean = 2.78 nm

30 20 10 0

1

20 nm


5

D

Frequency (%)

40 Pt/DWNT Mean = 2.72 nm

30 20 10 0 1

20 nm


FIGURE 1.18 TEM images and particle size distributions of (A) Pt/C, (B) Pt/SWNT, (C) Pt/MWNT, and (D) Pt/DWNT. All Pt loadings 30 wt% (Z. Chen et al., Electrochemical Society Transactions, 11, 1289 (2007). Reproduced by permission of The Electrochemical Society.)


41

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 current 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 socalled “Starbons” contain negligible microporosity, with average pore diameters 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 transporting 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 program 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 performance, 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.

42


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 underneath 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 development 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 decreasing CO coverage and increasing sites available for H 2 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 (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 solution has been the incorporation of water oxidation catalysts into the anode catalyst layer. Ralph, Hudson, and Wilkinson showed that the most effective


47

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 oxidation 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 pathways. In recent years, the deposition of PtRu particles onto novel carbon supports 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 spectroscopy 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 dissociation 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 promotes the reaction because it is able to produce OH species at lower potentials. This promotional effect is known as the “bifunctional” mechanism: CH3OH n CH3OHads CH3OHads n COads 4H 4e–

48


H2O n H2Oads H2Oads n OHads H e– COads OHads n 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 formation 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 reaction 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 overpotentials, the MeOH oxidation reaction was found to be zero order in MeOH concentration, 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 commercially 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 disagrees with the optimal ratios determined from bulk PtRu alloys.


49

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 60pC 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 characterized their surface composition by XPS.162 Similar values to the bulk compositions 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 complex.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 coverage 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 suggest that PtRu particles show Pt-segregated surfaces as predicted by theoretical 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 concentrations 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 thinfilm 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 activity enhancements for PtRuNi,170 PtRuFe,171 and PtRuIr.172 In addition, several reports have indicated that the addition of phosphorus to PtRu has

50


0.8 Pt Pt (50 at %) Ru (50 at %) Pt (65 at %) Ru (25 at %) Os (10 at %)

0.7

Voltage/V

0.6

(c)

0.5 (b) 0.4

(a)

0.3 0

100

200 300 400 Current Density/mA cm–2

500

600

FIGURE 1.21 Liquid-feed DMFC polarization curves with (a) Pt, (b) PtRu, and (c) PtRuOs anodes at 90pC, 0.5 MeOH, and O2 feeds. (K. L. Ley et al., Journal of the Electrochemical Society, 144, 1543 (1997). Reproduced by permission of The Electrochemical Society.)

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 developed 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 compared 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

51


Mass Current Density (mA/mg Pt)

800 PtPb/C PtRu/C

700 600 500 400 300 200 100 0

–100 0.0

0.2

0.4 0.6 E(V) vs Ag/AgCl

0.8

1.0

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

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 oxidation 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 profile 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 oxidation, 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 activity 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.

52


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 formulations can be considered as alternatives to the well-established PtRu system.

1.13 MeOH Oxidation Catalyst Stability

0.8

DMFC

40 mV Cathode Loss

0.7 0.6 0.5 0.4 Ru-free cathode Ru-contaminated cathode Pt-Ru cathode

0.3 0.2 0.00

0.05

0.10

0.15

Current Density (A cm–2)

0.20

iR-Corrected Cell Voltage (V)

iR-Corrected Cell Voltage (V)

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 activity. 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 operating 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 reduction of Pt oxides and restoring cathode activity.187 It is clear that to guard against DMFC performance degradation, improvements in PtRu stability are required. 0.8

DMFC (Extreme Ru Contamination) >200 mV Cathode Loss

0.7 0.6 0.5 0.4 0.3

Ru-free cathode Extremely Ru-contaminated cathode

0.2 0.00

0.02

0.04

0.06

0.08

0.10

Current Density (A cm–2)

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


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2 Catalyst Layers and Fabrication Zhong Xie, Chaojie Song, David P. Wilkinson, and Jiujun Zhang CONTENTS 2.1 Introduction .................................................................................................. 62 2.2 Catalyst Layer Components and Their Corresponding Functions ......64 2.2.1 Overview of Catalyst Layer Components and Functions ..........64 2.2.2 Properties of the Catalyst Layer..................................................... 66 2.2.2.1 Catalyst Loading and Catalyst Utilization .................... 66 2.2.2.2 Nafion Loading.................................................................. 68 2.2.2.3 Hydrophobicity and Hydrophilicity .............................. 68 2.2.2.4 Porosity ............................................................................... 69 2.2.2.5 Ionic (Proton) Conductivity and Electronic Conductivity ...................................................................... 70 2.3 Types of Catalyst Layers ............................................................................. 70 2.3.1 CCGDL .............................................................................................. 70 2.3.1.1 Uniform CCGDL ............................................................... 70 2.3.1.2 Gradient CCGDL ............................................................... 71 2.3.1.3 Dual-Bound Composite Catalyst Layer ......................... 75 2.3.2 CCM ................................................................................................... 76 2.3.2.1 Conventional CCM ........................................................... 76 2.3.2.2 Nanostructured Thin-Film Electrode ............................77 2.3.3 Novel Structural Catalyst Layer ....................................................77 2.3.3.1 CNT-Based Catalyst Layer ...............................................77 2.3.3.2 Columnar Oxide Supported Catalyst Layer..................77 2.3.3.3 Nanowire-Based Three-Dimensional Hierarchical Core/Shell Catalyst Layer ................................................ 79 2.3.3.4 Self-Supported Catalyst Layer.........................................80 2.3.3.5 Catalyst Layer with Additives .........................................80 2.3.3.6 Catalyst Layer with Novel Ionomers .............................. 81 2.4 Catalyst Layer Fabrication .......................................................................... 81 2.4.1 First-Generation Catalyst Layer Fabrication ................................ 81 2.4.1.1 Pt Black Catalyst Layer Fabrication ................................ 81 2.4.1.2 PTFE-Bound Catalyst Layer Fabrication Using Supported Catalyst ........................................................... 82

61

62


2.4.2

Thin-Film Catalyst Layer Fabrication ...........................................83 2.4.2.1 Ink-Based Catalyst Layer Fabrication.............................83 2.4.2.2 In Situ Catalyst Layer Fabrication ................................... 86 2.4.2.3 Other Methods .................................................................. 89 2.5 Catalyst Layer Optimization ...................................................................... 91 2.5.1 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 ............ 93 2.5.2 Catalyst Layer Microstructure Optimization .............................. 95 2.6 Prospects and Conclusion .......................................................................... 96 References............................................................................................................... 97

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 n 2H+ + 2e–

(2.1)

Cathode: O2 + 4H+ + 4e– n 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

63

Catalyst Layers and Fabrication

Electrically Conductive Fibers

Carbon Supported Catalyst

Proton Conducting Media

Carbon Supported Catalyst

Electrically Conductive Fibers e– H2O O2

H+

H2 H 2O GOL e–

Anode Catalyst Layer

PEM

Cathode Catalyst Layer

GOL

FIGURE 2.1 Schematic 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.)

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

64


(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: polytetrafluoroethylene (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-supported 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 conduction. 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/cm 2. Later it was found that the addition of Nafion ionomer (by brush coating or spraying) to the PTFEbound 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 performance. 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 reactions occur in the same layer. For this type of structure without a protonconducting 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 addition, 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.

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1 0.9

Cell Potential (V)

0.8 A

0.7 0.6

B C

0.5 0.4 0.3 0.2 0.1 0 0

20

40

60

80 100 120 140 Current Density (mA/cm2)

160

180

200

FIGURE 2.2 Polarization curves for H2/O2 fuel cells at 50pC, 1 atm pressure. Curve A: Nafion impregnated (brush coated) PTFE-bound electrode (0.35 mg/cm 2 Pt loading); curve B: PTFE-bound catalyst layer (Pt loading: 4 mg/cm 2); 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.)

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/cm 2 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 electrode structure, an electrode with a thin CL generally contains three layers:

66


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 commercialization. Therefore, one of the major goals in PEM fuel cell development 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: Ptutilization

ECAmeasured r 100% ECAcalculated

(2.3)

where ECAmeasured 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) measured from the cyclic voltammogram and the well-established standard value of 0.21 mC/cm2 for planar polycrystalline Pt [12,13] (here, cm 2 is the actual Pt surface), based on the following equation: ECAmeasured

Qdesp 0.21

r 1000

(2.4)

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30

20

Current Density (mA/cm2)

10

Qdesp

0 0 –10

0.2

0.4

0.6

0.8

1

Qadsp

–20

–30

–40

–50

Potential (V) vs NHE

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

where Qdesp is the charge under the hydrogen desorption peak (μC/cm 2) (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: ECAcalculated

6 r 10 4 r LPt dR

where LPt is the Pt loading in the catalyst layer (mg/cm 2) d is the Pt particle size (average particle diameter) in nanometers r is the density of Pt (g/cm3) the unit of ECAcalculated is also cm2/cm2.

(2.5)

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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 influencing 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]: LNafion 56 r

LPt PPt

(2.6)

where LPt is the Pt loading (mg/cm2) and PPt is the weight percent of Pt supported 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 disadvantages 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 microscope, 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.

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The capillary rise method, originally developed by Neumann and coworkers [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]: sinQ 1

$R gh2 2S

(2.7)

where q is the contact angle Δr is the difference between the densities of liquid and vapor g is the gravitational constant h is the meniscus height s is the liquid-gas surface tension of water. Therefore, the contact angle can be calculated from the experimentally measured meniscus height. Contact angle measurements on the CL may also be useful in the characterization 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 function 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% depending on the pore formers used. Unfortunately, cell performance is not proportional to catalyst layer porosity. 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%.

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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 usually determined by the Nafion ionomer loading. Proton conductivity in the CL can be measured using AC impedance spectroscopy. However, this measurement usually gives an average value over the whole CL, whereas, in reality, 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: PTFEbound 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 diffusion layer or microporous layer, and catalyst-coated membrane, in which the CL is directly coated on the proton exchange membrane. In the following 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.

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2.3.1.2 Gradient CCGDL A gradient CCGDL has a nonuniform distribution of catalyst (catalyst gradient 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 performance 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 corresponding 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, catalyst 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 electrolyte 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 2.4 Schematic diagram of the catalyst loading gradient: through plane (left) and in plane (right).

