University of Southern California Dissertations and Theses

Development of high-performance polymer electrolyte membranes for direct methanol fuel cells

(USC Thesis Other)

Development of high-performance polymer electrolyte membranes for direct methanol fuel cells

PDF

Download a page range

Download transcript

Copy asset link

Request this asset

Transcript (if available)

Content INFORMATION TO USERS
This manuscript has been reproduced from the microfilm master. UMI films
the text directly from the original or copy submitted. Thus, some thesis and
dissertation copies are in typewriter face, while others may be from any type of
computer printer.
The quality of this reproduction is dependent upon the quality of th e
copy subm itted. Broken or indistinct print, colored or poor quality illustrations
and photographs, print bleedthrough, substandard margins, and improper
alignment can adversely affect reproduction.
In the unlikely event that the author did not send UMI a complete manuscript
and there are missing pages, these will be noted. Also, if unauthorized
copyright material had to be removed, a note will indicate the deletion.
Oversize materials (e.g., maps, drawings, charts) are reproduced by
sectioning the original, beginning at the upper left-hand comer and continuing
from left to right in equal sections with small overlaps.
Photographs included in the original manuscript have been reproduced
xerographically in this copy. Higher quality 6" x 9” black and white
photographic prints are available for any photographs or illustrations appearing
in this copy for an additional charge. Contact UMI directly to order.
ProQuest Information and Learning
300 North Zeeb Road, Ann Arbor, M l 48106-1346 USA
800-521-0600
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
DEVELOPMENT OF HIGH PERFORMANCE
POLYMER ELECTROLYTE MEMBRANES
FOR DIRECT METHANOL FUEL CELLS
by
Anthony Richard Atti
A Dissertation Presented to the
FACULTY OF THE GRADUATE SCHOOL
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY
(CHEMISTRY)
August 2000
Copyright 2000 Anthony Richard Atti
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
UMI Number: 3042408
Copyright 2000 by
Atti, Anthony Richard
Ail rights reserved.
___ ®
UMI
UMI Microform 3042408
Copyright 2002 by ProQuest Information and Learning Company.
All rights reserved. This microform edition is protected against
unauthorized copying under Title 17, United States Code.
ProQuest Information and Learning Company
300 North Zeeb Road
P.O. Box 1346
Ann Arbor, Ml 48106-1346
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
UNIVERSITY OF SOUTHERN CALIFORNIA
THE GRADUATE SCHOOL
UNIVERSITY PARK
LOS ANGELES. CALIFORNIA 90007
This dissertation, written by
A nthony R. A t t i
under the direction of hiL s Dissertation
Committee, and approved by all its members,
has been presented to and accepted by The
Graduate School, in partial fulfillment of re
quirements for the degree of
DOCTOR OF PHILOSOPHY
Dean of Graduate Studies
Date 2000
DISSERTATION COMMITTEE
Chairperson
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
To my greatest heroes...
Mom and Dad
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Acknowledgements
I would like to first thank my advisor Professor G.K.S. Prakash for providing
me with guidance and support throughout my studies at USC. From the moment I
interviewed with him, his love and enthusiasm for chemistry was overwhelming and
inspirational Professor Prakash has a genuine interest in the well being of his
graduate students and is quick to praise hard work and offer encouragement when
research becomes frustrating. In my view Professor Prakash has redefined the
typical student-advisor relationship and I am honored to be his friend, as well as
colleague. I would also like to thank Professor George Olah for allowing me the
privilege o f working in a world-renowned institute where the focus is directed
toward learning and becoming better scientists. I am both grateful and humbled that
Dr. Prakash and Dr. Olah would entrust me with such an interesting, challenging and
important project that has allowed me to interact as a representative of USC with
researchers at the Jet Propulsion Laboratory. I would also like to acknowledge the
support of Professor William Weber whose encouragement and advice are typical of
the faculty at the Loker Hydrocarbon Institute, as well as the friendship and support
provided by Dr. Robert Aniszfeld.
This research project has given me the privilege o f enjoying a close working
relationship with researchers at the Jet Propulsion Laboratory. Therefore I would
like to thank Dr. Marshall Smart for his role as a mentor to me as I began my
graduate career. Dr. Smart provided me with guidance, encouragement and could
always be counted on for an idea helpful to our research. It was his work that
provided the foundation for this research project and I can only hope that I have
contributed as much to it as he has to me. I would also like to especially thank Dr.
S.R. Narayanan who has had a profound influence on my development as a scientist
introducing me to the fields o f electrochemistry and fuel cells while serving as a
second research advisor. Dr. Narayanan has always found time for an
electrochemistry lesson or discussions regarding our research and I will always
remember showing up unannounced at his office with notebook in hand and coming
iii
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
away with a better understanding o f the issue, as well as a sample of his lunch. In
addition, I would also like to thank Mr. Tom Valdez for his friendship, advice and
willingness to share his knowledge o f fuel cells with me. I would also like to
acknowledge Dr. Subbarao Surampudi, Group Supervisor of the Electrochemical
Technologies Group, for his support and acceptance o f me as a genuine member o f
the JPL family. Finally I would like to thank Dr. Robert Nowak at DARPA for
supporting our fuel cell efforts.
Several o f my colleagues at USC have also contributed to this research
project including Dr. Thomas Matthew and Mr. Michael Bradley with whom I first
prepared membrane samples. Mr. Robert Lieberman and Mr. Kyle Sherk have also
been extremely helpful in conducting methanol permeability studies and processing
precursor materials. Recently Dr. Akihisa Saitoh has joined our fuel cell team and
contributed to the processing o f PVDF precursors. I would also like to acknowledge
Ms. Virginie Pleynet, an employee o f E lf Atochem, N.A., who was instrumental in
coordinating our USC team efforts by providing me with necessary analysis data and
hydrolyzing membrane samples.
I would also like to express my gratitude to the faculty of the Department o f
Chemistry and Biochemistry at Ithaca College, Ithaca, NY. I was fortunate to not
only conduct research as an undergraduate but also did so under the supervision o f
Professor Anatol Eberhard, whose early guidance and support contributed greatly to
my scientific development. I would also like to thank Professors Heinz Koch and
Judith Koch for their support and interest in my degree progress.
My last academic thanks are directed to the faculty of West Seneca West
Senior High School, West Seneca, NY. I would like to thank Mr. John Ruh for
teaching me how to study and Mr. Paul Maciejewski for nurturing my early love o f
science.
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
On a more personal note, words do not adequately express my gratitude to
my family and friends. My wife Dana has supported me throughout my academic
career frequently putting her career on hold to further mine. It was her frith and love
that allowed me to move to California and pursue my goals believing that distance
would not drive us apart. Dana accepted the long weekends spent writing and
studying without ever complaining and I share this accomplishment with her, as she
is my best friend and partner. I look forward to moving into the next phase o f my
career with her by my side. I am also at a loss for words in thanking my Mom and
Dad, truly my greatest heroes. I have never met two people more willing to sacrifice
to ensure the well being o f their children. My parents always stressed the importance
of family and education while providing me with the foundation and support with
which to accomplish my goals. I realize now that the best way to repay them is to
make the same sacrifices and commitment some day for my children. I also owe a
great deal of thanks to my sisters Julie and Laura whose love and support is often
taken for granted. I hope that I have provided a good role model for them and I take
great pride in their achievements, as they are the true shining stars o f our family.
I also feel it is necessary to acknowledge the important role my extended
family has played throughout my education, including my Nani and Nano Carmina,
and Grandpa and Grandma Atti. I have been truly fortunate to have my grandparents
celebrate my achievements, as they have been a constant source o f support and
encouragement. I would also like to thank the Molino family for their support and
encouragement from an early stage in my education and treating me as a son, despite
taking their daughter across the country. Also, as is Italian custom, aunts, uncles,
cousins, club and friends are all an important part of my family and their support and
encouragement has been a constant reminder of my roots. I would also like to
express a special thanks to my California family, Denny and Peggy Boultinghouse.
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Finally I would like to acknowledge my friend Jay StruckhofF whose passing
was tragic and far too early. Jay had a bright career ahead o f him and it was he that
first welcomed me into Professor Prakash's research group. Jay's advice and
friendship came at a time when I was in a new city far from friends and family. It
was his council that helped me through the difficulties o f a cross-country relationship
as well as the rigors o f my academic program. There is not a day that goes by that I
do not remember Jay and miss him.
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Table of Contents
Dedication ii
Acknowledgements iii
List o f Figures xii
List of Tables xxi
Chapter 1: Introduction
Polymer Electrolyte Membranes: Applications in Fuel Cells
1.1 Historical Background 1
1.2 Solid Polymer Electrolytes 5
1.3 Direct Methanol Liquid-Feed Fuel Cells 8
1.3.1 Properties o f Perfluorosulfonate Ionomers 12
1.3.2 Water Permeability in Perfluorosulfonate Ionomers 16
1.3.3 Methanol Crossover in Perfluorosulfonate Ionomers 17
1.3.4 Methanol Permeability Mechanism(s) in Perfluorosulfonate 20
Ionomers
1.4 Alternative Polymer Electrolyte Membranes 25
1.4.1 PSSA Grafted Membranes 27
1.4.1.1 Water Management in PSSA Grafted Membranes 29
1.5 Interpenetrating Polymer Networks 34
1.5.1 Factors Affecting Polymer Miscibility in IPN Synthesis 37
1.5.1.1 Interfacial Tension in IPN Synthesis 38
1.5.1.2 Polymer Compatibility in IPN Synthesis 40
1.5.1.3 Polymer Composition in IPN Synthesis 41
1.5.1.4 Crosslinking Density in IPN Synthesis 42
vii
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
1.5.1.5 Temperature and Pressure in IPN Synthesis 44
1.6 PVDF-PSSA IPN Polymer Electrolyte Membranes 46
1.6.1 PVDF Synthesis 47
1.6.2 PVDF Crystallinity 48
1.6.2.1 Factors Influencing PVDF Crystallinity 50
1.6.3 Impact o f Crystallinity in IPN Synthesis 52
1.7 References 55
Chapter 2: Results and Discussion
PVDF-PSSA Polymer Electrolyte Membranes:
Fabrication and Analysis
2.1 PVDF-PSSA IPN Methodology 65
2.1.1 PVDF-PSSA Membrane Preparation 66
2.1.1.1 Kynar 740 Series: Membrane Morphology 67
2.1.1.2 Kynar 460 Series: Membrane Morphology 70
2.1.2 Thermal Analysis 72
2.1.2.1 PVDF Crystallinity 72
2.1.2.2 DSC Analysis: Polymer Interpenetration 81
2.1.2.2.1 PVDF-Polystyrene Interpenetration 82
2.1.2.2.2 PVDF-PSSA Interpenetration 85
2.1.2.3 Thermogravimetric Analysis 89
2.1.2.3.1 Thermogravimetric Analysis: Kynar 460 90
and 740 Precursors
2.1.2.3.2 Thermogravimetric Analysis: PVDF-PS 91
and PVDF-PSSA Membranes
2.1.3 Energy Dispersive X-Ray Analysis (EDAX) 95
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
2.1.4 Water Management: Membrane Characterization 100
2.1.4.1 Membrane Water Content 101
2.1.4.2 Proton Conductivity Measurements 103
2.1.4.3 Methanol Permeability 109
2.1.4.3.1 Methanol Permeability: NMR Analysis 112
2.1.5 Water Management: Membrane Analysis 116
2.2 References 120
Chapter 3: Results and Discussion
PVDF-PSSA Polymer Electrolyte Membranes:
Electrical Performance
3.1 Membrane Electrode Assembly: Fabrication Method I 122
3.1.1 Electrical Performance o f PVDF-PSSA Membranes: 1 " x l" 124
DMFCs
3.2 Membrane Electrode Assembly: Fabrication Method II 127
3.2.1 Electrical Performance o f PVDF-PSSA Membranes: 2"x2" 128
DMFCs
3.3 Membrane Electrode Assembly: Fabrication Method III 139
3.3.1 Electrical Performance o f PVDF-PSSA Membranes: 2"x2" 141
DMFCs
3.3.1.1 Impact of Methano 1 Concentration on Electrical 154
Performance and Crossover Rates
3.3.1.2 Impact of Methanol Crossover Rates on Fuel and 162
Fuel Cell Efficiency
ix
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
3.3.1.3 Parametric Stack and System Studies using 167
PVDF-PSSA ME As
3.3.1.4 Impact o f PSSA Uptake on ME A Properties and 172
Electrical Performance
3.3.1.5 Impact o f Membrane Thickness on MEA Properties 181
and Electrical Performance
3.4 Generation I 5-cell Stack using PVDF-PSSA Membranes 189
3.4.1 Electrical Performance o f PVDF-PSSA Membranes: 192
80 cm2 5-Cell DMFC Stack
3.5 Modified PVDF-PSSA Membrane Development 196
3.5.1 Electrical Performance o f PVDF-PSSA Membranes: 197
2"x2" DMFCs
3.5.1.1 Fuel and Fuel Cell Efficiency o f PVDF-PSSA 204
Membranes: 2"x2" DMFCs
3.6 Membrane Electrode Assembly: Fabrication Method IV 206
3.6.1 Electrical Performance o f PVDF-PSSA MEAs: 208
2"x2” DMFCs
3.7 Generation II 5-Cell Stack Development 216
3.7.1 Electrical Performance o f PVDF-PSSA MEAs: 217
80 cm2 Single Cell DMFC
3.8 Experimental Section 219
3.9 Conclusion 227
4.0 References 231
x
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
5.0 Bibliography
xi
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
List of Figures
Figure 1.1 Scheme o f Energy Conversion in a Fuel CelL
Figure 1.2 Principle Operation o f a H2/O2 Fuel Cell (Reprinted from
Reference 3).
Figure 1.3 Schematic representation of a Direct Methanol Liquid-
Feed Fuel Cell.
Figure 1.4 Structure o f Perfluorosulfonate Ionomer Membranes.
Figure 1.5 Proposed Mechanism o f Nafion® Formation.
Figure 1.6 Schematic representation of the molecular organization
o f a cluster in Nafion® (Reprinted from Reference 24).
Figure 1.7 Possible Arrangement of Nafion® showing an
intermediate ionic phase (Reprinted from Reference 25).
A. Fluorocarbon
B. Interfacial Zone
C. Ionic Clusters
Figure 1.8 Impact o f methanol crossover on the cathode potential in
a direct methanol fuel cell.
Figure 1.9 Structure o f Polystyrene-Sulfonic Acid Grafted onto
PVDF.
Figure 1.10 Common Crystalline Conformations o f PVDF.
Figure 2.1 DSC analysis o f a compression molded Kynar 460
precursor processed at USC.
Figure 2.2 DSC analysis o f the glass transition endotherm for a
compression molded Kynar 460 precursor processed at
USC.
Figure 2.3 DSC analysis o f a compression molded Kynar 460
precursor supplied by E lf Atochem, N. A.
Figure 2.4 DSC analysis o f a thermally extruded Kynar 740
precursor supplied by Westlake Plastics.
1
2
10
12
13
14
15
18
30
50
74
75
76
77
X ll
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.5
Figure 2.6
Figure 2.7
Figure 2.8
Figure 2.9
Figure 2.10
Figure 2.11
Figure 2.12
Figure 2.13
Figure 2.14
Figure 2.15
DSC analysis o f a thermally extruded Kynar 740 78
precursor supplied by Elf Atochem, N.A. The sample
was heat-treated at 200 °C for 15 minutes, cooled and
subjected to DSC analysis.
DSC analysis o f a PVDF-PS membrane sample prepared 83
using a compression molded Kynar 460 precursor
processed at USC.
DSC analysis o f the glass transition endotherm for a 84
PVDF-PS membrane sample prepared using a
compression molded Kynar 460 precursor processed at
USC.
DSC analysis o f a PVDF-PSSA membrane sample 86
prepared using a compression molded Kynar 460
precursor processed at USC.
DSC analysis o f the glass transition endotherm of a 87
PVDF-PSSA membrane sample prepared using a
compression molded Kynar 460 precursor processed at
USC.
TGA thermograph for Kynar 460 precursors processed at 90
USC and supplied by Elf Atochem, N.A. and Kynar 740
precursor samples supplied by Westlake Plastics.
TGA thermograph o f a PVDF-PS membrane sample 92
prepared using a compression molded Kynar 460
precursor processed at USC.
TGA thermograph o f a PVDF-PSSA membrane sample 94
prepared using a compression molded Kynar 460
precursor processed at USC.
EDAX plot o f a PVDF-PSSA membrane sample with 97
localized sulfur distribution along membrane edges.
EDAX plot o f a PVDF-PSSA membrane sample with 98
localized sulfur distribution within the bulk.
EDAX plot of a PVDF-PSSA membrane sample with 99
homogeneous sulfur distribution throughout the bulk.
xiii
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.16
Figure 2.17
Figure 2.18
Figure 2.19
Figure 2.20
Figure 2.21
Figure 2.22
Figure 2.23
Figure 2.24
Membrane water content for various PVDF-PSSA 102
membrane samples plotted vs. their corresponding PSSA
uptake values.
Specific conductivity analysis for a PVDF-PSSA 104
membrane sample using a 4-probe, D.C. apparatus. The
membrane sample was fully hydrated and tested at room
temperature.
Specific proton conductivity values for PVDF-PSSA 106
membrane samples plotted vs. their corresponding PSSA
uptake levels. Membrane samples were folly hydrated
and tested at room temperature.
Specific proton conductivity values for various PVDF- 107
PSSA membrane samples plotted vs. the corresponding
water content level for that specific PSSA uptake value.
Methanol permeability through Nafion®-117 and PVDF- 110
PSSA membrane samples over time. Tests were
conducted at room temperature using a 3.0 M methanol
standard.
Calculated methanol diffusion coefficients for various 111
membrane samples plotted vs. their corresponding
polystyrene weight uptake in grams impregnated per
sample.
Water-methanol partitioning coefficient determination 113
for various PVDF-PSSA membrane samples using NMR
analysis. Membranes were dried and equilibrated in
methanol solutions at room temperature for 24 hours.
Methanol diffusion coefficients for PVDF-PSSA 114
membrane samples as compared to Nafion®-l 17
determined using NMR analysis. Tests were performed
with methanol concentrations ranging from 1 - 5 M.
Water diffusion coefficients for PVDF-PSSA membrane 115
samples as compared to Nafion®-117 determined using
NMR analysis. Tests were performed with methanol
concentrations ranging from 1 - 5 M.
xiv
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.1
Figure 3.2
Figure 3.3
Figure 3.4
Figure 3.5
Figure 3.6
Figure 3.7
Figure 3.8
Figure 3.9
Figure 3.10
Performance o f a l"xl" PVDF-PSSA MEA in a direct
methanol fuel cell with 1.0 M methanol at 83 °C utilizing
pressurized oxygen at the cathode vented at various flow
rates.
Performance o f MEAs 97-30 and 97-33 in direct
methanol fuel cells at 60 °C with 1.0 M methanol
utilizing ambient oxygen at the cathode.
Effect of ambient oxygen flow rates on the performance
o f MEA 97-33 in a direct methanol fuel cell at 60 °C
with 1.0 M methanol.
Effect of pressurized oxygen and vented flow rate on the
performance o f MEA 97-33 in a direct methanol fuel cell
at 60 °C with 1.0 M methanol.
IR corrected anode performance o f MEA 97-33 in a
direct methanol fuel cell as compared to Nafion®-117 at
60 °C with 1.0 M methanol.
IR corrected cathode performance o f MEA 97-33 in a
direct methanol fuel cell at 60 °C with 1.0 M methanol
utilizing ambient and pressurized oxygen.
Effect of ambient air flow rates on the performance of
MEA 97-33 in a direct methanol fuel cell at 60 °C with
1.0 M methanol.
IR corrected cathode performance o f MEA 97-33 in a
direct methanol fuel cell at 60 °C with 1.0 M methanol
utilizing ambient air as compared to ambient oxygen.
Effect of cell conditioning (90 °C, pressurized oxygen)
on the performance of MEA 98-35 in a direct methanol
fuel cell as compared to the pre-conditioned MEA 98-35
and MEA 97-33 at 60 °C with 1.0 M methanol.
Effect of ambient oxygen flow rates on the performance
o f MEA 98-35 in a direct methanol fuel cell at 60 °C
with 1.0 M methanol.
126
130
131
132
133
135
136
138
143
144
xv
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.11
Figure 3.12
Figure 3.13
Figure 3.14
Figure 3.15
Figure 3.16
Figure 3.17
Figure 3.18
Figure 3.19
Figure 3.20
Figure 3.21
Effect o f pressurized oxygen and vented flow rate on the 145
performance of MEA 98-35 in a direct methanol fuel cell
at 60 °C with 1.0 M methanol.
Effect of ambient oxygen flow rates on the performance 146
o f MEA 98-35 in a direct methanol fuel cell at 90 °C
with 1.0 M methanol.
Effect of pressurized oxygen and vented flow rate on the 147
performance of MEA 98-35 in a direct methanol fuel cell
at 90 °C with 1.0 M methanol.
IR corrected anode performance o f MEA 98-35 in a 149
direct methanol fuel cell as compared to MEA 97-33 and
Nafion®-l 17 at 60 °C and 90 °C with 1.0 M methanol.
IR corrected cathode performance o f MEA 98-35 in a 150
direct methanol fuel cell at 60 °C with 1.0 M methanol
utilizing ambient oxygen.
IR corrected cathode performance o f MEA 98-35 in a 151
direct methanol fuel cell at 90 °C with 1.0 M methanol
utilizing ambient and pressurized oxygen.
Effect of ambient air flow rates on the performance of 152
MEA 98-35 in a direct methanol fuel cell at 60 °C with
1.0 M methanol.
IR corrected cathode performance o f MEA 98-35 in a 153
direct methanol fuel cell utilizing ambient air vs. ambient
oxygen flow rates at 60 °C with 1.0 M methanol
Effect of methanol concentration on the performance o f 155
MEA 98-35 in a direct methanol fuel cell at 60 °C
utilizing ambient oxygen at the cathode.
Impact o f methanol concentration on the IR corrected 157
anode performance o f MEA 98-35 in a direct methanol
fuel cell at 60 °C.
Impact o f methanol concentration on the IR corrected 158
cathode performance o f MEA 98-35 in a direct methanol
fuel cell at 60 °C utilizing ambient oxygen.
xvi
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.22
Figure 3.23
Figure 3.24
Figure 3.25
Figure 3.26
Figure 3.27
Figure 3.28
Figure 3.29
Figure 3.30
Figure 3.31
Figure 3.32
Effect of methanol concentration on the performance o f 161
MEA 98-35 in a direct methanol fuel cell at 60 °C
utilizing ambient air at the cathode.
Effect of methanol concentration on the fuel efficiency o f 164
MEA 98-35 in a direct methanol fuel cell at 60 °C.
Effect of methanol concentration on the fuel cell 166
efficiency o f MEA 98-35 in a direct methanol fuel cell at
60 °C utilizing ambient oxygen at the cathode.
Effect of methanol concentration on the fuel cell 167
efficiency o f MEA 98-35 in a direct methanol fuel cell at
60 °C utilizing ambient air at the cathode.
Performance o f MEA 98-35 in a direct methanol fiiel cell 169
as compared to Nafion®-117 at 55 °C with 0.50 M
methanol utilizing ambient air at the cathode.
Fuel efficiency o f MEA 98-35 in a direct methanol fuel 170
cell as compared to Nafion®-117 at 55 °C with 0.50 M
methanol.
Fuel cell efficiency of MEA 98-35 in a direct methanol 171
fuel cell as compared to Nafion®-117 at 55 °C with 0.50
M methanol utilizing ambient air at the cathode.
Effect of PSSA uptake on the measured cell resistivity o f 174
PVDF-PSSA MEAs in direct methanol fuel cells at 60 °C
and 90 °C.
Effect of PSSA uptake on the measured methanol 175
crossover rates of PVDF-PSSA MEAs in direct methanol
fuel cells at 60 °C and 90 °C with 1.0 M methanol.
Effect of PSSA uptake on the performance of PVDF- 176
PSSA MEAs in direct methanol fuel cells at 60 °C with
1.0 M methanol utilizing ambient cathode air.
Effect of PSSA uptake on the IR corrected anode 177
performance o f PVDF-PSSA MEAs in direct methanol
fuel cells at 60 °C with 1.0 M methanol.
xvii
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.33
Figure 3.34
Figure 3.35
Figure 3.36
Figure 3.37
Figure 3.38
Figure 3.39
Figure 3.40
Figure 3.41
Figure 3.42
Effect o f PSSA uptake on the IR corrected cathode 178
performance o f PVDF-PSSA MEAs in direct methanol
fuel cells at 60 °C with 1.0 M methanol utilizing ambient
air.
Effect o f ambient oxygen flow rates on the performance 182
o f MEA 98-72 in a direct methanol fuel cell at 60 °C
with 1.0 M methanol.
Impact o f ambient oxygen flow rates on the performance 183
o f MEA 98-72 in a direct methanol fuel cell as compared
to cell data measured at 60 °C and MEA 98-35 measured
at 90 °C, with 1.0 M methanol.
IR corrected anode performance of MEA 98-72 in a 184
direct methanol fuel cell as compared to MEA 98-35 at
60 °C and 90 °C with 1.0 M methanol.
IR corrected cathode performance o f MEA 98-72 in a 185
direct methanol fuel cell as compared to MEA 98-35 at
60 °C and 90 °C with 1.0 M methanol utilizing ambient
oxygen.
Effect o f ambient air flow rates on the performance o f 186
MEA 98-72 in a direct methanol fuel cell at 60 °C with
1.0 M methanol.
IR corrected cathode performance o f MEA 98-72 in a 188
direct methanol fuel cell as compared to MEA 98-35 at
60 °C with 1.0 M methanol utilizing ambient air.
Fuel cell performance o f cells 1-5 in a direct methanol 5- 194
cell stack as compared to MEA 98-35 at 55 °C with 0.50
M methanol utilizing ambient air at the cathode.
IR corrected anode performance of cells 1-4 in a direct 195
methanol 5-cell stack as compared to MEA 98-35 at 55
°C with 0.50 M methanol.
Effect o f ambient oxygen flow rates on the performance 199
o f MEA 98-95 in a direct methanol fuel cell at 60 °C
with 1.0 M methanol.
xviii
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.43
Figure 3.44
Figure 3.45
Figure 3.46
Figure 3.47
Figure 3.48
Figure 3.49
Figure 3.50
Figure 3.51
Figure 3.52
IR corrected anode performance o f MEA 98-95 in a 200
direct methanol fuel cell as compared to MEA 98-35 at
60 °C with 1.0 M methanol.
IR corrected cathode performance o f MEA 98-95 in a 201
direct methanol fuel cell as compared to MEAs 98-35
and 98-40 at 60 °C with 1.0 M methanol utilizing
ambient oxygen.
Performance o f MEA 98-95 in a direct methanol fuel cell 202
as compared to MEA 98-35 at 60 °C with 1.0 M
methanol utilizing ambient air at the cathode.
IR corrected cathode performance o f MEA 98-95 in a 203
direct methanol fuel cell as compared to MEA 98-35 at
60 °C with 1.0 M methanol utilizing ambient air.
Performance o f MEA 98-95 in a direct methanol fuel cell 204
as compared to MEA 98-35 and Nafion®-117 at 55 °C
with 0.50 M methanol utilizing ambient air at the
cathode.
Fuel efficiency o f MEA 98-95 in a direct methanol fixel 205
cell as compared to MEA 98-35 and Nafion®-117 at 55
°C with 0.50 M methanol.
Fuel cell efficiency o f MEA 98-95 in a direct methanol 206
fuel cell as compared to MEA 98-35 and Nafion®-117 at
55 °C with 0.50 M methanol utilizing ambient air at the
cathode.
Performance o f MEA 98-95B prepared using a non- 209
bonded MEA fabrication technique in a direct methanol
fuel cell as compared to MEA 98-95 at 60 °C with 1.00
M methanol utilizing ambient oxygen at the cathode.
IR corrected anode performance o f MEA 98-95B in a 210
direct methanol fuel cell as compared to MEA 98-95 at
60 °C with 1.0 M methanol.
IR corrected cathode performance o f MEA 98-95B in a 211
direct methanol fuel cell as compared to MEA 98-95 at
60 °C with 1.00 M methanol utilizing ambient oxygen.
xix
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.53
Figure 3.54
Figure 3.55
Figure 3.56
Figure 3.57
Figure 3.58
Figure 3.59
Figure 3.60
Performance o f MEA 98-95B prepared using a non
bonded MEA fabrication technique in a direct methanol
fuel cell as compared to MEA 98-95 at 60 °C with 0.50
M methanol utilizing ambient air at the cathode.
IR corrected anode performance o f MEA 98-95B in a
direct methanol fuel cell as compared to MEA 98-95 at
60 °C with 0.50 M methanol.
IR corrected cathode performance o f MEA 98-95B in a
direct methanol fuel cell as compared to MEA 98-95 at
60 °C with 0.50 M utilizing ambient air.
Performance o f MEAs Giner DE-6, 10 and 12 in a direct
methanol fuel cell at 60 °C with 0.50 M methanol
utilizing 0.10 L/min ambient air.
Performance o f MEA 99-18 (80 cm2) in a direct
methanol fuel cell as compared to MEAs 98-95 and 98-
95B at 60 °C with 0.50 M methanol utilizing ambient air
at the cathode.
Sulfonated polystyrene product from the 30% v/v
chlorosulfonic acid/chloroform bath.
Schematic of the apparatus used in the "four probe"
technique to measure the proton conductivity of
membrane samples (Reprinted from Reference 173).
Diagram of the apparatus used to evaluate methanol
permeability (Reprinted from Reference 75).
212
213
214
216
218
221
223
224
xx
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
List of Tables
Table 1.1
Table 1.2
Table 1.3
Table 2.1
Table 2.2
Table 2.3
Table 2.4
Table 2.5
Table 3.1
Table 3.2
Common Fuel Cell Types. 3
Properties o f Ion-Exchange Membranes (Reprinted from 7
Reference 3).
Direct Methanol Fuel Cells: Power Applications. 9
Membrane morphology resulting from the use o f 68
thermally extruded Kynar 740 precursors, supplied by
Westlake Plastics, impregnated using various
polymerization conditions. Membrane samples were
heated to the specific temperature and held for one hour
under the corresponding pressure.
Membrane morphology resulting from the use of 71
compression molded Kynar 460 precursors, processed at
USC and impregnated using various polymerization
conditions. Membrane samples were heated to each
specific temperature and held for one hour under the
corresponding pressure.
PVDF properties of the various Kynar grades tested and 79
their resulting degree of crystallinity as determined by
DSC analysis.
Specific proton conductivity values for PVDF-PSSA 105
membrane samples acquired from the 4-probe, D.C.
apparatus. Membrane samples were fully hydrated and
tested at room temperature.
Resistivity values for PVDF-PSSA membrane samples as 108
compared to Nafion-117 at 25 °C. Resistivity values
were determined using a 2-probe technique and
converted to conductivity values according to the
equation described in section 3.8.
PSSA uptake and related properties o f membranes 97-30 128
and 97-33 as compared to the same properties in
Nafion®-117.
PSSA uptake and related properties o f membrane 98-35 140
as compared to the same properties in Nafion®-117.
xxi
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Table 3.3
Table 3.4
Table 3.5
Table 3.6
Table 3.7
Table 3.8
Table 3.9
Table 3.10
Table 3.11
Table 3.12
Effect o f methanol concentration on the measured 159
methanol crossover rates o f MEA 98-35 in a direct
methanol fuel cell as compared to Nafion®-l 17 at 60 °C.
Effect o f methanol concentration on the measured 162
methanol crossover rates o f MEA 98-35 in a direct
methanol fuel cell at 60 °C as compared to Nafion®-117.
PSS A uptake and related properties o f membranes 98-36, 173
98-40 and 98-47A as compared to the same properties in
membrane 98-35 and Nafion®-117.
PSSA uptake and related properties o f membrane 98-72 180
as compared to the same properties in membrane 98-35
and Nafion®-117.
PSSA uptake and related properties o f potential stack 190
membranes as compared to the same properties in
membrane 98-35 and Nafion®-117.
Overall physical properties o f the potential stack 191
membranes as compared to membrane 98-35 and
Nafion®-117.
Measured resistivity values for cells 1-5 in a direct 193
methanol 5-cell stack as compared to MEA 98-35 at 60
°C.
PSSA uptake and related properties o f membrane 98-94, 197
98-95 and 98-96 as compared to the same properties of
membrane 98-35 and Nafion®-117.
PSSA uptake and related properties for potential stack 217
membranes as compared to the same properties in
membrane 98-95 and Nafion®-117.
Impact of AIBN initiator concentration on the 220
polystyrene uptake value(s) incorporated during
precursor impregnation.
xxii
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Chapter 1: Introduction
Polymer Electrolyte Membranes: Applications in Fuel Cells
1.1 Historical Background
The fuel cell concept actually dates back over 150 years to original work
done by Sir William Grove who found that oxidation o f hydrogen derived from
water electrolysis could generate low voltage direct current electricity.1 This early
precursor to the modem H2/O2 fuel cell used platinized platinum electrodes
immersed in sulfuric acid. Generally though, fuel cells are galvanostatic devices that
have their reactants continuously supplied from an external source. The fuel cell
converts the free chemical energy o f a fuel directly into electrical energy, as shown
in figure 1.1, and has the major advantage of being not limited by Carnot efficiencies
thus offering practical efficiency values of at least twice that o f the typical internal
combustion engine.2 This feature has been one of the driving forces behind fuel cells
as alternative power sources.
Figure 1.1 Scheme of Energy Conversion in a Fuel Cell.
chem ical ^eat -------► mechanical------- ^ electrical
energy energy energy
FUEL CELL
1
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Fuel cells consist o f a fuel electrode (anode) and an oxidant electrode
(cathode) separated by an ion-conducting electrolyte. Electrons generated by fuel
oxidation are harnessed to power an external load then returned to the cathode
completing the circuit. The electrolyte is necessary to transport ionic current
between electrode compartments. The operating principle o f a H2/O2 fuel cell was
demonstrated by Grove and is illustrated in figure 1.2.3
Figure 1.2 Principle Operation of a H2/O2 Fuel Cell (Reprinted from Reference
3)
Load
j K
F
Oxidant
/ He id
Porous ElectrolgtB
Anode
\
P o m s
Cathode
[hrera L I Cell
Anode: H, — ► 2H* + 2 « “
Cathode: £©2 2 H % 2 6 " H20
2
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
In the case o f pure H2 as the fuel, the fuel cell operates with zero emissions, a
fact that further enhances the environmental appeal. There are also a variety o f other
fuels available for fuel cell applications including N2H4, NH3, CH3OH, coal gas and
hydrocarbons. However, during the early development o f fuel cells up to the 1960’s,
it was determined that H2 was the only practical fuel since the listed alternative fuels
had insufficient reactivity. Air or pure oxygen is used typically as the oxidant in
hydrogen based fuel cells. The nature o f the electrolyte and temperature o f operation
are used to categorize fuel cells. The common electrolytes are either alkaline or
acidic and operate at low temperature (ambient — 100 °C) or high temperature (100 —
200 °C). The important fuel cell types are shown in table 1.1.3 ,4 The type o f fuel and
electrolyte used in each category typically sets restrictions on operating conditions as
well as practicality.
Table 1.1 Common Fuel Cell Types.