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Non-gradient 0.3 mg-Pt/cm2

0.4 mg-Pt/cm2

Inlet 5 cm

0.3 0.35

0.3

0.5

Outlet 5 cm 0.34 mg-Pt/cm2

0.3 mg-Pt/cm2

0.25 mg-Pt/cm2

0.3

0.2

0.15

0.35

0.3

0.25

0.35

0.35

0.3

(a) 1.0 Average Cathode Catalyst Loading 0.3 mg-Pt/cm2 (non-gradient) 0.4 mg-Pt/cm2 0.34 mg-Pt/cm2 0.3 mg-Pt/cm2 0.25 mg-Pt/cm2

0.9

Potential, V

0.8 0.7 0.6 0.5 0.4 0

200

400 600 800 Current Density, mA/cm2 (b)

1000

1200

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


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in a variety of fuel cell hardware, demonstrating that these performed better 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 catalyst loadings of 0.25 mg/cm2 and 0.3 mg/cm2, the gradient CCGDLs showed better performance than their nongradient counterparts. The gradient structure 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 gradient 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 content 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, theoretical studies and some experimental results have shown that gradient CLs with higher Nafion loading at the membrane/catalyst layer give higher performance 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) exhibited 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 distribution of Nafion ionomer inside the CL playing a role in extending the electrochemical 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 composite 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 loading and PTFE content were 0.29 and 0 mg/cm2, respectively; for the remaining third of the area close to the outlet, the Nafion loading and PTFE content were 0.6 and 0.37 mg/cm 2, 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

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10 0.9

Cell Voltage (V)

0.8 0.7 0.6 0.5 0.4

Surface 1.0 mg/cm2 - inside none Surface 1.0 mg/cm2 - inside 0.25 mg/cm2 Surface 1.0 mg/cm2 - inside 0.50 mg/cm2 Surface 1.0 mg/cm2 - inside 0.75 mg/cm2 Surface 1.0 mg/cm2 - inside 1.00 mg/cm2

0.3 0.2 0.1 0.0

0

200

400

600 800 1000 1200 Current Density (mA/cm2)

1400

1600

FIGURE 2.6 Cell voltage versus current density of the electrode with a Nafion content of 1.0 mg/cm 2 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.)

Sub-layer II

Sub-layer I

Dual Catalyst Layer Anode Membrane

Sub-layer I

Sub-layer II

Dual Catalyst Layer Anode

FIGURE 2.7 Schematic 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.)

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TABLE 2.1 Gradient Pt and Nafion Loadings in MEAs

Total Pt loading (mg) in catalyst layer (5 × 5 cm2) Nafion loading in sublayer (wt%)

Pristine MEA 10

33

MEA 1 Sublayer I

II

MEA 2 Sublayer I

MEA 3 Sublayer

MEA 4 Sublayer

II

I

II

I

II

7

3

7

3

7

3

7

3

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 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 performance than the nongradient CL. 2.3.1.3 Dual-Bound Composite Catalyst Layer Thin-film catalyst layers are usually hydrophilic, with no hydrophobic ingredients 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 1.1 Pristine MEA MEA 1 MEA 2 MEA 3 MEA 4

1.0 0.9 Voltage (V)

0.8 0.7 0.6 0.5 0.4 0.3 0.2 0

200

400

600 800 1000 1200 1400 1600 1800 Current Density (mA/cm2)

FIGURE 2.8 Polarization 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.)

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Cell Voltage (V)

1.0

PTFE-bound CL Ionomer-bound CL Dual-bound CL

0.8

0.6

0.4 0.0

0.2

0.4

0.6 0.8 1.0 1.2 Current Density (A cm–2)

1.4

1.6

FIGURE 2.9 Polarization 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.)

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 dualbound 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 catalystcoated 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 mixture bonded on a membrane. This was similar to the PTFE-bound catalyst


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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 conventional 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 catalyst ink was first applied to a Teflon blank and then transferred to the membrane 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 crystalline organic whiskers on which Pt had been deposited. The film was then transferred to the membrane surface using a decal method, and a nanostructured 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 attractive features of CNT-based CLs include improved thermal and charge transfer, 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 conducting ceramic columnar supports have also been proposed. Bonakdarpour et al. [48] characterized columnar titanium structures on smooth glassy

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FIGURE 2.10 Scanning 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.)

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

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Microporous Layer e– Carbon Backing Layer

Anode

PTFE+C PEM

Cathode

e–

H+ O2

H2

2e– 2H+

H2

Carbon Black Support

2e–

1 O 2 2

Nafion Pt Nanoparticle

Catalyst Layer

MWNT Support

H2O

FIGURE 2.11 Schematic 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.)

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 electrode, 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 heterostructure, 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 electrocatalysts for CLs in fuel cell applications.

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(a)

(b)

FIGURE 2.12 Scanning 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.)

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 temperature. 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 relative humidity and/or elevated temperature. Alternatively, proton conducting materials such as ZrP [55] or heteropoly acid HPA [56] have also been added


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into the CL to improve proton conductivity under low RH and/or elevated temperature. These types of CLs have the potential for applications in hightemperature 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 developed, 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 ionomers [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 performance 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 technology, 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/cm 2), large platinum agglomerates (~1 μm on average), lower electrochemical surface area (~25 m 2/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 following 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%.

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3. Add this diluted PTFE to the Pt/H2O slurry and mix them thoroughly. 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/cm 2 down to 0.4 mg/Pt cm 2 [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, followed 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/cm 2 [74]. However, the platinum utilization in this PTFE-bound catalyst layer still remains low: in the vicinity of 20% [75,76].


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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 drawbacks. 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 additive 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 blending 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 thermoplastic 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 improving 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 sacrificing electrode performance. In thin-film CL fabrication, the most common method is to prepare catalyst ink by mixing the Pt/C agglomerates with a solubilized 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 thinfilm 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).

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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 produces 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 160pC. 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], inkjet printing [91,92], and rolling [93]. Currently, screen printing and spray coating have become standard methods for conventional catalyst layer fabrication. Inkjet printing is also showing promise for fabricating low Pt loading 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 swelling 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 adhesion 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

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Squeegee Masked Silk Screen Substrate Film

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

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 diffusion layer or cast directly on the membrane. To prevent distortion and swelling 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 technology, 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 membrane to form a catalyst layer. The ink cartridges were filled with the catalyst inks. The membrane was secured to a cellulose acetate sheet and fed through the printer using the original paper feed platen. Computer software was used to control the print parameters, such as electrode dimensions, 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 catalyst layer fabrication, such as varied composition layer printing, and suggest

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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 question still remains of whether or not inkjet fabrication can offer any performance advantage and still lead to more efficient utilization of the Pt catalyst. Similarly, Taylor et al. [92] used inkjet printing to deposit catalyst materials onto GDLs. Their inkjet-printed catalyst layers with a catalyst loading of 0.020 mg Pt/cm2 showed high Pt utilizations. This research also demonstrated 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 fabrication 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 technology 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 altering 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 microstructure 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 Ng/cm2); 4. A simple preparation process that is easy to scale up; and 5. Adaptiveness to various substrates, such as GDL and membrane.


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FIGURE 2.14 Schematic 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.)

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 sputtered catalyst layers, such as low lateral electrical conductivity of the thin metallic films [96,108]. Furthermore, the smaller particle size of sputter-deposited 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 bombardment. 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/cm 2) 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

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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 catalyst layer because they do not meet the three-phase reaction requirements. In order to mitigate this drawback and increase Pt utilization, a pulse electrodeposition 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 protonation 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 threephase 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 membranes, to form CCMs [111]. The process involves dissolving a Pt precursor in appropriate solvents, followed by spraying the solution with an expansion 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 condensation 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 carbon 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 technology has the capacity and flexibility required to produce compositionally and


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structurally optimized CLs. The technology can be used to generate supported 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 platinum 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 turbomolecular 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 technique 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, achieving 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

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Roller

Membrane

Coating Nozzle

Powder Supporter Catalyst Additive

Nitrogen FIGURE 2.15 Schematic 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.)

in a knife mill. In order to obtain thin, homogeneous reactive catalyst layers, the materials are atomized and sprayed in a nitrogen stream through a slit nozzle directly onto the proton conducting membrane. Although adhesion 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 voltage (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 ultrasonic 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 improvements, which could contribute to better catalyst utilization than is achieved


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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 computercontrolled 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 successfully 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 solvent was then allowed to evaporate at room temperature. Pt loading was controlled 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 management. The composition and structure of a CL can affect all these functions

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to different extents. Catalyst layer optimization aims to satisfy these requirements, 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 interfaces) 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 ionomer 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 performance. There is a trade-off between ionomer and catalyst loadings for optimum 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 modeling, 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: r 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 dramatically. 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. r Song et al. [122] modeled optimal performance as a function of Nafion content as well as Pt loading. r Kamarajugadda and Mazumder [123] used numerical modeling and investigated the effects of ionomer (Nafion) loading, catalyst (Pt)


r

r

r

r

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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 fractions 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 performance 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 ionomer 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 particular 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 nanofabrication, 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 utilization to retain moisture, as well as to prevent membrane dehydration, especially at low current densities. Optimal Nafion content in the electrode is necessary to achieve high performance.

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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 content to be ~33 wt%. Pagainin, Ticianelli, and Gonzalez [83] confirmed that when the Nafion loading was increased from 0.87 to 1.75 mg/cm 2, performance could be significantly improved. If the Nafion loading was increased beyond 2.2 mg/cm 2 (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, better performance could be obtained in the high current density region. Xie et al. [32] prepared a three-sublayer CL containing a gradient distribution 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 proton transport at the catalyst layer/membrane interface, whereas at locations near the GDL interface, less ionomer is better to obtain more porous channels for gas flux. Apart from the optimization of a single component (Nafion), optimization of two components, Pt/C and Nafion, has been explored. Sasikumar, Ihm, and Ryu [131] examined the correlation between Pt loading and optimum Nafion content in the CL, finding that optimum Nafion content was dependent on Pt loading. If the Nafion content was increased, the Pt loading 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 performance was achieved with Pt loadings of 0.20 q 0.05 and 0.35 q 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;


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this demonstrated the possibility of reducing Pt loading using catalyst gradients 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 conductivity. However, if the ionomer loading is too high, transport of reactant gas to the reactive sites will be impeded, leading to significant mass transport loss. The optimal ionomer loading for a CL depends on what materials and fabrication methods are used. It is generally agreed that a Nafion ionomer 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 composition 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 formula. 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 carbonate, 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] proposed a colloidal ink fabrication procedure using low-dielectric-constant solvents to generate a good network and a uniformity of perfluorinated

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sulfone ionomers (PFSIs) on Pt particles in the CL. Wang et al. [146] optimized 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 electrochemically 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 technology. The strategies and methodologies to increase catalyst layer Pt utilization and reduce Pt loading constitute a major effort in fuel cell research and development. 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 better performance, we need to determine how various factors affect Pt utilization. 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 experimental 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 computation technology, computational modeling can provide a powerful complementary 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 performance, 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,