Oxidant
Polymer Eleetrolyte FC Air
Phosphorous Acid FC
Hydrogen
Direct Methanol
Fuel Cell
Methanol
Alkaline FC Hydrogen
Hydrocarbons
Molten Carbonate FC Air/COJ
Hydrocarbons 900 °C
3
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The electrolyte o f any fuel cell must meet a certain set o f requirements to
ensure adequate performance. The electrolyte serves as the ionic bridge that
separates the anode and cathode compartments; therefore the electrolyte must exhibit
high ionic conductivity to maintain minimal ohmic resistivity. The electrolyte must
be chemically, mechanically and thermally stable under operating conditions. The
electrolyte must also be electrochemically stable under highly reducing (anodic) and
oxidative (cathodic) atmospheres.
Alkaline fuel cells pursued under the U.S. Aerospace program were based on
high purity hydrogen and oxygen and used potassium hydroxide as the electrolyte.3
Fuel purity was necessary to avoid poisoning the anode and/or carbonization of the
aqueous KOH electrolyte by dissolved CO2 at the cathode. These requirements
limited fuel cell practicality for commercial or residential purposes and highlighted
the need for fuel cells containing components more tolerant o f diverse operating
conditions.
There has been substantial progress since the 1960’s in fuel cells based on
acid electrolytes and on fuels that can be processed into hydrogen. U.S. dependence
on foreign oil and environmental air pollution problems has driven this interest and
continued research has led to improvements in the electrolyte, fuels and catalysts
used in fuel cell processes. Early research in H2/O 2 fuel cells used concentrated
H2SO4 and H3PO4 acid electrolytes that resulted in promising electrical performance
despite the restrictive operating conditions.3 Acid electrolytes required cell hardware
that could withstand the corrosive nature o f the electrolyte. A key improvement in
4
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
this area was the development o f a solid polymer electrolyte as the ion-conducting
matrix.
1.2 Solid Polymer Electrolytes
Solid polymer electrolytes are polymer networks with attached functional
groups capable of exchanging cations or anions and are classified as ionomers or
ionomeric membranes. 5,6 Ionomers consist of a crosslinked organic polymer such as
polystyrene, polyethylene, polyvinyldifluoride (PVDF), tetrafluoroethylene (PTFE)
or PTFE analogs. Solid polymer electrolytes used in fuel cells are the cation-
exchange membranes that conduct hydronium ions using sulfonic acid groups bound
within the organic polymer matrix. 3 These materials are referred to as proton-
exchange membranes or polymer electrolyte membranes (PEM). The abbreviation
‘PEM’ used in this thesis can be used interchangeably.
The advantages o f using proton-exchange membranes are numerous. The
polymeric electrolyte serves as the electrode support(s) and is a cell structural
component. General Electric first used polymer electrolyte membranes in fuel cells
in the late I950’s during its program to develop membrane cells and electrolyzers. 7' 9
The earliest such membranes developed by GE were made by condensation o f
phenolsulfonic acid and formaldehyde. These membranes were brittle, cracked
when dried and rapidly hydrolyzed to sulfuric acid. Shortly thereafter Grubb
reported the use of Amberplex C-l resins for fuel cell applications. 10,11 These
membranes were based on partially sulfonated polystyrene held in an inert matrix
and had undesirable physical properties and poor lifetime stability o f only 2 0 0 hours
5
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
at 60 °C. Polymer degradation at or near the anode became evident using these early
hydrocarbon based ion-exchange resins and was attributed to radically induced
oxidative degradation at the weak alpha C-H bond o f polystyrene. GE then
developed the “D” series o f membranes made by crosslinking styrene-divinyl
benzene into an inert fluorocarbon matrix followed by partial sulfonation. 12 These
membranes had acceptable physical properties and were used in seven Gemini space
missions starting in 1964. Lifetime performance was still poor due to polymer
degradation but was slightly improved to 1000 hours at 60 °C as a result o f
optimizing the crosslinking density o f the membrane.
GE continued to develop ion-exchange resins and addressed the stability
issue with its “S” series of membranes prepared from homopolymers of
trifluorostyrene-sulfonic acid. 1 3 1 4 It was determined that these membranes had much
higher chemical stability when compared with the hydrocarbon based polystyrene-
sulfonic acid analogs tested years earlier. The improved stability was attributed to
the strength of the a C-F bond and its resistance to cleavage. However these
membranes had poor physical properties due to their high degree o f swelling.
Physical stability was improved by blending the ho mo polymer with PVDF using
triethyl phosphate as a plasticizer and resulted in lifetime performance o f 2 0 0 0 hours
at 80 °C. Grafting trifluorostyrene-sulfonic acid into an inert fluorocarbon matrix
further improved lifetime to about 3000 hours and was the foundation for future
work into grafted membrane systems.
6
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Research into novel ionomeric membranes was further enhanced by Dupont’s
development o f the Nafion® series in the late 1960’s. At the time these membranes
exhibited favorable properties as compared to the existing technologies, as shown in
table 1.2. 3 General Electric incorporated Nafion® membranes into fiiel cell
applications such as the Biosatellite missions in 1966, 1968 and used them
exclusively after 1969. Since then, Nafion® membranes have undergone continuous
improvement and achieved lifetime performance of greater than 40,000 hours at 80
°C.
Table 1.2 Properties of Ion-Exchange Membranes (Reprinted from Reference
3).
Properties of Ion Exchange Membranes
Property Desired
Value
Phenol
Sulfonic
Polystyrene
Sulfonic
Poly(trifluort>- Pcr-fluorinatcd
styrene sulfonic) Sulfonic (Nation)
Ionic Resistivity st 25 °C <20 100 < 20 < 20 <20
Ionic Resistance at 25 °C <0.7 10 <0.7 <0.7 <0.7
Hydrogen Permeability
at 25 "C
(cm1 m‘ h* arm 1 )
<21.5 < 21.5 <21.5 < 21.5 <21.5
Mullin Burst Strength
at 25 "C latm
> 3 4 10.2 2.0 - 2.7 5.1 -6.8
Tensile Strength at Break
at 25 "C (kg/cm*)
> 140 210 63 140-210
°o Elongation at 25 C > 100 • 100 40 » 100
Maximum Temperature of
Thermal-Hydrolytic Stability C C)
> 100 50 > 100 > 150 > 150
% Oxidative Degradation
After 18 b Exposure at 80 °C to 30% H ;0; *
20 ppm Fe2*
< I 100 100 < 1 < 1
Crease-Crack Resistance
(ability to withstand 5 folds
after drying at 100 °C)
pass fail pass fail pass
7
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Despite the successes o f ionomeric membranes such as Nafion® in H2/O2 fuel
cells, many problems still exist. Reforming o f hydrocarbon fuels to H2 gas brings up
several restrictive issues to fuel cell performance. Noble metals such as Pt have long
been used as catalysts for the electro-oxidation o f various fuels. 3 These catalysts are
very sensitive below 120 °C to ultra-low levels (< 100 PPM) of organic species such
as CO that are generated from incomplete fuel processing. These species bind
strongly to the catalyst and poison the surface. In-situ reforming o f fuel reduces
start-up time and also adds weight penalties, complexity and cost to the overall
system.
The use and storage of hydrogen gas also presents many logistical issues that
prove extremely problematic. Currently there lacks an existing infrastructure for
adequate H2 distribution and efforts to establish such a system would be expensive
and unrealized for years. Hydrogen storage requires gas compression to improve the
weight to volume ratio o f the gas in the storage cylinders thus requiring stringent
safety measures to prevent explosion. Recently the use o f metal hydrides for
hydrogen storage has been investigated but is currently weight inefficient. During
operation hydrogen fuel cells also require gas humidification to maintain adequate
proton conductivity within the polymer electrolyte membrane. This further increases
system complexity requirements in order to maintain adequate membrane hydration.
1.3 Direct Methanol Liquid-Feed Fuel Cells
Direct methanol fuel cells are attractive and promising alternatives to
hydrogen fueled polymer electrolyte systems. 15 The main benefit of the DMFC
8
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
concept is the liquid-feed delivery o f an aqueous solution of methanol as fuel. This
has tremendous logistical advantages over hydrogen and certain systems have
exhibited favorable electrical performance at temperatures ranging from ambient —
90 °C utilizing H2SO4 as the electrolyte. 3 The transition to a solid polymer
electrolyte as the conducting matrix has widened the scope o f potential power
applications, as shown in table 1.3.
Table 1.3 Direct Methanol Fuel Cells: Power Applications.
Applications for Direct Methanol Liquid-Feed Fuel Cells
Battery Replacement
Transportation
Light Duty Vehicles
Emergency Power
Consumer Tools
Marine Applications
Communications
Mobile Power
The current state o f art DMFC is based on metallic Pt-Ru catalyst at the
anode, metallic Pt at the cathode and a polymer electrolyte membrane separating the
electrode compartments. Catalyst materials are applied to porous gas diffusion
electrodes and hot pressed to the polymer electrolyte forming the membrane
electrode assembly (MEA). The schematic representation o f the direct methanol fuel
cell concept can be seen in figure 1.3.
9
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 1.3 Schematic representation of a Direct Methanol Liquid-Feed Fuel Cell.
Load
H:0
+ CO,
CH3 OH
+ H ,0
[ = £ > H 2 0
1 1 A ir
Pt-Ru
Catalyst Layer
Catalvst Layer
Porous Gas
Diffusion
Electrode
Polymer
Electrolyte
Membrane
Anodic Oxidation of Methanol
Anode: CH3OH + h 2 o
-------- ►
C 0 2 + 6 H+ + 6 e-
Cathode: 3/2 0 2 + 6 H+ + 6 e-
-------►
3 H2 0
Overall: CH3OH + 3/2 0 2 + H2 0
-------- ►
CO, + 3 H2 0
10
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
USC, JPL and Giner, Inc. demonstrated the direct methanol liquid-feed
concept in the early 1990’s and it was shown to have several advantages over state o f
art H2/O2 fuel cells. 15' 18 DMFC systems eliminate the need for gas humidification
and can operate at low temperatures. They also have fewer safety concerns since
handling and storage o f aqueous methanol is benign as compared to compressed
hydrogen. The viability o f this technology has since been independently verified by
a number o f other research teams. 19,20
Perfluorosulfonate ionomers such as Nafion®-117 have proven to be the
membrane o f choice for state of art DMFC systems. Structurally, Nafion® consists
o f a backbone o f tetrafluoroethylene with pendant side chains of perfluorinated vinyl
ether that terminate in sulfonic acid end groups, as shown in figure 1.4. 21 Nafion®
membranes have high proton conductivity and are attractive due to their ability to
attain high power densities under certain operating conditions. However Nafion® is
expensive ($900/m2) and its performance is characterized by high water and
methanol permeability rates from the anode compartment through the membrane to
the cathode. The "methanol crossover" results in a mixed cathode potential since Pt
will oxidize the fuel resulting in overall lowered cell voltages and operating
efficiencies. This phenomenon has often been referred to as a chemical “short-
circuit” o f the cell. The present study aims at developing a new membrane to
mitigate this problem.
11
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 1.4 Structure o f Perfluorosulfonate Ionomer Membranes.
Nafion
[(CF2CF2 )s----- (CF2CF)]x
o — c f 2c f c f 3
I
o — c f 2c f 2s o 3 h
Dow
[(C f 2c f 2 )„----- (CF2CF)]x
o — c f 2c f 2
I
so3 h
1.3.1 Properties o f Perfluorosulfonate Ionomers
Perfluorosulfonate ionomers such as Nafion® are synthesized according to
the scheme shown in figure 1.5. 22 Tetrafluoroethyene (1) is first reacted with SO3 (2)
forming a cyclic sultone (3). The sultone rearranges into a sulfonyl fluoride (4)
derivative that is then reacted with hexafluoropropylene oxide (5) and sodium
carbonate forming the corresponding sulfonyl fluoride vinyl ether (6 ). The vinyl
ether is then co-polymerized with tetrafluoroethylene (1) monomer to produce a XR
resin (7). This thermoplastic resin can then be processed into various forms,
including sheets and tubes using standard industrial fabrication techniques. Once
fabricated, the sulfonyl fluoride derivative is converted into the sulfonic acid,
Nafion-H® form. The polymer now has distinct two-phase character based on the
hydrophobic teflon-like backbone and its hydrophilic pendent side chain(s). It is
12
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
unlikely that Nafion® is a copolymer with blocks o f homopolymerized vinyl ether of
any significant size since the vinyl ether does not homo polymerize. The reactivity
ratio o f TFE (1) is also one hundred times greater than that of the vinyl ether
meaning a regular co-polymer can only be formed by means of a constant feed
polymerization with relatively low concentrations of TFE.
Figure 1.5 Proposed Mechanism o f Nafion® Formation.
F1 C =C F->
(i)
S O 3
(2)
F-)C C F ->
'I I ‘
O-----SOj
(3)
F: F
VN/VX (C Fi C C V N /N Z ‘
I
(7) 0(CF;>CF— 0)mCF;CF2 S0;X
CF3
X -F .O T C , OH
Copolymerize
F-.C^=CF->
( 1)
FSOiCFiCOF
(4)
Na,COi
Heat
, 0
/ \ (5)
c CF;
^C F,
F:C = C (OCF;CF)-OCF;CF:SO: F
C F 3 (6)
Although Nafion® is not covalently crosslinked the distinct bi-phase nature
results in a highly ordered structure. The sulfonic acid groups are far more
hydrophilic than the perfluorinated backbone and molecules arrange to maximize
interactions o f similar fragments. Eisenberg and Yeager first suggested organization
o f the ionizable sulfonic acid groups into clusters resulting in the formation of water-
containing pockets bound within the fluorocarbon backbone. 23 Since then there have
been many models proposed to describe the structural organization of Nafion®
13
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
and perfluorosulfonate ionomers like it. The first model illustrated in figure 1.6
suggests that all sulfonic acid groups are in an aqueous phase, consisting of well-
defined clusters and channels. 2 4 A second model proposed by Yeager et al.
postulates that there is a region intermediate between the perfluoro-backbone and
aqueous phase, as shown in figure 1.7. 25
Figure 1.6 Schematic representation of the molecular organization of a cluster in
Nafion® (Reprinted from Reference 24).
so:
so.
14
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 1.7 Possible Arrangement o f Nafion® showing an intermediate ionic
phase (Reprinted from Reference 25).
A) Fluorocarbon
B) Interfacial Zone
C) Ionic Clusters
TEM and SAXS analysis has been used to determine the cluster model most
likely associated with Nafion®. Ceynowa used TEM to determine the presence of
spherical ionic clusters 30-60 A in diameter. 26 The clusters are evenly dispersed and
there was no evidence o f regions with high or low concentration o f clusters. X-ray
experiments conducted by Gierke et al. also supported the presence of spherical
clusters with a pore breadth o f -50 A.2 4 Further studies o f the cluster size in Nafion®
indicate that it is highly dependent on membrane water content. Parthasarathy et al.
determined that the volume o f clusters increases as the membrane swells in
15
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
water. 2 7 The swelling properties o f Nafion® and other ionomeric membranes in
various solvents have been studied. 28 It was generally concluded that the swelling of
the membranes was strongly related to the type o f ion exchange group present and
not influenced by the perfluoro-carbon backbone. Therefore the grain-boundary
resistance o f the membrane is effectively lowered as small clusters swell, bear upon
other clusters forming large spherical domains. Though the actual structure o f
Nafion clusters remains an open question, the material is generally regarded as
microporous to both water and methanol.
1.3.2 Water Permeability in Perfluorosulfonate Ionomers
Water management is a major factor in the analysis o f direct methanol fuel
cells since adequate membrane hydration is critical to establish and maintain good
electrical performance. 2 9 ' 32 If the membrane is dry, conductivity falls resulting in
reduced cell performance through ohmic losses and poor utilization of the catalyst.
Zawodzinski et al. has reported proton conductivity for Nafion®-117 at various
degrees o f membrane hydration and determined that conductivity decreases roughly
linearly with decreasing water content. 33 Therefore membrane water content must be
raised above a certain threshold level o f 6-7 water molecules per S 03 ' for proton
transport to occur in Nafion® .34' 36 Falk has also used IR analysis to show that the
state o f water in Nafion of different equivalent weights (1400, 1200, 1100) differs
from that o f bulk water upon membrane hydration. 37,38 The water within Nafion®
becomes more liquid-like at EW = 1100 but still differs from the bulk state. 39
However the similarities to liquid-like water at lower equivalent weights has been
16
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
used to explain the corresponding increase in proton conductivity and methanol
permeability.
Water content and diffusion also has a direct impact on cathode performance
by permeating through Nafion® in two main ways: electro-osmotic drag o f water by
protons transported from anode to cathode and diffusion down concentration
gradients that build up. 33 Uribe et al. determined that higher water content within the
catalyst layer benefits electrode kinetics. 4 0 However, high water permeability rates
can create a flooding effect at the cathode whereby water build up can block oxygen
access to catalytic sites requiring higher flow rates to effectively remove water and
pressurized conditions to optimize electrode kinetics. Valdez et al. have recently
modeled the impact of high stoichiometric flow rates on an operating system and
determined that it increases ancillary power losses and inhibits the ability of the
system to maintain water balance. 41
1.3.3 Methanol Crossover in Perfluorosulfonate Ionomers
High methanol crossover rates have a negative impact on performance in
direct methanol liquid-feed fuel cells. During fuel cell operation the cathode
maintains two electrochemical processes, namely the reduction o f oxygen to water
and the oxidation of methanol to carbon dioxide. The cathodic current generated
from oxygen reduction that contributes to current/electricity generation is ic.
However, the process of methanol oxidation demands the flow o f additional current
in short circuited local galvanic cells. The rate o f methanol oxidation is equal to the
rate of oxygen reduction for this parasitic process. As a result the total observed
17
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
cathodic polarization [ic(Ec)] is due to the sum o f the current generating cathodic
process ic, and the parasitic process o f methanol oxidation, i«, and is illustrated in
figure 1.8. It has also been suggested that there is an additional poisoning o f the
cathode catalyst (Pt) as a result of methanol crossover further aggravating the mixed
potential by inhibiting the ability to effectively reduce oxgyen. 42
Figure 1.8 Impact o f methanol crossover on the cathode potential in a direct
methanol fuel cell.
I(EC ) = ic(Ec) + ia
Narayanan and coworkers first measured methanol crossover rates of
Nafion®-117 by measuring the CO2 content in the cathode exit stream using an in
line CO2 analyzer. 4 3 Narayanan et al. measured crossover rates at OCV and under an
applied current and observed that crossover rates decrease with increasing current
density and increase with temperature. The trend o f decreasing crossover with an
increase in applied current can be explained by the decrease of interfacial
concentrations at the anode at high current densities. Narayanan et al. reported the
methanol crossover rates as a parasitic crossover current density that at the time
corresponded to 200 mA/cm2 and 330 mA/cm2 at 60 and 90 °C, at an applied current
density o f 150 mA/cm2 with 1.0 M methanol. Researchers at the Los Alamos
National Laboratory later identified these crossover rates as well. 44 Since then
methanol crossover rates have been reduced considerably due to improvements in
electrode structure at the anode. Narayanan recently reported crossover rates of
18
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
90 mA/cm2 and 200 mA/cm2 at 60 and 90 °C, at the same concentration and applied
current density o f 150 mA/cm2.4 5 a,b
Several other research groups have since developed other techniques used to
study and measure methanol crossover rates in Nafion®-117. Wang et al. established
an in-line mass spectroscopic analysis o f the cathode exit stream to monitor COz
mass signals and determined that the methanol migrating over to the cathode was
oxidized almost quantitatively to COz with Pt acting as a heterogeneous catalyst. 4 2
Ravikumar et al. analyzed the relationship o f methanol concentration and its impact
on the measured cathode potential o f an operating cell. 4 6 It was determined that an
increase in concentration may benefit the anode, especially at high current densities
where mass transport becomes an issue. However there was a loss o f almost 50 — 80
mV at the cathode electrode due to increased methanol crossover.
Methanol crossover and the impact on cell performance were also studied
using perfluorosulfonate ionomers o f varied thickness. Narayanan et al. determined
that membrane thickness has a significant impact on crossover, corresponding to a
reduction of 40-50% from 5 mils to 14 mils in Nafion® .4 7 However there is a trade
off to using thicker ionomer membranes. Increased thickness has a negative impact
on the internal resistance o f the membrane causing depolarization from iR losses at
medium to high current densities. In similar studies conducted by Ren et al. they
also reported iR losses attributed to increased thickness in the analysis o f cell
performance with Nafion® 112 as compared to Nafion® 117.4 8
19
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Narayanan et al. also studied the relationship o f methanol permeability with
equivalent weight and its influence on cell performance. 49 Equivalent weight is a
measure of the amount o f polymeric material in grams, per molar amount of protons
within the membrane. Perfluorosulfonate ionomers o f high equivalent weight had
reduced methanol crossover and resulted in better performance at low current
densities where impact on the cathode is usually severe. Low equivalent weight
membranes performed better at medium to high current densities where the higher
rate o f proton transport reduced ohmic losses. However such improvement in
performance was neutralized by increased methanol crossover.
1.3.4 Methanol Permeability Mechanism(s) in Perfluorosulfonate Ionomers
Many groups have attempted to investigate the distribution of methanol
between the water and membrane phases in perfluorosulphonate ionomer
membranes. Verbrugge and co-workers used a radioactive tracer technique to
measure methanol diffusivity in Nafion® equilibrated with sulfuric acid at room
temperature and obtained a methanol-water partition coefficient o f 0 . 8 from these
measurements. 50 Ren et al. also investigated the methanol uptake o f Nafion®-117
equilibrated with a wide methanol concentration range of 1-8 M . 51 Proton NMR
experiments resulted in a partition coefficient o f 1 indicating Nafion shows no
preference for water or methanol sorption. They also found that Nafion® continues
to absorb methanol into the membrane in addition to the water already present. Skou
et al. has also used proton NMR experiments to determine the methanol-water ratio
in Nafion® 117.5 2 The mole fraction o f methanol in the membrane was plotted as
20
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
a function of the mole fraction o f methanol in the equilibrating liquid, revealing a
straight line. The slope o f this line gave a partition coefficient close to 1 indicating
that saturated Nafion® membranes show no preference for either methanol or water,
further supporting the earlier work o f Verbrugge50 and Ren51. However Nandon et
al. investigated the methanol uptake in Nafion® membranes equilibrated with pure
and 50/50 methanol-water mixtures and determined that Nafion® had a preferential
uptake of water, contrasting the earlier work mentioned. 53 They also found an
increased solvent uptake from the mixed solvent as compared to the uptake from
pure solvents. These discrepancies concerning solvent fractionation properties have
been attributed to the pretreatment o f the membranes.
The theoretical basis for the distribution and permeability o f methanol
through perfluorosulfonic acid ionomers is helpful to understand the mechanism(s)
o f transport. 52 The equilibrium conditions for a compound distributed between two
phases (a liquid phase I and a polymer phase p) is that the chemical potential (p,) of
the compound is identical in the two phases:
p°u + RT In aU I = p°i,p +RT In aiiP
In this expression, p° is the standard chemical potential and a is the activity o f the
component. The expression can then be rearranged to give the partition coefficient
Kf: 52
21
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
For a liquid mixture it is customary to use the mole fraction o f the component as a
measure o f its activity. This was applied to the various methanol-water mixtures
where gamma is an activity coefficient taking care o f deviations from ideality.
Xi.p
K; * gamma = — —
X -i.i
Therefore the mole fraction o f methanol in the polymer can be calculated as
the methanol mole fraction in the methanol-water mixture filling the ionic pores o f
Nafion®. Interactions between the water-methano 1 mixture and the polymer are not
calculated and are expressed in terms o f the activity coefficient gamma. The
partition coefficient o f a component can thus be determined from a plot o f the mole
fractions in the liquid phase and in the polymer phase
Studies o f the methanol-water partition coefficient in Nafion®-117 confirms
that the membrane does not discriminate in regards to methanol or water sorption.
Therefore the miscibilhy o f methanol in water suggests molecular diffusion for both
species should be related. Zawodzinski et al. studied and calculated the self-
diffusion coefficient o f water in Nafion®-117.3 3 Their results indicate that the self
diffusion coefficient o f membrane water (6.00 E"0 6 cm2/s) was roughly four times
less than that of the self-diffusion coefficient o f water in water (2.13 E" 0 5 cm2 /s at 25
22
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
°C). This was expected due to the increase in tortuosity o f the diffusion path within
the membrane since sidechain intrusions into the ion-lined aqueous pores, coupled
with the hydrophobic backbone, restricts the diffusive motion of water.
Verbrugge then calculated the self-diffusion coefficient for methanol in
Nafion®-117 and its relationship to membrane tortuosity. 5 4 His results indicate a
self-diffusion coefficient of 1.15 E-05 cm2/s, a value slightly lower than that o f
methanol in a dilute methanol/water solution standard (1.6 E"05 cm2/s). This high
diffusion rate suggests that sorbed methanol readily diffuses through the aqueous
phase of the polymer. Therefore Verbrugge determined methanol diffusion followed
Fick’s First Law and that membrane tortuosity explained the deviations in the
methanol diffusion coefficient within the membrane and solution standard. This data
suggests that methanol readily diffuses through the aqueous phase of the membrane
but not necessarily through the membrane itself.
Experiments using ionomers of varied equivalent weight were also tested in
hopes of further elucidating transport properties. As previously mentioned in section
1.3.3, Narayanan et al. correlated lower methanol permeability with higher
equivalent weight samples of Nafion®. This is attributed to the reduced number o f
sulfonic acid moieties in these samples limiting the aggregation of ionic side chains
into aqueous channels thus reducing water permeability. This suggests that overall
methanol permeability through the membrane is a direct result of water permeability.
In efforts to reduced methanol crossover Pu and co-workers attempted to
modify Nafion® using a methanol impermeable protonic conductor (MIPC)
23
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
composite electrolyte. 55 This consisted o f a thin palladium foil sandwiched between
two Nafion® membranes. These membranes were tested in H2/O2 fuel cells with
small amounts o f methanol added to the anode stream. The OCV was unaffected
indicating methanol did not permeate to the cathode where it could lower the
potential. Unfortunately these composites suffered from severe ohmic losses due to
the high contact resistance o f the Nafion®-Pd interface(s). These cells have not yet
been tested in direct methanol fuel cells.
Further modification o f Nafion® membranes involved doping phosphoric acid
into the polymer electrolyte. 56 The phosphoric acid acts as a Bronsted base, ionizing
the strong sulfonic acid groups and solvating the proton in the same manner as water.
These membranes are operated at temperatures greater than 100 °C in order to
minimize methanol crossover. At these conditions the membrane maintains proton
conductivity while the methanol permeability is reduced due to the lowered vapor
activity. Nafion®-H3P0 4 polymer electrolytes have been shown to yield sufficient
proton conductivity but have yet to be tested in an operating fuel cell.
The extensive amounts o f work done characterizing methanol and water
permeability in Nafion® reaffirm the cathode-related difficulties associated with
water removal and methanol crossover previously discussed in sections 1.3.2 and
1.3.3. Efforts to modify Nafion® in hopes o f reducing methanol crossover have met
with limited success. Therefore, much research in recent years has focused on the
development o f alternate polymer electrolyte membranes. They must exhibit the
same physical, thermal and chemical stability while also providing adequate
24
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
proton conductivity with reduced methanol crossover. Novel membranes must also
exhibit lower water permeability than Nafion® to achieve operation at lower
stoichiometries.
1.4 Alternative Polymer Electrolyte Membranes
Sulfonated polysulfones with various sulfonation levels have been prepared
and evaluated as solid polymer electrolytes. 5 7 Films were typically solvent or slurry
cast and sulfonated up to 90 mol% using chlorosulfonic acid. These films displayed
adequate proton conductivity but suffered mechanical degradation due to membrane
swelling and dissolution. Membrane stability was improved through the use of
crosslinking agents but at the expense o f reduced proton conductivity. Fuel cell
testing has been limited to H 2/O 2 cells due to the swelling issues that would arise in a
direct methanol liquid-feed system. The degree o f membrane sulfonation has also
been varied in other polymer electrolyte systems. Proton exchange membranes
derived from polyether-ether ketone and poly-4-phenoxybenzoyl-l,4-phenylene
polymers containing 65 mol% o f sulfonic acid, resulted in high proton conductivity
10* 2 - 1C T 1 S/cm at room temperature. 58 These membranes suffered mechanical
degradation from swelling and have been earmarked for high temperature (150 °C)
H2/O 2 cell testing.
Doping the polymer network with concentrated acids has been proposed to
avoid swelling arising from increased levels o f membrane sulfonation. Polymeric
electrolytes based on complexes o f polyethylene oxide-polymethyl methacrylate
doped with phosphoric acid have been evaluated for fuel cell applications. 59
25
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Membrane conductivity varied with acid concentration and peaked at 20 mol%
H3PO4. However, at ambient temperature the conductivity values were 1-2 orders of
magnitude lower than Nafion®-117. These membranes were tested in H 2/O2 fuel
cells and featured poor electrical performance attributed to poor electrode-electrolyte
contact and loss of acid from the polymeric blend. These membranes have not yet
been tested in direct methanol fuel cells.
Phosphoric acid-doped polybenzimidazole membranes have also been
proposed for fuel cell applications. 60 These membranes exhibit excellent thermal
stability but require operating temperatures greater than 130 °C to maintain adequate
proton conductivity. Although these membranes have demonstrated reduced
methanol permeability they are not adequate for low temperature liquid-feed
systems. A recent effort in the continued testing of high Tg , amorphous polymers
has been the development o f sulfonated and crosslinked polyphosphazene based
solid electrolytes. 61 These membranes have reported ion exchange capacities (1.4
mmol/g) greater than Nafion®-l 17 (0.91 mmol/g), swelled less than Nafion® and had
lower water and methanol diffusion rates. The membranes also have adequate
proton conductivity and have shown promise for fuel cell applications but there has
been no cell performance reported to date.
Recently there has been renewed interest in the use of polystyrene-sulfonic
acid based membranes. Grafting polystyrene into an inert fluorocarbon matrix
showed promise in the 1960’s but was quickly shelved when Nafion® was
introduced. The poor mechanical stability of polystyrene-sulfonic acid
26
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
membranes was improved upon blending or grafting to low Tg olefins. There has
since been an extensive amount o f work done in radiation grafting o f styrene to a
wide range o f perfluoro- and partially fluorinated polymers in hopes o f preparing
membranes with reduced methanol crossover.
1.4.1 PSSA Grafted Membranes
Graft co-polymerization is a convenient method used to modify the surface of
a polymer by introducing a second polymer graft. 6 2 Typical graft copolymers consist
of long segments o f one monomer (backbone polymer), with one or more branches
(grafts) consisting o f long segments o f a second monomer. These co-polymers have
been used extensively to design permselective membranes for a wide variety of
separation processes. One such method has used N-vmylpyrrolidone grafted onto a
wide variety o f films (polytetrafluoroethylene, polyethylene, polybutene and
polydimethylsiloxane) in pervaporation studies o f water-dioxane mixtures designed
to selectively remove organic toxins from water supplies. 63 Permselective
membranes based on polyvinyldifluoride-g-polyvinylalcohol have also been used for
alcohol separation processes in the beer industry. Permselective analysis o f both
membrane systems suggest that the active polymer need not be soluble in the solvent
to have a strong interaction with one o f the solution phases.
Sugiyama et al. introduced sulfonic acid groups onto various polymeric
substrates through the co-grafting o f sodium styrene-sulfonate and acrylic acid. 64
The substrates consisted o f various shapes o f polymers (tubes, films and fabrics)
made from polyethylene, polypropylene and polytetrafluoroethylene. They found
27
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
that copolymerization with acrylic acid enhanced the ability o f sodium styrene-
sulfonate to diffuse into the polymer matrix. These sulfonated films had ion
exchange capacity similar to commercially available cation-exchangers such as
Amberlite IR-24 (4.6 mol/kg). These materials have not yet been tested in direct
methanol fuel cells.
Radiation grafting has been used by a number o f other research groups to
prepare proton exchange membranes. Rouilly et al. prepared styrene-grafted films
using teflon-co-hexafluoropropylene as the substrate. 65 These membranes were
prepared by simultaneous radiation grafting o f styrene, a preferred technique since
the precursor film is irradiated while submerged in the monomer bath thus avoiding
radical coupling between individual teflon chains, a common phenomenon seen
using the pre-irradiation technique. Degrees o f grafting ranged from 13 - 52% and
FT-IR was used to confirm uniform styrene distribution. These membranes
exhibited very low specific resistivity values, especially at degrees of grafting >
30%. However these membranes swelled 80-90 vol% in water, two times the value
o f Nafion®-117. The increase in volume and dimension associated with membrane
swelling resulted in very poor physical properties.
Buchi et al. prepared similar teflon-co-FEP — g - polystyrene membranes but
used a combination of crosslinking agents to modify the material. 66 Divinyl benzene
(DVB) and triallyl cyanurate (TAC) was used as crosslinking agents and found to
have profound influence on physical properties relating to water uptake, thus
affecting specific resistance and conductivity. Their results indicated that DVB
28
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
interlinks the individual polystyrene chains rigidly, reducing their mobility
considerably. This resulted in lower water uptake and increased specific resistance.
TAC has longer side chains than DVB and allows greater chain mobility resulting in
membranes with higher swelling and lower specific resistance. Membranes prepared
with both crosslinking agents had higher power densities than Nafion® by 60% in
H 2/O 2 fuel cells at specific operating conditions. However, these membranes had
reduced overall performance due to poor electrode-electrolyte inter facial contact.
Research efforts using styrene-grafted membranes was found technologically
attractive because the resulting films combine the good physical stability and
chemical resistance o f the substrate with the favorable conductive properties of
sulfonated polystyrene. For these membranes to be used in direct methanol fuel cells
knowledge o f the water sorption properties and state o f membrane water is a
prerequisite for membrane characterization. Preliminary membrane studies have
shown a strong influence o f PSSA degree of grafting on water sorption properties
since increased styrene grafting resulted in low specific resistance and high
conductivity values.6 5 ’66 However the correlation between water transport and
methanol diffusion with PSSA degree o f grafting requires a better understanding of
water sorption and permeability in these materials.
1.4.1.1 Water Management in PSSA Grafted Membranes
The water sorption properties o f PSSA grafted membranes is strongly
dependent on the degree of grafting and is expected due to the increased amounts of
benzene rings available for sulfonation. Flint et al. studied the incorporation of
29
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
polystyrene grafts into a PVDF matrix and found that the introduction o f acid groups
into grafted membranes creates a pseudo two-phase hydrophobic, hydrophilic
polymer, as shown in figure 1.9.67 They also determined that membrane properties
such as water content and protonic conductivity are a direct result o f the aqueous
phase o f the polymer. ME As prepared from these grafted membranes exhibited
adequate cell resistivity in H2/O2 fuel cells but had low open circuit voltages as a
result o f poor interfacial contact at the electrode-electrolyte interface. This was
attributed to the distinct bi-phase nature o f the grafted membrane and is a significant
concern in these membrane systems since catalyst utilization is dependent upon
adequate interfacial contact. Hietala et al. have also studied PSSA grafted
membrane systems in order to understand the nature o f the two-phase hydrophobic,
hydrophilic behavior. 68 Based on thermal analysis o f PVDF-g-PSSA systems they
determined that the polystyrene grafts are incompatible with the PVDF matrix and
form phase separated microdomains within the amorphous regions o f the precursor.
Figure 1.9 Structure o f Polystyrene-Sulfonic Acid Grafted onto PVDF.
PVDF Base Polymer (F2C)— CH2)n
H
H,C-----Cv-w*
PVDF - g - Polystyrene
S 0 3-H+ (H20 )
PVDF - g - PSSA
30
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Ostrovskii et al. studied the state o f water and water sorption characteristics
of PVDF-g-PSSA membranes using Raman spectroscopy. 6 9 The investigation was
carried out on two different sets o f samples based on porous (Millipore) and non-
porous (Goodfellow) PVDF precursors. They found that the number o f water
molecules per sulfonic acid group, 11H 2O, is independent of the amount of sulfonate
groups in the material. Therefore the water sorption properties of the PVDF-g-PSSA
membranes, and thus, total amount of absorbed water are determined by the number
of sulfonic acid groups in the material.
Ostrovskii et al. also determined that the water content in the porous
membranes was significantly higher than that o f the non-porous films. This suggests
that the high degree o f porosity o f the initial PVDF matrix remains after styrene
grafting and facilitates water uptake. It was also determined that there is almost no
difference between the amount o f water taken up from the liquid phase or from
saturated water vapor for either membrane. This is different from the behavior
reported for Nafion® membranes for which the water uptake from water vapor is
considerably less than that from liquid water. 70*72 The difference in behavior of
water uptake between Nafion® and PVDF-g-PSSA membranes is explained by
studying the mechanism o f water sorption onto the sample surface. The hydrophobic
character o f the external fluorocarbon surface of Nafion® makes water condensation
difficult. In contrast, the surface o f the grafted membranes is a combination o f the
host matrix and that o f the grafted monomer (sulfonated polystyrene) that is
extremely hydrophilic and explains the indifference to liquid water or vapor. 70*72
31
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Raman spectroscopy has also been used to characterize the state of water
sorbed by the porous and non-porous membrane samples. The state o f water in the
porous PVDF-g-PSSA membranes is closer to bulk water whereas in the non-porous
material a significant part of the water molecules is confined to the backbone of the
polymer. This can be seen by determining the number o f water molecules per
sulfonic acid group in both materials. Membranes prepared using porous precursors
had 60 H2O per SO3' as opposed to only 24 H2O/SO3' in membranes prepared from
the non-porous precursors. This discrepancy has been used to explain the increased
water uptake measured in the porous membranes as compared to the non-porous
samples.
Molecular permeability through grafted membranes has also been
investigated. Tealdo et al. have characterized the permeability o f water-ethanol
mixtures through a number of PTFE grafted membranes. 73 They focused on the
impact of introducing low molecular weight styrene grafts to the PTFE backbone
using a grafting method developed by Stannet and co-workers. 74 The permeation
behavior depends on conditioning time due to rearrangement o f amorphous and
crystalline domains brought about by the permeating mixture. They determined that
the permeation rate (flux) increases with increasing degree o f grafting, however, the
selectivity factor is practically unaffected by styrene uptake.
The impact o f water uptake on permeability was further studied by Tealdo et
al.7 3 Highly grafted PSSA membranes can absorb much higher water content than
Nafion® resulting in superior conductivity values. However increased water content
32
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
with degree o f styrene grafting can result in unfavorable membrane swelling leading
to poor physical properties. Despite the importance o f swelling, there is no
relationship between permeation properties and swelling, a phenomenon well
documented. 75 Tealdo et al. concluded that permeation properties are mostly
dependent on the number o f active sites, much more so than with their distribution,
in agreement with the literature results. 73 Therefore it suggests that swelling may be
important to establish and maintain necessary water content but ultimately affects
membrane physical properties.
PSSA grafted membranes have shown potential to meet some o f the
necessary requirements for operation in direct methanol fuel cells. However the
micro structure o f the membranes suggest many similarities to the pore structure
proposed for Nafion®. The increase in bulk like water content with increasing PSSA
grafting is analogous to the impact o f decreasing equivalent weight in Nafion and
raises concerns regarding membrane stability considering the use o f a liquid-feed
system. Also agglomerated pockets o f PSSA grafts would most likely result in high
water and membrane permeability rates further inhibiting electrical performance.
Our research group at the University o f Southern California has proposed the
use o f ionomers prepared from interpenetrating polymer networks as novel polymer
electrolyte membranes for direct methanol fuel cell applications. 16 These materials
are different from co-polymers but similar in that they can mesh favorable individual
polymer properties into one component. 76 There is considerable activity in this area
since “new” products can be obtained by the physical mixing o f existing products.
33
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
1.5 Interpenetrating Polymer Networks
An interpenetrating polymer network is a blend o f two different polymer
systems without covalent bonds connecting the two. Millar7 7 , and later Shibayama et
a / . 78"81 first prepared IPNs in the early 1960’s and found that the technique provided
the only way o f achieving the equivalent o f a physical blend for systems containing
crosslinked polymers. Interpenetrating polymer networks are now possible for a pair
o f step polymerizations, a pair of chain polymerizations, or the combination of step
and chain polymerizations. IPNs have been used to modify the physical properties
o f commercially available plastics. The two individual polymer components form
continuous phases such that the phase o f the first formed network dominates physical
properties. For example, the inclusion o f a glassy polymer into a rubbery polymer
results in a toughened rubber. On the other hand, the inclusion o f a rubbery polymer
into a glassy polymer results in a high impact plastic. Some examples o f useful IPNs
are epoxy resin-poly sulfide, epoxy resin-polyester, epoxy resin-polyurethane and
polyurethane-polymethyl methacrylate. There are currently over 50 IPNs
commercially available.61
IPN systems are categorized according to the following methodologies:
simultaneous-IPN and sequential-IPNs. 82 Reacting a mixture of the monomers,
crosslinking reagents and catalysts for the two polymer systems produces a 'fully'
interpenetrating simultaneous-IPN. SimuItaneous-IPNs were first reported by Frisch
et al. in 1969 by mixing a urethane prepolymer with a low molecular weight epoxy
resin. 83 The mixture was then cured simultaneously via independent (non-
34
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
interacting) crosslinking reactions. In such cases where one of the polymers is not
crosslinked but the overall polymer system is still obtained in a one step process is
referred to as a semi-IPN. In certain publications these systems are also referred to
as pseudo-IPNs, a classification that will be used throughout this report.8 2
Reacting a mixture of one crosslinked, polymer precursor matrix and the
reagents for the other crosslinked polymer system produces a 'folly' interpenetrating
sequential-IPN. This methodology entails swelling a precursor film in a suitable
solvent or monomer and performing the polymerization in situ. These systems can
also be prepared by not crosslinking one o f the polymer networks and are referred to
as 'semi' interpenetrating sequential-IPNs. Sequential and je/w/'-sequential-IPNs
were first prepared and investigated by Sperling et al. and constitute a bulk of the
IPN work done in recent years.8 4 * 9 2 Incidentally the nomenclature used to
characterize interpenetrating polymer networks is often misleading so to avoid
confusion our proposed system is referred to as a se/m-sequential-IPN.
Adjusting the swelling o f the precursor polymer matrix according to solvent
and time can further modify sequential and j, ew/-sequential-IPNs. Sergeeva et al.
studied the impact of equilibrium and non-equilibrium precursor swelling on
polymer distribution in such systems.9 3 Gradient-IPNs are formed when non-
equilibrium swelling results in inhomogeneous distribution of the second polymer
phase whereby the amount o f material decreases from the exterior to the interior of
the matrix. The gradient technique is frequently used to adjust the physical
properties o f the resulting film Ensuring equilibrium swelling results in a sequential
35
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
or je/m-sequential-IPN since the monomer dissolves freely throughout the precursor
matrix.
Widmaier et al. further studied the effects o f precursor swelling on both
domain size and polymer distribution in sequential-IPN systems based on a
polystyrene-divinyl benzene matrix that exhibited heterogeneous distribution of
polybutadiene particles when the preformed network had a rather low degree of
swelling.9 4 When the precursor is swelled to a great degree, the polybutadiene
domains are larger and irregular but more uniformly distributed within the matrix.
This suggests that upon swelling the butadiene monomers self-polymerize within
preformed domains forming microgel particles that then agglomerate to form huge,
macromolecular clusters that then interconnect. Therefore the initial distribution of
domains within the precursor governs the distribution and morphology o f the final
network.
Miscibility o f the respective polymers also has a definitive impact on IPN
properties regardless o f the method(s) used to form the composite. Miscibility
though is not an essential requirement for a useful commercial polymer blend.
However, suitable mechanical properties are only achieved with favorable adhesion
between the respective polymer phases.9 5 " 9 7 IPN synthesis usually results in
multiphase morphology due to the well-known thermodynamic incompatibility of
polymers. IPNs exhibit various degrees o f phase separation depending on the
prehistory o f the system, interactions between the respective polymer phases and the
processing conditions used during system formation.
36
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
1.5.1 Factors Affecting Polymer Miscibility in IPN Synthesis
The important issues that determine morphology in IPN synthesis is the onset
o f phase separation, the rate o f phase separation and the time o f physical interlocking
o f the polymers.9 8 The onset point o f phase separation is the time when the Gibbs
free energy of mixing becomes positive due in part to the low entropy of mixing
between the individual polymer systems. The Gibbs free energy o f mixing is also
determined by the conversion (molecular weight), synthesis pressure(s) and
temperature(s) used in the polymerization. The rate of phase separation is related to
the mobility of the polymer chains and the medium viscosity o f the reaction mixture
and thus is indirectly related to the synthesis pressure and temperature. The time of
physical interlocking or network formation is the time when both component
polymers reach their sol gel transition, and thus the separated phase domain size
cannot increase much further beyond this point.
The relative rates o f the two competing kinetic processes o f phase separation
and network formation govern the morphology of the final product. When the
interlocking of the two component polymers occurs before the onset of phase
separation, the interlocked state prohibits further phase separation, and the resulting
simultaneous-IPN is a homogeneous mixture. When the interlocking occurs after the
onset of phase separation, simultaneous-IPNs with heterogeneous morphology are
obtained where only partial interpenetration exists. Sequential- and semz'-sequential-
IPNs have been prepared in part to reduce the kinetic contribution from phase
separation since they involve the use o f a preformed polymer network. The final
37
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
IPN is formed when the absorbed monomer polymerizes in situ and reaches its sol
gel point within the matrix. However in these systems the physical structure o f the
precursor and its interactions with the added polymer control the final morphology.
There are several additional factors that affect the relative rates o f phase
separation and network formation depending on the IPN methodology used. The
degree o f interfacial tension between polymer systems is a strong indicator as to the
level of miscibility between blends at a molecular scale and can be modified to
improve polymer mixing. Varying the blend composition can then optimize
favorable interactions between polymer components. The level (or lack of
crosslinking) in each system can also be adjusted to enhance the rates of network
formation in otherwise immiscible systems thus minimizing phase separation. The
synthesis pressure(s) and temperature(s) used are also factors that can influence
strongly the relative kinetic rates affecting IPN synthesis.
1.5.1.1 Interfacial Tension in IPN Synthesis
The high inter facial tension between polymer chains is a consequence of
specific intermolecular interactions such as hydrogen bonding, ionic interactions and
n-Pi complex formation.9 9 ,1 0 0 The sum of these interactions dictates the respective
miscibility between polymer systems and controls the degree o f phase separation
during IPN synthesis. Reducing the interfacial tension between immiscible polymer
systems is desirable to maintain a macroscopically “homogeneous” material despite
thermodynamically unfavorable interactions at the molecular level.
38
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Eisenberg et al. studied specific ion-ion, ion-dipole, and ion-pair-ion
electrostatic interactions between polymer blends during IPN synthesis.1 0 1 These
interactions typically apply only to ionomeric blends or polymers containing
ionizable functional groups. Therefore Eisenberg et al. analyzed polymer systems
with and without the presence o f such ionizable species in order to gauge their
influence on polymer miscibility. Despite their low interaction energy, PMMA and
polystyrene are not miscible and typically phase separate upon network formation
during IPN synthesis. However, copolymerization o f MMA with 4-vinyl pyridine
and styrene with styrene-sulfonic acid results in more intimate mixes of the
respective polymers. Eisenberg determined proton transfer from the styrene-sulfonic
acid regions to the vinyl-pyridine reduced segment mobility due to favorable specific
ionic interactions between the polymer chains.
Siqueira et al. also investigated the specific polymer-polymer interactions in
immiscible polymer networks mixed with copolymers.9 5 They prepared polymer
blends of PVDF and polystyrene that resulted in a very rough morphology with large
phase domains due to the high interfacial tension between the two systems. From the
Flory-Huggins interaction parameter X’ ^ d the solubility parameters 6, for
respective polymer systems, Siqueira et al. were able to calculate the interaction
energy B, for the mixing segments o f polymers. The interaction energy reflects the
enthalpy of mixing between polymer systems. Immiscible blends such as
PVDF/poIystyrene have an endothermic interaction energy, B = 18.4 E6 J/m*3.
However polymer blends of PVDF/PMMA have an exothermic interaction energy =
39
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
-17.3 E6 J/m'3 and result in polymers with well dispersed, uniform domains. The
highly exothermic enthalpy o f mixing in these blends is due to hydrogen bonding
interactions. Siquiera et al. determined that the domain size within
PVDF/polystyrene blends decreased with increasing PMMA content. This was
explained by the relatively low interaction energy (B = 0.2 E6 J/m*3) between
polystyrene and PMMA. Thus PMMA acts as an effective compatibilizer that
reduces interfacial tension between the two relatively immiscible polymer systems.
1.5.1.2 Polymer Compatibility in IPN Synthesis
It is widely known that the presence of certain polymeric species, usually
suitably chosen block or graft copolymers, can compatibilize immiscible systems by
reducing the degree o f interfacial tension between the two. The segments o f these
copolymers can be chemically identical with those in the respective phases1 0 2 * 1 0 7 or
miscible with or adhered to one of the phases.1 0 8 ' 11 0 Paul et al. pointed out that this
type o f interaction should 1) reduce the interfacial tension between the respective
phases, 2) permit a finer dispersion during mixing, 3) provide a measure of stability
against gross segregation, and 4) result in improved interfacial adhesion.1 1 1 In most
cases each block is at least partially miscible with one of the blend components and
the copolymer is located at the interfeces between the immiscible phases, increasing
the adhesion between them
Huelk et al. studied polymer compatibility in a sequential-IPN based on a
crosslinked, polyethyl acrylate precursor matrix impregnated with polystyrene-dvb
and found distinct two-phase character.1 1 2 The compatibility was improved by co-
40
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
polymerization o f methyl acrylate with styrene since PEA and MMA are chemically
isomeric and expected to be much more compatible toward each other than the PEA-
polystyrene pairs. Incorporation of the PMMA blocks into the blend resulted in
disappearance of the polystyrene cellular envelopes and retention of the fine 100 A
domain structure within the PEA precursor. Anastasiadis et al. further studied the
compatibilizing effect o f block copolymer addition to an immiscible IPN of
polystyrene/l,2-polybutadiene.1 1 3 They found a sharp decrease in inter fecial tension
between the two homopolymers with the addition o f polystyrene-block- 1,2-butadiene
to the blend. The few percent of block copolymer additive was required to
essentially saturate the interface and reach the limiting interfacial tension.
1.5.1.3 Polymer Composition in IPN Synthesis
Adjusting the polymer composition o f the respective IPN can influence the
relative rates of phase separation and network formation during IPN synthesis.
Widmaier studied the effects o f polymer composition on IPN morphology in
polyurethane-poly styrene pseudo-IPNs.1 1 4 The concentration of linear polystyrene
polymerizing simultaneously with crosslinked polyurethane determines the degree of
phase separation between the two systems. Polymerization of polystyrene with
molecular weights up to about 40,000 does not induce phase separation before the
gel point o f polyurethane. For higher molecular weights of polystyrene, phase
separation begins not only before gelation o f polyurethane, but continues to develop
during the formation o f polystyrene. The phase separation at high molecular weights
is caused by the mutual incompatibility o f the polymer systems. Network formation
41
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
prior to phase separation resulted in homogeneous films whereas phase separation
prior to gelation resulted in a fine dispersement o f the components and a gross phase
separation o f polystyrene moduli surrounded by a polyurethane rich shell.
The kinetics of network formation vs. phase separation was further studied by
Zhou et al. in simultaneous-IPNs based on polycarbonate urethane (PCU) and
polystyrene that exhibited a single-phase morphology when PCU levels reached 50%
or more o f the IPN composition.1 1 5 This morphology was achieved despite the
relative incompatibility o f the polymers. However pseudo-IPNs based on linear
polystyrene and/or polycarbonate urethane exhibited distinct two-phase morphology
with phase separated domains ranging from 1500-2500 A regardless o f PCU
composition. These results indicate that the crosslinking density o f the two systems
was required to establish network formation prior to phase separation considering the
two phases are immiscible.
1.5.1.4 Crosslinking Density in IPN Synthesis
In sequential-IPN synthesis the precursor network is the dominant,
continuous phase that dictates membrane morphology. Crosslinking density is the
controlling factor in determining the size and distribution o f domains within the
precursor matrix. Sperling et al. studied the impact o f crosslinking density on
domain formation within the precursor and their effects on the distribution o f the
second polymer system 1 1 6 Sequential-IPNs o f polystyrene polymerized within a
polystyrene-co-butadiene precursor had domain sizes o f 800 - 1000 A when the
matrix was crosslinked with 0.2% dicumyl peroxide, compared to a domain size o f
42
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
1600 - 1800 A when crosslinked with 0.1% o f the crosslinking agent. Later work
done by Donatelli et al. using similar materials confirmed that domain size is
controlled primarily by the crosslinking density o f the first formed network.1 1 7
Therefore the greater the degree of crosslinking, the finer the domain size within the
precursor resulting in a more uniform distribution of the impregnated polymer within
the host matrix.
Kim et al. also studied the effects o f crosslinking density in simultaneous-
IPNs o f polyurethane and polymethyl methacrylate that had fine domain sizes (500 -
2000 A ) due to the more favorable solubility parameters o f the monomers.1 1 8
However, pseudo-IPNs consisting of a linear and crosslinked polymer(s) exhibited
bi-phase morphology due to phase separation prior to network formation evident by
phase separated domain sizes of 2000 - 5000 A in diameter. This supports the
contention that the interlocking of chains during polymerization and crosslinking o f
each polymer component prevents, to a certain extent, the demixing o f the polymer
chains, thus increasing the degree of mixing o f the immiscible component networks.
In addition to its impact on polymer mixing, crosslinking can also affect the
physical properties o f an IPN. Allen et al. characterized the physical and mechanical
properties o f a simultaneous-IPN of polyurethane-polymethyl methacrylate.1 1 9 They
found that both the modulus and impact strength were strongly dependent on
crosslinking density, Me. When the network was highly crosslinked (low Me) the
composite had a high modulus but was relatively brittle. Conversely for lightly
crosslinked networks (high Me) the composite IPNs had high impact strength at the
43
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
expense o f modulus. Jasso et al. also investigated the physical characteristics of a
se/m-sequential-IPN comprised of a slightly crosslinked polystyrene matrix
impregnated with butyl acrylate monomer.1 2 0 The swelling o f the polystyrene
precursor was varied in order to form gradient-IPNs for physical comparison to the
se/w'-sequential-IPN systems. The physical properties o f the resulting films varied
with polybutylacrylate distribution such that gradient-IPNs with localized
polybutylacrylate at the surface had a toughening affect on the polymer composite.
IPN analysis confirmed the formation o f a distinct two-phase morphology thus
supporting the belief that the polymers need not be miscible to impart favorable
physical properties.
1.5.1.5 Temperature and Pressure in IPN Synthesis
In simultaneous-IPNs the synthesis pressure and temperature influence the
competing rates of network formation and phase separation. Lee et al. reported the
ability to increase the degree o f interpenetration between incompatible polymers
during the polymerization o f simultaneous-IPNs o f polyurethane-polymethyl
methacrylate and polyurethane-po lystyrene by applying pressure during the
polymerization process.1 2 1 In a separate report they indicated that the polyurethane
domain size(s) decreased from about 300 to 30 A with increasing synthesis pressure
from 1250 to 20,000 kg/cm2.1 2 2 At high synthesis pressure, the onset point o f phase
separation moves toward higher conversion and the mixture stays homogeneous at
high molecular weight. The synthesis pressure also affects the rate of phase
separation by reducing the mobility o f the polymer chains by reducing their free
44
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
volume. Lee et al. has since confirmed this relationship in several other analogous
simultaneous-IPN systems.1 2 3 Hourston et al. also studied the effects o f synthesis
pressure in sequential-IPNs o f polyurethane-polymethyl methacrylate.1 2 4 Increased
synthesis pressure resulted in improved mixing of the two networks but was not the
significant factor determining membrane morphology. The domain structure of the
precursor held precedence over all other synthetic factors relating to membrane
morphology, in accordance with previous results utilizing the sequential-IPN
methodology.
The synthesis temperature can also affect the miscibility o f multi-component
polymer systems. The temperature has both direct and indirect effects on the onset
point o f phase separation, the rate o f phase separation and the time o f interlocking.
Kim et al. studied the effects of synthesis temperature in .ye/m'-sequential-IPNs based
on a polyurethane matrix impregnated with polystyrene.9 8 The materials are
relatively immiscible but still showed a decrease in polystyrene domain size o f 2000
A to 300 A with a decrease in temperature from 40 to 0 °C. This was explained by
the feet that at low synthesis temperatures the chain mobility o f polystyrene is
reduced due to increased medium viscosity, thereby reducing the rate o f phase
separation. Despite the fact that the rate of network formation is also slightly
decreased at low temperatures the reduced ratio of the rate o f phase separation to the
rate o f network formation increases the degree of mixing o f the two component
polymers at the time of interlocking.
45
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
1.6 PVDF-PSSA IPN Polymer Electrolyte Membranes
Our research team has proposed polymer electrolyte membranes based on a
■se/m-sequential-IPN o f PVDF-polystyrene and subsequent sulfonation to PVDF-
polystyrene-sulfonic acid, for application in direct methanol fuel cells.1 2 5 The
physical stability, chemical resistance and methanol rejecting properties o f PVDF
make it an attractive precursor material for IPN synthesis. The incorporation of
polystyrene-sulfonic acid into the matrix provides the means for the water uptake,
permeability and proton conductivity necessary for cell performance. The use of an
IPN system is proposed to ensure uniform distribution o f PSSA throughout the
PVDF matrix.
The degree o f phase separation in PVDF-PSSA IPN membranes has a direct
correlation to water/methanol permeability and thus DMFC performance. The
presence o f localized regions o f PSSA results in the formation o f hydrophilic pores
analogous to the two-phase morphology o f Nafion® and PSSA grafted systems
leading to high rates o f water and methanol permeability. The distinct two-phase
morphology can also have an inhibitory affect on catalyst utilization at the electrode
electrolyte interface(s), especially in PSSA grafted systems. Catalytic particles need
to be in direct contact with the sulfonic acid moieties o f the membrane to allow
proton transfer to the bulk, thus maintaining a low interfacial resistance.
Macroscopic phase separation can result in irregular PSSA distribution whereby
many catalyst particles contact the PVDF matrix, not PSSA thus rendering them
inactive. The PVDF-PSSA membranes must maintain a uniform distribution of
46
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
PSSA throughout the matrix to ensure optimum catalyst utilization at the interface
and minimize macroscopic PSSA domains within the bulk.
Our research group has focused on many o f the issues that relate to phase
separation and thus the degree o f homogeneity in PVDF-PSSA sem/'-sequential-
IPNs. The type o f PVDF precursor selected governs the ultimate morphology o f the
system since it is the preformed polymer network. However, most sem/-sequential-
IPNs utilize a crosslinked precursor matrix with a linear additive polymer. Our
system proposes the reverse whereby the polymer additive is crosslinked and
polymerized within the precursor matrix. Therefore crystalline regions within the
PVDF precursor serve as pseudo-crosslinks determining domain size and subsequent
distribution o f impregnated styrene-dvb. The overall PVDF molecular weight, its
distribution and processing conditions influence the degree o f crystallinity and its
impact on styrene-dvb homogeneity. A thorough study o f the issues that affected
precursor characteristics and their relationship to IPN synthetic conditions was
initiated in hopes o f preparing uniform polymer electrolyte membranes suitable for
DMFC testing.
1.6.1 PVDF Synthesis
Polyvinyldifluoride is the addition polymer o f 1,1 difluoroethane (vinylidene
fluoride, VF2) and is polymerized readily by free radical initiators to form a high
molecular weight, partially crystalline polymer.1 2 6 The spatial symmetrical
disposition o f the hydrogen and fluorine atoms along the polymer chain gives rise to
unique polarity influences that affect solubility, dielectric properties and crystal
47
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
morphology. PVDF is readily melt processed by standard methods o f molding or
extrusion whereas coatings or porous membranes are prepared from solution using
polar solvents.
Vinylidene fluoride can be polymerized using a variety o f well-known
methods. Emulsion, suspension and solution polymerization can all be used and
typically employ chain transfer reagents to moderate the molecular weight and its
distribution. As is the case during conventional polymerization o f vinyl
monomers1 2 7 , reverse addition o f VF2 to the growing chain end may occur; the extent
of which is influenced by temperature.1 2 8 The incidence o f these defects referred to
as head-to-head and tail-to-tail units has been confirmed by a number of
spectroscopic techniques.1 2 9 ' 1 3 1 The polymerization conditions employed in
conjunction with the number o f defects can strongly influence PVDF properties such
as crystallinity and mechanical strength.
1.6.2 PVDF Crystallinity
The most profound impact on the morphology o f the resulting PVDF-PSSA
membranes is precursor crystallinity. The crystalline forms of PVDF involve
lamellar and spherulitic structures whose size, morphology and distribution strongly
influence precursor characteristics and thus IPN properties. PVDF crystallinity can
vary from about 35% - 70% depending on the method o f preparation,
thermomechanical history and processing condition(s). The degree o f crystallinity is
also an important factor influencing physical properties such as toughness,
mechanical strength and the high impact resistance o f the polymer. PVDF exhibits a
48
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
complex crystalline polymorphism absent in most other synthetic polymers. There
are 4 distinct crystal forms with a corresponding chain conformation: a (TGTGT’), P
(TTTT’), y (TTTGTTTG’) and 5 (the chain conformation for the 8 form has yet to be
confirmed).1 3 2 The polymorphs are present to different proportions in samples and
depend on the following processing conditions: pressure, intense electric fields,
controlled melt crystallization, precipitation from different solvents or seeding
crystallization.
The a form usually arises during melt fabrication processing and has a chain
conformation that is trans-gauche, placing the hydrogen and fluorine atoms
alternately on each side o f the chain, as shown in figure 1.10.1 3 3 The a form is
thermodynamically most stable and is most readily obtained from the melt under a
variety o f conditions. The a chain conformation is polar with dipole moment
contributions both normal and parallel to the chain. However, the unit cell is non
polar because the intrinsic structural polarity is neutralized by an anti-parallel chain-
packing configuration.
The P form consists o f an all trans chain conformation placing the fluorine
atoms on one side and the hydrogen atoms on the other, as shown in figure 1.10. 1 3 3
The P phase develops during thermal extrusion, calendering or compression molding
o f the melt processed polymer. It can also be directly obtained through the
mechanical drawing o f a phase films.1 3 4 ' 1 3 5 Mechanical deformation o f the melt
forming the P phase results in highly oriented films with exceptional mechanical
strength. Recent work suggests the P phase can also be obtained from slow
49
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
heating of a DMF solution o f PVDF and annealing at a specific temperature for an
extended period of time.1 3 6 The other two crystalline forms o f PVDF are usually
present to a very small degree and rely on unusual processing conditions. The y
crystals are only obtained from annealing o f very high molecular weight PVDF
blends while 8 crystals are formed by a distortion of one o f the other phases under a
high electrical field.1 3 7 138
Figure 1.10 Common Crystalline Conformations of PVDF.
Alpha Phase Beta Phase
1.6.2.1 Factors Influencing PVDF Crystallinity
The study of the crystallization behavior of PVDF is voluminous yet
important to understand the behavior of the polymer as a suitable precursor in IPN
synthesis. Factors such as molecular weight and processing conditions strongly
influence the degree o f crystallinity and the corresponding crystalline morphology.
Research into the use o f PVDF as a precursor in IPN synthesis suggests the
50
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
crystalline domains behave as pseudo crosslinks. Therefore the degree and order o f
crystalline domains could influence the distribution o f the impregnated polymer thus
controlling the morphology o f the blend.
Gregorio et al. studied the effects o f temperature on the crystallization
behavior of solution cast PVDF films.1 3 6 Solutions o f PVDF/DMA cast and heated
at temperatures below 160 °C and then quenched showed only the a phase.
Annealing the films at 130 °C resulted in the highest degree o f the a phase indicating
the maximum crystallization rate occurred here. Samples evaporated at temperatures
greater than 160 °C and then subsequently cooled resulted in predominately a phase
with small amounts o f y crystals. However, films prepared from the melt at
temperatures greater than 185 °C resulted in only the a phase since quenching passes
through a temperature that favors crystallization o f only this particular phase.
Since PVDF obtained from the melt results exclusively in a phase;
mechanical stretching is used to induce the formation o f P phase crystals. However
this process still results in films with a considerable amount o f the a phase. Several
other studies indicate that rapid quenching of the melt at low temperatures ( 0 - 3 0
°C) can result primarily in the formation of P crystals.1 3 3 ’1 3 9 The use of various
solvents to cast PVDF films can also induce crystallinity changes.1 4 0 PVDF has the
a morphology when cast from acetone or cyclohexanone and the p morphology
when cast from hexamethylphosphoramide. The difference in the types o f crystals
formed in the cast samples results from the various specific interactions between the
solvent and PVDF.
51
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Kim et al. further studied the factors that determine the formation o f the p
crystalline phase o f PVDF and PVDF/PMMA blends.1 3 9 The inclusion o f 20 wt%
PMMA to the melt resulted in films with maximum p phase content. Since the
content of the P phase depends strongly on the quenching temperature and rate, they
concluded that the inclusion o f PMMA drastically reduced the rate and temperature
o f crystallization. This is mainly due to the increase in the glass transition
temperature and is recommended as one o f the best ways to enhance the efficiency of
quenching.
1.6.3 Impact o f Crystallinity in IPN Synthesis
The use of PVDF in grafted polymer systems indicates immiscibility of
crystalline domains with the graft monomer results in polymerization within the
amorphous regions o f the precursor.6 7 This relationship is much less understood in
the realm o f IPN synthesis since polymer blends containing PVDF are usually
grafted systems or obtained from the melt. An understanding o f the impact of
precursor crystallinity on the uptake and distribution o f monomer during semi-
sequential-IPN synthesis would be helpful to better design more miscible polymer
blends. PVDF is among the few partially crystalline polymers that exhibit
thermodynamic compatibility with other polymers1 4 1 , in particular acrylic or
methacrylic resins.1 4 2 Strong dipolar interactions are important to achieve miscibility
with PVDF, as suggested by the observation that polyvinylfluoride is incompatible
with PVDF.1 4 3 The earlier work reviewed in section 1.5.1.1 based on polymer
blends o f PVDF and polystyrene indicated that the addition o f compatibilizers
52
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
such as PMMA were necessary to avoid macroscopic phase separation between the
PVDF and PSSA phases. PMMA acts as an effective compatibilizer for such
systems since it is widely considered to be compatible with PVDF due to favorable
hydrogen bonding interactions between the two polymers.9 5 However, the inclusion
of a non-crystalline polymer such as polystyrene into the melt o f a crystalline matrix
can impact spherulitic growth. An understanding o f this interaction and its impact
on crystallinity and thus the distribution o f the impregnated polymer can be
tremendously helpful in designing homogeneous IPN systems.
Saito et al. studied the intraspherulitic structure o f PVDF upon addition of
PMMA using polarized microscopy and small-angle X-ray scattering.1 4 4 Optical
micrographs o f PVDF crystallized from the melt show very compact, closed
spherulites having a clear Maltese cross pattern. However, inclusion of PMMA to
the melt resulted in PVDF spherulites with a very course, or open structure
suggesting PMMA hinders the radial growth of the crystalline lamellae, as suggested
by Morra and Stein.1 4 5 Saito et al. then specifically studied the distribution of
PMMA within the PVDF matrix. They found that the thickness o f the amorphous
region increased upon addition o f non-crystalline PMMA. This suggests that the
non-crystalline polymer is incorporated between the crystalline lamellae, not within
the lamellar bundles. This phenomenon has also been verified by a number o f other
research groups.1 4 6 "1 5 0
The work by Saito et al. and others with miscible PVDF based polymer
systems indicates that the second polymer diffuses into the amorphous regions o f the
53
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
crystalline polymer matrix. This is somewhat analogous to a sequential-IPN
whereby the second monomer polymerizes within pre-existing domains formed as a
result of precursor crosslinking. Since crystalline regions are thought to behave as
pseudo-crosslinks it is important to understand their impact on distribution of the in
situ polymerized polymer.
The je/m'-sequential-IPN methodology has been proposed since the
characteristics of the PVDF matrix, and thus final network can be predetermined.
Therefore the degree and distribution of crystalline regions that have a significant
impact on the distribution o f impregnated polystyrene-sulfonic acid can be evaluated
prior to membrane fabrication. Our laboratory has since conducted an exhaustive
analysis of suitable PVDF precursors and their impact on membrane morphology
upon in situ polymerization o f styrene-dvb.
54
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
1.7 References
1. Grove, W.R., Phil. Mag. 14, 127 (1839).
2. Scherer, G.G, Solid State Ionics 94, 249-257 (1997).
3. Appleby, A.J., Foulkes, F.R., "Fuel Cell Handbook", Krieger Publishing Co.,
Malabar, Florida, 1993.
4. K. Kordesch, G. Simader, "Fuel Cells and Their Applications", VCH
Verlagsgesellschaft mbH, Weinheim, Germany, 1996.
5. Eisenberg, A., King, M., "Ion Containing Polymers" Academic, New York,
1977.
6. Holiday, L., (ed.), "Ionic Polymers", Halstead, New York, 1975.
7. Harrison, J.W., (General Electric Company), the Potential of Fuel Cells for
Generation o f Electricity from Hydrogen, G.E. Direct Energy Conversion o f
Electricity from Hydrogen, G.E. Direct Energy Conversion Programs, Rep.
(undated).
8. LaConti, A.B., (General Electric Company), Introduction to SPE Coal
Technology, G.E. Direct Energy Conversion Programs, Rep. (undated).
9. McElroy, J.F., (General Electric Company), Status of Solid Polymer
Electrolyte Fuel Cell Technology, G.E. Direct Energy Conversion Program,
Rep. (undated).
10. Grubb, W.T., J. Electrochem. Soc. 106, 275-278 (1959).
11. Grubb, W.T., Niedrach, L.W., J. Electrochem. Soc. 107, 131-134 (1960).
12. Appleby, A.J., Yeager, E.B., Energy 11, 137 (1986).
13. Chapman, L.E., Proc. 7th Intersoc. Energy Conv. Eng. Conf., 466 (1972).
14. Hodgdon, R.B., Enos, J.F., Aiken, E.J., (U.S.P. 3,341,366; September 12,
1967; assigned to General Electric Company).
15. Surampudi, S., Narayanan, S.R., Vamos, E., Frank, H., Halpert, G., LaConti,
A., Kosek, J., Prakash, G.K.S., Olah, G.A., J. Power Sources 47, 377 (1994).
16. Smart, M., "Chemical and Electrochemical Oxidation of Small Organic
Molecules", UMI Ann Arbor, MI, 1998. «
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
17. Narayanan, S.R., Kindler, A., Jeffries-Nakamura, B., Chun. W., Frank, H.,
Smart, M., Surampudi, S., Halpert, G., "Performance o f PEM Liquid Feed
Direct Methanol-Air Fuel Cells", Electrochemical Society Proceedings
Volume 95-23, 261 (1995).