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

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3 Proton Exchange Membranes Timothy J. Peckham, Yunsong Yang, and Steven Holdcroft CONTENTS 3.1 Introduction ................................................................................................ 108 3.2 PEM Properties and Structure–Property Relationships ...................... 108 3.2.1 Proton Conduction......................................................................... 108 3.2.2 Oxygen Permeability and Methanol Crossover ........................ 119 3.2.2.1 Oxygen Permeability ...................................................... 119 3.2.2.2 Methanol Crossover........................................................ 122 3.2.3 Water Transport ............................................................................. 127 3.2.4 Mechanical Properties and Chemical Stability of PEMs ......... 129 3.2.4.1 Mechanical Properties .................................................... 129 3.2.4.2 Chemical Stability ........................................................... 131 3.3 Brief Overview of Existing PEM Materials ............................................ 137 3.3.1 Statistical Copolymer PEMs ......................................................... 137 3.3.1.1 Perfluorinated and Partially Fluorinated .................... 137 3.3.1.2 Polyarylenes ..................................................................... 142 3.3.1.3 Miscellaneous Monolithic, Statistically Sulfonated Copolymer PEMs ........................................ 149 3.3.2 Block and Graft Copolymer PEMs .............................................. 150 3.3.2.1 Block Copolymers ........................................................... 151 3.3.2.2 Graft Copolymers............................................................ 155 3.3.3 Polymer Blends and Composite PEMs........................................ 159 3.3.3.1 Polymer Blends ................................................................ 161 3.3.3.2 Ionomer-Filled Porous Substrates and Reinforced PEMs ............................................................. 165 3.3.3.3 Composite PEMs for High-Temperature Operation and Alternative Proton Conductors .......... 166 3.4 Future Directions ....................................................................................... 170 References............................................................................................................. 171

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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 academic 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 academic 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: r possess high proton conductivity (Section 3.2.1); r be impermeable to gases (specifically oxygen) and/or fuel (e.g., methanol) (Section 3.2.2); r achieve balanced water transport (Section 3.2.3); r possess high thermomechanical and chemical stability to fuel cell conditions (Section 3.2.4); and r be 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 membrane, 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.

109

Proton Exchange Membranes

+

Breaking of Hydrogen Bond

+

+ Formation of Hydrogen Bond

+

+

+

H5O2

H9O4

H5O2

(Zundel-ion)

(Eigen-ion)

(Zundel-ion)

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

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 follows 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 sulfonate 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 counter-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 demonstrated by Kreuer5 and will be discussed later in this section.

110


(a)

(b)

FIGURE 3.2 Connectivity 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.)

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 schematically 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, s H , can be related through the Nernst–Einstein relationship to the activity of protons (aH ) in the membrane as well as to the mobility (mH ) of those protons:

S H FaH MH

–

– –

– –

– –

–

– – –

(a)

(3.1)

–

–

–

– –

– –

– –

–

(b) – = –SO3H

FIGURE 3.3 Proximity 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.)

111


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 polyarylenebased 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 content must be higher (l s 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 ionexchange capacity (IEC) of the membrane. As can be seen in the figure, proton conductivity is strongly dependent upon acid content for all the PEMs. This is not surprising, given that conductivity is dependent upon proton 0.25

Conductivity (S/cm)

0.2

Nafion SPEEK BAM membrane ETFE-g-PSSA

0.15

0.1

0.05

0 0.00

0.50

1.00

1.50 2.00 IEC (meq/g)

2.50

3.00

3.50

FIGURE 3.4 Proton conductivity as a function of IEC for ETFE-g-PSSA polyethylenetetrafluoroethylenegraft-polystyrene sulfonic acid, BAM membrane substituted poly(trifluorostyrene) sulfonic 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.)

112


.'& %"%(&" #

'&!* )$+$), %

- /

FIGURE 3.5 Proton 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.)

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 conductivity. 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 increasing 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 intermolecular proton transfer. However, at low water contents, the degree of dissociation of the proton from the sulfonate groups decreases, thus leading

113


1e–3

1e–3 Dσ 1e–4

NAFION

DH2O Dσ

D/cm2s–1

DH2O 1e–5

Dσ DH2O

Pure Water 1e–4

PEEKK (70% sulfonated)

1e–5

T = 300 K

1e–6

1e–6

1e–7

1e–7

1e–8

1e–8 0.1 Water Volume fraction XV

1

FIGURE 3.6 Proton mobility (Ds) and water self-diffusion coefficient (D H2O) as a function of the water volume 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.)

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 concentration. 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 increasing water content. However, in the case of BAM¢ membrane, the increase in water content is much higher for similar acid contents versus either ETFEg-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 content 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,

114


'#" )++*$"

$#%#"#""%$%!#"

& (

FIGURE 3.7 Proton 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.)

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, proton 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 separation 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,


115

resistance against dissolution. Studies by small-angle x-ray scattering (SAXS) and small-angle neutron scattering (SANS)—with many of the primary studies 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 microstructure of the polymer with a range of polymer-to-solvent ratios in addition 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 clusters (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 isolated 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 n 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 separated hydrophilic and hydrophobic microdomains, corresponding to a larger average separation between neighboring acid sites. A schematic illustration 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 channels 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

116


Dry Membrane Perfluorinated Matrix

0 Swollen Membrane Ionic Domains

Percolation 0.25

*Structure Inversion*

0.50

Connected Network of Polymers Rods

0.75 Colloidal Dispersion of Rod Like Particles Solution

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

117


Nafion

Sulfonated Polyetherketone (PEEKK) O

— CF2—CF2—CF—CF2— n

O

OCF2CF—O(CF 2)2SO3H m O

CF3

O

SO3H

1 nm

– – : -SO3

+ : Protonic charge carrier : H 2O

Wide channels More separated Less branched Good connectivity Small –SO–3 /-SO–3 separation pKa ~ –6

Narrow channels Less separated Highly branched Dead-end channels – – Large –SO3 /-SO3 separation pKa ~ –1

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

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 scattered 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

118


(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 morphology 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 completed, 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 hydrophilic 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 conductivity have also been carried out on PEMs other than Nafion and sulfonated poly(ether ketone). These include studies in which phenomenological examinations 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

A

Cathode: O2 + 4H+ + 4e–

2H2O

Pt-AFM Tip U

Membrane

Electrode

Anode 2H2O

O2 + 4H+ + 4e–

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

119


"4*)

5

3

*&2#(5

5 !

5

,

5

(

5

"!+

5

+*-*)#*)$.#-'/'-1 0'$%'*)#*)$.#-'/'-1

5 5

)

5

FIGURE 3.11 Conductivity 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.)

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 proton 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 membrane 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

120


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 t 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 membranes have been studied. It has been found that the diffusion coefficients of all the studied membranes decrease with increasing EW, while the solubility correspondingly increases. These trends are the same as found in


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previous research for perfluorinated membranes.26,27 The oxygen diffusion coefficient is affected to a greater extent by increasing EW than is solubility, 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 content. However, it must be clarified that the water contents were measured at 298 K and represent “wet” (i.e., fully saturated) membranes. Thus, the correlation 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 solubility 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 oxygen 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 corresponding 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 inspection 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.

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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 permeability was also found for SPEEK (at 30pC, 8.7 t 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 oxygen 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 treatment is above the Tg of Nafion film. 3.2.2.2 Methanol Crossover Sluggish methanol (MeOH) oxidation reaction kinetics and MeOH crossover 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 compartment 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 liquid 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 membrane, 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


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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 sufficiently 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 membranes 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 diffusion 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 gradient (PFG) NMR techniques. The MeOH self-diffusion coefficient was 6 t 10–6 cm2 S –1 and roughly independent of concentration over the range of 0.5–8.0 M at 30pC. A similar diffusion coefficient was obtained for Nafion 117 at 22pC by Hietala, Maunu, and Sundholm with the same technique.46 Kauranen and Skou determined the MeOH diffusion coefficient of 4.9 t 10–6 cm2 S –1 for Nafion 117 in the presence of 2.0 M H2SO4 solutions at 60pC by using an electrochemical 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-butanol (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 permeability ranges from 1.0 t 10 –6 cm2 S –1 to 1.49 t 10 –6 cm2 S –1 were determined for Nafion 117 membranes at 22pC depending on different pretreatment protocol. 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 permeability for a series of S-SEBS membranes.49 Both proton conductivity and MeOH

124


Δ × 106 (Ω–1cm–3s)

6

4

2 Nafion 117 0 0

10 20 30 40 Sulfonation Degree (mol%)

50

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

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 permeability for Nafion 117 was measured for comparison reasons and found to be 2.6 t 10 –6 cm2 s–1. When comparing this value with the MeOH permeability (1.2 t 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 indicates 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 membrane 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

125

Methanol to Water Uptake Ratio


9

6

3

0 0

10 20 30 Sulfonation Degree (mol%)

40

FIGURE 3.13 Ratios 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.)

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 10 ns) and length (>10 nm) scales. Mesoscale models are needed to bridge the gap between the chemical structure of the polymer 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 associated with water molecules. In agreement with experiments,73 Khalatur and colleagues found a linear dependency of microscopic swelling on l, attributed 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 constructed 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 distribution 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 atomistic 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

Physical Modeling of Materials for PEFCs

363

to structural reorganization of the phase-separated morphology.89,90 Simulations 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 evolution 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 withl. 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 content 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 morphology of the membrane shows a percolation-type transition from isolated hydrophilic clusters to the three-dimensional network of randomly interconnected water channels. Very recently, Wu et al.118 applied extensive DPD simulations to study the morphologies of Nafion, SSC, and 3M PFSA membranes at various hydration levels and ionomer equivalent weights. These DPD simulations 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 correlation functions (e.g., diffusion), making them less applicable to extract structure-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

364


FIGURE 6.5 Snapshots 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.)

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 energies between the distinct beads are specified. In hydrated ionomer membranes, 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

365


FIGURE 6.6 Coarse-grained model of Nafion with 20-unit oligomer (length ~ 30 nm).

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 nm 3) of a four-monomeric unit of polytetrafluoroethylene PTFE (-[-CF2-CF2-CF2CF2-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 nm 3. 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 simulations are modeled by the Lennard–Jones (LJ) potential: §¤ A ³ 12 ¤ A ³ 6 ¶ ij ij U LJ (r ) 4 C ij ¨¥ ´ ¥ ´ · , ¨¦ r µ ¦ r µ · © ¸

(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 el (r )

£ i j

qi q j r

.

(6.3)

366


Interactions between chemically bonded beads in ionomer chains are modeled by harmonic potentials for the bond length and bond angle: Vbond (r )

1 K (r r0 )2 2 bond

(6.4)

1 Vangle (r ) Kangle [cos(Q ) cos(Q 0 )]2 , 2

where the force constants are Kbond 1,250 kJ mol–1 nm–2 and Kangle 25 kJ mol–1, respectively.117,119 r0 and q 0 are the equilibrium bond length and angle, respectively. 119 Membrane simulations were performed with l 4, 9, and 15.117 The mesoscopic structure of the hydrated membrane is visualized in Figure 6.7, revealing 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 indicating 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 6.7 Snapshots of the final microstructure in hydrated Nafion membrane at different water contents. Hydrophilic domains (water, hydronium, and side chains) are shown in gray, while hydrophobic domains are shown in black.

367


Atomistic-MD CG-MD

g (r)

20

10

0

0

1 r (nm)

2

FIGURE 6.8 Site–site W–W radial distribution function obtained from CGMD simulation and compared with that of the atomistic MD simulation using the force-matching procedure.

So far, CG approaches offer the most viable route to the molecular modeling 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 models 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 sliding 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 conductivity studies.24 This value is similar to the theoretical percolation threshold for bond percolation on a face-centered cubic lattice. It indicates a highly interconnected network of water nanochannels. Notably, this percolation threshold is markedly smaller, and thus more realistic, than those found in atomistic simulations, which were not able to reproduce experimental values.

368


The ultimate goals of molecular modeling studies for PEM materials based on fully atomistic or CG models are to develop predictive models 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 properties, for the sake of material design we must understand how the chemical 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 utilized 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 appropriate choices in defining the bead structure and the interaction parameters between beads. The requirements for self-consistent approaches in molecular modeling and computational materials science are (1) an appropriate structural representation 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 simplistic (often, too short) representations of the ionomer chains could lead to largely inaccurate predictions of structure and properties. Atomistic models often fail in reproducing sizes and shapes of water clusters and polymer 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 modeling 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 description is able to capture essential parameters in synthesis, characterization, and development of advanced membrane materials for PEFCs at the relevant time and length scales.


369

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 understanding the structure and distribution of water in PEMs has stimulated many efforts in experiment and theory. Major classifications of water in PEMs distinguish (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 polymeric 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 conducting elements with strongly restricted, surface-like mobility of protons control membrane conductivity at low l, while proton current at large l will be carried 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 conductivity 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 determine 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 ~20pC. Notably, a freezing point suppression has also been observed by Cappadonia et al.132 and Cappadonia, Erning, and Stimming133 in Arrhenius

370


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 liquid water inside PEMs is unphysical and misleading. It is guided by empirical 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 frequently cited Schroeder’s paradox,136,137 which indeed refers to a rather logical 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 capillary 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 usually 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 consequence 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 90p.146 These assumptions are, however, largely uncorroborated. There is evidence neither for a significant hydrophobic gas porosity nor for the existence


371

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 predominantly 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 SchmidtRohr80 show no indication of hydrophobic pores; these models correspond 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 protogenic 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 interfacial 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 > l s, equilibrium water uptake is controlled by capillary forces. Deviations from capillary equilibrium arising at l < l s 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

372


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 temperatures (Ta,Tc), vapor pressures (Pav,Pcv), and total gas pressures (Pal, 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 considered conditions of membrane operation; this implies uniform temperature 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,

MwPEM (L ) M wext ( a v ) RT ln a v , a v P v/P s ,

(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: Pc Pg Pl .

(6.6)

In this description, the local state of water in the membrane is thus determined by two independent variables, m wPEM 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 Pcv – Pav y 0, or in total gas pressures, ΔPg Pcg – Pag y 0, between cathode and anode compartments will generate a finite water flux, jw y 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

373


Under fuel cell operation, a finite proton current density, jp y 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 membrane 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 polymer elasticity. In order to justify a scenario based on capillary condensation, isopiestic vapor sorption isotherms for Nafion138 in Figure 6.9(a) are compared with data on pore size distributions in Figure 6.9(b) obtained by standard porosimetry.142–144 In Figure 6.9(a), a simple fit function, ¤ Pv ³ L 3.0 ¥ s ´ ¦P µ

0.2

4

¤ Pv ³ 11.0 ¥ s ´ , ¦P µ

(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 distribution, similar to the function suggested in Eikerling et al.7:

L L max ,

¯

rc

0

¤ ¤ log(r/r ) ³ 2 ³ m dr exp ¥ ¥ ´, ¥¦ ¦ log s ´µ ´µ

¤ ¤ log(r/r ) ³ 2 ³ m with , dr exp ¥ ¥ ´, ¥¦ ¦ log s ´µ ´µ 0

¯

c

(6.8) where rm 0.75 nm, s 0.15, and l max 14. We can compute the Gibbs free energy of water sorption139 using the relationship $G s $G w RT ln

Pv , Ps

(6.9)

For water uptake by capillary condensation, the Gibbs free energy of water in pores is given by $G c $G w P cVw ,

(6.10)

374


#!$# % "##

P v/P s

#!$# % "##

%

"

FIGURE 6.9 Water 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) capillary isotherms (extracted from J. Divisek et al., Journal of the Electrochemical Society 145 (1998) 2677–2683) fitted by Equation (6.8).

with the capillary pressure Pc

2S cos Q , rc

(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

375


# !

!

!

&

&

&

&

%G &

%G &

&

& &

& &

&

&

$$"

$$"

FIGURE 6.10 Gibbs 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). v

experimental relations L f ( PPs ) and l g(rc) gives two expressions of the Gibbs free energy of water sorption as functions of l: $G s (L ) $G w RT ln f 1 (L ) and $G c (L ) $G w

2S Vw cos Q . (6.12) g 1 (L )

In Figure 6.10 we compare these two expressions obtained from the independent sets of experimental data. At l/l max > 0.2, the two graphs are indistinguishable—ΔG s(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/l max < 0.2, where the steeply increasing strength of water binding, seen in |ΔG s(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 mechanical 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 c $G w P cVw

2S Vw cos Q ¤ Pv ³ RT ln ¥ s ´ . rc ¦P µ

(6.13)

376


For given external vapor pressure and total gas pressure, the liquid pressure (or swelling pressure) in the membrane is obtained from Equations (6.6) and (6.13): Pl Pg

RT ¤ P v ³ ln . Vw ¥¦ P s ´µ

(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 l Pel .

(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 ionexchange resins,127 including the statistical mechanical theory of Flory and Rehner154 and the more recent model of Freger.155 Flory–Rehner theory provides the following relation for the elastic (or swelling) pressure: ¤ c ³ Pel G ¥ Fp1 3 Fp ´ , 2 µ ¦

(6.16)

where G is the shear modulus of the polymer matrix156 and f p is the volume fraction of polymer. The volume fraction of excess water, which corresponds to l b l – l s, is

Fw 1 Fp

NL b . NL b R

Here it is assumed that only excess water causes swelling. The parameter R Vp/Vw 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 Vp Mp/Rp 0.55 L mol 1 , with the molar mass of a monomer unit Mp 1,100 g mol–1 (i.e., the equivalent weight) and the density of the dry polymer r p 2.0 g cm–3. We thus obtain r 30.6, where we have used the value Vw 0.018 L mol 1 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 f p n 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

377


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, Pel

2 G ` Fp1 3 Fp7 3 3

(6.17)

with G`

N `kBT , V0

where V0 is the volume of the dry resin and Nb 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 Pel and f p is sensitive to the structural model of the membrane and the mechanism of swelling. Equation (6.17) may be appropriate 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 cylindrical domains of polymer fibrils or water channels, will lead to modified relations of Pel versus f p. 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 (Pg and Pv) and l: 13 7 3 § ¤ R ³ ¶ ¤ v 2 ¨¤ R ³ · P g RT ln P ³ .

G` ¥ ´ ¥ ´ Vw ¥¦ P s ´µ 3 ¨¦ NL b R µ ¦ NL b R µ · © ¸

(6.18)

This relation can adequately reproduce the shape of experimental vapor sorption isotherms in the regime corresponding to capillary condensation. The approach to the limit of l b n 0 (only surface water remaining) is given by

Lb

3 R ª g RT ¤ P v ³ ¹ ln º. «P 4 NG ` ¬ Vw ¥¦ P s ´µ »

(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).

378


The corresponding capillary pressure, Pc, Equation (6.11), and the external gas pressure, Pg, 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 relation between P el and the volume fraction of polymer f p given by Equation (6.16) or (6.17). f p can be related to the water uptake by

Lb

R 1 Fp , N Fp

where

R Vp Vw 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 1.4 103 atm. This implies that vapor equilibration of PEMs corresponds to large negative liquid pressures inside the membrane or that l b 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 l b 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 structural 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 n 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 selforganized polymer is an effectively cross-linked elastic medium. Under liquid-equilibrated conditions, swelling is not controlled by external vapor

379


pressure because capillary condensation does not occur at the membrane surfaces, but rather by the elastic pressure of the electrostatically crosslinked polymer. The equilibrium condition reduces to a purely mechanical balance of pressures: 13 7 3 ¤ R ³ ¶ 2 §¨¤ R ³ · Pg ,

¥ G` ´ 3 ¨¥¦ NL b R ´µ ¦ NL b R µ · © ¸

(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 maximum 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 observation of different water uptake by vapor-equilibrated and by liquid waterequilibrated 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 external membrane interfaces determines gradients in water activity or concentration; 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 condensation kinetics at membrane–vapor interfaces is jw

F k P v,eq (Lint ) P v , 2 RT v

[

]

(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 corresponds to the membrane water content, l int, at the interface. The relation

380


Pv,eq(l int) is the vapor sorption isotherm, as depicted in Figure 6.9(a), applied at the membrane–gas interface. Equally, we could write jw

¤ $G s (Lint ) ³ P v ¹ F ª k v P s «exp ¥

º 2 RT ¦ RT ´µ P s » ¬

(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 permeabilities 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

Nv Ps L , k v RTDeff cmax PEM

(6.23)

with Pv 2 if the PEM is in contact with vapor on both sides and Pv 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 steadystate 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): LcPEM

N v RTDeff cmax kv Ps

(6.24)

This parameter helps distinguishing the relative importance of interfacial kinetics and bulk transport. For LPEM < LPEMc, water transport through the PEM is dominated by interfacial water exchange, whereas for LPEM > LPEMc, bulk permeation of water prevails. The data obtained in Monroe et al.153 yield LPEMc _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


381

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 pathways. 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 mesoscopic scale, as well as the statistical geometry and percolation effects of the phase-segregated random domains of polymer and water at the macroscopic scale. A question of utmost interest is whether high proton mobility in aqueousbased PEMs is possible under conditions of minimal hydration and at elevated 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 encouraging insights for this endeavor. These experiments suggest that lateral proton 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 protogenic 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-functionalized 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 current 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 detrimental effects of membrane dehydration are multifold. It triggers morphological transitions that have been studied recently in experiment51,74–76 and theory.24,129,161,162 At water contents below the percolation threshold, the wellhydrated 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.