18. Narayanan, S.R., Kindler, A., Jeflries-Nakamura, B., Chun. W., Frank, H.,
Smart, M., Valdez, T., Surampudi, S., Halpert, G., Kosek, J., Cropley, C.,
"Recent Advances in PEM Liquid-Feed Direct Methanol Fuel Cells",
Proceedings o f the IEEE 11th Annual Batteiy Conference on Applications and
Advances, p.l 13, Long Beach, CA (Jan. 1996).
19. (a) Ren, X., Wilson, M.S., Gottesfeld, S., J. Electrochem. Soc. 143, L13
(1996). (b) Ren X., Wilson, M.S., Gottesfeld, S., Electrochem. Soc. Proc.,
PV 95-23 (1995) The Electrochemical Society, Pennington, NJ.
20. Shukla, A.K., Christensen, P.A., Hamnett, M., Hogarth, M., J. Power Sources
55, 87 (1997).
21. Moore, R.B., Martin, C.R., Macromolecules 21, 1334-1339 (1998).
22. Sondheimer, S.J., Bunce, N.J., Fyfe, C.A., JMS-Rev. Macromol. Chem. Phys.
C26(3), 353-413 (1986).
23. Eisenberg, A., Yeager, H.L., (eds.), "Perfuorinated Ionomer Membranes"
ACS Symposium Series 180, American Chemical Society, Washington, D.C.,
1982.
24. Gierke, T.D., Munn, G.E., Wilson, F.C., J. Polym. Sci., Polym. Phys. Ed. 19,
1687(1981).
25. Yeager, H.L., Chap. 4 in Reference 22.
26. Ceynowa, J., Polymer 19, 73 (1978).
27. Parthasarathy, A., Dave, B., Srinivasan, S., Appleby, A.J., Martin, C.R., J.
Electrochem. Soc. 139, 1634-1640 (1992).
28. Yeo, R.S., Chan, S.F., Lee, J., Journal o f Membrane Science 9 273-283
(1981).
29. Zawodzinski, T.A., Springer, T., Davey, J., J. Electrochem. Soc. 140, 1981
(1993).
30. Springer, T., Zawodzinski, T.A., Gottesfeld, S., J. Electrochem. Soc. 140,
1981 (1993).
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
31. Fuller, T.F., Newman, J., J. Electrochem. Soc. 140, 1218 (1993).
32. Nguyen, T.V., White, R.E., J.Electrochem.Soc. 140,2178 (1993).
33. Zawodzinski, T.A., Derouin, C., Radzinski, S., Sherman, R.J., Smith, V.T.,
Springer, T.E., Gottesfeld, S., J. Electrochem. Soc. 140, 1041 (1993).
34. Yeo, R.S., J. Electrochem. Soc. 130, 533 (1983).
35. Pourcelly, G., Oikonomou, A., Hurwitz, H.D., Gavach, C., J. Electroanal.
Chem. 43, 287 (1990).
36. Nandan, D., Pushpa, K., Kartha, P.K., Wahi, P.K., Iyer, R.M., Indian J.
Chem. 33A, 395 (1994).
37. Falk, M., in Perfluorinated Ionomer Membranes, ACS Series, 1982, Ch.8,
p.139.
38. Falk, M., Can. J. Chem. 58, 1495 (1980).
39. Ostrovskii, D., Brodin, A., Torell, L.M., Solid State Ionics 85, 323 (1996).
40. Uribe, F.A., Springer, T.E., Gottesfeld, S., J. Electrochem. Soc. 139 765
(1992).
41. Valdez, T i., Narayanan, S.R., Rohatgi, N., "Development o f a 150-W Direct
Methanol Fuel Cell System" Proceedings of the 15th Annual Battery
Conference on Applications and Advances, Long Beach, CA., Jan 11-14,
2000. IEEE pg. 37-40, Eds. Das, R.S.L., Frank, H.
42. Wang, J.T., Wasmus, S., Savinell, R.F., J. Electrochem. Soc. 143, 4 (1996).
43. Narayanan, S.R., Kindler, A., Jeffries-Nakamura, B., Chun. W., Frank, H.,
Smart, M., Surampudi, S., Halpert, G., "Performance o f PEM Liquid Feed
Direct Methanol-Air Fuel Cells", Electrochemical Society Proceedings
Volume 95-23,261 (1995).
44. Maricle, D.L., Murach, B.L., Van Dine, L.L., Abstract 35, The
Electrochemical Society Extended Abstracts, Vol. 94-1, San Francisco, CA,
meeting, May 22-27, 1994.
57
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
45. (a) Narayanan, S.R., Proceedings o f the 11th Annual Battery Conference on
Applications and Advances, Long Beach, CA., Jan 9-12, 1996. IEEE pg. 113-
122.
(b) Valdez, T.I., Proceedings o f Proton Conducting Membrane Fuel Cells II,
ECS Proceeding Series, Pennington, NJ, 1998 pg. 380-387.
46. Ravikumar, M.K., Shukla, A.K., J. Electrochem. Soc. 143, 8 (1996).
47. Narayanan, S.R., Frank, H., Jeffries-Nakamura, B., Smart, M., Chun. W.,
Halpert, G., Kosek, J., Cropley, C., "Studies on Methanol Crossover in
Liquid-Feed Direct Methanol PEM Fuel Cells", Electrochemical Proceedings
Volume 95-23, 278 (1995).
48. Ren, X., Henderson, W., Gottesfeld, S., J. Electrochem. Soc. 144, 9 (1997).
49. Narayanan, S.R., Vamos, E., Surampudi, S., Frank, H., Halpert, G., Abstract
73, p. 126, The Electrochemical Society Extended Abstracts, Vol 93-2, New
Orleans, LA., Meeting, October, 1993.
50. Verbrugge, M.R., J. Electrochem. Soc. 136, 417 (1989).
51. Ren, X., Zawodzinski, T.A., Uribe, F., Dai. H., Gottesfeld, S., in Proc.First
Int.Symp.Proton Conducting Membrane Fuel Cells I, Gottesfeld, S., Halpert,
G., Landgrebe, A., (Eds.) The Electrochemical Society, 1995.
52. Skou, E., Kauranen, P., Hentschel, J., Solid State Ionics 97 333-337 (1997).
53. Nandon, D., Mohan, H., Iyer, R.M., J. Membrane Sci. 71, 69 (1992).
54. Verbrugge, M.R., J. Electrochem. Soc. 136, 2 (1989).
55. Pu, C., Huang, W., Ley, K.L., Smotkin, E.S., J. Electrochem. Soc. 142, 7
(1995).
56. Savinell, R., Yeager, E., Tryk, D., Landau, U., Wainright, J., Weng, D., Lux,
K., Litt, M., Rogers, C., J. Electrochem. Soc. 141,4 (1994).
57. Nolte, R., Ledjeff, K., Bauer, M., Mulhaupt, R., Journal o f Membrane
Science 83, 211-220 (1993).
58. Kobayashi, T., Rikukawa, M., Sanui, K., Ogata, N., Solid State Ionics 106,
219-225 (1998).
59. Przyluski, J., Dabrowski, A., Stys, S., Wieczorek, W., Solid State Ionics 60,
141-146 (1993).
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
60. Bouchet, R., Siebert, E., Solid State Ionics 118, 287-299 (1999).
61. Guo, Q., Pintauro, P.N., Tang, H., O'Connor, S., Journal o f Membrane
Science 154, 175-181 (1999).
62. Odian, G., "Principles o f Polymerization", 3rd Edition, Wiley-Interscience,
1991.
63. Morel, G., Jozefbnvicz, J., Aptel, P., Journal o f Applied Polymer Science 23,
2397-2407 (1979).
64. Sugiyama, S., Tsuneda, S., Saito, K., Furusaki, S., Sugo, T., Makuuchi, K.,
React Polym. 21, 187-191 (1993).
65. Rouilly, M.V., Kotz, E.R., Haas, O., Schere, G.G., Chapiro, A., J. Membrane
Sci. 81, 89-95 (1993).
66. Buchi, F.N., Gupta, B., Haas, O., Scherer, G.G., J. Electrochem. Soc. 142,
3044 (1995).
67. Flint, S.D., Slade, R.C.T., Solid State Ionics 97, 299-307 (1997).
68. Hietala, S., KoeL, M., Skou, E., Elomaa, M., Sundholm, F., J. Mater. Chem.
8(5), 1127-1132(1998).
69. Ostrovskii, D.I., Torell, L.M., Paronen, M., Hietala, S., Sundholm, F., Solid
State Ionics 97, (1997) 315-321.
70. Hinatsu, J.T., Mizuhata, M., Tokenaka, H., J. Electrochem. Soc. 141 (1994)
1493.
71. Zawodzinski, T.A., Dorouin, S., Radzinski, S., Sherman, R.J., Smith, V.T.,
Springer, T.E., Gottesfeld, J. Electrochem. Soc., 140, (1993) 1041.
72. Zawodzinski, T.A., Gottesfeld, S., Shoichet, S., McCarthy, J., J. Appl.
Electrochem. 23 (1993) 86.
73. Tialdo, G.C., Canepa, P., Munari, S., Journal o f Membrane Science 9, (1981)
191-196.
74. Bentvelzen, J.H., Kimura-Yeh, F., Hopfenberg, H.B., Stannet, V., J.
Appl.Polym.Sci. 17 (1973) 809.
75. Michaels, A., Baddour, R.F., Bixter, H.J., Choo, C.Y.,
Ind.Eng.Chem.Process.Des.Dev. 1 (1962) 14.
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
76. Paul, D.R., Barlow, J.W., Keskkula, H., in "Polymer Blends", pp.399-461 in
"Encyclopedia o f Polymer Science and Engineering," Vol. 12, Mark, H.F.,
Bikales, N.M., Overberger, C.G., Menges, G., Eds., Wiley-Interscience, New
York, 1988.
77. Millar, J.R., J. Chem. Soc. 1311 (1960).
78. Shibayama, K., Suzuki, Y., Kobunshi Kagaku 23,24 (1966).
79. Shibayama, K., Zairyo 12, 362 (1963).
80. Shibayama, K., Kobunshi Kagaku 19, 219 (1962).
81. Shibayama, K., Kobunshi Kagaku 20, 221 (1963).
82. Sperling, L.H., "Interpenetrating Polymer Networks and Related Materials",
Plenum Press, New York, 1981.
83. Frisch, K.C., Klempner, D., Mukherjee, S.K., Frisch, H.L., J. Appl. Polym.
Sci. 18, 689 (1974).
84. Sperling, L.H., Friedman, D.W., J. Polym. Sci., Part A-2 7, 425 (1969).
85. Sperling, L.H., Taylor, D.W., Kirkpatrick, M.L., George, H.F., Bardman,
D.R., J. Appl. Polym. Sci. 14, 73 (1970).
86. Sperling, L.H., George, H.F., Huelck, V., Thomas, D.A., J. Appl. Sci. 14,
2815(1970).
87. Sperling, L.H., Thomas, D.A., Covitch, M.J., Curtius, A.J., Polym. Eng. Sci.
12, 101 (1972).
88. Sperling, L.H., Chiu, T.W., Hartman, C.P., Thoma, D.A., Int. J. Polym.
Mater. 1,331 (1972).
89. Sperling, L.H., Thomas, D.A., Huelk, V., Macromolecules 5, 340 (1972).
90. Sperling, L.H., Sarge, H.D., J. Appl.Polym.Sci. 16, 3041 (1972).
91. Sperling, L.H., Mihalakis, E.N., J. Appl.Polym.Sci. 17, 3811 (1973).
92. Sperling, L.H., Chiu, T.W., Thomas, D.A., J. Appl.Poly.Sci. 17, 2443 (1973).
60
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
93. Sergeeva, L.M., Gorbach, L.A., Russian Chemical Reviews 65, (4) 345-354
(1996).
94. Widmaier, J.M., Tabka, M.T., Eur. Polym. J. 28, 5 (1992).
95. Siquiera, D.F., Galembeck, F., Nunes, S.P., Polymer 32, 6 (1991).
96. Barlow, J.W., Paul, D.R., Polym. Eng. Sci. 27, 1482 (1987).
97. Barlow, J.W., Paul, D.R., Polym. Eng. Sci.24, 525 (1984).
98. Kim, B.S., Lee, D.S., Kim, S.C., Macromolecules 19, 2589-2593 (1986).
99. Cowie, J.M.G., in "Integration o f Fundamental Polymer Science and
Technology", (Eds. Kleintjens, L.A., Lemstra, P.J.), Elsevier, London, 1986,
p. 138.
100. Walsh, D.J., in "Integration o f Fundamental Polymer Science and
Technology", (Eds. Kleintjens, L.A., Lemstra, P.J.), Elsevier, London, 1986,
p. 119.
101. Natansohn, A., Murali, R., Eisenberg, A., Chemtech 6, 1990.
102. Riess, G., Jolivet, Y., in "Copolymers, Polyblends, and Composites", Platzer,
N.A.J., Ed.; Advances in Chemistry Series 142; American Chemical Society:
Washington, D.C., 1975; p. 243.
103. Fayt, R., Jerome, R., Teyssie, Ph., J. Polym. Sci. Polym. Lett. Ed. 24, 25
(1986).
104. Ramos, A.R., Cohen, R.E., Polym. Eng. Sci. 17, 639 (1977).
105. Gailard, P., Ossenbach-Sauter, M., Riess, G., Makromol. Chem., Rapid
Commun. 1, 771 (1980).
106. Marie, P., Selb, J., Rameau, A., Duplessix, R., Gallot, Y., in "Polymer Blends
and Mixtures", Walsh, , D.J., Higgins, J.S., Maconnachie, A., Eds.; NATO
Advanced Science Institute Series; Mortinus NijhofF: Dordrecht, The
Netherlands, 1985; p. 449.
107. Barensen, W.M., Heikens, D., Piet, P., Polymer 15, 119 (1974).
108. Ouhadi, T., Fayt, R., Jerome, R., Teyssie, Ph., Polym. Commun. 27, 212
(1986).
61
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
109. Ouhadi, T., Fayt, R., Jerome, R., Teyssie, Ph., J. Polym. Sci., Polym. Phys.
Ed. 24, 973 (1986).
110. Ouhadi, T., Fayt, R., Jerome, R., Teyssie, Ph., J. Appl. Polym. Sci. 32, 5647
(1986).
111. Paul, D.R., in "Polymer Blends", Paul, D.R., Newman, S., Eds., Academic
Press, New York, 1978, 2, p.35.
112. Huelk, V., Thomas, D.A., Sperling, L.H., Macromolecules 5, 348 (1972).
113. Anastasiadis, S.H., Gancarz, I., Koberstein, J.T., Macromolecules 22, 1449-
1453 (1989).
114. Widmaier, J., Macromol. Symp. 93, 179-186 (1995).
115. Zhou, P., Frisch, H.L., Journal o f Polymer Science: Part A: Polymer
Chemistry 30, 887-893 (1992).
116. Sperling, L.H., Donatelli, A.A., Thomas, D.A., in "Recent Advances in
Polymer Blends, Grafts and Blocks", Plenum Press, New York, 1974, p. 375.
117. Donatelli, A.A., Sperling, L.H., Thomas, D.A., Macromolecules 9, 671
(1976).
118. Kim, S.C., Klempner, D., Frisch, K.C., Radigan, W., Frisch, H.L.,
Macromolecules 2, 258 (1976).
119. Allen, G., Bowden, M.J., Blundell, D.J., Hutchinson, F.G., JefEs, G.M.,
Vyvoda, J., Polymer 14, 597 (1973).
120. Jasso, C.F., Martinez, J.J., Mendizabal, E., Laguna, O., Journal o f Applied
Polymer Science, 58, 2207 (1995).
121. Lee, D.S., Kim, S.C., Macromolecules 18, 2173-2179 (1985).
122. Lee, D.S., Kim, S.C., Macromolecules 17, 268-272 (1984).
123. Lee, J.H., Kim, S.C., Macromolecules 19, 644-648 (1986).
124. Hourston, D.J., Huson, M.G., Journal o f Applied Polymer Science 46, 973-
979 (1992).
62
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
125. Prakash, G.K.S., Smart, M.C., Wang, Q.J., Atti, A., Olah, G.A., Narayanan,
S.R., Chun, W., Valdez, T., Surampudi, S., J. Electrochem. Soc. Paper to be
submitted.
126. Encyclopedia o f Polymer Science and Engineering, 17, 2n d Ed., p. 532, John
Wiley & Sons, Inc., 1989.
127. Marvel, C.S., Denon, C.E., J. Am. Chem. Soc., 60, 280 (1938).
128. Gorlitz, M., Minke, R., Trautvetter, W., Weisberger, G., Angew. Makromol.
Chem. 29/30, 137 (1973).
129. Ferguson, R.C., Baume, E.G., J. Phys. Chem. 83, 1379 (1979).
130. Bachmann, M.A., Gordon, W., Koenig, J.L., Lando, J.B., J. Appl. Phys. 50,
6106(1979).
131. Mattem, D.E., Fu-Tyan, L., Hercules, D.M., Anal. Chem. 56, 2762-2769
(1984).
132. Benedetti, E., Catanorchi, S., D'Alessio, A., Moggi, G., Vergamini, P.,
Pracella, M., Ciardelli, F., Polymer International 41, 35-41 (1996).
133. Ludwig, B. W., Urban, M. W., Polymer 39, 24 (1998).
134. Yang, D., Thomas, E.L., J. Mater. Sci. Lett. 3, 929 (1984).
135. Petermann, J., Gohill, R.M., J. Mater. Sci. 14, 2260, (1979).
136. Gregorio, R., Cestari, M., Journal o f Polymer Science: Part B: Polymer
Physics, 32, 859-870 (1994).
137. Weinhold, S., Litt, M.H., Lando, J.B., J. Polym. Sci. Polym. Lett. Ed. 17, 585
(1979).
138. Litt, M.H., Lando, J.B., J. Polym. Sci. Polym. Phys. Ed. 20, 535-552 (1982).
139. Kim, K.J., Cho, Y.J., Kim, Y.H., Vibrational Spectroscopy, 9, (1995) 147-
159.
140. Tazaki, M., Wada, R., Okabe, M., Homma, T., Kobunshi Ronbunshu, 50, 533
(1993).
141. Paul, D.R., Barlow, J.W., J. Macromol. Sci. Rev. Macromol. Chem. C18, 109
(1980). ^
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
142. Mijovic, J., Luo, H.L., Han, C.D., Polym. Eng. Sci. 22(4), 234 (1982).
143. Guerra, G., Karasz, F.E., MacKnight, W.J., Macromolecules 19, 1935 (1986).
144. Saito, H., Stuhn, B., Macromolecules 27, 216-218 (1994).
145. Morra, B.S., Stein, R.S., J. Polym. Sci., Polym. Phys. Ed. 20, 2261 (1982).
146. Keith, H.D., Padden, F.J., J. Appl. Phys. 34, 2409 (1963).
147. Keith, H.D., Padden, F.J., J. Appl. Phys. 35, 1270 (1964).
148. Warner, F.P., Macknight, W.J., Stein, R.S., J. Polym. Sci. Polym. Phys. Ed.
15,2113(1977).
149. Stein, R.S., Khambatta, F.B., Warner, F.P., Russell, T., Escala, A., Balizer,
E., J. Polym. Sci., Polym. Symp. 63, 313 (1978).
150. Hudson, S.D., Davis, D.D., Lovinger, A.J., Macromolecules 25, 1759 (1992).
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Chapter 2: Results and Discussion
PVDF-PSSA Polymer Electrolyte Membranes:
Fabrication and Analysis
2.1 PVDF-PSSA IPN Methodology
The precursor polymer matrix dictates membrane morphology according to
the sequential-IPN methodology by determining domain size and thus distribution of
the impregnated polymer. The precursor must be capable o f solvent and/or monomer
swelling without membrane dissolution. Kynar (PVDF) studies indicate that the
local dipole moment generated by chain packing allows for favorable interactions
with various polar aprotic solvents such as acetone that swells PVDF up to 28% by
weight with negligible dissolution o f the film over a 24-hour period. In preparation
o f membrane samples various Kynar precursors were first swollen in acetone and
then immersed in a styrene-dvb-initiator solution and subsequently treated in a hot
press for polymerization. Membrane samples were then weighed to determine
polystyrene-dvb uptake according to the following equation:
Weight (g) (PVDF-PS(DVB)) - Weight (g) PVDF
----------------------- X 100%
Weight (g) PVDF
Membrane samples were then either impregnated a second time or sulfonated in a
chlorosulfonic acid bath and converted to the PSSA form by water hydrolysis. A
more detailed account o f membrane preparation can be found in experimental section
3.8.
65
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Our first qualitative analysis o f the PVDF-PSSA membranes focused on their
physical morphology post-polymerization. Surface morphology is a very important
criterion in selecting membranes for MEA fabrication and fuel cell testing. The
MEA fabrication procedure involves direct catalyst deposition followed by hot
pressing catalyzed electrodes to the membrane. Membranes with irregular surface
morphology and/or polystyrene-dvb distribution can lead to poor interfacial contact
at the electrode-electrolyte interface. This can result in a high contact resistance
featuring poor charge transfer kinetics that reduce catalyst utilization at the anode
and cathode.
Phase separation in IPN films has been shown to manifest itself as
macroscopic morphological irregularities as compared to that o f the original
precursor. One such notable example in our work was severe membrane wrinkling
and brittleness resulting from styrene-dvb polymerization that was further
accentuated after sulfonation and hydrolysis. Our results indicated that the degree o f
surface irregularities was a function o f the PVDF precursor selected. We also found
that precursor swelling in addition to polymerization conditions contributed to the
degree o f phase separation. Exhaustive efforts optimizing the PVDF precursor in
addition to preparatory conditions were made to control surface irregularities as well
as to determine the nature o f physical deformations.
2.1.1 PVDF-PSSA Membrane Preparation
Much o f the original membrane work conducted in our laboratory featured
the je/w/'-sequential-IPN methodology utilizing PVDF powder (Aldrich) compression
66
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
molded into films (Q. Wang and M. Smart).1 5 1 These precursors ultimately led to the
fabrication o f PVDF-PSSA membranes with suitable physical morphologies that
were tested in direct methanol fuel cells. Despite the promising electrical
performance o f these membranes, to be discussed in section 3.1, precursor
fabrication was a difficult and time-consuming process since compression molding
of PVDF powder requires high temperature and pressure and typically resulted in
sheets with irregular thickness. Therefore efforts were directed toward the use of
commercially available grades o f PVDF, sold under the trademark Kynar. As
mentioned earlier, PVDF is readily available as thermally extruded, uniform films
with a wide variety o f thickness. Commercial precursor samples were provided by
Westlake Plastics and Elf Atochem, N.A. and incorporated into our se/m'-sequential-
IPN methodology.
2.1.1.1 Kynar 740 Series: Membrane Morphology
Westlake Plastics provided us with thermally extruded, Kynar 740 precursors
with thickness ranging from 5- 10 mils per sample. PVDF-PSSA membranes
prepared from Kynar 740 sheets were extremely wrinkled and brittle. Swollen
membranes were subjected to polymerization using a wide range o f temperatures and
pressures yet still exhibited very poor surface morphology. We also varied precursor
swelling in order to assure equilibrium swelling and thus homogeneous styrene
sorption throughout the bulk. However, even equilibrium membrane swelling
coupled with modifications to the polymerization conditions foiled to alleviate the
wrinkling problem. Polystyrene-dvb uptake values ranged from 6 - 8% per
67
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
impregnation while values dropped with higher polymerization temperatures
presumably due to styrene loss from evaporation. The data in table 2.1 gives some
of the polymerization conditions used and the effective surface morphology o f the
resulting films.
Table 2.1 Membrane morphology resulting from the use o f thermally extruded
Kynar 740 precursors, supplied by Westlake Plastics, impregnated using various
polymerization conditions. Membrane samples were heated to the specific
temperature and held for one hour under the corresponding pressure.
Kynar 740 Temperature Pressure Surface Morphology
Membrane: 97-3 150 °C 300 psi Wrinkled
Membrane: 97-5 150 °C 600 psi Wrinkled
Membrane: 97-8 120 °C 600 psi Wrinkled
Membrane: 97-9 135 °C 600 psi Wrinkled
Membrane: 97-11 130°C 1200 psi Wrinkled
Membrane: 97-14 171 °C 600 psi Wrinkled
One theory as to the cause of membrane wrinkling focused on the type of
spherulitic crystalline phase(s) present in the PVDF precursor. As previously
mentioned in section 1.6.2, PVDF is known to crystallize in several conformations
depending on precursor properties and/or fabrication procedures. Thermal extrusion
is used to preferentially form ( 3 crystallites as a result o f chain orientation from
stretching during PVDF processing.1 5 2 Kynar 740 is prepared in this manner
suggesting a high degree of p character that may influence membrane morphology.
However, Benedetti et al. reported that heating PVDF above a certain threshold
temperature would erase the thermal history o f the polymer thus eliminating the P
phase.1 5 3 Therefore, Kynar 740 precursors were heat treated in order to destroy p
68
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
character and quenched at temperatures favoring crystallization o f a phase
spherulites. Incidentally, Kynar 740 precursors that were heat treated in our
laboratory and provided by Elf Atochem, N.A. and processed using a wide variety of
swelling and polymerization conditions still exhibited extreme membrane wrinkling.
This suggested that the crystal morphology does not solely influence membrane
morphology, but is instead determined by contributions from a number o f other
factors.
The degree and distribution o f crystalline regions, not necessarily the
crystalline spherullitic phase, may play a larger role in the degree o f miscibility
between PVDF and polystyrene. The orientation o f crystalline spherulites within the
Kynar 740 precursor could localize styrene sorption and thus polymerization to
amorphous regions within the matrix. Since Kynar 740 is a thermally extruded film
a significant amount o f orientation can occur during precursor processing. Therefore
membrane wrinkling may be influenced by the presence o f phase-separated
polystyrene domains localized in highly ordered amorphous pockets heterogeneously
dispersed throughout the highly ordered, crystalline PVDF matrix.
We obtained KYNAR 740 powder in hopes o f preparing compression
molded precursor samples thus avoiding chain orientation common from thermal
extrusion. Membranes prepared from these precursors had improved physical
morphologies but still exhibited broad, irregular wrinkles throughout the membrane
surface. Polystyrene uptake values were very similar to the commercial sheets
suggesting these membranes incorporated the same amount o f monomer. At this
69
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
juncture we returned to compression molding Kynar 460 PVDF powder despite the
inconvenience and problems associated with precursor thickness. Over time though,
optimization o f our precursor fabrication methodology resulted in fairly uniform
films ranging from 8-13 mils in thickness.
2.1.1.2 Kynar 460 Series: Membrane Morphology
Precursors prepared from compression molding Kynar 460 powder (Aldrich)
had resulted in adequate PVDF-PSSA polymer electrolyte membranes as previously
mentioned in section 2.1.1. The unfavorable morphology issues associated with
Kynar 740 precursors, coupled with improvements to our precursor fabrication
process allowed us to return to the use o f Kynar 460 samples. Kynar 460
membranes prepared using the same polymerization temperatures and pressures as
Kynar 740 (table 2.1) had much improved membrane morphology. Membrane
samples were wrinkled but featured very broad, thick irregularities as opposed to the
distinct, oriented wrinkles of Kynar 740 based membranes. Polystyrene uptake
values were also very similar to that of Kynar 740 suggesting monomer sorption was
similar and not an issue.
As we began to understand the impact o f temperature and pressure on phase
separation between polymers we attempted to further modify our polymerization
conditions. Membrane morphology improved dramatically with reduced temperature
and pressure and polystyrene uptake levels increased due to increased initiator
utilization and reduced evaporative losses. Polystyrene uptake levels ranged from 10
- 15% depending on acetone swelling and initiator concentration. Therefore, we
70
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
found that precursors swollen to equilibrium in acetone coupled with low
polymerization temperature and pressure produced membranes with preferred
surface morphologies, as shown in table 2.2. Incidentally membranes prepared with
no polymerization pressure had very severe surface irregularities caused by the
polymerization o f pools o f agglomerated styrene localized on the surface. We found
a minimum threshold pressure helped to ensure uniform styrene polymerization on
the surface and presumably in the bulk.
Table 2.2 Membrane morphology resulting from the use o f compression molded
Kynar 460 precursors, processed at USC and impregnated using various
polymerization conditions. Membrane samples were heated to each specific
temperature and held for one hour under the corresponding pressure.
Kynar 460 Temperature Pressure Surface Morphology
Membrane: 97-18 120 °C 300 psi Smooth broad wrinkles
Membrane: 97-19 90 °C 160 psi Smooth
Membrane: 97-20 90 °C 160 psi Smooth
Membrane: 97-25 60 °C - 90 °C (1 Hour Each) 160 psi Smooth
Membrane: 97-26 60 °C - 90 °C (1 Hour Each) 160 psi Smooth
N 'fem brane: 97-30 60 °C - 90 °C (1 Hour Each) 160 psi Smooth
The switch to Kynar 460 compression molded samples resulted in a dramatic
improvement in the surface morphology o f our membranes. However, it was still
unclear what the main factor was that contributed to membrane wrinkling since our
optimized polymerization conditions still resulted in extreme wrinkling using Kynar
740 precursors. Our work with Kynar 740 and 460 suggested that the degree and
distribution o f crystalline domains was the dominant force influencing membrane
morphology, not necessarily the type o f crystalline phase. A closer look at the
71
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
PVDF materials indicates that Kynar 460 has a higher Mw than Kynar 740, a factor
known to influence degree o f crystallinity. Also there are significant differences
regarding precursor preparation (compression molding vs. thermal extrusion) that
could result in discrepancies regarding the degree and distribution o f crystalline
domains within the two materials.
Thermal analysis was used to investigate the nature o f the PVDF precursor(s)
specifically relating to the degree of crystallinity, miscibility between the respective
polymer systems and thermal stability of the membranes. These tests were intended
to elucidate the nature and influence of the PVDF precursor on the ultimate
morphology o f the prepared polymer electrolyte membranes.
2.1.2 Thermal Analysis
Thermal analysis is a technique widely used to study thermal transitions in
polymeric materials.1 5 4 Specifically, Differential Scanning Calorimetry is used to
understand thermal transitions associated with polymer mobility (Tg) and
crystallinity (Tm ). The corresponding heat capacities for these thermal transitions
can be used to elucidate polymer characteristics such as degree o f crystallinity, rate
of chain crystallization, molecular weight distribution as well as degree of
interpenetration between polymer blends.
2.1.2.1 PVDF Crystallinity
DSC analysis o f Kynar 740 and 460 PVDF precursors was proposed to
reveal the underlying differences between the materials resulting in such stark
contrasts in membrane surface morphology. DSC was ideally suited to specifically
72
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
address the degree and distribution o f crystalline domains between the two
precursors and thus their impact on membrane morphology. Tazaki et al. calculated
the degree of crystallinity o f various PVDF samples from their respective DSC
curves using the following equation:1 5 5
Xc = [Hsm /H*m ] X 100%
where Hsm and H*m are the heats o f fusion o f the PVDF sample and o f perfectly
crystalline PVDF, respectively. Rosenberg and associates previously reported an
H*m value = 104.7 J/g used by Tazaki et al. in their measurements.1 5 6
Membrane morphology as a function o f PVDF precursor suggests the degree
and distribution o f crystalline domains is a dominant factor. DSC analysis was
conducted on both Kynar 460 and Kynar 740 samples in hopes of quantifying this
relationship. All the Kynar 460 samples evaluated were compression molded and
either prepared in our laboratory or provided by Elf Atochem, N.A. The DSC
analysis o f a compression molded Kynar 460 precursor, prepared at USC, in figure
2.1 reports a melt temperature o f 158.14 °C with a corresponding endothermic heat
capacity o f 45.05 J/g. This is almost completely inline with corresponding values
found in the literature for a phase crystals and equate to a degree of crystallinity =
43%.1 5 7 It is interesting to point out that the onset o f melting is difficult to determine
and the melt endotherm is relatively broad with a small shoulder at approximately
160 °C. This qualitatively suggests a somewhat broad molecular weight distribution
while the shoulder is attributed to the presence o f y crystals known to crystallize
from the melt at temperatures close to that o f a crystals.1 5 8 The thermal transition
73
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
at — 60 °C is unclear and thought to be a crystallization phase change since PVDF is
known to undergo a p - a transition at temperatures ranging from 50 °C - 150 °C.1 3 9
Figure 2.1 DSC analysis o f a compression molded Kynar 460 precursor
processed at USC.
w P A R T N O . 9 9 6 3 8 1 .9 0 1
S o a p le : PVQF
S lz a : 1 9 .7 9 0 0 ag
Xotnocfc 5*C/m in
c a u n t : 5*C/MIN A r PURGE 100CC/HIN
0 .1 -1 --------------------------------------------------------
0 . 0 -
• T il
DSC
F I la : SRNDSC.008
O p e ra to r: G ary P l e t t
Run D ate: I S - J u n - 9 8 14: 39
- 0 . 1 -
2 - 0 . 2 -
- 0 . 3 -
-0.4-
156.14*C
-200 - 1 0 0 0 100
T e m p e ra tu re (*C)
300 300
DSC V 4.0B O uP ont 2000
A second DSC scan concentrating in the low temperature region indicates a
glass transition temperature o f -40.64 °C, as shown in figure 2.2. This is also in
agreement with PVDF Tg values obtained from the literature for Kynar 460.1 5 2
74
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.2 DSC analysis o f the glass transition endotherm for a compression
molded Kynar 460 precursor processed at USC.
V P A R T M O . 9 0 4 3 6 1 S O I
S a o p le : PVOF
S iz e : 1 9 .7 9 0 0 eg
M ethod: 5 * C /« in
C o a a e n t: 5*C/MIN a t PURGE 100CC/MIN
F i l e : SfiNDSC.000
O p e r a to r : G ary P l a t t
Run D ate: l5 - J u n - 9 8 1 4 :3 9
a . 02
o.ao-
-4 0 .64*C (I)
o - 0 . 0 2 -
- 0 .0 4
-100 -120 -20 40 -00 2 0 - 4 0 -6 0
T e e o e r a tu re (°C) OSC V4.0B D uPont 2000
Compression molded Kynar 460 samples were also prepared for us by Elf
Atochem, N.A. in attempts to incorporate commercial precursors into our membrane
fabrication process. Prior to membrane fabrication, E lf Atochem PVDF precursors
were also subjected to the same DSC analysis as our in-house materials to maintain
precursor reproducibility. These precursors resulted in the same preferred membrane
morphologies as our in-house materials and were later incorporated into ME As
(section 3.3.1.5).
75
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
DSC analysis o f compression molded Kynar 460 samples supplied by Elf
Atochem, N.A. indicates similar glass transition and melt temperatures as compared
to the samples prepared at USC, discussed in figures 2.1 and 2.2. As illustrated in
figure 2.3, the corresponding melt heat capacity o f 48.24 J/g is very similar to that of
the in-house sample and corresponds to a degree o f crystallinity = 46%. It is
important to note that the DSC scan rate was slower for this sample and resulted in a
more pronounced shoulder on the melt endotherm at 160 °C confirming the presence
o f a small amount o f y crystals that co-crystallize with a crystals upon cooling the
PVDF powder from the melt. The data also indicates a phase transition at — 50 °C
thought to be consistent with the (3 - a phase transition discussed earlier in figure 2.1.