382


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 seminal publication in 1806—long before the discovery of the proton by Ernest Rutherford, which can be dated around 1918—the ingenious physical chemist Freiherr Theodore von Grotthuss introduced the rudimentary concept of structural diffusion of positively charged moieties (protons) in acidic solutions.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 proton 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, H 5O2 . Transformations H9O4 k H5O2 are triggered by hydrogen bond fluctuations 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 k 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 nanochannels, electrostatic effects at interfaces, and desolvation phenomena.129,160,170–174 In the case of cellular membranes and lipid monolayers, systematic experimental 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 90pC extinguishes the most favorable mechanism of proton transport through bulk water. Inevitably, aqueous-based PEMs for operation at higher temperatures (120–200pC) have to attain high rates of proton transport with a minimal amount of tightly bound surface water.


383

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 morphological 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 bulklike 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 parameterizations 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 solution 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 powerful 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 confinement in nanometer-sized pores and the role of acid-functionalized polymer walls on solvation and transport of protons in PEMs. The calculations by the Voth group revealed an inhibiting effect of sulfonate ions on proton motions. The EVB model by Kornyshev, Spohr, and Walbran88,94,186 was specifically designed to study effects on proton mobility due to charge delocalization within SO3– groups, side chain packing density, and fluctuations

384


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 conductivity 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 contributions 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,

£

L x p L

z0

¯ dzM (z)R (z),

(6.25)

z0

where L is the length of the pore and qz0 denotes the positions of the interfacial 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 contribution. 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


385

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 interactions 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 interfacial 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 > 100pC. 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 realistic 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 proton exchange between water of hydration and surface groups. The detailed calculations by Paddison and Elliott of hydrated polymeric fragments, including several side chains attached to a single polymer backbone, are insightful 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

386


O4 S1

H2 O3

S2

2.5 Å

4.5 Å O2 H1 2.4 Å O1

(a) Native Crystal

(b) Intermediate State

FIGURE 6.11 Ab 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 conformation of the native crystal (left).

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 proton complexes, which are formed almost simultaneously, stabilize the intermediate 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 conditions 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 composition and water content are fixed. Only minimal hydration could be considered. 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

387


Side chain “length” Fixed Carbon (a)

dCC (b)

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.)

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 mesoscopic 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 hydrophobic 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 concerted dynamics of polymer side chains, water, and protons at polymeric aggregates in PEMs. The approach implies that the effect of polymer dynamics on processes inside pores is primarily due to variations in chemical architecture, 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 c dCC c 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

388

Formation Energy (eV)


0.0 –0.5

–1.5 Point c Partially Dissociated Point b Point a

–2.0 –2.5 –3.0 5.0

Formation Energy (eV)

Point e Point d

–1.0

6.0

7.0

8.0

9.0

CF3CF2SO3H

10.0 11.0 12.0 C-C Distance (Å)

13.0

14.0

15.0

16.0

0.0 –0.5 –1.0 –1.5

Point b

Point c

–2.0

Partially Dissociated

–2.5 –3.0 5.0

Point e

Point d

CF3OCF2CF2SO3H

Point a 6.0

7.0

8.0

9.0

10.0 11.0 12.0 C-C Distance (Å)

13.0

14.0

15.0

16.0

FIGURE 6.13 The formation energy (Efuc) 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.)

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: c E uc f Etotal (dCC ) Etotal ,

(6.26)

c 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 dissociation 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 ( 1 mm. Ultrathin CLs with LCL _ 100–200 nm, on the other hand, can operate well without these additional components, based on sufficiently 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) consist 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 agglomerates. 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 volume fractions of primary and secondary pores are Xm and XM, respectively. The primary optimization target of CLs is the effectiveness factor of Pt utilization, (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 0 2 103 [mPt ] j 0* ' stat , with ' stat E S/V ' a g(Sr )

f (X PtC, X el ) X PtC

.

(6.58)


405

The structure-related statistical factors include the surface-to-volume atom ratio of Pt nanoparticles, e S/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,” ECL.7 The criterion for uniform reaction rate distributions is that the reaction penetration depth is comparable to or larger than the thickness of the CL (i.e., d CL > 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, d CL. 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 determine the irreversible voltage loss incurred by the CCL, h 0. Therefore, relations 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 h 0. We can express the corresponding current density as a two-variable function, j0 f(j0, d CL), where the reaction penetration depth, d CL, depends on h 0. This function can be used to determine the effectiveness factor, (CL. In the case of severely limited oxygen diffusion, the following relations for local oxygen partial pressure and current density can be obtained: § L ¤ z ³¶ p( z) exp ¨ CL ¥ 1 ·, LCL ´µ ·¸ ¨© D CL ¦ ª § LCL ¤ z ³ ¶ ¹ 1

·º «1 exp ¨ ¥ LCL ´µ ·¸ ¨© D CL ¦ » ¬

j( z) I

LCL D CL

D CL LCL

¤ H ³ j0 exp ¥ 0 ´ , I ¦ 2b µ

(6.59)

406


where z 0 is the PEM–CCL boundary z LCL is the CCL–GDL boundary p PO2 /PO2 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

4 FPO2 Do RTLCL

is a characteristic current density of diffusive flux through the layer (with a unit of amperes per square meter). For a fixed h 0, the overall effectiveness of Pt utilization can be defined by ' CL ' stat 'D

j0 2 103 [mPt ] j 0*

j0 , j0id

(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, j0id, 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

D CL LCL

¤ LCL ³ ¹ ª «1 exp ¥ ´ º. ¦ D CL µ » ¬

(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 d CL LCL. If we replace d CL by h 0, using Equation (6.59), we obtain an explicit expression for (CL as a function of the catalyst layer voltage loss: ' CL ' stat

§ ¤ H ³ ¶ ¹ ¤ H ³ ª j0 j0 exp ¥ 0 ´ «1 exp ¨ exp ¥ 0 ´ · º . I I ¦ 2b µ ·¸ » ¦ 2b µ ¨© ¬

(6.62)

Equations (6.59)–(6.61) represent a highly simplified scheme for evaluating various catalyst layer designs. Refinements of this crude framework for evaluating 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


407

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 ionomer is to be found in solubilized, colloidal, or precipitated forms. This influences 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 characterize 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 effective medium), and (2) relating effective properties to performance. The main structural effects in CLs occur at well-separated scales: at catalyst nanoparticles (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 electrocatalytic activity is critical in view of the design of highly performing catalyst systems.223–225 Obviously, a reduction in particle size improves the surface-tovolume ratio,e S/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 modeling of reactivities based on Monte Carlo simulations or mean field methods 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.

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The d-band model of Hammer and Norskov229–231 has been successful in relating trends in chemisorption energies for various adsorbates on transition 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, systematic 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 chemisorption 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 components 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 properties can be evaluated in view of capabilities for controlling sizes of carbon/Pt agglomerates, ionomer domains, and the resulting pore network topology. At macroscopic level, the overall relations between structure and performance are strongly affected by the formation of liquid water. Solution of such a model that accounts for these effects provides full relations among structure, properties, and performance, which in turn allow predicting architectures 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 catalyst 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, condensation); and


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terms that account for the transport of species—that is, the migration of electrons and protons in conduction media, the diffusion of dissolved 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 solvent 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 protons (H) and water (W) behave similarly. The correlation between hydrophilic species (H and W) and ionomer (N) is significantly stronger than that between those species and carbon (C). The autocorrelation functions, gSS and g HH, exhibit a similar structure as g WW. This indicates a strong clustering 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 suppressed compared to the primary peak in g WW due to electrostatic repulsion between these charged beads. g BB 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 backbone clusters (~2–3 nm). More detailed analysis of RDFs revealed a strong correlation of carbon particles and polymer backbones (gCB). This suggests that polymer backbones

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FIGURE 6.16 Equilibrium 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.)

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 provide valuable structural information at the nanometer scale that allows refining the picture of the phase-segregated catalyst layer morphology. Ionomer 10 W-W S-S H-H

2 1 0

W-W H-H

1

6 4

1.5

C-B

1

C-C S-S

0

C-B C-S

2 g (r)

5 4 3

2.5 B-B C-C

8 g (r)

g (r)

8 7 6

2 3 r (nm)

C-S

2 4

0

0.5 B-B

0

1

2 3 r (nm)

4

0

0

1

2 3 r (nm)

4

FIGURE 6.17 Site–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.)

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backbones form clusters or fibers that are attached to the carbon agglomerates. 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 structural correlations.22 Overall, the magnitude of the RDF peaks decreases from low to high polar solvent. Low F 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 connectivity 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 ionomer aggregates. Figure 6.18 depicts size distributions of ionomer domains in CLs obtained with differente. These coarse-grained MD calculations helped consolidate the main features of microstructure formation in CLs of PEFCs. They showed that the final microstructure depends on carbon particle choices and ionomer–carbon 8000 Dielectric = 2 Dielectric = 20 Dielectric = 80

a.u.

6000

4000

2000

0

0

5 10 Ionomer Cluster Size (nm)

15

FIGURE 6.18 Effect 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 permission from ACS.)

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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 ionomer, 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 catalyst 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 assumption 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.


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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 highfrequency limit. Three regimes of current density exist in electrode polarization curves: 1. a kinetic regime at small current densities, j0 I, in which all oxygen is consumed in a sublayer of thickness d CL, 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 distributions 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 thickness. 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:

D CL

I L . j0 CL

(6.64)

In the oxygen depletion regime, j0 >> I, only a thin sublayer with thickness d CL 0, these effective properties also vary spatially 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

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equation, which relates capillary pressure, pc, and capillary radius, rc, to local gas pressure, pg, and liquid pressure, pl: pc

2S cos(Q ) pg pl . rc

(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 agglomerates. The wettability of these pores is a vital property that controls the water formation in CCLs. Assuming that these pores are partially hydrophilic (typical wetting angles of carbonaceous materials are in the range of q ~ 80–90p), the liquid saturation is 1 S Xp

rc

¯ dr ' 0

dX p (r `) dr `

.

(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 vv tR(r , t) v j (r , t) Q(r , t) tt

(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 representation of the pore space by a bimodal E-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 E-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 x lv { 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


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current density are high (g { 0.8–0.9); the surface area for evaporation, x lv, 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-tovapor 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 parameters (10 mm thickness, 5 atm cathodic gas pressure, 89p 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 context, 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, bistability) 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, Dflo 2.0 t 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 transition between two principal states of operation occurs, as illustrated in

418


1 Primary Pores Filled

0.9

E/V

0.8 0.7

Transition Region

0.6 Fully Flooded

0.5

0

0.2

0.4 0.6 j0/A cm–2

0.8

1

FIGURE 6.19 Plot 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.

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 transition between the two states of operation can occur monotonously or can involve bistability as a signature of nonlinear coupling of liquid water accumulation, 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 current 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 parameters 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 management in CCLs is to push back capillary equilibrium to small enough pores so that liquid water formation cannot block gaseous transport in secondary

419


!