Figure 2.3 DSC analysis o f a compression molded Kynar 460 precursor supplied
by E lf Atochem, N.A.
S a m p le 4 6 0 E P V D F Sim
S ta r 1 9 .9 1 0 0 m g
M a tfio d P V D F
Convnant SCArmtoOC 1 C/Wnto 225
C
DSC
F ie C .A T A U 3a W O S C W *O G 2
O p e ra to r G a ry P la n
R u n Da 16>M ar-99 1 7 0 1
000
•31. i ■ 0 02-
0 1
£
I
m
9
-0 .0 6 -
I«.4(rd«2-7V C
•100 9 0 1 00
Temperature (*C)
190 •90 200 2 90
76
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
DSC analysis o f thermally extruded Kynar 740 films, commercially available
from Westlake Plastics, signifies a dramatic change from that o f compression
molded Kynar 460. The resulting melt endotherm in figure 2.4 has two distinct
peaks at 165 °C and 171 °C corresponding to significant amounts o f both a and y
spherulites. The increase in melt temperature also suggests an increase in
crystallinity that is supported by a higher heat capacity for the combined melt region
corresponding to a degree o f crystallinity = 55%. This is only 10% higher than
Kynar 460, however the sharp, distinct melt endotherm suggests a much narrower
molecular weight distribution. The DSC plot does however maintain the (3 - a
transition at — 50 °C suggesting this phase transition is inherent to PVDF. The data
also indicates a glass transition temperature = -41.9 °C consistent with Kynar 460.
Figure 2.4 DSC analysis o f a thermally extruded Kynar 740 precursor supplied
by Westlake Plastics.
SOT0K 740 W PVDF %n
$ tc a 1 3 .0 8 0 0
MtftoOrPVDF
Oar m w r S C M n B C C 1 C A I W 3 2 5
DSC
F Ik C .V T A 'fia te C S O P W f 0C 1
Oom*or. S s y R *
W in C a te IS -U a r-8 9 10 M
N.. •«iyc .
s jr c '--J
S C O
1
u.
1
• a o 200 2 80
IV2JCTM
77
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
In our efforts to improve membrane morphology using commercial materials
we also evaluated thermally extruded Kynar 740 precursors as supplied directly from
Elf Atochem, N .A As previously discussed in section 2.1.1.1, these precursor
samples were heat treated to destroy the thermal history o f the polymer yet still
exhibited poor membrane morphologies even after using our improved
polymerization conditions. The resulting melt endotherm in figure 2.5 indicates two
sharp, distinct peaks with identical melt temperatures (165, 171 °C) as the sample
from Westlake (Fig. 2.4) with a corresponding degree o f crystallinity = 55%. The
glass transition temperature = -42.52 °C and is consistent with the Kynar 460 and
Kynar 740 precursors previously tested.
Figure 2.5 DSC analysis o f a thermally extruded Kynar 740 precursor supplied
by E lf Atochem, N .A The sample was heat-treated at 200 °C for 15 minutes, cooled
and subjected to DSC analysis.
S a m p tK 7 4 0 E P V D F fMm
S h e 1 8 .1 9 0 0 m g
M e ix itP V P F
Commorc 5 CAran so 0 C 1 C/mtnto225
DSC
Fie. C.\TA\DaortDSOPvdf 003
O p a ra k v -G a ry PlaQ
R i* D a i*. 1 7 -U a r-M 11 C3
•0.02 ~
«
-0.10
• 0 0 90
Temperature (*C )
100 150 200
78
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The DSC analysis o f the PVDF precursors summarized in table 2.3 supports
many o f our premonitions regarding the causes o f poor membrane morphology. It is
clear that Kynar 740 either thermally extruded or heat-treated has a higher degree o f
crystallinity than compression molded Kynar 460 precursors. We investigated the
factors involved in determining precursor crystallinity in order to better understand
subsequent issues regarding polymer miscibility.
Table 2.3 PVDF properties o f the various Kynar grades tested and their
resulting degree o f crystallinity as determined by DSC analysis.
PVDF Sample Processing Mw Tm°C % Crystallinity
use (Aldrich)
Kynar 460
Compression
Molding
540,000 160 43%
Elf Atochem
Kynar 460
Compression
Molding
540,000 160 46%
Westlake
Kynar 740
Thermal
Extrusion
370,000 166,171 55%
Elf Atochem
Kynar 740
Post - Heat Treated
Thermal
Extrusion
370,000 166,171 55%
Molecular weight has been shown to have a strong influence on the degree of
crystallinity in semi-crystalline polymers.1 5 4 Increasing Mw results in a decrease in
the degree o f crystallinity in most polymer systems and is due to a greater number of
chain entanglements that interfere with the crystallization process. This is supported
by the data in table 2.3 indicating that Kynar 740 has a Mw o f 370,000 as compared
to a Mw = 540,000 for Kynar 460. Since Kynar 460 is the highest molecular weight
PVDF polymer available, we proposed using copolymers such as PVDF-HFP to
79
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
further reduce crystallinity and its negative impact on membrane morphology.
Unfortunately these precursors completely dissolved during acetone swelling.
The actual PVDF polymerization process can also affect the degree of
crystallinity and is closely associated with molecular weight. Kynar 460 is
polymerized using a proprietary process designed to achieve high molecular weights
and is known to introduce a large number o f head-head and tail-tail chain defects
into the polymer.1 6 0 These defects are known to interfere with the chain packing
mechanism necessary for crystallization. Kynar 460 has been determined to have 1
defect per 8-10 monomer units while Kynar 740 has only 1 defect per 20 monomer
units. The higher rate of defects in Kynar 460 supports the lower degree of
crystallinity measured for these samples.1 5 2
The Kynar 740 polymerization process is also proprietary and results in lower
molecular weight as well as a narrow molecular weight distribution.1 6 0 The polymer
is thermally extruded into sheets and films according to a process that further
enhances orientation o f crystalline domains since a strong sheer force is exerted on
the polymer. By contrast Kynar 460 has a broad molecular weight distribution that
results in a very broad melting endotherm. The relationship between molecular
weight distribution and the shape o f the melting endotherm has been confirmed by
discussions with E lf Atochem, N.A.1 6 0 and other literature sources.1 5 4
Compression molding o f a high molecular weight polymer such as Kynar 460
is very difficult since the polymer has a high melt viscosity and requires high
processing temperature and pressure in order to form sheets o f uniform thickness.
80
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
However, the favorable membrane morphology resulting from the use of Kynar 460
as compared to Kynar 740 indicates that it is the superior precursor. DSC analysis
suggests that the high degree and ordered distribution o f crystalline domains within
the Kynar 740 precursors enhance phase separation since the polystyrene rich
amorphous domains can not interconnect and are thus heterogeneously dispersed.
The reduced order and degree o f crystalline domains in the Kynar 460 precursors
allows for a more homogeneous distribution o f the impregnated polymer since
microgels o f polystyrene are allowed to polymerize and interconnect thus realizing
interpenetration. DSC analysis was further used to assess the level of
interpenetration between the PVDF matrix and impregnated polystyrene-dvb and
polystyrene(dvb)-sulfonic acid.
2.1.2.2 DSC Analysis: Polymer Interpenetration
DSC analysis was also proposed to characterize the degree o f interpenetration
between PVDF and polystyrene during membrane processing. It has long been
understood that completely miscible polymer blends will exhibit a single Tg , usually
somewhere intermediate between the Tg's o f the two respective polymer
systems.1 6 ,1 1 6 2 For incompatible blends, two Tg's are detected which approach one
another as a function of the degree of component compatibility. As previously
discussed in section 1.5.1.2, Zhou et al. used DSC analysis to characterize the phase
morphology o f simultaneous-IPNs consisting o f polycarbonate urethane/polystyrene
and determined that polymers containing greater than 50 wt% PCU exhibited a
single Tg indicating the existence of a single polymer phase.1 6 3
81
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Del Rio et al. further used DSC analysis to calculate the Flory-Huggins
interaction parameter for PVDF/PS compatibilized blends.1 6 4 The Tg values
determined from DSC analysis were fit into the Fox equation used to calculate the
volume fractions o f individual polymers within blended systems. The derived
interaction parameters indicated that polystyrene dissolves into PVDF only in the
presence o f compatibilizers such as carbon black or polyvinyl alcohoL PVDF does
not, however, dissolve into polystyrene under any circumstances.
2.1.2.2.1 PVDF-Polystyrene Interpenetration
Membrane morphological analysis of PVDF-PSSA membranes based on
Kynar 460 and Kynar 740 precursors suggested the former resulted in membranes
with adequate polymer miscibility. This was a strictly qualitative analysis based on
the physical morphology o f the membranes prepared from both precursors under a
variety o f polymerization conditions. We attempted to further study the degree of
polymer miscibility in membranes prepared from Kynar 460 precursors.
DSC analysis probing polymer miscibility was first conducted on a PVDF-PS
membrane sample prior to sulfonation. This was proposed to gauge the impact o f
sulfonation and corresponding water uptake on polymer interactions. The melting
endotherm in figure 2.6 at 159.5 °C has a heat capacity o f 46.4 J/g resulting in a
degree o f crystallinity = 44%. These values are almost identical to that o f the Kynar
460 precursor discussed in figure 2.1. This suggests that the incorporation of
polystyrene-dvb does not interfere with the crystallization process in PVDF.
However, the P - a phase transition identified in the Kynar 460 analysis in figure 2.1,
82
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
has shifted from 60 to 75 °C and may be attributed to the presence o f polystyrene-
dvb interfering with the phase transition. However, the relatively small shift in
temperature coupled with the already broad, (3 - a transition temperature range,
makes this conclusion uncertain. One interesting note is the presence o f a thermal
transition previously unseen in PVDF at ~130 °C and attributed to the polystyrene-
dvb glass transition temperature. Literature reports indicate a normal Tg — ~ 1 10 °C
for linear polystyrene and the presence o f 10 mol% dvb may account for the shift to
higher temperatures as mobility is further restricted.1 5 4
Figure 2.6 DSC analysis o f a PVDF-PS membrane sample prepared using a
compression molded Kynar 460 precursor processed at USC.
S am ple: PVOF-PS F ilm I — ] Q p F i l e : SRNOSC.0O6
S iz e : 1 9 .0 6 0 0 mg I J O O p e ra to r: 6 a ry P l e t t
M ethod: 10*C /m ln Run O ate: 5 - J u n - 9 8 09: 33
C om nant: 10*C/MIN Xr PURGE 100CC/MIN
•C . 2 -
e - 0 .4 -
9
X
-0 . 6 -
-0.0
-SO -1 5 0 100 200 2 5 0
OSC V4.O0 D uP ont 2000
-1 0 0 150
83
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The Tg o f PVDF in the PVDF-PS sample was also expanded in figure 2.7
since the incorporation of a second polymer into an IPN is known to impact the
precursor Tg. The analysis indicates a PVDF glass transition temperature = -40.87
°C that is similar to the values reported in literature and confirmed by our previous
precursor analysis.1 5 7 This suggests that the incorporation o f polystyrene does not
affect the Tg o f PVDF upon monomer uptake or in situ polymerization.
Figure 2.7 DSC analysis o f the glass transition endotherm for a PVDF-PS
membrane sample prepared using a compression molded Kynar 460 precursor
processed at USC.
Sam ola: PVCF-PS F ilm O O F i l e : SRNCSC.006
S iz e : 19.0600 *9 LJ O L-* O p e ra to r: Gary P l a t t
Metnoo: 10*C/m in Run D ata: S -J u n -9 6 09:33
Co am ent: 10*C/HIN Ar =»unG E 100CC/MIN
0.02
3 . 0 0 -
- 4 0 . 5 7 * c ( i ;
- 0 . 10
-150 -1 0 0 5 0 100 -50
T em p eratu re (*C) 0SC V4.0S DuPont 2000
84
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
2.1.2.2.2 PVDF-PSSA Interpenetration
DSC interpenetration analysis was then conducted on sulfonated samples of
our membranes. The data in figure 2.8 shows a melt endotherm at 158 °C with a
corresponding heat capacity = 37.88 J/g resulting in a calculated degree of
crystallinity = 36%, well below that o f our previous samples. However, due to the
broad melt endotherm it was difficult to identify where the onset o f melting takes
place and thus the heat capacity measurement may have a significant degree o f error.
It was also difficult to identify the p - a transition and polystyrene Tg from the plot.
The broad a transition suggests it may be hidden in the PVDF melt endotherm while
the presence o f sulfonic acid residues may further increase the Tg of polystyrene-dvb
since water loss during analysis results in extreme rigidity of the polystyrene-
sulfonic acid moieties thus masking or even eliminating the glass transition
temperature.
85
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.8 DSC analysis o f a PVDF-PSSA membrane sample prepared using a
compression molded Kynar 460 precursor processed at USC.
w PART NO. 994361.901
S a a o le : P o ly s ty re n e Film
S iie : 19.1 8 0 0 0 9
Method: 1 0 * C /ain
C onsent: 10*C/MIN 4r PURGE IOGCC/mZN
F ile : SRNCSC.007
O p e ra to r: Gary P l e t t
Run D ate: 5 - Ju n -9 6 13: * * S
x
x
a
a
o
- 0 . 5 —
- 1 5 0 -1 0 0 -5 0 50
" e s p e r s t o r e
20 0 250
OSC 7 4 .0 3 O u P o n t 2000
The Tg o f PVDF in the PVDF-PSSA membrane sample was again identified
to be at -42 °C according to figure 2.9. This remains consistent with our previous
samples again suggesting a lack o f influence from polystyrene-sulfonic acid on the
segmental motion o f PVDF.
86
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.9 DSC analysis o f the glass transition endotherm o f a PVDF-PSSA
membrane sample prepared using a compression molded Kynar 460 precursor
processed at USC.
PART NO. 906361 901
S a a o le : P o ly s ty r e n e F ilm
S iz e : 1 9 .1 8 0 0 mg
Metnoo: io * c /rn m
Comment: 10*C/MIN Ar PURGE 100CC/MIN
0 .05 - --------------------------------------------------------
DSC
F ile : SRNOSC.Q07
O p e ra to r: Gary P l a t t
Run G ate: 5— Ju n -9 0 13: 45
-42.26*C (U
-200 -150 - lOC -50 0 50 100
Temperature CC2 DSC V4.0S OuPcnt 2000
DSC analysis o f the PVDF-PS and PVDF-PSSA samples foiled to provide
conclusive data regarding polymer miscibility. However, this does not suggest that
there are no interactions between the respective polymer systems. The level o f
*
polymer interaction between immiscible systems such as PVDF and polystyrene may
lie at the molecular level out o f the boundaries o f DSC resolution. Also the lack of
influence, o f impregnated polystyrene on PVDF crystallinity or Tg values further
support the conclusion that styrene polymerizes in the amorphous domains within the
precursor matrix. The incorporation o f polystyrene-sulfonic acid within amorphous
domains o f PVDF may not specifically affect thermal transitions such as
87
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
crystalline melting or long-range segmental motion o f PVDF. Thus, the presence of
polystyrene in amorphous regions would preclude its influence on those transitions
since they are typically associated with crystalline segments of a polymer.
Polymer miscibility within sequential-IPN systems indicates that the nature
o f the precursor and its interactions with a second polymer system lie at the
microscopic level. The re/m-sequential-IPN methodology further suggests that
polystyrene-sulfonic acid is anchored in amorphous domains encapsulated by PVDF
and thus the distribution o f these interconnected amorphous domains, with respect to
the crystalline phases, dictates membrane morphology. However, the lack of
conclusive miscibility data from DSC analysis suggests other analytical methods
may be necessary to truly understand the nature of polymer interactions at the
molecular level.
Recently, solid state NMR has been used to study the morphologies of
polymer blends based on both miscible and immiscible systems. Papavoine et al.
used cross-polarization magic angle spinning (CPMAS) experiments to study
PMMA/PVF2 blends.1 6 5 The CPMAS technique involves transfer o f the
magnetization from lH protons o f PMMA to l9F fluorines on PVF2 . This technique
discriminates between PMMA and PVF2 segments that are close to or remote from
each other. They determined that the CPMAS experiment is sensitive to small
separations between polymer domains on the order o f 2-3 nm and therefore provides
unique information on miscibility at a molecular scale. Their results indicated that
phase separation increased with crystallization of PVF2 in the polymer blend since
88
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
magnetization from PMMA was lost. This supports our conclusions regarding the
exclusion o f PSSA from crystalline regions o f PVDF into amorphous domains, a
phenomenon only aggravated by the use of highly crystalline precursors such as
Kynar 740.
The CPMAS technique is a very promising method used to study polymer
miscibility at the molecular level with much better resolution than DSC. We have
not yet conducted these studies on our materials but are currently devising a set of
experiments based on this procedure. However, other thermal analytic techniques
such as thermogravimetric analysis can also be used to probe polymer miscibility.
The degree o f interpenetration o f the respective polymer systems can result in single
phase polymer degradation suggesting uniform distribution o f polystyrene domains
throughout the PVDF matrix. The presence o f a multi-phase degradative mechanism
suggests macroscopic phase separation and a truly immiscible system.
2.1.2.3 Thermogravimetric Analysis
Thermogravimetric analysis was proposed to assess the thermal stability o f
the PVDF precursors), PVDF-PS and PVDF-PSSA membranes. TGA is useful to
study the weight loss o f polymeric material from degradation as a function o f
increasing temperature and was used to substantiate our claims that these membranes
could tolerate proposed operating conditions ranging from ambient to 100 °C. It was
also used to identify thermal limits for hot pressing catalyzed electrodes to the
membrane relating to the membrane electrode assembly fabrication procedures to be
discussed in chapter 3.
89
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
2.1.2.3.1 Thermogravimetric Analysis: Kynar 460 and 740 Precursors
PVDF is a physically tough, thermally stable polymer melt processed at
temperatures ranging from 200 - 280 °C.1 5 2 Under these conditions there is no
reported risk o f decomposition except in the presence o f contaminants. TGA
analysis o f compression molded Kynar 460 as well as thermally extruded Kynar 740
provided by Westlake and Elf Atochem, N.A. indicates thermal stability up to 350 °C
as shown in figure 2.10. This is consistent with values found in the literature that
state decomposition resulting in HF formation from chain depolymerization at
temperatures greater than 350 °C.1 5 2 The TGA plot also suggests that the different
precursor processing conditions do not contribute to or influence degradation
temperatures in any significant way.
Figure 2.10 TGA thermograph for Kynar 460 precursors processed at USC and
supplied by E lf Atochem, N.A. and Kynar 740 precursor samples supplied by
Westlake Plastics.
TGA Analysis: PVDF 460 & 740
1002
740 V f
740 E -
« Q E
1 0 0 -
*
2S0 300 4 00 360 100 160
90
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
2.1.2.3.2 Thermogravimetric Analysis: PVDF-PS and PVDF-PSSA Membranes
TGA has previously been used to characterize PSSA based polymer
electrolyte membranes. Hietala et al. used TGA to study the thermal stability and
miscibility o f styrene grafted (PVDF-g-PS) and sulfonated (PVDF-g-PSSA) polymer
electrolyte membranes.1 6 6 They were able to specifically identify degradative
polymeric products from TGA using inline mass spectroscopic analysis. Their work
indicated that the degradation of the PVDF-PS samples is a two-step process starting
at -340 °C featuring degradative products typical from polystyrene. Thermal
degradation o f the PVDF backbone occurred at temperatures - 430 °C. This
suggests that the polystyrene grafts do not alter the inherent decomposition o f the
PVDF backbone and further supports their conclusion that polystyrene grafts are
incompatible with the PVDF matrix. Therefore, membranes prepared using this
grafting technique resulted in the formation of phase separated PSSA microdomains
within the PVDF matrix. Thermal degradation of the PVDF-PS(DVB) crosslinked
films occurred at 10 - 15 °C lower temperatures than the non-crosslinked
membranes. This surprising relationship is in contrast to the thermal stabilization
crosslinking agents provide to most polymeric materials.1 6 7 ,1 6 8 Hietala et al.
suggested the preferential reactivity o f dvb at grafting sites restricted polymer
mobility upon heating and resulted in fragmentation o f crosslinked polystyrene
segments.
91
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
We also conducted TGA studies on our PVDF-PS(DVB) membranes,
prepared using Kynar 460, to determine if improved polymer miscibility afiected
thermal stability. The TGA results in figure 2.11 indicate membrane decomposition
begins at — 325 °C and continues with a sharp drop in membrane weight at
temperatures exceeding 360 °C. This suggests improved miscibility o f PVDF-PS
using the se/n/'-sequential-IPN methodology as opposed to styrene grafting since the
TGA indicates a single degradative phase.
Figure 2.11 TGA thermograph o f a PVDF-PS membrane sample prepared using a
compression molded Kynar 460 precursor processed at USC.
Sam ple: p v d f-p s F ilm
S iz e : 2 5 .7 3 5 0 tag
Method: 5 * C /a ln to 400 *C
Comaent: 5 * C /a in to 4C0*C N2 p u rg e t0 0 c c /® m
102
GA
1 0 0 -
9 0 -
F l l e : SPN98.004
O p e ra to r: Sary P l a t t
Pun D ate: tS-M ay-98 10: 58
2 6 1 0 4 * C
i
,'L ,
\
v
\
9 6 -
9 4
9 2 - j -
0 150 200 250
T em p eratu ra !*C)
200 250 400
G eneral V4. tc DuPont 2000
Hietala et al. also used their TGA-Mass spectroscopic analysis to study the
degradative products o f crosslinked PVDF-g-PSSA membranes.1 6 6 Their results
indicated a mass loss o f ~10% between 100 - 180 °C that was attributed to water loss
92
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
from the membrane. Massive membrane degradation began at 220 °C and
corresponded to the mass peaks o f SO2 and SO3, respectively. Weight loss from
desulfonation was even seen as low as 200 °C and reached a maximum at 320 °C.
The next degradative weight loss corresponded to polystyrene grafts at 380-400 °C
and finally the PVDF backbone at 430 °C. Membranes prepared without the use of
crosslinking agents had improved stability by 10-15 °C but followed the same
degradative mechanism.
We conducted the same TGA studies on our crosslinked PVDF-PSSA
membranes to determine the relationship between sulfonation and reduced thermal
stability. The TGA plot in figure 2.12 indicates initial mass loss o f 20% at 75-100
°C that we attributed to membrane water loss. The next substantial weight loss
begins at ~200 °C and reaches a maximum at 300 °C. This correlates to the values
reported by Hietala et al. and is most likely the result of desulfonation. The final
degradative loss begins at 350 °C and corresponds to uniform membrane degradation
as seen in our analysis o f the PVDF-PS(DVB) samples (figure 2.11) prior to
sulfonation.
93
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.12 TGA thermograph o f a PVDF-PSSA membrane sample prepared
using a compression molded Kynar 460 precursor processed at USC.
Saapls: P o ly s ty re n e F i l s (Fuel C ell) “|“ A F ile : C :SPN 98.00l
Size: 84.7000 sg * w M. O p e ra to r: 9 ery P l e t t
Method: 5 * C /sm to 400*C Pun D ate: 9 -* ay -9 0 11: 40
Cossent: 5 * C /sln to 4Q0*C ns purge lO O cc/sin
1 0 0 C .4
-C .3
o
a
x
152.3‘
60
150 250 50 1 0 0 300 350 400
T6A V5.1A DuPont 2000
The TGA studies suggest that the incorporation o f polystyrene-sulfonic acid
into PVDF using the remz-sequential-IPN methodology result in the formation of
uniform polymer electrolyte membranes. The lack o f a multi-step degradative
mechanism as seen in PVDF-g-PSSA grafted systems indicates an improved level o f
polymer miscibility. We have, however, researched other techniques to characterize
polystyrene-sulfonic acid distribution within the PVDF precursor. This is especially
important since homogeneous sulfur distribution is critical to water sorption
characteristics and thus membrane electrochemical performance.
94
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
2.1.3 Energy Dispersive X-Ray Analysis (EDAX)
Membrane morphology specifically relating to polystyrene-sulfonic acid
distribution is a major parameter in designing novel polymer electrolyte membranes.
Infrared spectroscopy has typically been used to characterize membrane uniformity
in a number o f polymer electrolyte membranes. Bouchet et al. used IR analysis to
characterize nitrogen protonation throughout acid-doped polybenzimidazole
membranes.1 6 9 Flint et al. have also used this technique to monitor surface
polystyrene-sulfonic acid uptake and distribution in PVDF-g-PSSA membranes.1 7 0
IR spectroscopy can be qualitatively used to characterize the surface of a membrane
but is unable to monitor the distribution o f conductive moieties throughout the bulk.
This is critical since adequate proton conductivity is necessary along both the lateral
and transverse axis o f the membrane in order to maintain performance.
Energy Dispersive X-Ray Analysis (EDAX) is a technique we used to
measure the homogeneity o f polystyrene-sulfonic acid distribution along the cross-
section o f a membrane. Samples are placed in liquid nitrogen, fractured and coated
with an electronically conductive layer o f mixed gold and platinum. The sample is
then mounted in a holder and pulsed with a 20 kev beam in a scanning electron
microscope. A detector then measures the elemental X-rays emitted from the sample
after electron pulsing. The pulse is then applied throughout the cross-section in 10
micron intervals to give a representative sample of X-Ray emission. The detector
can be fine-tuned to measure certain elemental X-Ray emissions from carbon,
fluorine and sulfur. Sulfur X-Rays can only be emitted by the sulfonic acid group on
95
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
polystyrene and give a qualitative analysis o f polystyrene-sulfonic acid distribution
throughout the bulk.
EDAX proved to be a helpful tool in not only gauging sulfur distribution but
also modifying our sem/-sequential-IPN methodology. We initially varied the
acetone swelling times in order to understand conditions conducive to equilibrium
swelling. These times typically ranged from 1.5, 4, 8 and 12 hour swelling periods.
The resulting EDAX map seen in figure 2.13 shows the sulfur distribution for
membranes prepared using the 1.5 and 4 hour swelling period(s). The X-Ray map
indicates uniform fluorine distribution (top) but a staggered sulfur distribution, rich
at the edges but blank in the middle (bottom). This suggests that the precursor
acetone swelling step was insufficient since styrene polymerized at the edges but
never diffused into the bulk. Based on the sulfur distribution in figure 2.13 a
standard 12 hour heated acetone bath was incorporated into our membrane
fabrication methodology in order to ensure uniform styrene sorption and thus
homogeneous distribution upon in-situ polymerization
96
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.13 EDAX plot of a PVDF-PSSA membrane sample with localized sulfur
distribution along membrane edges.
EDAX was also used to identify problems in membrane handling prior to
hot-press polymerization. In our initial attempts to solve membrane morphology
related issues the membrane was typically wiped down prior to insertion into the hot
press. This was done under the assumption that agglomerated surface polystyrene
contributed to poor membrane morphology. Membranes prepared using this
technique had adequate surface morphology but poor sulfur distribution. The EDAX
map in figure 2.14 indicates a strong fluorine presence at the surface o f the sample
attributed to PVDF. However, the sulfur map indicates a lack of sulfur at the edges
with a stronger concentration in the middle. The X-ray map suggests the wiping
97
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
procedure prior to polymerization removed available monomer from the surface
since the bulk distribution o f polystyrene-sulfonic acid is adequate.
Figure 2.14 EDAX plot o f a PVDF-PSSA membrane sample with localized sulfur
distribution within the bulk.
EDAX analysis o f our membrane samples suggested a minimum acetone
swelling time was required for adequate styrene sorption as well as minimal handling
o f the membrane to avoid monomer loss. We modified our precursor swelling and
monomer incorporation accordingly to prepare membranes with desirable surface
morphology and a homogeneous distribution of polystyrene-sulfonic acid, as seen in
figure 2.15. In this analysis the corresponding fluorine and sulfur X-ray maps are
uniform throughout the sample cross-section.
98
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.15 EDAX plot o f a PVDF-PSSA membrane sample with homogeneous
sulfur distribution throughout the bulk.
EDAX analysis proved to be a highly effective tool used to characterize
sulfonic acid distribution in our membrane samples. In addition to validating our
je/w/'-sequentiai-IPN methodology, it is now used as a quality control measure to
effectively screen membranes for uniform sulfonic acid distribution. The ability to
adequately characterize the sulfonic acid distribution within our polymer electrolyte
membranes allowed us to move into the next phase o f membrane characterization
relating to water management properties.
99
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
2.1.4 Water Management: Membrane Characterization
Many important factors relating to fuel cell performance depend on the water
management characteristics o f the membrane. Membrane parameters such as water
content and proton conductivity are two such items that can impact cell performance
in several areas. For example, a highly conductive surface is necessary to maintain
protonic transport at the electrode-electrolyte interface influencing catalyst
utilization since a low inter fecial resistance is necessary to maintain adequate charge
transfer kinetics. Also the water and methanol permeability characteristics o f the
membrane impact cell performance issues specifically relating to the cathode.
Sufficient water permeability through the bulk maintains a low grain-boundary
resistance necessary to provide the cathode with the protonic current required
sustaining oxygen reduction. However, increased rates o f water and methanol
permeability can result in cathode flooding requiring higher stoichiometric flow rates
and mixed cathodic potentials that reduce overall cell performance.
In the analysis of PSSA grafted systems it has been shown that the degree of
grafting, and thus sulfonic acid content o f the membrane, dictates membrane water
management properties. Therefore the homogeneous uptake and distribution of
conductive polystyrene-sulfonic acid moieties in our je/w/'-sequential-IPN system
would be expected to behave accordingly. We performed a series of experiments
designed to better understand the relationship between PSSA uptake and the many
water management issues affecting cell performance.
100
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
2.1.4.1 Membrane Water Content
Membrane water content is an important parameter that has a cascading
influence on other water-related issues. As previously mentioned in section 1.3.2,
minimum membrane water content is necessary to maintain proton conductivity in
Nafion® and PSSA grafted systems. However, membrane water content also has an
impact on the physical integrity o f the polymer electrolyte membrane. PSSA grafted
membranes with high degree o f grafting typically display high proton conductivity
but lack stability due to unrestricted membrane swelling. The introduction of
crosslinking agents reduce swelling but do so at the expense of reduced water
content resulting in lower conductivity values.
The use of the ^ew/'-sequential-IPN methodology is attractive since the PSSA
phase is distributed throughout the PVDF precursor resulting in uniform water
sorption within the stable, physical confines of the PVDF matrix. We investigated
the influence of PSSA uptake and related crosslinking density on membrane water
content in a wide variety o f membrane samples measured as water weight gain
according to the following equation:
W eight (g) H20 - Weight (g) Dry
^ J X 100%
Weight (g) Dry
The measured membrane water content per sample was then plotted vs. the
corresponding PSSA uptake as shown in figure 2.16. The data indicates that the
membrane water content increases linearly with PSSA uptake and even exceeds that
101
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
o f Nafion®-117 above certain PSSA uptake levels. This relationship would be
expected since increasing polystyrene uptake would make available more aromatic
rings for sulfonation. This trend has been previously established in PSSA grafted
systems and fits well with our se/m'-sequential-IPN methodology.
Figure 2.16 Membrane water content for various PVDF-PSSA membrane samples
plotted vs. their corresponding PSSA uptake values.
50%
Water Uptake vs. PSSA Content
45%
40%
35%
■ 10% X-Unk
• 6 % X-Unk
30%
Nafion 117 = -3 0 %
25%
« 20%
1 5 %
10%
5%
0 %
10 20 25 0 5 15
PSSA Content %
The data in figure 2.16 also suggests that polystyrene crosslinking density
does not have a significant impact on membrane water sorption. Membrane samples
crosslinked with 6 mol% DVB exhibited slightly higher water content levels than
membranes crosslinked with 10 mol% DVB at comparable PSSA uptake values.
This is caused by the increased rigidity o f membranes with higher crosslinking
102
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
density limiting swelling and thus effective water uptake. One interesting note is the
impact o f crosslinking density on PSSA uptake observed during styrene
impregnation. Membrane samples crosslinked with 10 mol% DVB typically
required two polystyrene impregnations to achieve PSSA uptake levels similar to
that o f 6 mol% crosslinked samples impregnated only once. This is presumably due
to the increased reactivity o f DVB at higher crosslinking densities limiting available
monomer for polymerization. However, the data indicates that total PSSA uptake,
not the number o f styrene impregnations, governs membrane water content.
One important observation included in figure 2.16 is that the water uptake
measurements performed on Nafion®-117 result in water content = 30 wt%, in
accordance with published data. Therefore the results in figure 2.16 indicate that
membranes with a minimum PSSA uptake of 15 wt% exhibit water weight gain
similar to that of Nafion®-117. Since Nafion® is known to exhibit adequate proton
conductivity with sufficient water content we anticipated similar proton conductivity
values in the PVDF-PSSA membranes with water content values analogous to
Nafion®-117.
2.1.4.2 Proton Conductivity Measurements
The specific proton conductivity o f membrane samples was tested using a
D.C., 4-probe apparatus developed in our laboratory. The technique was
standardized using fully hydrated samples of Nafion®-117 since it has been
previously reported in section 1.3.1 that conductivity values increase and reach a
steady state value when the membrane is fully hydrated. This test resulted in specific
103
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
proton conductivity values = 70-75 mS/cm for Nafion®-117, in accordance with
published data. The technique was then used to determine the specific conductivity
value for a fully hydrated PVDF-PSSA membrane sample as shown in figure 2.17.
Figure 2.17 Specific conductivity analysis for a PVDF-PSSA membrane sample
using a 4-probe, D.C. apparatus. The membrane sample was fully hydrated and
tested at room temperature.
0.1
SpeciGc Conductivity Determination
USC PVDF-PSSA 98-35 0.09
0.06
0.07
0.06
0.05
Slope - 1971 Ohms
Thickness * 10 mils
W i d t h - 172 mOs
Conductivity — 47 mS/cm
O. 0.04
0.03
00 2
0 01
0.Q2 0.05 0.005 001 0.015 0.025 0.035 0.04 0.045 0 003
Current (mA)
Specific conductivity analysis was applied to a wide range of PVDF-PSSA
membrane samples o f varied PSSA uptake and crosslinking densities. The
conductivity measurements o f these samples are included in table 2.4 and range from
20-90 mS/cm, well within and even exceeding experimental values of Nafion®-117.
104
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Table 2.4 Specific proton conductivity values for PVDF-PSSA membrane
samples acquired from the 4-probe, D.C. apparatus. Membrane samples were frilly
hydrated and tested at room temperature.
Membrane Sample Specific Conductivity Average Specific Conductivity
Nation 1) 75 mS/cm
2) 70 mS/cm
72.5 mS/cm
97-33 1) 57 mS/cm
2) 60 mS/cm
58.5 mS/cm
98-35 1) 47 mS/cm
2) 50 mS/cm
48.5 mS/cm
98-36 1) 25 mS/cm
2) 34 mS/cm
29.5 mS/cm
98-37 1) 51 mS/cm
2) 57 mS/cm
54.0 mS/cm
98-38 1) 47 mS/cm
2) 53 mS/cm
50.0 mS/cm
98-40 1) 73 mS/cm
2) 75 mS/cm
74.0 mS/cm
98-41 1) 81 mS/cm
2) 90 mS/cm
85.5 mS/cm
98-42 1) 75 mS/cm
2) 81 mS/cm
78.0 mS/cm
A closer analysis of the conductivity values in table 2.4 suggests a strong
correlation with PSSA uptake supported by the accompanying plot in figure 2.18 that
indicates an increase in conductivity with membrane PSSA uptake. The data also
points out that the degree o f polystyrene crosslinking at either 10 or 6 mol% DVB
had no influence on proton conductivity values. This observation suggests that
PSSA uptake levels dictate proton conductivity while degree of crosslinking has a
greater impact on membrane water sorption, as previously shown in figure 2.16.
105
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.18 Specific proton conductivity values for PVDF-PSSA membrane
samples plotted vs. their corresponding PSSA uptake levels. Membrane samples
were fully hydrated and tested at room temperature.
140
Membrane Conductivity vs. PSSA Uptake
4-Probe, D.C. Analysis
120
100
S
cn
JL 8 0
u 80
3
e
o
U
20
25 5 10 20 30 35 0 15
PSSA Uptake %
We further investigated the influence o f water content on proton conductivity
in PVDF-PSSA samples o f varying PSSA uptake, as shown in figure 2.19. The
membranes were hydrated at varied time intervals from the dry state until they
exhibited a consistent water uptake over time. Conductivity tests performed before
attainment o f steady state water content resulted in very low values that subsequently
increased with complete hydration and exhibited a linear dependence on membrane
water content as would be expected. This observation confirms the necessity of
minimum membrane water content to establish and maintain consistent proton
conductivity.
106
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.19 Specific proton conductivity values for various PVDF-PSSA
membrane samples plotted vs. the corresponding water content level for that specific
PSSA uptake value.
90
Proton Conductivity vs. Water Content
50
G 30
20
1 5 * 20% 35% 45% 50% 40% 0% 5%
W ater Content %
The data in figure 2.19 also suggests that once minimum membrane water
content is achieved further membrane swelling does not improve proton
conductivity. Conductivity values for membrane samples crosslinked with 10 mol%
DVB exhibited the same conductivity as membrane samples crosslinked with 6
mol% DVB at the same PSSA uptake value. This was observed despite slightly
higher membrane water content in membranes o f lower crosslinking density, as
shown in the upper right quadrant of the plot in figure 2.19. This observation
supports our earlier results discussed in figure 2.18 regarding the impact of
crosslinking density or lack thereof on proton conductivity. The data ultimately
107
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
indicates that a membrane water content = ~30 wt % is sufficient for proton
conductivity values on the order of 50-60 mS/cm, data similar to the measured
membrane water content and proton conductivity values o f Nafion®-117.
Our laboratory also developed the use o f a 2-probe screening technique to
further test resistivity in membrane samples. The data in table 2.5 shows 2-probe
resistance values for a number o f membrane samples, including a Nafion®-117
standard, that exhibit the same linear correlation to PSSA uptake as the 4-probe
results in figure 2.18. However, the 2-probe results are purely qualitative and result
in conductivity values an order of magnitude lower than the 4-probe technique. The
decrease in conductivity is most likely associated with the high contact resistance at
the membrane-electrode interface inherent with the experiment.
Table 2.5 Resistivity values for PVDF-PSSA membrane samples as compared
to Nafion-117 at 25 °C. Resistivity values were determined using a 2-probe
technique and converted to conductivity values according to the equation described
in section 3.8.
Membrane Sample Resistivity Average Conductivity
Nafion 0.88 Ohms 10 mS/cm
97-33 1.34 Ohms 6.6 mS/cm
98-35 1.40 Ohms 6.3 mS/cm
98-36 3.80 Ohms 2.3 mS/cm
98-37 1.35 Ohms 6.5 mS/cm
98-38 6.00 Ohms 1.5 mS/cm
98-40 1.12 Ohms 8.0 mS/cm
98-41 1.08 Ohms 8.1 mS/cm
98-42 1.10 Ohms 8.0 mS/cm
108
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The 2-probe technique was used to assess conductivity along the transverse
axis o f a membrane sample since the 4-probe technique is only able to discriminate
PSSA distribution along the lateral axis of the membrane. For example, samples
with localized surface or bulk PSSA that exhibited adequate 4-probe conductivity
had high 2-probe resistivity values. Discrepancies between the two techniques
suggested irregular, localized distribution o f PSSA and prompted development o f the
ED AX technique described in section 2.1.3. The use of ED AX analysis proved
invaluable in elucidating the nature o f the discrepancies between the two techniques.
Therefore the 4-probe technique coupled with ED AX analysis offered insight into
conductivity along all axis o f the membrane. These two techniques are helpful
screening methods used to determine effective PSSA uptake levels and their relative
distribution throughout the PVDF matrix.
2.1.4.3 Methanol Permeability
The unique morphology o f the se/m-sequential-IPN system was proposed to
avoid the distinct two-phase character of Nafion® and PSSA grafted membranes,
thus reducing methanol crossover. We prepared a series of experiments that were
designed to monitor methanol diffusion through membrane samples along a
concentration gradient. PVDF-PSSA membrane samples were sandwiched between
a standard solution o f methanol and de-ionized water and aliquots of water were
analyzed over time using a gas chromatograph to monitor methanol diffusion
through the membrane. The data in figure 2.20 indicates that the methanol diffusion
rates in a variety o f membrane samples are roughly a third of Nafion®-117.
109
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.20 Methanol permeability through Nafion®-117 and PVDF-PSSA
membrane samples over time. Tests were conducted at room temperature using a 3.0
M methanol standard.
0 009
Methanol Permeability
3.0 M Methanol Standard, Room Temperature 0 008
0 007
Nation
0 008
S 0003
u s e P V D F -PSSA
Membranes
98-46N
98-45
97-33
98-37
98-35
0 001
98-36
210 00 120 150 180 240 0 30 80
Time (Minutes)
The slope of the diffusion rate(s) from figure 2.20 were then used to calculate
methanol diffusion coefficients and plotted vs. the amount o f polystyrene
impregnated per sample, as shown in figure 2.21. The plot indicates a linear increase
in the methanol diffusion coefificient(s) with increasing weight of polystyrene uptake
per sample. This trend would be expected since membrane water content increases
linearly with PSSA uptake and would provide a more aqueous environment for
methanol permeability. Despite this relationship membrane samples with PSSA
uptake = 20-25 wt% still exhibited methanol diffusion coefficients half that o f
Nafion®-117. Incidentally methanol was found to be completely impermeable
through the PVDF precursor prior to styrene impregnation indicating that methanol
permeability was solely attributed to the PSSA phase of the membrane.
110
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.21 Calculated methanol diffusion coefficients for various membrane
samples plotted vs. their corresponding polystyrene weight uptake in grams
impregnated per sample.
2.506-07
Methanol Permeability
•5 1.506-07
S 1.006-07
5.00E-06
3.5 2.5 0 0.5 1.5 2 3 1 4
Styrene Uptake (grmms)
The methanol permeability studies conducted on the PVDF-PSSA samples
revealed a significant reduction in methanol permeability as compared to Nafion®-
117. However, the water/methanol sorption and permeability mechanism(s) within
the membrane is not obvious since certain membrane samples exhibited adequate
water content and proton conductivity values as compared to Nafion®-117 despite
reduced methanol permeability rates. This suggests that the mechanism(s) o f proton
conductivity and methanol permeability in PVDF-PSSA membranes may not be
analogous to Nafion® or PSSA grafted systems and may instead rely on contributions
from a number o f other factors. In order to better understand the dynamics o f
111
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
methanol diffusion we developed a collaborative effort with Dr. Steve Greenbaum at
Hunter College using ‘H NMR studies to characterize methanol sorption and
permeability.
2.1.4.3.1 Methanol Permeability: NMR Analysis
Greenbaum et al. used ‘H NMR analysis to determine the methanol-water
partitioning coefficient for membrane samples equilibrated in concentrated solutions
o f methanol.1 7 1 The data in figure 2.22 indicates that membranes equilibrated at low
methanol concentrations (below 3M) preferentially sorbed methanol vs. water. The
data points that fall below the theoretical methanol-water partitioning coefficient
curve for that specific concentration confirm this. However, at concentrations
greater than 5 M methanol the membrane samples preferentially sorbed water since
the data points lie above the theoretical partition coefficient curve. This is in contrast
to the data presented for Nafion®-117 in section 1.3.4 that indicated a partition
coefficient = 1 meaning no preference to methanol or water sorption at
concentrations up to 15 M. Greenbaum et al. plotted sorption data from Nafion®-117
in figure 2.22 and it follows the theoretical partition coefficient value of 1 at most
concentrations.
112
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.22 Water-methanol partitioning coefficient determination for various
PVDF-PSSA membrane samples using NMR analysis. Membranes were dried and
equilibrated in methanol solutions at room temperature for 24 hours.
Membrane-CH3 0H-H2 0 solution presoaked for 48 hours
use 97-18
USC 97-33
USC 97-06
Nafion-117
USC 97-30
USC 97-33
USC 98-45
USC 9 8 -4 7 B
USC 98-52
USC 98-53
USC 98-57
USC 9 8 - 5 8 K
CH3 0H-H20 concentration (M)
Greenbaum et al. were also able to calculate methanol diffusion coefficients
over a wide concentration range for both PVDF-PSSA samples and Nafion®-117.
The results in figure 2.23 indicate that methanol diffusion coefficients are reduced in
the PVDF-PSSA samples as compared to Nafion®-! 17 at all concentrations tested.
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.23 Methanol diffusion coefficients for PVDF-PSSA membrane samples
as compared to Nafion®-117 determined using NMR analysis. Tests were performed
with methanol concentrations ranging from 1 - 5 M.
JL 1 0 *
c
c
£
3 10“
+■ ♦ +
Q
ia7
CH3 peak
*
*
5 10
CRjOH-HjO concne* ration (M)
H.O
+ CHjOH-HjO
-
USC 97-18
USC 97-33
▼
USC 97-06
Nnfion-117
15
Greenbaum et al. then calculated diffusion coefficients for water in PVDF-
PSSA samples as well as Nafion®-117 over the same concentration range as the
methanol experiments. The results in figure 2.24 show a reduction of water
permeability in the PVDF-PSSA samples as compared to Nafion®-117. The
difference between the calculated diffusion coefficients o f the PVDF-PSSA samples
and that o f Nafion® were o f the same order o f magnitude as the calculations
performed on the methanol experiment, previously discussed in figure 2.23.
114
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 2.24 Water diffusion coefficients for PVDF-PSSA membrane samples as
compared to Nafion®-117 determined using NMR analysis. Tests were performed
with methanol concentrations ranging from 1 - 5 M.
io- 1
OH peak
lo -5
H.O
+
CH^OH-HjO
c
USC 97-18
USC 97-33
T
USC 97-06
Naficn-I 17
c
«
o
£
£
a
1 0 *
■ t o - 7
10 15
CH 3O H-H 2O Concentration (M)
NMR analysis o f methanol permeability in PVDF-PSSA membrane samples
indicates that the methanol diffusion coefficients are reduced despite the preference
for methanol sorption at low concentrations. The NMR technique also indicated that
the water diffusion coefficients are reduced as well. This is an important observation
since it supports the belief that methanol permeability is closely associated with the
aqueous phase o f the polymer. The methanol sorption results suggest that despite
preferential sorption o f methanol the lack o f aqueous channels reduce methanol
permeability. However, membranes still exhibited adequate proton conductivity at
specific PSSA uptake levels despite the reduced water and methanol permeability
115
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
rates. This phenomenon suggests that the mechanism(s) of proton transport and
methanol permeability may be independent variables.
2.1.5 Water Management: Membrane Analysis
The mechanism(s) of methanol permeability and proton transport in polymer
electrolyte membranes has been an area o f significant research interest in recent
years. As previously discussed in section 1.3.3, methanol is believed to permeate
through aqueous channels present in Nafion® and PSSA grafted materials. The
formation o f these channels is thought to depend on the agglomeration of ionic
clusters formed during membrane hydration. The mechanism(s) of proton transport
is also strongly dependent on membrane water content as well as the state of the
bound water within the membrane. In order to better understand transport
phenomenon in our materials it is important to study these concepts at a more
fundamental level.
Kreur et al. has recently published a comprehensive review detailing the
mechanism(s) of proton conductivity in a wide range of materials. 172 He found that
the most trivial case is the assistance o f proton migration by the translational
dynamics o f bigger species referred to as the vehicle mechanism since the proton
diffuses together with a vehicular species like water, as H3O V 73 The
counterdiffusion o f unprotonated vehicles such as water allows for the net transport
o f protons. The relevant rate for the observed conductivity is that o f vehicle
molecular diffusion throughout the medium. A second conductivity principle
involves the transfer o f protons from one vehicle to another within hydrogen bonds.
116
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Additional reorganization o f the proton environment then results in the formation o f
an uninterrupted trajectory for proton migration. This mechanism is frequently
termed the Grotthuss mechanism and relies on the relevant rates of proton transfer
and the reorganization o f its solvent environment.1 7 4
Proton conductivity has been widely studied in Nafion and PSSA grafted
systems (reviewed in sections 1.3 and 1.4) in order to determine the relative proton
transport mechanism(s). Zawodzinski et al. used ‘H NMR analysis to calculate FT
and H2O diffusion coefficients in Nafion®-! 17 at various degrees o f membrane
hydration.1 7 5 They concluded that at low membrane water content, FT and H 2O most
likely diffuse according to the vehicle mechanism driven by a concentration gradient
from anode to cathode. However the FT and H 2O diffusion coefficients differ greatly
with increasing membrane water content. From this behavior Zawodzinski et al.
determined that the transport o f FT- ions by a Grotthus mechanism probably becomes
significant with high membrane water content. This is presumably due to the
increasing presence o f more bulk-like water that allows for effective reorganization
o f the solvent environment.
Ostrovskii et al. studied the state o f water within PVDF-g-PSSA grafted
membranes using both porous and non-porous PVDF precursors.1 7 6 They reported
that significant amounts o f water within the non-porous PVDF-g-PSSA grafted
membranes are confined to the PVDF backbone and differ from that o f bulk water.
This was further supported by the use o f Raman spectroscopy to calculate the
number o f water molecules per sulfonate group in both the porous and non-porous
117
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
samples. Membranes prepared using non-porous precursors had -2 0 water
molecules per sulfonate group whereby porous precursors had — 60 water molecules
per sulfonate group. This determination supported the increased water content o f the
porous membranes as compared to the non-porous membranes. However, increasing
the degree o f grafting ultimately resulted in a distinct, 2-phase morphology with
water increasingly bulk like in nature, especially within the porous precursors. This
is analogous to the increased presence o f bulk like water in Nafion® with decreasing
equivalent weight.
The results o f the water management experiments conducted on our PVDF-
PSSA membranes suggest that the transport mechanism(s) are fairly unique as
compared to Nafion® and PSSA grafted systems. The semz-sequential-IPN
morphology o f the PVDF-PSSA membranes likely minimizes the formation of
aqueous channels, thus explaining the reduced water and methanol permeability
rates, and may be attributed to the use o f a non-porous precursor confining a
significant portion o f the water along the PVDF backbone. Therefore, despite
membrane water content similar to that o f Nafion®-117 at specific PSSA uptake
values the lack of bulk like water may explain the reduced permeability rates
measured.
The ability to maintain adequate proton conductivity in PVDF-PSSA
membranes may be attributed to the number and distribution o f charge carriers
within the membrane. The reduced water diffusion coefficients determined from ‘H
NMR analysis suggests that these charge carriers may not aggregate thereby
118
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
regulating proton transport via the vehicle mechanism since bulk like water content
is thought to be low. Therefore the vehicle (HsO*) would be forced to follow the
same diffusive pathway as unprotonated vehicles (H2O). However, membrane water
content and proton conductivity values similar to Nafion®-117 suggest that Grottuss
transport may contribute to proton migration since this is thought to be the diffusive
mechanism in Nafion. If this were the case we would see higher water and methanol
permeability rates in our membranes but instead they are consistently lower than
Nafion® at all PSSA uptake levels. This situation could be resolved by determining
the diffusion coefficient for H+ as compared to water and evaluating any
discrepancies. We have not yet conducted these experiments.
The ability to prepare PVDF-PSSA membranes with adequate proton
conductivity and reduced methanol permeability was and continues to be a
significant challenge. However, understanding the dynamics of transport
phenomenon is crucial to optimize membrane parameters relating to PSSA uptake
and crosslinking density. The development o f membranes exhibiting favorable
characteristics ultimately led to the fabrication o f membrane electrode assemblies for
further electrical fuel cell testing.
119
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
2.2 References
151. S m art, M., "Chemical and Electrochemical Oxidation o f Small Organic
Molecules", UMI Ann Arbor, MI, 1998.
152. Encyclopedia of Polymer Science and Engineering, 17, 2n d Ed., p. 532, John
Wiley & Sons, Inc., 1989.
153. Benedetti, E., Catanorchi, S., D'Alessio, A., Moggi, G., Vergamini, P.,
Pracella, M., Ciardelli, F., Polymer International 41, 35-41 (1996).
154. Odian, G., "Principles o f Polymerization", 3rd Edition, Wiley-Interscience,
1991.
155. Tazaki, M., Wada, R., Okabe, M., Homma, T., Kobunshi Ronbunshu, 50, 533
(1993).
156. Rosenberg, Y., Sigmann, A., Narkis, M., Shkolnik, S., J. Appl. Polym. Sci.
43, 535(1991).
157. Lee, J.H., Kim, S.C., Macromolecules 19, 644-648 (1986).
158. Stein, R.S., Khambatta, F.B., Warner, F.P., Russell, T., Escala, A., Balizer,
E., J. Polym. Sci., Polym. Symp. 63, 313 (1978).
159. Liu, Z., Marechal, P., Jerome, R., Polymer 38, 19 (1997), 4925-4929.
160. Discussions with Elf Atochem, N.A.
161. Olabisi, O., Robeson, L.M., Shaw, M.T., in "Polymer-Polymer Miscibility",
Chpt. 3, Academic, New York (1979).
162. Paul, D.R., Newman, S., in "Polymer Blends", Chpt. 5, Academic, New York
(1978).
163. Anastasiadis, S.H., Gancarz, I., Koberstein, J.T., Macromolecules 22, 1449-
1453 (1989).
164. Del Rio, C., Acosta, J.L., Eur. Polym. J. 32, 7 (1996) 913-917.
165. Papavoine, C.H.M., Maas, W.E.J.R-, Veeman, W.S., Buning, G.H.W.,
Vankan, J.M.J., Macromolecules 26, 6611-6616 (1993).
166. Hietala, S., Koel, M., Skou, E., Elomaa, M., Sundholm, F., J. Mater. Chem.
8(5), 1127-1132(1998). 120
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
167. Buchi, F.N., Gupta, B., Haas, O., Scherer, G.G., J. Electrochem. Soc. 142,
3044 (1995).
168. Buchi, F.N., Gupta, B., Haas, O., Scherer, G.G., Electrochim. Acta 40, 345
(1995).
169. Bouchet, R-, Siebert, E., Solid State Ionics 118, 287-299 (1999).
170. Flint, S.D., Slade, R.C.T., Solid State Ionics 97, 299-307 (1997).
171. Greenbaum, S.G., Work in Progress.
172. Kreuer, K.D., Chem. Mater. 8, 610-641 (1996).
173. Kreuer, K.D., Weppner, W., Rabenau, A., Angew. Chem., Int. Ed Engl. 21,
208 (1982).
174. Grotthuss, C.J.D., Ann. Chim. 58, 54 (1806).
175. Zawodzinski, T.A., Derouin, C., Radzinski, S., Sherman, R.J., Smith, V.T.,
Springer, T.E., Gottesfeld, S., J. Electrochem. Soc. 140, 1041 (1993).
176. Ostrovskii, D.I., Torell, L.M., Paronen, M., Hietala, S., Sundholm, F., Solid
State Ionics 97, (1997) 315-321.
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Chapter 3: Results and Discussion
PVDF-PSSA Polymer Electrolyte Membranes:
Electrical Performance
3.1 Membrane Electrode Assembly: Fabrication Method I
The fabrication o f membrane electrode assemblies continues to evolve with
developments in direct methanol fuel cells. The desire to reduce catalyst loading(s)
thus improving the cost efficiency o f fuel cells has been a driving force in this aspect
of the technology. One early approach used in the development of Nafion®-based
MEAs for H 2/O2 fuel cells involved hot pressing a mixture o f Pt powder and PTFE
on both sides o f a proton exchange membrane.1 7 7 This method resulted in H2/O2 fuel
cells capable o f high power densities at specific operating conditions. Significant
improvements in MEA fabrication and the corresponding cell electrical performance
have since been made by the incorporation o f Nafion®-H ionomer into the catalyst
mix either replacing or in addition to PTFE. This catalyst mix is then applied to gas
diffusion electrodes and hot-pressed onto the membrane resulting in enhanced cell
performance with Pt loadings as low as 0.35 mg/cm2.1 7 8
Gottesfeld and co-workers have developed a similar technique capable of
applying "thin-film" catalyst layers directly to the membrane surface.1 7 9 By virtue of
their thinness and the high ionomer contents achievable with these catalyst layers,
high catalyst utilization is obtained since the continuity and integrity of the catalyst
layer-membrane interface is greatly improved. This application technique has
frequently been referred to as the "decal" process whereby the catalyst ink is cast
122
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
onto Teflon blanks for transfer to the membrane by hot-pressing and has since been
used in the testing o f Nafion®-based direct methanol fuel cells. 1 8 0 ‘1 8 2 Membrane
electrode assemblies were prepared by the application o f anode (Pt-Ru) and cathode
(Pt) inks to Teflon blanks that were subsequently hot-pressed to the membrane at 125
°C and 105 atm for 120s. The resulting MEAs had catalyst loading(s) o f
approximately 2.2 mg/cm2 and were inserted into fuel cell hardware and sandwiched
between two 5.5 cm2 uncatalyzed, carbon-cloth gas diffusion backings from E-Tek.
MEAs prepared using this technique exhibited high cell voltages (0.5 V) at a high
corresponding current density (670 mA/cm2) under operating conditions o f high
temperature with pressurized oxygen at the cathode.
Much o f the early electrical testing o f PVDF-PSSA membranes (prepared by
M. Smart and Q. Wang) was based on MEA fabrication procedures developed by our
research team.1 8 3 This process involved hot-pressing catalyzed, Teflon-impregnated
porous carbon electrodes to a catalyzed membrane surface at 112 °C and 1000-3000
psi for 15 minutes. The preparation o f catalyzed substrates was achieved by
preparing an ink consisting of anode (Pt-Ru) or cathode (Pt) catalyst, Nafion®-H
ionomer and an aqueous solution o f PTFE. The ink was painted onto the Teflon-
impregnated electrodes and onto the membrane surface. This process was used to
prepare r'x l" (6.45 cm2) and 2”x2" (25 cm2) MEAs with an anode/cathode catalyst
loading o f ~4 mg/cm2.
123
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
3.1.1 Electrical Performance o f PVDF-PSSA Membranes: 1 "xl" DMFCs
The l"xl" PVDF-PSSA membrane electrode assembly was inserted into the
necessary cell hardware and evaluated for electrical performance. The cell was
tested with 1.0 M methanol at 83 °C utilizing pressurized oxygen at the cathode
vented at various flow rates. As illustrated in figure 3.1, the l"x l" cell exhibited
moderate cell voltages at medium range current densities and also had excellent
OCV methanol crossover rates o f 18 mA/cm2 at 60 °C and 40 rnA/cm2 at 83 °C,
respectively. This compares quite favorably to Nafion®-117 OCV methanol
crossover values o f 135 mA/cm2 at 60 °C and 300 mA/cm2 at 90 °C representing a
significant reduction. The electrical performance of the 1 "x l" cell exhibited a strong
dependence on the flow rate o f oxygen vented from the pressurized cathode, as
shown in figure 3.1. The best overall cell performance was observed under
conditions of high oxygen pressure with low flow rate as indicated by optimum cell
performance at 30 PSIG and 0.0 L/min oxygen. In this mode the cell is considered
‘dead-ended’ whereby a minimal oxygen flow rate (< 0.1 L/min) is passed across the
cathode.
The ability o f the cell to operate at these cathode conditions is very unique
considering the inability o f Nafion®-117 based systems to operate in such a manner
since the high water permeability rates via diffusion and osmotic drag flood the
cathode, requiring high flow rates to alleviate this condition. The decreased water
permeability rates within the PVDF-PSSA membranes (supported by the NMR
analysis discussed previously in section 2.1.4.3.1) restrict cathode water content thus
124
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
avoiding the flooding phenomenon- In fact, increasing the cathode flow rate while
maintaining pressurization resulted in reduced cell performance and increased
resistivity attributed to a drying effect at the cathode caused by water removal not
compensated for by water flux from diffusion or osmotic drag. The lowest OCV cell
resistance measured was 180 mOhms observed during dead-ended testing at 83 °C
and increased with oxygen flow rate. Incidentally this value is substantially higher
than that of Nafion®-l 17 (6 mOhms) at similar test temperatures and was attributed
to the poor interfacial contact between the membrane surface and the catalyzed
electrodes as a direct result of the conditions used during MEA fabrication. This
discrepancy in resistivity contributed to the significant reduction in electrical
performance as compared to Nafion®-! 17.
125
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.1 Performance o f a l"x l" PVDF-PSSA MEA in a direct methanol fuel
cell with 1.0 M methanol at 83 °C utilizing pressurized oxygen at the cathode vented
at various flow rates.
Evaluation of PS/PVdF Membranes
in Direct Methanol Fuel Cells
1.0
0 3
0.7-
2 0 p s ig O x y - 2 .2 U m n . F la w R a t*
2 0 p & g O xy. - 0 l3 L . ‘m in. R o w R ata
2 0 psig O xy. - Q. '< L 'm n F lo w R a t*
2 0 p sig O xy. • 0 .0 0 3 U 'irin . F lo w H ata *
30 ps*g O x y - QO L/n-un. B o w Rat• '
a t
0 3 -
0.4-
oj -
0.0
100 250 300 150 0 200 5 0
Current Density (mA/cmA 2)
The initial electrical results using the PVDF-PSSA membranes were very
promising considering this was a novel system. The significant decrease in methanol
crossover as compared to Nafion®-117 also suggests higher fuel cell efficiencies are
attainable. However practical, portable power applications require a system capable
o f operation utilizing low flow rate ambient air, not pressurized oxygen since the use
o f pressure and flow rate increase parasitic losses drawn from the system. Also the
ability to operate at temperatures lower than 83 °C reduce the need for large
126
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
start-up heaters that further decrease system efficiency. Therefore modifications to
the MEA fabrication procedure, in addition to membrane optimization according to
PSSA uptake and morphology, was seen as an effective way to improve electrical
performance.
3.2 Membrane Electrode Assembly: Fabrication Method II
The ability to improve the melt flow characteristics o f the membrane, without
polymer degradation, is vital to establish intimate contact at the catalyst layer-
membrane interface. Therefore the two most important components involved in
MEA fabrication are processing temperature and pressure. The processing pressure
is a fixed variable since the electrode papers have a predetermined pressure above
that the papers will crush. The initial MEA fabrication process previously discussed
(method I) operated within that threshold (1000-3000 PSI) depending on the size o f
the electrode papers (l"x l" or 2”x2”). However, the hot-press temperature and its
impact on interfacial contact was an area of intense investigation. It was generally
agreed upon that increasing MEA fabrication hot-press temperatures would have an
advantageous affect on interfacial bonding at the catalyst layer-membrane and
electrode-electrolyte interfeces. As we improved membrane preparation in regards
to PSSA uptake and surface morphology we prepared two membranes for MEA
fabrication that exhibited favorable conductivity as compared to Nation®-117, as
shown in table 3.1. It is important to note that at the time these membranes were
prepared our methanol diffusivity experiment was not yet setup so proton
conductivity was our major screening criterion.
127
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Table 3.1 PSSA uptake and related properties o f membranes 97-30 and 97-33 as
compared to the same properties in Nafion®-! 17.
Membrane PSSA Uptake Crosslinking Density Thickness Proton Conductivity
Nafion-117 N/A N/A 7 mils 60/65 mS/cm
97-30 33% 6% 10 mils 120/127 mS/cm
97-33 24% 10% 10 mils 57/60 mS/cm
The MEAs were prepared using a procedure similar to that of our initial
P x l" MEA discussed in section 3.1. However, the pressing temperature was raised
from 112 °C to 150 °C. The anode/cathode catalyst ink was prepared and applied to
the graphite electrodes and to the membrane surface resulting in a catalyst loading in
the range o f 8-12 mg/cm2 at both electrodes. Once fabricated the MEAs were
assembled into the cell hardware and evaluated for electrical performance.
3.2.1 Electrical Performance o f PVDF-PSSA Membranes: 2"x2" DMFCs
The temperature modification to the MEA fabrication process had an
immediate impact on the two MEAs since the measured resistivity was considerably
lower for both samples than that reported for our initial l"xl" cell tests discussed in
section 3.1.1. MEA 97-30 had an OCV resistance o f 17 mOhms at 60 °C while
MEA 97-33 had an OCV resistance o f 20 mOhms at the same temperature. These
measurements were still considerably higher than that o f Nafion-117 (9 mOhms)
despite the similar conductivity values reported in table 3.1 and suggested that the
increased membrane thickness (10 mils) could have contributed to the disparity in
128
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
resistivity values. Membrane thickness and its impact on resistance, methanol
crossover and cell performance will be evaluated later in section 3.3.1.4.
Incidentally the measured OCV methanol crossover rates for the two MEAs
at 60 °C with 1.0 M methanol were considerably lower than that of Nafion®-117
(135 mA/cm2) yet higher than that o f our initial 1 "x l" MEA (18 mA/cm2). MEA 97-
30 had a parasitic loss o f 60 mA/cm2 at 60 °C while MEA 97-33 had a loss o f 50
mA/cm2, also at 60 °C. The importance of methanol crossover and the overall
impact on cell performance, fuel utilization and fuel cell efficiency will be addressed
further in sections 3.3.1.1 and 3.3.1.2.
The electrical behavior o f the two MEAs was evaluated with particular
interest paid to ambient performance. The results in figure 3.2 indicate that the cell
performance o f MEA 97-33 far exceeds that of MEA 97-30 achieving a cell voltage
o f 0.344 V at 160 mA/cm2 at 60 °C with 1.0 M methanol utilizing 0.10 L/min
ambient oxygen at the cathode. This represented a significant improvement as
compared to our initial r ’xl" results (figure 3.1) yet were still inferior to Nafion -
117. Incidentally it was observed from post-testing analysis that the graphite
electrodes o f MEA 97-30 had delaminated in some spots and shredded in others
explaining the poor performance. It was determined that excessive membrane
swelling (as a result o f the high PSSA uptake and low cross-linking density) had
destroyed the integrity o f the papers. It was for this reason that further testing o f
MEA 97-30 was limited.
129
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.2 Performance o f MEAs 97-30 and 97-33 in direct methanol fuel cells
at 60 °C with 1.0 M methanol utilizing ambient oxygen at the cathode.
Electrical Performance
0.8
use PVDF-PSSA MEAs 97-30 & 97-33
2"x2" MEA (25 cm2 )
1.0 M Methanol, 60 °C, Ambient Oxygen
0 .7
0.6
0 .5
u
oa
I 0 4
>
s
0 .3
MEA 97-33:0.10 L/nan AnfcetS Oxygen
MEA 97-30:0.10 L /n n A irb c rt Oxygen
0.2
0.1
330 360 380 270 300 30 6 0 90 120 150 180 210 240 0
Current Density (mA/cm2 )
Further analysis o f the cell performance o f MEA 97-33 illustrated in figure
3.3 indicates that there is a strong dependence on the flow rate of oxygen supplied to
the cathode, a phenomenon previously discussed in figure 3.1. The cell exhibited a
cell voltage o f 0.344 V at 160 mA/cm2 utilizing 0.10 L/min ambient oxygen that
decreased considerably with increasing flow rate. Moreover, the corresponding cell
resistivity increased with flow rate and was attributed to cathode drying as a result of
water removal. This effect was especially severe at low current densities where
water flux via osmotic drag is minimal as a result o f reduced anode catalyst
utilization leaving simple diffusion o f water as the only method available to hydrate
130
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
the cathode. However, the impact o f flow rate on cell resistance decreased at high
current densities since anode catalyst utilization is greater thus increasing the
osmotic drag associated with proton transport through the bulk.
Figure 3.3 Effect o f ambient oxygen flow rates on the performance o f MEA 97-
33 in a direct methanol fuel cell at 60 °C with 1.0 M methanoL
0.8
0.7
0.6
0.5
£
j “
3
0.3
0.2
0.1
0
0 30 60 90 120 150 180 210 240 270 300 330 360 390
Current Density (mA/cmJ)
Despite the ability to perform utilizing ambient cathode flow rates, optimum
cell performance of MEA 97-33 was observed utilizing pressurized oxygen at the
cathode and was attributed to the increase in oxygen solubility associated with
pressure. The data in figure 3.4 indicates that MEA 97-33 achieved a voltage of
0.406 V at 160 mA/cm2 at 20 PSIG with 0.00 L/min oxygen, a benchmark that
decreased with increasing cathode flow rates. The dependence o f cell performance
on flow rate even under pressurized conditions was similar to the performance trends
exhibited by our initial l"x l" MEA, discussed in section 3.1.1. This suggested
131
Electrical Performance
PVDF-PSSA MEA 97-33
2"x2" MEA (25 cm2 )
1.0 M M ethanol, 60 °C, Ambient Oxygen
0.10 L/min Ambient Oxygen
O.SO L/min Ambient Oxygen
1.00 L/min Ambient Oxygen
1.50 L/min Ambient Oxygen
2.00 L/min Ambient Oxygen
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
that despite the reduced cell resistivity o f MEA 97-33 as compared to our earlier
experiments, adequate cathode hydration continued to be a problem.
Figure 3.4 Effect o f pressurized oxygen and vented flow rate on the performance
o f MEA 97-33 in a direct methanol fuel cell at 60 °C with 1.0 M methanol.
Electrical Performance
0.8
u s e PVDF-PSSA M EA 97-33
2"x2" MEA (25 cm2 )
1.0 M M ethanol, 60 °C, 20 PSIG Oxygen
0.7
0.8
20 PSIG Oxygen, 0.00 L/min 0.5
>
20 PSIG Oxygen, 1.00 L/min
0
1 04
o
>
X
20 PSIG Oxygen, 3.00 L/min
w
0.3
0.2
0.1
360 390 90 120 150 240 270 300 330 420 0 30 60 180 210
Current Density (mA/cm2 )
In order to better understand the impact o f MEA fabrication modifications on
cell performance we designed an experiment that could measure the individual
potential o f the two electrodes. The anode polarization experiment effectively
separates the cell into two half-cell components utilizing the cathode as a pseudo-
reference, or normal hydrogen electrode. The anode voltage is then measured upon
cell polarization and corrected for contributions from the bulk (grain-boundary)
resistance to give a true representation o f the anode performance, as illustrated in
figure 3.5. The data indicates that despite the low cell resistance and improved
132
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
electrical performance o f MEA 97-33 as compared to our initial r 'x l ” MEA, the
anode suffers polarization losses much greater than that o f Nafion®-l 17. The
decrease in anode performance as compared to Nafion®-117 may be explained by
poor catalyst utilization as a result o f poor electrode-electrolyte interfacial contact,
despite the modified MEA fabrication procedure used (method II).