!

!

FIGURE 6.20 The stability diagram displays the effect of the wetting angle on CCL operation, distinguishing 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.)

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 90p, high total gas pressure, and high temperature of operation.

6.10 Concluding Remarks Numerous demonstrations in recent years have shown that the level of performance of present-day polymer electrolyte fuel cells can compete with current energy conversion technologies in power densities and energy efficiencies. However, for large-scale commercialization in automobile and portable 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 breakthroughs, 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

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concepts of proton-conducting membranes and catalyst layers. Due to recent detailed accounts of the status of catalyst layer research, the present 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 systematic optimization of performance. The grand challenge is to understand relations of primary chemical structure of materials, composition of heterogeneous 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 determined 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 benchmark 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 experimental studies exist for alternative membranes, Nafion provides the most complete 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 channels 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 primary chemical architecture of the ionomer and allow predicting the morphology, 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.


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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 statistical sampling of structural configurations or elementary transport processes. 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 forcematching 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 selforganized, phase-segregated morphology. Structural features of domains of polymer and water and of the interfaces between them are in good agreement 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 percolation threshold is markedly smaller than those found in other simulations that employed significantly shorter representations of the ionomer. This finding 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 transport properties at acid-functionalized polymer–water interfaces, as discussed in Section 6.7.3. They provide insights into the structure of surface water, correlation effects between surface groups, and fundamental transport mechanisms 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 systematic 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

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bottom-up design of advanced polymer materials that are optimized for operation under minimal hydration—that is, low relative humidity and/or elevated temperature (>100pC). 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 molecules correspond to capillary condensation of bulk-like water in pores. Based on this mechanism of water uptake, the thermodynamic model of water sorption 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, Pc, 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 Pel and the volume fraction of polymer f p. 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 properties, which can be included in the model of membrane operation in PEFCs. As discussed in Section 6.8, hydraulic permeation transpires as the dominant 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 permeation 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 existing 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 modeling 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;


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understanding the coupling of water distribution and fluxes in the membrane 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 parasitic voltage losses, balancing of water distribution and fluxes, and optimization 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 operation. In spite of the structural complexity of catalyst layers, existing tools in molecular modeling and physical theory have contributed to the fundamental 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 minimization 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 comparing 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; and identify 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 vapor cmax: 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)

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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) 0 j : 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) kPm: 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 flux Mp: molar mass of polymer (g mol–1) Nb: effective number of polymer chains in resin r N l : molar flux of liquid water in the membrane n0: number of SO3– groups in the dry membrane nd: electro-osmotic drag coefficient in PEM nt: number of electrons transferred in the overall fuel cell reaction P: 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)


425

rO rM: radius of primary and secondary pores in CCL (nm) ΔS: reaction entropy (J mol–1 K–1) T: temperature (K) Sr: liquid saturation U(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 a ij: effective bead radius (nm) (CL: effectiveness factor of CL d CL: reaction penetration depth (mm) e fc: fuel cell efficiency e th: ideal theoretical efficiency e V: voltage efficiency hCCL: voltage losses incurred in the CCL (V) h other: parasitic voltage losses incurred in components other than PEM and CCL (V) q: contact angle (p) l: water content in PEM, number of moles of water molecules per moles of acid head groups l b: water content corresponding to bulk-like water l s: percolation threshold of water content l s: water content corresponding to surface water m wPEM: chemical potential of water (kJ mol–1) n0: average pore volume in the dry membrane r(r): density distribution (nm–3) r p: density of polymer (g cm–3) s: surface tension (N m–1) s bb: percolation bond conductance (S) s p: proton conductivity of PEM (S cm–1) f p: volume fraction of polymer f w: volume fraction of water Dij0: depth of the LJ potential well (kJ mol–1 K–1)

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Index 2050-A, 199 2050-HF, 199 2050-L, 199 A Acetylene black, 247 Aciplex®, 138, 353 AFM. See Atomic force microscopy (AFM) Aluminum, 335 Aluminum-magnesium alloy, 335 Anode catalyst(s), 4 loadings, 4 reformate-tolerant, 41–45 improved, 43–45 mechanistic studies, 42–43 Atom-transfer radical polymerization (ATRP), 154 Atomic force microscopy (AFM), 117–118, 126, 145, 153 Atomistic simulations, 359–362 ATRP. See Atom-transfer radical polymerization (ATRP) Avcarb 1071HCB, 215 B Bac2 Conductive Composites, 322 BAM®, 112 Nafion vs., 114 oxygen reduction and transport characteristics, 120 proton mobility, 113 SPEEKK vs., 114 Bipolar plate(s) based composite, 316 basic structure, 309–311 comparative costs, 332 competition 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, 315 from lab-scale to large quantity production, 336–337 materials with better performance and durability, 334–335 comparison of major properties and cost of, 338 composite, 316–325 carbon/carbon, 317–319 filler, 316 matrix, 316 thermoplastic-based, 321–325 thermoset-based, 319–321 metal, 325–333 key fabrication processing of, 328–333 materials used for, 326–327 welding, 332–333 progress and challenges in development of, 315–333 traditionally used, 314–315 new activities and development trends, 333–337 new manufacturing processes, 335–336 performance targets of, 312 requirements, 313 schematic indication in simplified PEM fuel cells, 310 SS316 vs. graphite for, 315 stack cost, 315 transportation applications, 311, 312, 313 Borides, 328, 334 Boronization, 330, 331 BP2000, 34 BPSH. See Disulfonated biphenol poly(arylene ether sulfone) (BPSH)

435

436

C Cabot Fuel Cells, 12 Capillary flow porometry, 259 Capillary pressure, 259 Capillary rise method, 69 Carbides, 36 Carbon(s) PTFE and, 223 synthetic mesoporous, 39, 41 Carbon black(s), 37 conventional, 37–38 modification of, 38 Carbon cloth, 207–208 fabrication, 207–209, 210 SEM of, 208–209 Carbon corrosion, 30 measurements, 34 mechanism, 33–34 Pt promoted, 36 Carbon fiber paper, 196–207 with aerogels, 206–207 large-scale production, 204–205 manufacturing, 204–206 SEM picture, 202–203 Carbon nanofibers, 38 Carbon nanotube (CNT), 38 catalyst layer and, 77 Carbon support materials, 32, 37–41 synthetic, 38–41 mesoporous, 39, 41 Catalyst-coated gas diffusion electrode (CCGDL), 70–75 CCM vs., 77 gradient, 71–75 catalyst, 71–73 dual, 73–75 Nafion, 73 uniform, 70 Catalyst-coated membrane (CCM), 76–77 CCGDL vs., 77 conventional, 76–77 nanostructural thin-film electrode, 77 Catalyst layer (CL), 62 with additives, 80–81 CNT-based, 77 columnar oxide supported, 77–78

Index

components and their corresponding function, 64–70 dual-bound composite, 75–76 fabrication, 81–91 first generation, 81–82 Pt black catalyst, 81–82 PTFE-bound, 82 gas porosity, 404 ionic conductivity, 70 ionomer in, 403–419 challenges for design and operation, 404–406 macrohomogenous, 412–414 mesoscale simulations of selforganization in, 409–412 multiscale modeling scheme, 406–409 water management in, 414–419 location of, 62 Nafion loading, 68 nanowire-based three-dimensional hierarchical core/shell, 79 with novel ionomers, 81 novel structural, 77–81 optimization, 91–96 composition, 92–95 experimental studies on, 93–95 microstructure, 95–96 modeling and simulation in, 92–93 porosity, 69 properties, 66–70 prospects and conclusions, 96–97 Pt utilization and, 66–68 PTFE-bound, 64 self-supported, 80 thin-film, 64, 70 dual-bound composite, 75–76 fabrication, 83–91 DLR process in, 89–90 dual ion-beam assisted deposition in, 87–88 electrodeposition in, 88 electrophoretic deposition in, 91 electrospray technique in, 90–91 ink-based, 83–86 inkjet printing in, 85–86 pulsed laser deposition, 89

437

Index

RSDT in, 88–89 screen printing in, 84–85 in situ, 86–89 sol-gel Pt application in, 91 spray coating in, 85 sputter deposition in, 86–87 nanostructured, 77, 78 preparation of, 65 types of, 70–76 CCGDL, 70–75 (See also Catalyst-coated gas diffusion electrode (CCGDL)) CCM, 76–77 (See also Catalystcoated membrane (CCM)) types of, 70–81 wetting property, 68–69 Catalyst performance targets, 4 Cathode catalyst(s) enhanced activity, 14–29 Pt alloy, 14–20 loadings, 4 stability, 29–37 alternative supports for, 35–37 high voltage and, 32–34 potential cycling and, 31, 32 Pt electrochemical area loss and, 29–30 toward high-voltage excursions, 34–35 Cathode loadings, 4 CCGDL. See Catalyst-coated gas diffusion electrode (CCGDL) CCM. See Catalyst-coated membrane (CCM) Cell reversal tolerance, stability and, 46–47 Chemical vapor deposition, 223 Chemical vapor-infiltrated graphite carbon, 317 Chromium, 330 CL. See Catalyst layer (CL) Cladding process, 331 CNT. See Carbon nanotube (CNT) CO2 poisoning, 42 Coarse-grained molecular dynamics (CGMD) simulations, 363–367, 409

Coating compounds, 328 Cold-start tests, 279 Contact angle of moving droplet, 254 Contact angle tests, 254, 281 Copolymerization, 144 Copper diffusion layers, 220–221 Corrosion studies, 279–280 Cross-link density, 163 Cross-linkers, 150, 156, 163 Cyclic voltammogram, 67 D Darcy’s coefficient, 399 Darcy’s law, 261 Density functional theory (DFT), 351, 408, 421 calculations of water binding, 371 hybrid, 360 VASP and, 387 Desulfonation, 136 Diffusion layer (DL), 192–288 capillary pressure, 259 cold-start tests, 279 contact angle of moving droplet, 254 contact angle tests, 281 corrosion studies, 279 costs, 194 direct visualization in DMFC, 267 electricity conductivity, 273–274 in-plane, 273 through-plane, 273–274 engineered, 215–221 flow field interaction, 282–286 future direction, 286–288 gas transport properties, 260–267 general transport properties, 255–260 hydrophobic treatment, 227–233 (See also Hydrophobic treatments) hydrophobicity and hydrophilicitiy, 251–255 in-plane permeability, 261–264 internal contact angle, 254–255 liquid transport properties, 267–272 permeability, 270–271 mechanical properties, 276–278

438

performance, 224–227 in PEMFCs, 224–226 polarization curves showing effect of, 225 pore size distribution, 256–259 porosity, 255–256 pressure drop measurements, 260 pressure drop tests, 282–284 properties and measurements, 248–286 sessile drop for, 251 silicon-based, 221, 223 thermal conductivity, 274–276 thickness, 249–251 through-plane permeability, 264–266 transport, 255–272 treatment and coatings, 227–248 hydrophilic, 233–234 hydrophobic, 227–233 microporous layers, 234–248 types, 196–227 carbon cloth commercially available, 210 fabrication, 207–208 SEM of, 208 carbon cloth fabrication, 207–209, 210 SEM of, 208, 209 carbon fiber paper commercially available, 198–201 fabrication, 196–207 SEM of, 202–203 engineered, 215–221 metal, 209, 211–221 foams, 215 meshes, 211 micromachined, 214 SEM of, 212 sintered, 213–214 silicon-based materials, 221–223 visualization techniques, 284–286 water balance analysis, 271–272 Wilhelmy method for, 252–254 Diffusion medium. See Diffusion layer (DL) Direct ethanol fuel cell, 211

Index

Direct membrane fuel cells (DMFCs), 4, 28, 39, 52, 120 cell configurations, 29 liquid-feed polarization curve, 50 methanol crossover, 122 performance degradation, 52 PtRu particles for, 39, 52 reviews of technology, 47 Direct methanol fuel cell, 211, 215 cathode loadings, 4 diffusion layers (See Diffusion layer (DL)) performance comparison of, 226–227 hydrophilic treatment, 234 hydrophobic treatment, 232–233 microporous layers, 246–248 visualization of gas bubbles in, 267 Dissipative particle dynamics (DPD), 363 Disulfonated biphenol poly(arylene ether sulfone) (BPSH), 120, 144–145, 153 mechanical properties, 130 MeOh diffusion coefficients, 126 Nafion 117 vs., 121 Divinylbenzene (DVB), 131, 156, 158, 159 DLR process, in thin-film CL fabrication, 89–90 DMFC. See Direct membrane fuel cell (DMFC) DMFCs. See Direct methanol fuel cells (DMFCs) DoE Hydrogen Program, 318 Double-walled nanotubes (DWNTs), 39 DPD. 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, 35 Electricity conductivity, 273–274 in-plane, 273 through-plane, 273–274

439

Index

Electro-osmotic drag, 122–123, 394–397 Electrocatalyst(s), 5–52 anode loadings, 4 carbon support materials, 37–41 cathode loadings, 4 computation studies, 7–9 discovery, 5–9 electrochemical screening of, 8 preparation, 9–13 colloidal routes for, 10–11 conventional routes for, 9–10 molecular precursor routes for, 11–12 vapor phase routes, 12–13 scanning electron micrographs, 6 testing, 13–14 Electrochemical impedance spectroscopy, 355 Electrodeposition, in thin-film CL fabrication, 88 Electron micrograph, restored highresolution transmission, 4 Electron microscopy, 355 Electrophoretic deposition, in thin-film CL fabrication, 91 Electrospray technique, in thin-film CL fabrication, 90–91 Enhancement layer, 214 EP40, 198 EP40T, 198 ETFE-g-PSSA. See Ethylene-tetrafluoroethylenecopolymer (ETFE-g-PSSA) Ethylene-tetrafluoroethylenecopolymer (ETFE-g-PSSA), 125 BAM vs., 114 methanol uptake, 125 Nafion vs., 114, 125 oxygen permeability, 121 oxygen reduction and transport characteristics, 120 proton conductivity, 111, 112 comparative studies of, 113, 114, 120, 121 water content and, 112 radiation grafted, 157 SPEEKK vs., 114 Ex situ visualization techniques, 270

F Fe-and Co-based catalysts, 25–27 Fe-Cr-V alloy, 334 FEA. See Finite element analysis (FEA) Fenton test, 134 FEP. See Fluoroethylenepropylene (FEP) Finite element analysis (FEA), 307 Flemion®, 353 Flow field interaction, 282 Flow field plate. See Bipolar plate(s) Fluoride evolution rate (FER), 134–135 Fluorinated ethylene-propylene copolymer, 130–131 Fluoroethylenepropylene (FEP), 227 Forchheimer equation, 261 Fortron PPS, 322 Fourier transform infrared spectroscopy, 355 Furnace blacks, 37 G Gas diffusion layer, 62, 306, 309. See also Diffusion layer (DL) Gas-permeable elastomer, 223 Gas transport properties, 260–267 in-plane permeability, 261–264 relative permeability, 266 through-plane permeability, 264–266 GD05505G, 200 GD07508G, 200 GD12012G, 200 GD65055G, 200 GD05505T, 200 GD07508T, 200 GD12012T, 200 GD65055T, 200 GDL10BA, 200 GDL24BA, 201 GDL25BA, 201 GDL34BA, 201 GDL10BB, 201 GDL10BC, 201 GDL24BC, 201 GDL25BC, 201 GDL34BC, 201 GDS1120, 199 GDS2120, 199

440

GDS3215, 198 GDS22100, 199 Gebel’s calculations, 115 GLAD. See Glancing angle deposition (GLAD) Glancing angle deposition (GLAD), 78 GRAFCELL, 320, 321 Grafcell flexible graphite diffusion layers, 216–220 Graphite, 323, 328 alternative materials to, 315–316 bulk electrical conductivity, 315 chemical structure, 314 cost, 315, 319, 330 fillers, 316, 317 functional requirement of plates and, 320 ratio between matrix and, 320 high-density, 314 layered metals and thermal expanded, 316 metal vs., 333, 335 natural, 314, 319 artificial vs., 320 nonporous, 314 normal carbon vs., 314 physical and chemical properties, 314 reinforced, 324 shortcomings, 314 SS316L vs., 315, 326, 327, 333 synthesized, 319 synthetic, 314 thermoplastic-based composite, 316 thermoset-based composite, 317, 319, 321 Grotthuss mechanism, 109 H H1-Pt catalysts, 11 Hafnium, 326 Hagen–Poiseuille–Kozeny equation, 399 Hexafluoropropylene, 228 High-throughput screening, 5–7 High-velocity electromagnetic forming process (HVEF), 335–336

Index

HVEF. See High-velocity electromagnetic forming process (HVEF) Hydrophilic treatment(s), 233–234 measurement of, 68–69, 251–255 for PEMFC, 233–234 Hydrophobic treatment(s), 227–233 for DMFCs, 232––233 effect, 229–232 fabrication processes and procedures, 227–229 gas permeability and, 229 measurement of, 251–255 thermal conductivity and, 231 water flooding and, 230 Hyflon® Ion E83, 353 I Imidazole, proton conduction and, 169 In-plane conductivity, 273 In-plane permeability, 261–264 Infrared spectroscopy (IR), 355 Inkjet printing, in thin-film CL fabrication, 85–86 Internal contact angle, 254–255 Ionomer(s), 165–166 in catalyst layer, 403–419 challenges for design and operation, 404–406 macrohomogenous, 412–414 mesoscale simulations of selforganization, 409–412 multiscale modeling scheme, 406–409 water management in, 414–419 membranes molecular modeling of selforganization in, 359–368 perfluorinated sulfone, 96 perfluorosulfonic acid, 353 status and directions in research, 353–355 structural organization and dynamic properties of, 352–368 structure and dynamics in, 355–359

441

Index

K Ketjen, 32, 34 Key fabrication processing, 328–333 coating materials, 328–330 coating process, 330–332 forming process, 328 L Lattice–Boltzmann simulations, 363 Lennard–Jones potential, 365 Liquid crystal polymer (LCP), 322 Liquid transport properties, 267–272 permeability, 270–271 Long term durability tests, 169 LT1200-N, 199 LT1300-N, 199 M MEA. See Membrane electrode assembly (MEA) Mechanical properties, 276–278 Membrane electrode assembly (MEA), 4, 62, 63 constituents, 192, 193 cyclic voltammogram for cathode of, 67 description, 308 targets, 4, 5 durability of, 6 performance of, 5 MeOH oxidation catalyst(s), 43, 47–52 mechanistic advancements, 47–48 MeOH-tolerant oxygen reduction catalysts, 27–29 Mercury porosimetry, 256 Mesoscale simulations coarse-grained, 362–368 DPD in, 363 Lattice–Boltzmann, 363 SCMF theory in, 362 of self-organization, 409–412 Metal(s) for bipolar plates, 325–333 for DL, 209, 211–221 foams, 215 meshes, 211 micromachined, 214

SEM of, 212 sintered, 213–214 graphite vs., 333, 335 welding, 332–333 Methanol crossover, 122–126 diffusion mode vs. electro-osmotic drag, 122–123 electro-osmotic drag vs.diffusion mode, 122–123 sulfonation and, 123–124 Method of standard porosimetry, 257–258 Microelectromechanical system fabrication techniques, 221 Micromachined meshes, 214 Microporous layers, 234–248 in DLFC, 246–248 effect on fuel cell performance, 237–239 fabrication processes, 236–237 hydrophobic content, 240, 242 multilayered, 244–246 parameters affecting, 239–244 pore formers in, 244 thickness and carbon loading, 239–240 types of carbon particles, 242–244 Molybdenum, 330 Monte Carlo simulation, 407 Multiple-walled nanotubes (MWNTs), 39 MWNTs. See Multiple-walled nanotubes (MWNTs) N Nafion®, 64, 65, 138 coarse-grained model, 365 degradation mechanism, 134 MeOH diffusion flow, 123, 126 microphase separation, 117 microstructure, 115, 116 in hydrated membranes at different water contents, 366 morphology, 118 N117, 117 BPSH vs., 121 MeOH permeability, 123 S-SEBS vs., 124

442

oxygen reduction and transport characteristics, 120 phase separation, 114, 117 proton conduction, 109 comparative studies of, 119 proton mobility, 113 PVDF and, 161 QENS data for hydrated, 357 source of protons, 111 SPEEKK vs., 113, 114, 115, 117 water absorption, 123 Navier–Stokes equation, 395 Nb-doped TiO2, 36 NEDO. See New Energy and Industrial Technology Development Organization (NEDO) Nernst–Einstein relationship, 110 Neutron imaging in fuel cells, 268–269 New Energy and Industrial Technology Development Organization (NEDO), 4 Niobium, 326 Nitridation, 330 Nitrides, 36, 328, 334 NMR. See Nuclear magnetic resonance (NMR) NMR relaxometry, 357 Non-Pt catalysts, 24–29 Nonconductive whiskers, 36–37 Nuclear magnetic resonance (NMR), 355 O ONIOM method, 361 Optical contact measurement system, 68 Oxides, 35–36, 328 Oxygen permeability, 119–122 Oxygen reduction reaction, 238, 242 P P50, 198 P75, 198 PAN. See Polyacrylonitrile (PAN) PBI. See Poly(benzimidazole) (PBI) Pd-based catalysts, 25