Figure 3.5 IR corrected anode performance o f MEA 97-33 in a direct methanol
fuel cell as compared to Nafion®-! 17 at 60 °C with 1.0 M methanoL
Anode Perform ance
0 .6
use PVDF-PSSA M EA 97-33 vs. Nafion-117
2 "x 2 " M EA (25 cm2)
1.0 M M ethanol, 60 °C
0.5
0.4
MEA 97-33: Anode Performance
Nafion: Anode Performance
0.1
1 0 1 0 0 1000 1
Current Density (mA/cm2 )
133
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Based on the overall cell and anode performance, the cathode potential can
then be calculated by rearranging the following equation:
Ecdl — Eeathode " Eanode
Eeathode — EcellE an ode
Once the respective electrodes were corrected for iR contributions the
cathode potentials were calculated and plotted vs. current density, as shown in figure
3.6. The variations in the cathode performance o f MEA 97-33 utilizing ambient vs.
pressurized oxygen illustrates the impact o f pressure and flow rate on cathode
performance. Optimum cathode performance was achieved utilizing pressurized
oxygen since the concentration of reactant at the catalyst layer is increased while the
minimal flow associated with dead-ended testing allowed for adequate hydration
thus maintaining reduction kinetics. The relatively linear slope o f the polarization
plot suggests oxygen content is sufficient and maintained even at high current
densities where mass transfer of oxygen becomes a limiting issue. The downward
arc o f the ambient cathode potentials at high current densities supported this since
the reduction in reactant solubility contributes to polarization losses not seen with
pressure.
134
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The ambient cathode polarization losses also increased slightly with
increasing flow rate up to current densities o f 100 mA/cm2 due to the inhibitory
impact on cell resistance caused by uncompensated water removal. However, at
higher current densities the cathode performance benefits from increased flow rate
since water content is greater due to increased osmotic drag and effectively removed
exposing sites for oxygen diffusion. The benefit o f flow rate was only realized up to
1.00 L/min beyond that the cathode exhibited severe drying and potentials dropped
accordingly.
Figure 3.6 IR corrected cathode performance o f MEA 97-33 in a direct methanol
fuel cell at 60 °C with 1.0 M methanol utilizing ambient and pressurized oxygen.
Cathode Perform ance
1 .2
use PVDF-PSSA MEA 97-33
2"x2” MEA (25 cm2)
1.0 M M ethanol, 60 °C, Oxygen
0 .8
z
5
>
20 PSIG. 0.00 L/min Oxygen
m
0.10 L/min Ambient Oxygen
o
a.
0.4 0 50 L/irnn Ambient Oxygen
1.00 L/min Ambient Oxygen
1.50 L/min Ambient Oxygen 0 .2
1000 1 0 0 1 1 0
Current Density (mA/cm2 )
135
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Once cell performance had been characterized utilizing oxygen at the cathode
we evaluated MEA 97-33 supplied with ambient air. As illustrated in figure 3.7, the
cell exhibited moderate cell voltages at relatively high current densities reaching a
voltage o f 0.318 V at 100 mA/cm2 at 60 °C with 1.0 M methanol utilizing 0.10 L/min
ambient air. Cell performance was unaffected by flow rates as high as 0.50 L/min
ambient air but dropped considerably above 1.00 L/min flow. This is presumably
due to the effective removal o f water from the cathode making available more sites
for oxygen diffusion that reached a limit at 0.50 L/min, above that the cathode began
to dry out and resistivity increased.
Figure 3.7 Effect o f ambient air flow rates on the performance o f MEA 97-33 in
a direct methanol fuel cell at 60 °C with 1.0 M methanol.
Electrical Performance
0.7
use PVDF-PSSA M EA 97-33
2"x2" M EA (25 cm2 )
1.0 M M ethanol, 60 °C, A m bient Air
0 .6
0.5 -
0.10 L min Ambient Air
0.50 L/min Ambient Air
1.00 L/min Ambient Air
1.S0 L/min Ambient Air
2.00 L/min Ambient Air
0 30 60 90 1 2 0 150 2 1 0 180 240
Current Density (mA/cm2 )
136
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The cathode performance o f MEA 97-33 at 60 °C with 1.0 M methanol
utilizing ambient air at various flow rates was then calculated and plotted in figure
3.8. The cathode performance utilizing ambient oxygen at 0.10 L/min has also been
included as a reference. The significant polarization losses evident at the cathode
utilizing ambient air as compared to oxygen is a result o f the mass transfer
limitations o f reactant to the active catalyst layer due to the reduced concentration o f
oxygen in air. Increasing cathode flow rate in order to remove product water and
thus make available more sites for oxygen diffusion is limited at low current
densities since cathode water content is low. However, cathode potentials improve
slightly with flow rate at medium to high current densities (> 100 mA/cm2) since
water content from diffusion and osmotic drag is higher and needs to be removed in
order to maintain oxygen diffusion. Incidentally the cathode potentials utilizing
ambient air, regardless o f flow rate, have a distinct downward slope as compared to
the plot for ambient oxygen. This is attributed to the mass transfer limitations
associated with ambient air operation coupled with the inhibitory affect o f methanol
crossover, an issue to be discussed in section 3.3.1.1.
137
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.8 IR corrected cathode performance o f MEA 97-33 in a direct methanol
fuel cell at 60 °C with 1.0 M methanol utilizing ambient air as compared to ambient
oxygen.
Cathode Performance
1 2
use PVDF-PSSA MEA 97-33
2"x2" MEA (25 cm2)
1.0 M Methanol, 60 °C, A ir & Oxygen
0 8
Z
m
>
> 0 .6
e
u 0.10 L/min Ambient Oxygen
0.10 L/min Ambient Air
0.4
0.50 L/min Ambient Air
1.00 L/min Ambient Air
0 2
1.50 L/min Ambient Air
1 0 0 1000 1 0 1
Current Density (mA/cm2 )
The promising cell performance o f MEA 97-33 at low temperatures utilizing
low flow rate ambient cathode conditions, as compared to our initial l ”xl" MEA,
suggested that the modifications made to the MEA fabrication procedure benefited
cell performance. However, the disparity in cell resistivity and anode performance
as compared to Nafion®-117 suggests poor interfacial contact and thus poor catalyst
utilization. The inhibitory impact o f increasing cathode flow rates on cell resistivity
and electrical performance further supports this belief. One remedy to reduce the
grain-boundary resistance thus minimizing bulk polarization losses and maintaining
cathode hydration involved increasing the PSSA uptake value o f the membranes.
Unfortunately the physical degradative problems associated with MEA 97-30
138
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
and increased methanol crossover rates correlating to high PSSA uptake values
complicated this option. Therefore, we continued to investigate methods to improve
MEA fabrication thus enhancing cell performance.
3.3 Membrane Electrode Assembly: Fabrication Method in
The dramatic improvement in cell performance following our initial
modification (method II) to the original MEA fabrication procedure (method I)
prompted continued research into this area.1 8 4 The ability to further improve the
melt flow characteristics of the membrane would allow for intimate contact between
the applied catalyst and conductive moieties o f the membrane. We initiated an
exhaustive investigation into the thermal stability o f the membrane since hot-press
temperatures influence melt flow. The TGA analysis discussed in section 2.1.2.3.2
indicated that the membrane is stable at up to 220 °C before degradation begins.
Therefore, we proposed further increasing the hot-press temperature to 180 °C.
In addition to hot-press temperature we also investigated the use o f a co
solvent capable o f improving the melt flow characteristics o f the membrane. As
discussed in section 1.6.2.1, PVDF is known to swell and even dissolve in a number
o f polar aprotic solvents. In early MEA work performed by M. Smart dimethyl
acetamide was one such solvent tested in the fabrication o f MEAs.1 8 3 Dimethyl
acetamide is a solvent used to solution-cast films o f PVDF powder and is highly
soluble in water. This is a favorable characteristic since it can be readily dissolved
into our pre-existing catalyst ink mixture.
139
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
In conjunction with our evaluation o f MEA fabrication procedure(s)
membrane development was modified to optimize favorable characteristics. The
difficulties in MEA handling and increased methanol crossover rates associated with
high PSSA uptake samples o f low crosslinking density (MEA 97-30) prompted us to
regulate the PSSA uptake level o f potential membranes. An effort was made to
fabricate membranes with the minimum PSSA uptake necessary to establish
conductivity values comparable to Nafion-117 while minimizing methanol
permeability. One such membrane (98-35) was prepared by reducing the PSSA
uptake and establishing a standard crosslinking density o f 10% dvb. The PSSA
related properties were evaluated according to our screening techniques discussed
earlier in section 2.1.4 and shown in table 3.2. The membrane was then fabricated
into an MEA using our modified (method HI) procedure.
Table 3.2 PSSA uptake and related properties o f membrane 98-35 as compared
to the same properties in Nafion®-! 17.
Membrane PSSA
Uptake
Thickness Proton
Conductivity
CH3OH
Diffusion Coefficient
MEA
Thickness
Nafion-117 N/A 7 mils 60/65 mS/cm 3.60 E-07 cm2/s 20 mils
98-35 15% 13 mils 45/50 mS/cm 6.40 E-08 cm2/s 26 mils
The MEA was prepared using the same catalyst application technique
discussed earlier in sections 3.1 and 3.2 with the exception o f the addition o f DMA
to the ink composition. The membrane was also sanded to roughen the surface in
hopes o f promoting improved contact between the catalyst ink and membrane
140
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
surface. The anode and cathode catalyst ink was applied to the electrode papers and
membrane surface resulting in a loading o f 8-12 mg/cm2 at both electrodes. The hot-
press temperature was then raised from 150 °C to 180-185 °C and then assembled
into the cell hardware and flushed with water for 48 hours to hydrate the membrane
and remove any residual DMA.
3.3.1 Electrical Performance o f PVDF-PSSA Membranes: 2"x2" DMFCs
The benefits o f the revised MEA fabrication methodology were not
immediately apparent in our evaluation o f MEA 98-35. The initial OCV resistance
measured at 60 °C was 23 mOhms, a value comparable to that o f MEA 97-33 (18
mOhms) and not unexpected since MEA 98-35 had substantially lower PSSA uptake
(15%) than that of MEA 97-33 (24%). However, the initial cell performance of
MEA 98-35 exhibited erratic cell voltages characterized by large polarization losses
at medium current densities (100-200 mA/cm2) that improved slowly with time. The
cell originally achieved a potential o f 0.370 V at 160 mA/cm2 at 60 °C with 1.0 M
methanol utilizing 0.10 L/min ambient oxygen, as illustrated in figure 3.9. However,
it was very difficult to maintain steady state resistivity values and/or cell voltage(s).
Therefore prior to continued cell testing at 60 °C, MEA 98-35 was heated to
90 °C with pressurized oxygen at the cathode and slowly polarized to higher current
densities. This method o f testing conditioned the MEA since the resistance slowly
decreased while the cell voltage slowly increased all while holding the cell steady at
a specific current density. Once steady state resistance and voltage values were
achieved the current density was stepped up and the process repeated. The ability to
141
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
extend the cell to high current densities allowed for adequate hydration o f all ionic
moieties due to the diffusion o f water associated with osmotic drag. Cell
conditioning was repeated over and over until the cell reached a steady state OCV
resistance (12 mOhms at 90 °C) with reproducible cell voltages.
Due to the beneficial effect o f pressurized conditioning at 90 °C, MEA 98-35
was re-evaluated at 60 °C and immediately exhibited a lower OCV resistance (17
mOhms) as compared to the pre-conditioned value o f 23 mOhms. The measured
resistivity was still high as compared to Nafion®-117 (9 mOhms) despite the similar
conductivity values reported in table 3.2 prior to MEA fabrication. This once again
suggested that membrane thickness could have contributed to the discrepancy and
was further supported by the reduced methanol diffusion coefficients reported in
table 3.2 that corresponded to an OCV methanol crossover loss = 30 mA/cm2 at 60
°C with 1.0 M methanol. This represented a substantial reduction in crossover as
compared to MEA 97-33 (50 mA/cm2) and Nafion®-117 (135 mA/cm2). The results
in figure 3.9 also indicate an improved and stable post-conditioned cell voltage of
0.397 V at 160 mA/cm2 at 60 °C with 1.0 M methanol utilizing 0.10 L/min ambient
oxygen, a significant improvement as compared to the pre-conditioned cell.
142
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.9 Effect o f cell conditioning (90 °C, pressurized oxygen) on the
performance o f MEA 98-35 in a direct methanol fuel cell as compared to the pre
conditioned MEA 98-35 and MEA 97-33 at 60 °C with 1.0 M methanol.
Electrical Perform ance
0 .8
u s e PVDF-PSSA M EA 97-33 & 98-35
2"x2" M EA (25 cm2 )
1.0 M M ethanol, 60 °C, Oxygen
0.7
0 .6
MEA 9£-35 Post Cond: 0.10 L/min Ambient Oxygen
0.5
MEA 98-35 Pre-Co nd: 0.10 L/min Ambient Oxygen
u
I 04
o
>
MEA 97-33: 0.10 L/min Ambient Oxygen
0.3
0 .2
240 270 300 330 360 390 1 2 0 150 180 2 1 0 0 30 60 90
Current Density (mA/cm2 )
The data in figure 3.9 also indicates that the conditioned cell performance of
MEA 98-35 was improved at most current densities as compared to our earlier test
results from MEA 97-33, fabricated using method II, at similar test conditions.
However, at medium current densities (> 240 mA/cm2 ) the performance o f MEA 97-
33 is superior to that o f MEA 98-35. This was surprising since polarization losses at
these current densities are usually attributed to poor bulk resistivity limiting proton
flux thus starving and drying the cathode. This does not coincide with the measured
resistivity values for MEA 97-33 that are higher than MEA 98-35 at all current
densities, despite increased PSSA uptake. Therefore the resistive losses of MEA 97-
143
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
33 were not necessarily grain-boundary related but instead resulted from poor
interfacial contact at the electrode-electrolyte interface. This suggests that the
modified MEA fabrication procedure (method HI) resulted in improved interfacial
contact in MEA 98-35 explaining the low resistivity values measured despite the low
PSSA uptake as compared to MEA 97-33.
The cell performance o f MEA 98-35 was also evaluated at 60 °C with
ambient oxygen supplied to the cathode at various flow rates, as illustrated in figure
3.10. Despite the benefits of cell conditioning, cell performance still exhibited a
strong dependence on cathode flow rate dropping considerably with slight flow rate
increases while cell resistivity increased accordingly, especially at flow rates greater
than 0.5 L/min. The cell ultimately achieved a voltage o f 0.397 V at 160 mA/cm2 at
60 °C with 1.0 M methanol utilizing 0.10 L/min ambient oxygen.
Figure 3.10 Effect o f ambient oxygen flow rates on the performance o f MEA 98-
35 in a direct methanol fuel cell at 60 °C with 1.0 M methanol.
Electrical Performance
0.9
use PVDF-PSSA M EA 98-35
2"x2" MEA (25 cm2 )
1.0 M M ethanol, 60 °C, Ambient Oxygen
0 .8
0 7
♦ 0 10 L/min Ambient Oxygen
0.50 L/min Ambient Oxygen
— 1.00 L/min Ambient Oxygen
_ 0.4
0.3
0 .2
O .t
0 30 60 90 1 2 0 150 180 2 1 0 270 300 240
C u rren t Denaicy (mA/cm3)
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
MEA 98-35 was further tested at 60 °C utilizing pressurized oxygen at the
cathode, as shown in figure 3.11, and was able to reach high current densities and
preferred operation at minimum flow rates. The cell exhibited a voltage o f 0.457 V
at 160 mA/cm2 measured at 20 PSIG with 0.0 L/min oxygen representing a dramatic
improvement in cell performance as compared to similar testing o f MEA 97-33
(0.406 V at 100 mA/cm2). The MEA was able to operate at these dead-ended
conditions repeatedly with only slight purging necessary to remove product water.
This is supported by the slight increase in voltage with increasing flow rate observed
at current densities greater than 270 mA/cm2.
Figure 3.11 Effect o f pressurized oxygen and vented flow rate on the performance
o f MEA 98-35 in a direct methanol fuel cell at 60 °C with 1.0 M methanol.
Electrical Perform ance
0.9
use PVDF-PSSA MEA 98-35
2"x2" M EA (25 cm2 )
60 °C, 1.0 M M ethanol, 20 PSIG Oxygen
0 .8
0.7
0.6
>
I
o
>
X
u
0.0 L/min
0.3
0 .2
0 .1
0 160 2 1 0 30 1 2 0 180 300 330 60 90 240 270 360 390 420
Current Density (mA/cm1 )
145
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
MEA 98-35 was also tested at 90 °C based partly on the favorable impact
increasing temperature had on cell resistance and performance during conditioning.
However, the data in figure 3.12 indicates a drop in performance at 90 °C utilizing
ambient oxygen as compared to ambient oxygen test data measured at 60 °C (figure
3.10). This was surprising since temperature is known to benefit diffusion and
reaction kinetics in analogous systems using Nafion®. The strong dependence o f cell
performance on flow rate suggested that the cathode suffered from severe drying.
However the steady state OCV resistance at 90 °C was 12 mOhm, identical to that
observed during conditioning, and only increased with extremely high flow rates (>
1.0 L/min). This suggests that membrane water content is sufficient to minimize
bulk resistivity yet the increased vapor pressure o f water at 90 °C results in severe
evaporative losses at the cathode aggravated by flow rate.
Figure 3.12 Effect o f ambient oxygen flow rates on the performance of MEA 98-
35 in a direct methanol fuel cell at 90 °C with 1.0 M methanol.
Electrical Performance
0.8
use PVDF-PSSA MEA 98-35
2"x2” MEA (25 cm2 )
1.0 M Methanol, 90 °C, Ambient Oxygen
0.7
0.6
0.5
v
M — •—0.10 L/min Ambient Oxygen
- • —0.50 L/min Ambient Oxygen
0.2
0.1
0
0 30 90 180 240 60 120 150 210 270 330
Current Density (mA/cm3)
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Further cell testing o f MEA 98-35 at 90 °C utilizing pressurized oxygen at the
cathode resulted in a dramatic improvement in cell performance, as indicated in
figure 3.13. The cell exhibited a cell voltage o f 0.505 volts at 160 mA/cm2 at 20
PSIG 0.0 L/min oxygen (dead-ended) and was able to operate in such a manner for
over an hour before slight purging was necessary, consistent with our earlier
pressurized testing at 60 °C. The dramatic effects of cathode pressure, especially at
90 °C, suggested that the increased oxygen solubility associated with pressure and
the minimal flow rate of dead-ended operation allowed for hydration o f the catalyst
layer necessary to maintain reaction kinetics. The effects of temperature and
pressure on cathode performance will be addressed further in the analysis of figure
3.16.
Figure 3.13 Effect of pressurized oxygen and vented flow rate on the performance
o f MEA 98-35 in a direct methanol fuel cell at 90 °C with 1.0 M methanol.
Electrical Performance
use PVDF-PSSA MEA 98-35
2"x2" MEA (25 cm 2 )
1.0 M M ethanol, 90 °C, 20 PSIG Oxygen
0 9
0.7
20 PSIG Oxygen, 0.00 L/min
20 PSIG Oxygen, 0.10 Lm m
20 PSIG Oxygen, 0.50 L/min
03
0 2
0 30 60 90 120 150 180 210 240 270 300 330 360 390 420 450 480 510 540
Current Density (mA/cm2 )
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The reduction in cell resistivity and improved electrical performance o f MEA
98-35, as compared to our previous MEAs, suggested that our MEA fabrication
modifications (method IE) further enhanced interfacial contact at the electrode
electrolyte interface. Therefore the anode performance o f MEA 98-35 was evaluated
in order to quantitatively study the respective anode and cathode polarization
characteristics. The data in figure 3.14 indicates a dramatic improvement in the
anode performance o f MEA 98-35 at 60 °C with 1.0 M methanol as compared to
MEA 97-33 and comparable performance to Nafion®-117. Incidentally the anode
performance o f MEA 98-35 at 90 °C has also been included illustrating the beneficial
affect o f temperature on anode kinetics. The overall anode analysis at both 60 °C
and 90 °C suggests that the interfecial bonding, specifically at the anode, benefited
from our MEA fabrication adjustments (method HI).
148
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.14 IR corrected anode performance o f MEA 98-35 in a direct methanol
fuel cell as compared to MEA 97-33 and Nafion®-117 at 60 °C and 90 °C with 1.0 M
methanol.
Anode Performance
use PVDF-PSSA MEA 9S-35
2"x2" MEA (25 cm1 )
1.0 M MeOEL, 60 °C
0.8
U
= 0.8
MEA 97-33: Anode Performance 60 oC
— • — MEA 98-35: Anode Pcrfonitmcc 60 oC
— a— Nafion: Anode Performance 60 oC
MEA 98-35: Anode Performance 90 oC
0.2
1000 100 1 0 1
Current Density (mA/cm2)
Since the anode performance o f MEA 98-35 at 60 °C was favorable and even
consistent with that of Nafion®-117, it seemed evident that variations in cell
performance were a direct result o f cathode-related polarization losses. The cathode
performance of MEA 98-35 at 60 °C with 1.0 M methanol utilizing ambient oxygen
at various flow rates was calculated and plotted in figure 3.15. The data indicates
that the cathode potentials decrease with flow rate due to the inhibitory impact on
cell resistivity associated with water removal from the cathode. The downward slope
usually associated with mass transfer limitations is somewhat unusual since oxygen
supply is plentiful suggesting polarization losses may be attributed to inadequate
cathode hydration as a result o f bulk resistivity losses or poor oxygen solubility.
149
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.15 IR corrected cathode performance o f MEA 98-35 in a direct methanol
fuel cell at 60 °C with 1.0 M methanol utilizing ambient oxygen.
Cathode Performance
use PVDF-PSSA MEA 98-35,
2"x2" MEA (25 cm2)
1.0 M Methanol, 60 °C, Ambient Oxygen
0.8
Z 0 6
i
0.10 L/min Ambient Oxygen
>
0.50 IVmm Ambient Oxygen
1
c
a o.4
o
a.
1.00 L/mm Ambient Oxygen
0.2
1000 100 10 1
Current Density (tnA/cm3)
The dramatic improvement in the cell performance o f MEA 98-35 at 90 °C
utilizing pressurized oxygen as compared to ambient flow rates can be explained by
studying the beneficial affects o f pressure on cathode performance, as illustrated in
figure 3.16. The poor ambient performance suggests that evaporative water losses
severely inhibit reduction kinetics and are aggravated by increasing flow rate.
However, the ability to adequately hydrate the cathode under dead-ended conditions
alleviates cathode drying and thus improves reduction kinetics while the associated
increase in oxygen solubility as a result o f pressure increases availability o f reactant.
It is interesting to note that the cathode potentials do not significantly improve as
pressure is increased from 10-20 PSIG, suggesting that despite increased oxygen
solubility, kinetics are limited by utilization o f available catalyst.
150
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.16 IR corrected cathode performance o f MEA 98-35 in a direct methanol
fuel cell at 90 °C with 1.0 M methanol utilizing ambient and pressurized oxygen.
Cathode Perform ance
1
u se PVDF-PSSA M EA 98-35
2"x2" MEA (25 cm1 )
1.0 M M ethanol, 90 °C, Oxygen
0.8
0.6
20 PSIG, 0.00 L/mm
10 PSIG, 0 00 L/min
0.10 L/min Ambient Oxygen
0.50 L/mm Ambient Oxygen
0.4
0.2
0
1000 100 1 0 1
Current Density (mA/cm2 )
The promising performance of MEA 98-35 utilizing oxygen, especially at
low temperatures and ambient flow rates, was an important step on our road to
developing a practical DMFC system. However, the performance must be validated
utilizing ambient air since the use of oxygen places certain restrictions on an
operating system. Therefore we aggressively evaluated the ambient air performance
of MEA 98-35, as illustrated in figure 3.17. The cell achieved a voltage o f 0.406 V
at 100 mA/cm2 at 60 °C with 1.0 M methanol utilizing 0.10 L/min o f ambient air that
decreased with flow rate, further suggesting that cathode hydration is a sensitive
issue. However, the ambient air performance o f MEA 98-35 exceeded that o f
Nafion®-l 17 further highlighting the ability o f the membrane to operate at low flow
151
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
rates due to the lack o f accumulated water that blocks oxygen access in Nafion-based
systems. The significant reduction o f methanol crossover in the PVDF-PSSA MEA
also minimizes cathode polarization losses already complicated by air operation.
Figure 3.17 Effect o f ambient air flow rates on the performance o f MEA 98-35 in
a direct methanol fuel cell as compared to Nafion®-117, at 60 °C with 1.0 M
methanol.
Electrical Perform ance
0.8
use PVDF-PSSA M EA 98-35 & Nafion-117
2"x2" MEA (25 cm 2)
1.0 M M ethanol, 60 °C, A m bient Air
0.7
0.6 ■
MEA 98-35: 0.10 L/min Ambient Air
0.5
MEA 98-35: 0.50 L/min Ambient Air
I 0.4
Nafion-117: 0.10 L 'raia Ambient Air
0.3
0.2
0.1
0 30 60 150 180 120 210 240 90
Current Density (mA/cm2 )
Based on the anode polarization data of MEA 98-35 discussed previously in
figure 3.14, the cathode performance utilizing ambient air at 60 °C with 1.0 M
methanol was calculated and plotted in figure 3.18. The data indicates that the
cathode potentials decrease with increasing flow rate corresponding to non
compensated water removal drying the cathode. The downward slope o f the cathode
plot utilizing ambient air and the 85 mV reduction in the potential at 100
152
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
mA/cm2, as compared to utilizing oxygen, highlight the mass transfer issues
associated with the reduced concentration o f oxygen in air at the catalyst su rface.
Efforts to alleviate this situation by increasing flow rate in order to remove product
water, thus exposing sites for oxygen diffusion, reduced cell performance as a result
o f cathode drying.
Figure 3.18 IR corrected cathode performance o f MEA 98-35 in a direct methanol
fuel cell utilizing ambient air vs. ambient oxygen flow rates at 60 °C with 1.0 M
methanol.
Cathode Performance
1
u se PVDF-PSSA MEA 98-35,
2"x2" MEA (25 cm2 )
1.0 M M ethanol, 60 °C
Am bient Oxygen/A ir
08
0.6
60 oC: 0.10 L/min Ambient Oxygen
60 oC: 0.10 L/min Ambient Air
0.4
60 oC: 0.50 L/min Ambient Air
0.2
0
100 10 1 1000
Current Density (mA/cm2 )
153
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
3.3.1.1 Impact o f Methanol Concentration on Electrical Performance and Crossover
Rates
In our continued efforts to improve cell performance at low temperatures
utilizing low flow rate ambient air and oxygen we investigated the impact of
methanol concentration on cell performance and methanol crossover rates. During
the initial cell testing o f the l"xl" MEA reported in figure 3.1, the methanol
crossover rates measured at 1.0 M methanol were extremely low but reasonable
performance was only observed at high temperatures utilizing pressurized oxygen.
The development and testing o f MEAs 97-30 and 97-33 represented an improvement
in cell performance, specifically at low temperature ambient flow rate testing, but
corresponded to crossover rates only half that o f Nafion®-l 17 using 1.0 M methanol.
However MEA 98-35 exhibited superior ambient flow rate performance with
accompanying crossover rates only 25% that o f Nafion®-117 at similar test
conditions. The ability o f MEA 98-35 to achieve promising cell performance and
low methanol crossover rates with 1.0 M methanol prompted us to investigate the
effects o f methanol concentration on cell performance in hopes of improving cell
voltages while maintaining low methanol crossover rates.
The cell performance o f MEA 98-35 at methanol concentrations o f 1.0, 1.5
and 2.0 M was measured at 60 °C utilizing 0.10 L/min o f ambient oxygen, as
illustrated in figure 3.19. The data indicates that at low to medium current densities
(0-150 mA/cm2) 1.0 and 1.5 M methanol exhibited roughly the same performance
while 2.0 M methanol is consistently inferior by 20-25 mV throughout the course of
the experiment. However, at current densities greater than 150 mA/cm2 the
154
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
performance o f 1.5 and 2.0 M methanol is superior to that o f 1.0 M methanoL This
suggests that methanol concentration can benefit or inhibit cell performance
depending on the operating current density.
Figure 3.19 Effect o f methanol concentration on the performance o f MEA 98-35
in a direct methanol fuel cell at 60 °C utilizing ambient oxygen at the cathode.
Electrical Perform ance
0.8
use PVDF-PSSA MEA 98-35
2"x2" MEA (25 cm2 )
60 °C, 0.10 L/min Ambient Oxygen
0.7
0.6
0.5
1.0 M Methanol
= 0.4
1.5 M Methanol
2.0 M Methanol
0.3
0.2
0.1
50 100 200 250 300 350 450 0 150 400
Current Density (mA/cm2 )
Increasing methanol concentration has been shown to directly benefit anode
performance, specifically at high current densities where methanol mass transfer
limitations are alleviated. In order to verify this relationship we conducted anode
polarization experiments on MEA 98-35 using methanol concentrations o f 1.0, 1.5
and 2.0 M at 60 °C, as illustrated in figure 3.20. The data indicates that at current
densities up to 100 mA/cm2 the anode is slightly less polarized with 1.5 M methanol
155
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
than the other concentrations tested. The increase in polarization loss with 1.0 M
methanol at the same concentration range can be attributed to mass transfer limits of
methanol to the catalyst layer. However, the increase in polarization loss with 2.0 M
methanol suggests that the fuel is not consumed due to catalyst utilization limits at
such low overpotentials. The anode potentials begin to differentiate though beyond
100 mA/cm2 exhibiting optimum performance with 2.0 M methanol attributed to the
reduction in mass transfer limits of methanol to the now active catalyst layer. The
increase in anode performance with increasing concentration beyond 100 mA/cm2
correlates to the increased fuel cell performance at the same current density range
utilizing 1.5 and 2.0 M, as compared to 1.0 M, as reported previously in figure 3.19.
This suggests that beyond 100 mA/cm2 increasing concentration alleviates mass
transfer limitations at the anode thus sustaining methanol oxidation and improving
cell performance.
156
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.20 Effect o f methanol concentration on the ER corrected anode
performance o f MEA 98-35 in a direct methanol fuel cell at 60 °C.
Anode Performance
0.7
u s e PVDF-PSSA MEA 98-35
2"x2" MEA (25 cm2 )
Anode Performance 60 °C
0.6
0.5
Z
0.4
1.00 M Methanol
1.50 M Methanol
2.00 M Methanol
0.1
1000 100 1 0 i
Current Density (mA/cm2 )
We then calculated the cathode performance of MEA 98-35 utilizing 0.10
L/min ambient oxygen at 60 °C based on the cell and anode performance per
concentration tested. We decided to evaluate cathode performance utilizing oxygen
in order to reduce the mass transfer losses associated with air operation allowing us
to study the true impact o f methanol concentration on cathode potentials more
closely. The data in figure 3.21 indicates that cathode polarization losses at 60 °C
are similar for 1.00 and 1.50 M methanol yet is significantly increased utilizing 2.0
M methanol. This suggests that despite the beneficial impact on anode performance,
increasing concentration results in cathode polarization losses presumably due to the
associated increase in methanol crossover rates.
157
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.21 Effect o f methanol concentration on the IR corrected cathode
performance of MEA 98-35 in a direct methanol fuel cell at 60 °C utilizing ambient
oxygen.
Cathode Performance
1
u s e PVDF-PSSA MEA 98-35
2"x2" MEA (25 cm2 )
60 °C, 0.10 L/min Ambient Oxygen
0.8
1.00 M Methanol 0.6
I.SO M Methanol
0.4
0.2
0
1000 1 0 100 1
Current Density (mA/cm2 )
The reduction in cathode potentials with increasing concentration can be
further explained by measuring the methanol crossover rates per concentration
tested. The data in table 3.3 reports an increase in methanol crossover with
concentration, as would be expected. This indicates that the beneficial affects o f
increasing concentration on anode performance (figure 3.20) are compensated by the
inhibitory affect o f methanol crossover rates on cathode potentials (figure 3.21)
especially at concentrations greater than 1.5 M methanol.
158
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Table 3.3 Effect o f methanol concentration on the measured methanol crossover
rates of MEA 98-35 in a direct methanol fuel cell as compared to Nafion®-l 17 at 60
°C.
M eOH [ ] M EA 98-35
OCV X-Over Rates
Nafion-117
OCV X-Over Rates
1.0 M
1.5 M
2.0 M
33 mA/cm2
53 mA/cm2
63 mA/cm2
135 mA/cm2
158 mA/cm2
N/A
Our analysis o f MEA 98-35 utilizing ambient oxygen indicated that
increasing methanol concentration had a beneficial impact on cell performance up to
certain current densities. The improvement in anode performance with increasing
methanol concentration enhanced cell performance at high current densities (> 150
mA/cm2) where mass transfer issues at the anode are alleviated. However, this
benefit was only realized up to a concentration o f 1.50 M methanol since the
inhibitoiy crossover rates exhibited at 2.0 M methanol compromised any
improvement at the anode. Therefore the benefits o f concentration on cell
performance must also be weighed against the negative impact o f methanol
crossover on cathode potentials. Incidentally cell performances at low current
densities (< 150 mA/cm2) preferred reduced methanol concentration since anode
catalyst utilization at these current densities is poor and methanol crossover can have
a particularly inhibitory impact on the cathode, and thus cell performance.
159
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Fuel cell operation utilizing ambient air further compounds cathode
performance since mass transfer o f oxygen to the catalyst layer, coupled with
methanol crossover, becomes a significant issue. Therefore in our efforts to optimize
fuel cell performance we investigated the impact of reducing methanol concentration
below 1.0 M methanol anticipating a beneficial impact on the cathode. We re
evaluated fuel cell performance with 0.5, 0.75 and 1.0 M methanol at 60 °C utilizing
0.10 L/min ambient air, as illustrated in figure 3.22, and the effects o f reduced
concentration were immediately apparent. Cell performance at 0.5 M methanol
reached 0.428 V at 100 mA/cm2 as compared to 0.4 V using 0.75 or 1.0 M methanol
at the same operating current density. The fact that the cell exhibits optimum
performance with 0.5 M methanol utilizing ambient air is contrary to the data
observed utilizing ambient oxygen (figure 3.19) whereby optimum performance was
achieved with 1.5 M methanol. This suggests that the cathode utilizing ambient air
is highly sensitive to methanol crossover and cathode polarization losses outweigh
any potential gain at the anode. This is presumably due to the mass transfer
limitations o f oxygen in air coupled with the poisoning effect o f methanol crossover.
160
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.22 Impact o f methanol concentration on the performance of MEA 98-35
in a direct methanol fuel cell at 60 °C utilizing ambient air at the cathode.
Electrical Performance
u se PVDF-PSSA MEA 98-35
2"x2" MEA (25 cm2)
60 °C, 0.10 L/min Ambient A ir
0.9
0.8
0.7
0.50 M Methanol
0.75 M Methanol
1.00 M Methanol
j u
U 0.4
0.3
0.2
0.1
200 150 250 100 50 0
Current Density (mATtnC)
The favorable cell performance in figure 3.22 with 0.50 M methanol
suggested that the cathode benefited from the reduced methanol crossover rates
associated with decreasing concentration. The measured OCV methanol crossover
rates illustrated in table 3.4 further support this since the lowest parasitic losses were
measured at 0.50 M methanol, corresponding to the concentration o f optimum cell
performance. Unfortunately, we were unable to calculate cathode potentials since
anode polarization was only performed with 1.00 M methanol. However, the data in
table 3.4 qualitatively supports our hypothesis regarding crossover and cathode
performance utilizing ambient air.
161
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Table 3.4 Effect o f methanol concentration on the measured methanol crossover
rates o f MEA 98-35 in a direct methanol fuel cell at 60 °C as compared to NaJfion®-
117.