Index

PEM. See Proton exchange membrane (PEM) PEMFC. See Polymer electrolyte membrane fuel cell (PEMFC); Proton exchange membrane fuel cell (PEMFC) Perfluorinated PEM, 140, 141 Perfluorinated sulfone ionomers (PFSIs), 96 Perfluoroalkoxy polymer, 130 Perfluorosulfonic acid ionomer membranes, 353, 360, 363, 371 effect of hydration on local structure, 361 enhancing proton conduction with, 404 transport properties, 359 Perforated pores, 217–219 Perylene red, 37 PFSA. See Perfluorosulfonic acid ionomer membranes PFSIs. See Perfluorinated sulfone ionomers (PFSIs) Phase separation, 114 Platinum alternative promoters to, 24 electrochemical 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, 15–17 nanoparticles and particle size effect, 17–20 Platinum catalyst(s) alternative, 50–52 stabilization, 31, 34–35 layer for, 65 Platinum core-shell catalysts, 20–24 Polphosphazenes, 149 Poly(4-vinylpyridine) (P4VP), 163 Polyacrylonitrile (PAN), 196, 197, 207 carbon fiber manufacturing from, 204 Poly(aryl ether ketones), 143 Polyarylenes, 142–149

Index

Poly(benzimidazole) (PBI), 143, 144, 163, 165, 169, 355 Poly(dimethylsiloxane), 223 Poly(ether sulfone) (PES), 161, 163, 164 Polyethylene substrates, 125 Poly(ethyleneimine) (PEI), 163, 164 Polyimides, 143, 144 Polymer electrolyte fuel cell (PEFC) catalyst layers, 348 challenges for materials and operation, 346–347 energy conversion, 344–346 hierarchy y of scales, 351–352 membrane, 348 atomistic simulations of, 359–362 conductivity, network model of, 390–393 in fuel cell modeling, 397–403 mesoscale coarse-grained simulation, 362–368 molecular modeling of selforganization of, 359–368 structural evolution of, 354 structural 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, 369–381 seven-layer structure and basic processes in, 345 water sorption in, 369–381 phenomenology of, 370–371 thermodynamic model of, 371–381 Polymer electrolyte membrane. See Polymer electrolyte fuel cell (PEFC), membrane Polymer electrolyte membrane fuel cell (PEMFC), 192 cost component distribution, 195 POLYMET®, 224 Polyol method, 11 Polyphenylene oxide (PPO), 143, 162 Polyphenylene sulfide (PPS), 322, 324

443

Polypropylene (PP), 321 Polysiloxane, sulfonated, 150 Polytetrafluoroethylene (PTFE), 192, 233 carbon and, 223, 236 catalyst, 64 fabrication process and procedure, 227–228 effects of, 229–231 thermal conductivity and, 231 Poly(vinyl imidazole), 169 Poly(vinyl phosphanate)-b-polystyrene, 162 Poly(vinyl triazole), 169 Poly(vinylidene difluoride-cohexafluoropropylene), 154 Poly(vinylidene fluoride) (PVDF), 125, 131, 322 as base substrates, 156 dehydrofluorination, 161 Nafion and, 161 radiation grafted, 157 SEBS and, 162 sulfonated, 153 Pore size distribution, 256–259 Porosity, 255–256 Positron annihilation spectroscopy, 355 PP. See Polypropylene (PP) PPO. See Polyphenylene oxide (PPO) PPS. See Polyphenylene sulfide (PPS) Pressure drop tests, 260, 282–284 Proton conduction, 108–119 acid content and, 111 AFM in study of, 118 bulk water for, 354 connectivity of aqueous domains, 110 distance between acid groups, 110 enhancing, 404 imidazole and, 169 measurement, 118 morphology and, 114 Nernst–Einstein relationship and, 110 polymer microstructure and, 114 pore-scale models of, 383–385 of SPEKK/PEI blends, 164 sulfonation and, 123–124 triazole and, 169 water and, 109, 112–113, 354

444

Proton exchange membrane (PEM), 62 B,C C-trifluorostyrene-based, 139 block copolymer, 151–155 chemical stability, 131–136 composite, 165–166 cross-linked,sulfonicacid-substituted, polyphosphazene-based, 150 development of new, 354 fluorinated block copolymer, 154 fluorosulfonic acid-based, 139 fuel cell (PEMFC) air-breathing, 231 cost component distribution, 195 diffusion layers (See Diffusion layer (DL)) performance comparison of, 224–226 direct hydrogen, 195 factors affecting power and density, 120 hydrophobic treatment for, 227–232 polarization curves with different carbon loadings, 241 water flooding, 230–231 future directions, 170 Gierke model, 355 graft copolymer, 155–159 for 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, 137–169 mechanical properties, 129–131 methanol crossover, 122–126 network model of conductivity, 390–393 oxygen permeability, 119–122 perfluorinated, 140, 141 polyarylene-based, 143 polymer blends, 161–164

Index

properties and structure-property relationships, 108–136 proton conduction, 108–119 (See also Proton conduction) QENS studies, 357, 358 radiation-grafted, 156 random network model, 355 simple water channel models, 356 sites for radical attack in, 134 stability, 356 statistical copolymers, 137–150 structure of water in, 369–370 synthesis, reviews of, 355 water sorption in, 369–381 phenomenology of, 370–371 thermodynamic model of, 371–381 water transport, 127–129 Proton mobility, near polymer-water interface, 385–390 Proton transport, 381–397 near polymer-water interface, 385–390 Ohm’s law of, 396 pore-scale models of, 383–385 P50T, 198 P75T, 198 PTFE. See Polytetrafluoroethylene (PTFE) PtMo catalyst(s), 44–45 durability, 46 PtRu catalyst(s), 38, 47, 52 durability, 46 layer for, 46 preparation, 12 variants, 43–44, 48–50 Pulsed field gradient-NMR experiments, 357, 358 Pulsed laser deposition, 4, 89 PVDF. See Poly(vinylidene fluoride) (PVDF) P4VP. See Poly(4-vinylpyridine) (P4VP) PWB-3, 224 PyroCell, 314 Q Quasi-elastic neutron scattering (QENS), 355

Index

R Radial flow permeability testing apparatus, 263 Raman spectroscopy, 355 Reactive spray depositon technology (RSDT), 88–89 Reference interaction site model (RISM), 362 Reformate-tolerant catalyst(s), 41–47 anode, 41–45 improved, 43–45 mechanistic studies, 42–43 stability, 45–47 RISM. See Reference interaction site model (RISM) Rotating disk electrode, 7 RSDT. 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, 355 SCMF. See Self-consistent mean field theory Screen printing, in thin-film CL fabrication, 84–85 SDAPP. See Sulfonated Diels-Alder poly(phenylene) SECM. See Scanning electrochemical microscopy (SECM) Self-consistent mean field theory, 362 SENS. See Small-angle neutron scattering (SENS) Separation plate. See Bipolar plate(s) Sessile drop method, 68, 251 Shutdown-start-up tests, 169 SIGRAFLEX, 324 Silicon diffusion layers, 220–221 Single-walled nanotubes (SWNTs), 39

445

Sintered metals, 213–214 Small-angle neutron scattering (SANS), 115, 140, 355 Small-angle x-ray scattering (SAXS), 115, 140, 355 Sol-gel Pt application, in thin-film CL fabrication, 91 SPAEKs. See Sulfonated poly(arylene ether)s SPAES. See Sulfonated poly(aryl ether sulfone)s SPEEKK. See Sulfonated poly(ether ether ketone ketone) (SPEEKK) SPES. See Sulfonated poly(ether sulfone) Spray coating, in thin-film CL fabrication, 85 Spray pyrolysis, 12 sPSO2, 145 Sputter deposition, in thin-film CL fabrication, 86–87 SS316, 315, 326, 327 SS430 core, 331 S-SEBS. See Sulfonated polystyrene-bpoly(ethylene-r-butylene)b-styrene (S-SEBS) S-SIBS. See Sulfonated polystyrene-b(isobutylene)-b-sulfonated polystyrene (S-SIBS) SS316L, 333 Stack cost, 315 Stack targets, 5 Starbons, 41 Sulfonated Diels–Alder poly(phenylene), 147 Sulfonated poly(aryl ether ketones), 144 Sulfonated poly(aryl ether sulfone)s, 143–144, 162 Sulfonated poly(arylene ether)s, 142, 143 Sulfonated poly(ether ether ketone ketone) (SPEEKK), 112 BAM vs., 114 ETFE-g-PSSA, 114 microphase separation, 117 Nafion vs., 113, 114, 115, 117 SAXS analyses, 115

446

Sulfonated poly(ether sulfone), 161 Sulfonated polysiloxane, 150 Sulfonated polystyrene-b-(isobutylene)b-sulfonated polystyrene (S-SIBS), 124 casting solvent, 153 Sulfonated polystyrene-bpoly(ethylene-r-butylene)b-styrene (S-SEBS), 120 Nafion 117 vs., 124 ratios of MeOH to water uptakes for, 125 ratios of proton conductivity to MeOH permeability, 124 SuPAES. See Sulfonated poly(aryl ether sulfone)s Superior MicroPowder, 12 Superprotonic conductor, 355 SWNTs. See Single-walled nanotubes (SWNTs) T Tantalum, 326 TEM. See Transmission electron microscopy (TEM) TG-090, 224 TGP-H030, 199 TGP-H060, 199 TGP-H090, 199, 228 TGP-H120, 199 Thermal conductivity, 231, 274–276 Through-plane permeability, 264–266 TiO2, 35 Ti4O7, 35 Titanium, 326 Transmission electron microscopy (TEM), 115 Transparent fuel cells, 267–268 Transportation fuel cells, 312 Triallyl cyanuarate, 156 Triazole, as proton conductor, 169 Tungsten carbide, 36 Tungsten nitride, 36

Index

U U. S. Department of Energy (DoE), 4 Ultrasmall-angle x-ray scattering (USAXS), 356 USAXS. See Ultrasmall-angle x-ray scattering (USAXS) V Vanadium, 330, 334 VASP. See Vienna Ab Initio Simulation Package (VASP) Vectra LCP, 322 Vienna Ab Initio Simulation Package (VASP), 386 Visualization techniques, 284–286 Vulcan XC-72R, 247 W Water balance analysis, 271–272 Water channel models, 356 Water flooding, 230–231 Water management layer, 243 Water transfer region, 245 Water transport, 127–129 WAXS. See Wide-angle x-ray scattering (WAXS) Wide-angle x-ray scattering (WAXS), 355 Wilhelmy method, 68, 252–254 Wilhelmy plate gravimetric technique, 69 Williamson ether synthesis, 153 X X-ray radiography, 269–270 XC72R, 32, 34 Y YLP-100 single-mode ytterbium fiber laser, 332 Z Zirconium, 326

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