MeOH [ ] MEA 98-35
OCV X-Over Rates
Nafion-117
OCV X-Over Rates
0.50 M
0.75 M
l.OOM
17 mA/cm2
24 mA/cm2
33 mA/cm2
60 mA/cm2
N/A
135 mA/cm2
Our analysis of MEA 98-35 utilizing ambient air indicated that increasing
methanol concentration had a negative impact on cell performance at all current
densities. This was a direct result of the reduction in cathode potentials with
increasing concentration that offset any marginal improvements at the anode. This
phenomenon highlights the difficulties associated with ambient air operation since
methanol crossover, and its poisoning affect on the cathode catalyst, further
compounds the mass transfer limitations o f oxygen in air. However, the ambient air
cell performance with 0.50 M methanol was highly competitive to that of Nafion®-
117 with 25% the methanol crossover rate suggesting higher fuel and fuel cell
efficiencies are attainable.
3.3.1.2 Impact o f Methanol Crossover Rates on Fuel and Fuel Cell Efficiency
The most immediate consequence o f methanol crossover in direct methanol
fuel cells prior to the inhibitory impact on the cathode is the reduction in fuel
utilization at the anode. Therefore the reduced methanol crossover rales o f the
162
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
PVDF-PSSA membranes translate into increased fuel efficiency values, especially as
compared to Nafion®-based systems. The fuel efficiency o f an operating cell can be
calculated according to the following equation:
^Load
Fuel Efficiency % =
^Load + ^Crossover
The fuel efficiency of MEA 98-35 was determined from a variety of
methanol concentrations at 60 °C, as shown in figure 3.23. The data indicates that
the fuel efficiency values increase with decreasing molarity as a result of the
reduction in methanol crossover rates, trends previously reported in tables 3.3 and
3.4. MEA 98-35 exhibited a fuel efficiency o f 84.7% at 100 mA/cm2 utilizing 0.5 M
methanol representing a 15% increase as compared to the fuel efficiency o f Nafion®-
117 (73.8%) evaluated at the same concentration and current density. It is important
to note that the fuel efficiency values of MEA 98-35 may actually be slightly higher
than reported since methanol crossover rates are known to decrease with current
density as a result o f increased anode utilization thus realizing enhanced fuel
efficiency values. Unfortunately we were unable to measure crossover rates with
current density for our PVDF-PSSA MEAs since our CO 2 analyzer is only capable of
monitoring evolved CO2 in a gas stream of at least 2.0 L/min, an operating condition
unsuitable for our MEAs.
163
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.23 Effect o f methanol concentration on the fuel efficiency o f MEA 98-35
in a direct methanol fuel cell at 60 °C.
ICO
90
80
70
£ 6 0
8
2 30
E
t o !
Z
■ 1 4 0
30
2 0
1 0
0
0 50 100 150 200 250 300 350 400 450
Current Density (mA/cm*)
The data in figure 3.23 indicates that reducing methanol concentration has a
favorable impact on fuel utilization yet this only represents a part o f the overall
picture when evaluating efficiencies using the PVDF-PSSA membrane. The overall
cell performance as compared to the theoretical potential o f 1.24 V must be factored
into the equation in order to truly understand the impact o f methanol crossover on all
facets o f fiiel cell operation. The fuel cell efficiency o f an operating system factors
in both fuel efficiency and voltage efficiency according to the following equation:
Fuel Cell
Efficiency %
1.24 V
164
Operating Voltage
------------------------ X Fuel Eflficiency%
Fuel Efficieucy%
-0.50 M Methanol
- 0.75 M Methanol
-1.00 M Methanol
1.50 M Methanol
- 2.00 M Methanol
- Nafion-117:0.50 M Methanol
u s e PVDF-PSSA MEA 98-35
2 ”X 2 ” MEA (25 cm2 )
Fuel Efficiency % , 60 °C
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The fuel cell efficiency o f MEA 98-35 was calculated based on the fuel
utilization and cell performance with 1.0, 1.5 and 2.0 M methanol at 60 °C utilizing
0.10 L/min o f ambient oxygen, as illustrated in figure 3.24. The data indicates that a
fuel cell efficiency o f 30.3% is observed at 100 mA/cm2 with 1.0 M methanol that
decreases with increasing concentration up to a current density o f 200 mA/cm2. This
is due to the decrease in fuel utilization with increasing concentration, and thus
methanol crossover, that effectively limits any benefit concentration may have on the
voltage efficiency. However, at current densities beyond 200 mA/cm2 the fiiel cell
efficiency with 1.5 and 2.0 M methanol is superior to that o f 1.0 M methanol and is
attributed to the increase in voltage efficiency compensating losses in fuel utilization.
This phenomenon is only seen at medium to high current densities where methanol
crossover is reduced on account o f the increase in anode catalyst utilization.
Unfortunately the voltage efficiencies at these current densities is low resulting in
poor fuel cell efficiencies and power densities and also require higher stoichiometric
rates to sustain performance, unfavorable from a systems viewpoint.
165
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.24 Effect o f methanol concentration on the fuel cell efficiency of MEA
98-35 in a direct methanol fuel cell at 60 °C utilizing ambient oxygen at the cathode.
Fuel Cell Efficiency %
35
use PVDF-PSSA MEA 98-35
2"x2" MEA (25 cm2 )
60 °C, 0.10 L/min Ambient Oxygen
2 0
1.00 M Methanol
I SO M Methanol
ZOO M Methanol
a.
350 150 250 300 400 450 100 0 50
Current Density (mA/cm1 )
The dependence o f fuel cell efficiency on concentration suggested we might
be able to realize favorable efficiencies utilizing ambient air since optimum cell
voltages were achieved with reduced methanol molarity (figure 3.22). Therefore we
calculated the fuel cell efficiency of MEA 98-35 with 0.5, 1.0 and 1.5 M methanol
utilizing ambient air at 60 °C, as illustrated in figure 3.25. The data indicates that a
fuel cell efficiency o f 29.5% is observed at 100 mA/cm2 utilizing 0.50 M methanol
and is consistently superior to that of the other concentrations tested and Nafion®-
117, at all current densities. The ability to obtain similar fuel cell efficiency values
at 0.50 M utilizing ambient air as compared to 1.0 M methanol utilizing ambient
oxygen, despite the dramatic discrepancy in voltage efficiency, confirms the
166
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
inhibitory impact o f methanol crossover on fuel utilization and fuel cell efficiency.
Since fuel utilization is such a significant contributor to fiiel cell efficiency at low
current densities the benefits o f utilizing 0.50 M, despite the problems associated
with air operation, are readily apparent.
Figure 3.25 Effect o f methanol concentration on the fuel cell efficiency o f MEA
98-35 in a direct methanol fuel cell at 60 °C utilizing ambient air at the cathode.
Fuel Cell Efficiency %
use PVDF-PSSA MEA 98-35
2"x2” MEA (25 cm2 )
60 °C, 0.10 L/min Air
25
0.50 M Methanol
0.75 M Methanol
1.00 M Methanol
Nation* 117: 0.50 M Methanol
50 100 150 250 0 200
C urrent Density (mA/cm2)
3.3.1.3 Parametric Stack and System Studies using PVDF-PSSA Membranes
The favorable cell performance of MEA 98-35 at low temperatures utilizing
ambient air prompted our JPL team to study the feasibility o f incorporating the
membrane into an operating system.1 7 1 Parametric studies conducted by our JPL
team established a minimum set of operational parameters necessary for an MEA to
provide the required power density to operate a stack and subsequent system: 0.400
167
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
V at 100 mA/cm2 at 60 °C with 0.50 M methanol utilizing 0.10 L/min o f ambient air.
The ability o f a membrane to meet the target performance with minimal methanol
crossover would translate favorably into the design o f a lightweight (12 kg) 150 W
portable power system.
In a Nafion®-based system, operation at such a low current density (100
mA/cm2) maintains the required power density but exhibits poor fuel and fuel cell
efficiency values as a result o f high methanol crossover rates. Adjusting the target
performance to higher current densities where methanol crossover is reduced due to
improved anode catalyst utilization, is limited by reduced cell voltages and thus
decreased power density. Also operation at higher current densities requires much
greater air stoichiometry rates necessary to remove product water and maintain
adequate reactant supply. Unfortunately the parasitic losses associated with
increased air stoichiometry rates, coupled with reduced power density further
complicates system design.
The ability o f MEA 98-35 to meet the desired target cell performance at low
current densities with crossover rates 25% o f Nafion®-l 17 translated into greater fuel
and fuel cell efficiency values as compared to Nafion®-l 17, as previously reported in
figures 3.23, 3.24 and 3.25. This was particularly exciting since it presented an
opportunity to construct an efficient 150 W portable power system utilizing a
membrane other than Nafion®-! 17. MEA 98-35 was exhaustively tested at a variety
o f methanol concentrations and operating temperatures in an attempt to optimize
operating stack parameters. Our JPL team concluded that stack and system
168
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
complexity could be further reduced by maintaining our target performance of 0.400
V at 100 mA/cm2 while operating at 55 °C with 0.50 M methanol utilizing ambient
air at 0.10 L/min. As illustrated in figure 3.26, MEA 98-35 was evaluated at these
conditions and exhibited a cell voltage o f 0.401 volts at 100 mA/cm2, highly
competitive to that ofNafion®-l 17.
Figure 3.26 Performance o f MEA 98-35 in a direct methanol fuel cell as
compared to Nafion®-l 17 at 55 °C with 0.50 M methanol utilizing ambient air at the
cathode.
Electrical Performance
0.9
0.8
0.7
0.6
0.5
0.4
0.3
0.2
u se PVDF-PSSA MEA 98-35 & Nafion-117
2"x2" MEA (25 cm2)
55 “C, 0.50 M M ethanol
0.10 L/min Ambient A ir
MEA 98-35
0 50 100 150 200 250
Current Density (mA/cm)
The benefit of utilizing MEA 98-35 at the proposed operating conditions is
readily apparent when studying fuel efficiency as compared to Nafion®-117, as
shown in figure 3.27. The OCV parasitic loss o f 17 mA/cm2 at 55 °C with 0.5 M
methanol is roughly 25% o f Nafion®-l 17 and corresponds to a fuel efficiency of
169
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
85.4% at 100 mA/cm2, as compared to a value o f 74% for Nafion®-l 17. The ability
to improve fuel efficiency while maintaining the target performance o f 0.400 V at
100 mA/cm2 highlights the benefits o f the PVDF-PSSA membrane developed at
USC, especially while operating at such low current densities.
Figure 3.27 Fuel efficiency of MEA 98-35 in a direct methanol fuel cell as
compared to Nafion®-! 17 at 55 °C with 0.50 M methanol.
Fuel Efficiency %
100
90
80
70
60
USC PVDF-PSSA MEA 98-35 & Nafion-117
2"x2" MEA (25 cm2)
0.50 M M ethanol, 55 °C
50
40
30
2 0 MEA 98-35
1 0
Nafion-117
0
1 0 0 150 200 250 0 50
C urrent Density (mA/cm2 )
The corresponding fuel cell efficiency values were also calculated for MEA
98-35 and Nafion®-117 at 55 °C with 0.50 M methanol and a cathode flow rate of
0.10 L/min ambient air, as illustrated in figure 3.28. The data indicates that despite
the slightly improved voltage efficiency o f Nafion-117 (figure 3.26) MEA 98-35
exhibited a higher fuel cell efficiency (27.8%) at 100 mA/cm2 attributed to the low
methanol crossover rates. The ability o f MEA 98-35 to at least maintain the
170
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
desired stack performance o f 0.400 V utilizing low flow rate ambient air, while
maintaining superior fuel utilization and fuel cell efficiency suggests the USC
membrane had proven to be a capable membrane alternative.
Figure 3.28 Fuel cell efficiency o f MEA 98-35 in a direct methanol fuel cell as
compared to Nafion®-117 at 55 °C with 0.50 M methanol utilizing ambient air at the
cathode.
Fuel Cell Efficiency %
USC PVDF-PSSA MEA 98-35 & Naflon-117
2"x2" MEA (25 cm2 )
0.50 M M ethanol, 55 °C
0.10 L/min Ambient Air
1 0 ■
MEA 98-35
0 50 100 2 0 0 250 150
C urrent Density (mA/cm2)
The development o f modified MEA fabrication techniques leading to the
improved cell performance and efficiency values o f MEA 98-35 provided us with a
baseline membrane comparable to Nafion®-117. However we decided to investigate
the impact o f varying membrane parameters on fuel cell performance now that
method III had been established as our preferred MEA fabrication technique.
Therefore we prepared membrane samples of varied PSSA uptake and thickness in
171
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
hopes o f continuing the improvement in cell performance, especially as compared to
Nation®-117, while maintaining low methanol crossover rates.
3.3.1.4 Impact of PSSA Uptake on MEA Properties and Electrical Performance
In our development o f membranes and subsequent electrical cell testing of
MEAs it was clear that many fuel cell related properties like membrane water
content, proton conductivity and methanol permeability, discussed in detail in
section 2.1.4, were dependent upon PSSA uptake. Membrane 98-35 had displayed
comparable proton conductivity, as compared to Nafion®-117, based on a PSSA
uptake o f — 15% and exhibited promising electrical performance at low temperatures
utilizing ambient air. In hopes o f further improving cell performance we also
prepared membranes with PSSA uptake values above and below 15% in order to
probe the impact on fuel cell related properties. As mentioned earlier, PSSA uptake
was regulated according to the number o f styrene impregnations conducted per
membrane and the following samples were prepared and characterized, as shown in
table 3.5, then fabricated into MEAs using the methodology established for MEA 98-
35.
172
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Table 3.5 PSSA uptake and related properties o f membranes 98-36, 98-40 and
98-47A as compared to the same properties in membrane 98-35 and Nafion®-! 17.
Membrane PSSA
Uptake
Thickness Proton
Conductivity
c h 3 o h
Diffusion Coefficient
Nafion-117 N/A 7 mils 60/65 mS/cm 3.60 E-07 cm2 /s
98-35 15% 1 3 mils 45/50 mS/cm 6.40 E-08 cm2 /s
98-36 1 0% 13.5 mils 25/34 mS/cm
5.6 E-08 cm2 /s
98-40 19% 11.5 mils 73/75 mS/cm N/A
98-47A 12% 12 mils 45/46 mS/cm 6.42 E-08 cm2 /s
The membranes were assembled into the necessary test hardware and
evaluated for cell resistance only after an exhaustive conditioning regimen,
previously used in the testing of MEA 98-35. The steady state OCV resistance
values for each cell was taken at 60 °C and 90 °C and plotted vs. PSSA uptake, as
shown in figure 3.29. The data indicates that there is a linear relationship o f
decreasing cell resistance with increasing PSSA uptake; an expected trend based on
the increasing membrane water content associated with increasing PSSA uptake
lowering the grain-boundary resistivity.
173
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.29 Effect o f PSSA uptake on the measured cell resistivity o f PVDF-
PSSA MEAs in direct methanol fuel cells at 60 °C and 90 °C.
Cell Resistance of USC PVDF-PSSA MEAs
25
L re a r(6 0 o O
MEA
98-36
— Linear (90 o O
2 0
15
MEA
98-47A
M EA
98-35
MEA
98-40
1 0
5
0
18 2 0 1 0 1 2 8 16 14
PSSA Uptake %
The various MEAs were then evaluated for parasitic losses as a result of
methanol crossover, and plotted vs. PSSA uptake as shown in figure 3.30. The data
indicates that there is a linear increase in methanol crossover with increasing PSSA
uptake at both 60 °C and 90 °C, respectively. This would be expected since the
membrane screening and analysis conducted in section 2.1.4 indicated that the
membrane water content and associated methanol diffusion coefficients increase
with PSSA uptake. However, even at PSSA uptake values o f 20% (MEA 98-40)
methanol crossover values at 1.0 M methanol were still substantially lower than
Nafion®-117.
174
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.30 Effect o f PSSA uptake on the measured methanol crossover rates of
PVDF-PSSA MEAs in direct methanol fuel cells at 60 °C and 90 °C with 1.0 M
methanol.
M ethanol C rossover in PVDF-PSSA MEAs
6 0
5 0
1
s
MEA
98-40
w
b
>
98-35
s
2
MEA
98-47A
— Linear (90oC)
>
8
Linear (60 oC)
MEA
98-36
16% 2 0 % 16% 6 % 1 2 % 14%
PSSA Uptake %
Once the respective MEAs were adequately conditioned, cell performance
was evaluated at a variety of test parameters. The data in figure 3.31 presents fuel
cell performance at 60 °C with 1.0 M methanol utilizing 0.10 L/min o f ambient air at
the cathode. The performance o f MEA 98-35 has also been included as a reference
to our baseline PSSA uptake value and its corresponding MEA cell performance.
The data indicates that despite variations above and below the 15% PSSA uptake
baseline, MEA 98-35 still exhibited superior cell voltages as compared to the other
MEAs. The fact that the cell performance for MEAs 98-36 (10% PSSA) and 98-40
(19% PSSA) is similar is quite surprising considering the difference in PSSA
175
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
uptake (table 3.5), steady state resistivity values (figure 3.29) and methanol
crossover rates (figure 3.30) between the two MEAs. The poor cell performance o f
MEA 98-47A (12% PSSA), especially at current densities greater than 100 mA/cm2,
was surprising since it exhibited resistivity values (18 mOhms) comparable to MEA
98-35 (17 mOhms) suggesting polarization losses were not necessarily grain-
boundary related.
Figure 3.31 Effect o f PSSA uptake on the performance of PVDF-PSSA MEAs in
direct methanol fuel cells at 60 °C with 1.0 M methanol utilizing ambient cathode
air.
Electrical Performance
0.9
use PVDF-PSSA MEAs
2"X2" MEA (25 cm2 )
1.0 M M ethanol, 60 °C, 0.10 L/min A m bient A ir
o .g
0 .6
MEA 98-35: 15% PSSA Uptake
MEA 98-36: 10% PSSA Uptake
MEA 98-47A. 12% PSSA Uptake
MEA 98-40: 19% PSSA Uptake
0.3
0 .2
50 200 1 0 0 250 0 150
Current Density (mA/cm1 )
We performed anode polarization analysis on each MEA in order to probe the
reproducibility o f our MEA fabrication technique on anode performance, as well as
determine discrepancies in cell voltages. The resulting anode polarization plots
shown in figure 3.32 were measured at 60 °C with 1.0 M methanol and are fairly
176
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
reproducible to that measured for MEA 98-35, with the notable exception o f MEA
98-47A However, upon post-testing analysis it was evident that MEA 98-47A had
significant electrode damage at the anode as a result of excessive torque applied to
the MEA in the cell hardware thus explaining the poor anode and overall cell
performance. Therefore the anode polarization data suggested that our MEA
fabrication procedure results in adequate anode catalyst utilization, regardless o f
PSSA uptake.
Figure 3.32 Effect o f PSSA uptake on the IR corrected anode performance o f
PVDF-PSSA MEAs in direct methanol fuel cells at 60 °C with 1.0 M methanol.
Anode Performance
0.7
use PVDF-PSSA MEAs
2"X2" MEA (25cm2 )
1.0 M M ethanol, 60 °C
0 6
0.5
Z 0.4
.2 0.3
— MEA 98-35: 15% PSSA Uptake
MEA 98-40: 19% PSSA Uptake
MEA 98-47A 12% PSSA Uptake
— Nafion-117
MEA 98-36: 10% PSSA Uptake
0 2
1 1 0 100 1000
C urrent Density (mA/cm2)
The resulting similarities in the anode polarization data suggest disparities in
overall cell performance are a direct result o f cathode-related issues. This would be
expected since the variations in PSSA uptake led to a broad range o f resistivity
177
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
values (figure 3.29) and methanol crossover rates (figure 3.30), factors known to
influence cathode potentials. The cathode performance per MEA was calculated
based on the anode and overall cell performance obtained at 60 °C with 1.0 M
methanol utilizing 0.10 L/min ambient air, as illustrated in figure 3.33.
Figure 3.33 Effect o f PSSA uptake on the IR corrected cathode performance of
PVDF-PSSA MEAs in direct methanol fuel cells at 60 °C with 1.0 M methanol
utilizing ambient air.
Cathode Perform ance
u s e PVDF-PSSA MEAs
2"X 2" MEA (25cm2 )
1.0 M M ethanol, 60 °C
0.9
0.8
0.7
Z
3
|
o MEA 98-35: 15% PSSA Uptake
0.3
MEA 98-36: 10% PSSA Uptake
MEA 98-40: 19% PSSA Uptake 0.2
MEA 98-47A 12% PSSA Uptake
0.1
1000 100 1 1 0
Current Density (mA/cm2)
The cathode polarization data from the various MEAs illustrated in figure
3.33 can be used to determine the impact o f PSSA uptake on cathode related issues
such as proton flux and methanol crossover. Increasing the PSSA uptake beyond
15% reduced resistivity but increased methanol crossover having a negative impact
on cathode potentials, as evident by the performance o f MEA 98-40 (19%
178
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
PSSA). The cathode exhibited severe polarization losses attributed mainly to
methanol crossover and cathode flooding supported by independent cell testing that
indicated optimum cell performance at medium to high flow rates (0.50 - 1.00 L/min
ambient air) where effective water removal improved oxygen access and diffusion.
Cathode analysis of low PSSA uptake samples such as MEA 98-36 (10%)
and MEA 98-47A (12%) indicated that despite the reduced methanol crossover rates
(figure 3.30) cathode polarization losses were severe and similar to that o f the high
PSSA uptake sample (98-40). The overall cell performance of both MEAs 98-36
and 98-47A was extremely sensitive to flow rate suggesting inadequate cathode
water content was a contributing factor to the reduced cathode potentials. In the case
o f MEA 98-36 this is supported by the high grain-boundary resistivity value o f 23
mOhms, substantially higher than MEA 98-35 (15 mOhms). Since MEA 98-47B
had a slightly higher PSSA uptake (12%) than MEA 98-36 it would be expected to
have slightly improved cathode potentials since the corresponding cell resistivity was
18 mOhms, somewhere between that o f MEAs 98-35 and 98-36. This is supported
by the data in figure 3.33 indicating that the cathode polarization losses o f MEA 98-
47A are slightly reduced as compared to MEA 98-36 suggesting that the poor cell
performance was attributed to the severe anode polarization losses resulting from the
aforementioned electrode damage.
The overall analysis o f the various MEAs suggested that the PSSA uptake o f
MEA 98-35 was sufficient to provide the cathode adequate proton flux and thus
water content while minimizing excessive methanol crossover. Increasing the PSSA
179
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
uptake above the 15% baseline resulted in cathode flooding while reducing it
inhibited performance due to resistive losses. Therefore we established and
maintained a baseline PSSA uptake value o f 15% and attempted to prepare multiple
membrane samples according to this benchmark. Unfortunately much o f our
membrane work was hampered by irregularities in precursor thickness as well as
difficulties in consistently attaining a 15% PSSA uptake value per styrene
impregnation. However, we prepared one such membrane (98-72) with a PSSA
uptake o f 14.6% that exhibited desirable PSSA related screening characteristics, as
compared to membrane 98-35, with the physical exception o f decreased membrane
thickness, as shown in table 3.6. Despite the thickness differential the membrane
was fabricated into an MEA and tested in hopes o f determining the impact o f
membrane thickness on fuel cell performance.
Table 3.6 PSSA uptake and related properties o f membrane 98-72 as compared
to the same properties in membrane 98-35 and Nafion®-! 17.
Membrane PSSA
Uptake
Thickness Proton
Conductivity
CHjOH
Diffusion Coefficient
MEA
Thickness
Nafion-l 17
98-35
98-72
N/A
15%
14.60%
7 mils
13 mils
11 mils
60/65 mS/cm
45/50 mS/cm
56/58 mS/cm
3.60 E-07 cm2/s
6.40 E-08 cm2/s
6.70 E-08 cm2/s
20 mils
26 mils
23 mils
180
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
3.3.1.5 Impact of Membrane Thickness on MEA Properties and Electrical
Performance
The immediate impact o f membrane thickness was readily apparent based on
the steady state resistivity values o f MEA 98-72 measured at 60 °C (12 mOhms) and
90 °C (10 mOhms), as compared to 17 and 14 mOhms for MEA 98-35, respectively.
The differences in resistivity between the two MEAs were not too significant but did
translate into a slight increase in methanol crossover rates. MEA 98-72 had an OCV
parasitic loss of 43 mA/cm2 at 60 °C with 1.0 M methanol as compared to 33
mA/cm2 for MEA 98-35. The disparity was not unexpected considering the reduced
MEA thickness and corresponding increase in the methanol diffusion coefficient of
membrane 98-72 illustrated previously in table 3.6. Once conditioned, MEA 98-72
was evaluated at 60 °C with 1.0 M methanol utilizing ambient oxygen at the cathode.
The data in figure 3.34 reports a cell voltage o f 0.411 volts at 160 mA/cm2 at 0.10
L/min oxygen, a slight improvement over MEA 98-35 (0.397 V at 160 mA/cm2 ).
However, the cell performance was relatively independent of flow rate up to 1.00
L/min, contrary to the data observed for MEA 98-35 (figure 3.10) that exhibited a
decrease in cell voltages with even the slightest increase o f flow rate.
181
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.34 Effect o f ambient oxygen flow rates on the performance o f MEA 98-
72 in a direct methanol fuel cell at 60 °C with 1.0 M methanol.
Electrical Performance
u s e PVDF-PSSA MEA 98-72
2"x2" MEA (25 cm2 )
1.0 M Methanol, 60 °C, Ambient Oxygen
0.9
0.8
0.7
0.6
>
0.10 L/min Ambient Oxygen
O
«a
■
0.5
e
>
0.50 L/min Ambient Oxygen
0.4
1.00 L/min Ambient Oxygen
0.3
0.1
420 450 330 360 390 90 120 150 300 0 30 60 180 270 210 240
Current Density (mA/cm2 )
Since the cell performance of MEA 98-72 was relatively tolerant to
increasing flow rate at 60 °C, exhibiting no increase in resistivity, the cell was heated
to 90 °C and evaluated paying special attention to ambient performance. The data in
figure 3.35 indicates an improvement in the cell performance at 90 °C with 1.0 M
methanol utilizing 0.10 L/min ambient oxygen, even as compared to similar cell
testing at 60 °C (figure 3.34). This is in stark contrast to similar testing of MEA 98-
35 that exhibited poor cell performance at 90 °C (figure 3.12) utilizing ambient
cathode flow rates, especially as compared to testing at 60 °C (figure 3.10).
182
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.35 Effect o f ambient oxygen flow rates on the performance o f MEA 98-
72 in a direct methanol fuel cell as compared to cell data measured at 60 °C and
MEA 98-35 measured at 90 °C, with 1.0M methanol.
Electrical Performance
USC PVDF-PSSA MEA 98-72 & 98-35
2"x2" MEA (25 cm2)
1.0 M M ethanol, 60 °C, Ambient Oxygen
0.9
0.8
0.7
MEA 98-72: 90 oC, 0.10 L/min Ambient Oxygen 0.6
v
“ 0 5
o
>
X
C J 0.4
MEA 98-72: 60 oC, 0.10 L/min Ambient Oxygen
MEA 98-35: 90 oC, 0.10 L/min Oxygen
0.3
02
0.1
120 150 180 210 240 270 300 330 360 390 420 450 480 510 540 0 30 60 60
C u rren t Density (mA/cm2)
We performed anode polarization analysis on MEA 98-72 in an attempt to
understand the discrepancies in cell performance as compared with MEA 98-35
utilizing ambient oxygen flow rates. The data in figure 3.36 indicates that the
standard high temperature MEA fabrication methodology used for MEA 98-72
resulted in favorable anode performance consistent with that of MEA 98-35 at both
60 °C and 90 °C. Therefore variations in overall cell performance were a direct
result o f cathode-related polarization losses.
183
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.36 IR corrected anode performance o f MEA 98-72 in a direct methanol
fuel cell as compared to MEA 98-35 at 60 °C and 90 °C with 1.0 M methanol.
Anode Performance
0.7
u se PVDF-PSSA MEA 98-72 & 98-35
2"x2" MEA (25 cm2)
1.0 M M ethanol, 60/90 °C
0.6
0.5
>
u
0.3 s
« »
MEA 98-72: 60 oC
o
e.
MEA 98-35: 60 oC
0.2
MEA 98-72: 90oC
MEA 98-35: 90 oC
0.1
1000 100 10 1
C urrent Deiuity (mA/cm2)
The reduction in cell resistivity and associated increase in methanol crossover
had an impact on the cathode performance o f MEA 98-72, as compared to MEA 98-
35, at both 60 and 90 °C utilizing ambient oxygen. The cathode potentials for the
two MEAs at 60 °C are almost identical up to a current density o f 120 mA/cm2
where they then begin to diverge. The cathode potentials of MEA 98-72 increase
dramatically as compared to MEA 98-35 at higher current densities where the lower
cell resistivity reduces bulk polarization losses thereby maintaining proton flux and
thus adequate cathode hydration. This is further supported by the ability o f the MEA
to operate with increasing cathode flow rates without enduring cathode dryout
(figure 3.34).
184
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The disparity in the cathode potentials between the two MEAs is much more
dramatic at 90 °C where the performance o f MEA 98-72 is superior at almost all
current densities. The data suggests that despite the increased vapor pressure o f
water at 90 °C the cathode water content in MEA 98-72 is sufficient and maintained
presumably due to the increased rates o f proton and water permeability. The linear
cathode plot for MEA 98-72 at both 60 and 90 °C also suggests mass transfer o f
oxygen is not a concern, contrary to the performance o f MEA 98-35. The slight
increase in methanol crossover o f MEA 98-72 was not sufficient to significantly
polarize the cathode especially considering the reduced cell resistivity and
availability o f oxygen. However, the cathode potentials for each MEA were reduced
at 90 °C as compared to 60 °C, as a result o f increased methanol crossover,
suggesting cell performance at 90 °C is governed by polarization gains at the anode.
Figure 3.37 IR corrected cathode performance o f MEA 98-72 in a direct methanol
fuel cell as compared to MEA 98-35 at 60 °C and 90 °C with 1.0 M methanol
utilizing ambient oxygen.
Cathode Performance
0.9
0.8
use PVDF-PSSA MEA 98-72 & 98-35
2"x2" MEA (25 cm2 )
1.0 M Methanol, 0.10 L/min Ambient Oxygen
0.7
W
2 ao
i
S 0.5
l 0.4
o
a
MEA 98-72:90 oC
MEA 98-35:90 oC
0.2
MEA 98-35:60 oC
0.1
1 1 0 100
Current Density (mA/cms)
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The favorable cell performance o f MEA 98-72 utilizing ambient oxygen,
especially as compared to MEA 98-35, prompted us to evaluate cell performance
utilizing ambient air. The data in figure 3.38 represents cell performance o f MEA
98-72 at 60 °C with 1.0 M methanol utilizing ambient cathode air at a variety o f flow
rates. The data indicates that a cell voltage o f 0.388 V at 100 mA/cm2 is obtained at
an ambient cathode flow rate o f 0.10 L/min, an 18 mV drop as compared to MEA
98-35 at the same operating parameters (figure 3.10). However, cell performance
improved with increasing flow rate attaining a cell voltage of 0.407 volts at 100
mA/cm2 with 0.50 L/min ambient air. The ability o f MEA 98-72 to tolerate
increasing cathode flow rates was first observed during testing with ambient oxygen
and was again contrary to the data previously obtained from MEA 98-35 that
exhibited a strong, negative dependence on increasing cathode flow rates.
Figure 3.38 Effect o f ambient air flow rates on the performance o f MEA 98-72 in
a direct methanol fuel cell at 60 °C with 1.0 M methanol.
Electrical Performance
0.8
u s e PVDF-PSSA MEA 98-72
2"x2" MEA (25 cm2 )
1.0 M Methanol, 60 °C, Ambient A ir
0.7
06
0.5
0.10 L/min Ambient Air
0.50 L/min Ambient Air
SJ
03
1.00 L/min Ambient Air
0.2
0.1
0 60 90 30 120 210 240 270 150 180 300
C urrent Density (mA/c«nJ) 1 $ 6
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Based on the anode performance previously reported in figure 3.36, we also
calculated the cathode potentials o f MEA 98-72 utilizing ambient air, as illustrated in
figure 3.39. We were particularly interested in the fact that the cell performance
utilizing ambient oxygen at 0.10 L/min was superior to MEA 98-35 yet similar
testing utilizing ambient air was inferior. The data in figure 3.39 indicates that
ambient air operation at 0.10 L/min lowers the cathode potential o f MEA 98-72 by
37 mV as compared to MEA 98-35 at 100 mA/cm2 and the gap widens further with
flow rate. However, at current densities beyond 160 mA/cm2 the cathode
performance o f MEA 98-72 surpasses that of MEA 98-35 and is attributed to the
removal of product water thus maintaining oxygen diffusion. The slight increase in
the methanol crossover rates o f MEA 98-72 further aggravates the situation,
particularly at low current densities, highlighting the difficulties associated with air
operation. Incidentally the impact of methanol crossover at high current densities is
less severe due to the increase in catalyst utilization at the anode.
187
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.39 IR corrected cathode performance o f MEA 98-72 in a direct methanol
fuel cell as compared to MEA 98-35 at 60 °C with 1.0 M methanol utilizing ambient
air.
Cathode Performance
0.9
0 7 use PVDF-PSSA MEA 98-72 & 9
2”x2” MEA (25 cm2 )
0 6 1.0 M Methanol, 60 °C, Ambient
08
Z
— MEA 98-72: 0.10 L/mm Ambient Air
MEA 98-72: 0.50 L/min Ambient Air
o
— M EA 98-72: 1.00 L/min Ambient Air
— MEA 98-35: 0.10 L/min Ambient Air
0 2
0.1
0
1 0 100 1000
C urrent Density (mA/cm3)
The impact o f membrane thickness on cell performance resulted in reduced
cell resistivity accompanied by a slight increase in methanol crossover in membranes
of comparable PSSA uptake. The reduction in cell resistance at both 60 °C and 90
°C for MEA 98-72 translated into superior cell performance utilizing ambient oxygen
since the reduction in the grain-boundary resistivity reduced ohmic losses throughout
the bulk. Therefore cathode kinetics benefited from the improved proton flux and
availability o f oxygen despite the slight increase in methanol crossover.
Cell testing utilizing ambient air did not necessarily benefit from the
reduction in bulk resistivity since the mass transfer issues associated with ambient air
testing required higher flow rates necessary to provide for effective water removal
thus sustaining oxygen diffiisivity. Unfortunately the need for increased flow rates
188
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
necessary to maintain a sufficient level of performance ultimately places further
restrictions on an operating system thereby increasing parasitic power losses. Based
on the cell analysis o f MEAs 98-35 and 98-72 we determined that precursor
thickness would be an important screening criterion in reproducing our 15% PSSA
uptake baseline membranes. We then concentrated efforts on the preparation of
reproducible baseline membranes as per MEA 98-35 for incorporation into a 5-cell
stack.
3.4 Generation I 5-Cell Stack using PVDF-PSSA Membranes
The decision to incorporate PVDF-PSSA membranes into a 5-cell stack was
based on the performance parameters established by our parametric studies and met
by MEA 98-35, as discussed previously in section 3.3.1.3. However, continued
parametric system analysis called for an 80 cm2 MEA, as opposed to 25 cm2,
operating at our original target performance of 0.400 V at 100 mA/cm2 at 55 °C with
0.50 M methanol utilizing 0.10 L/min of ambient air. The increase in effective
active area would result in the attainment of higher power output per cell thereby
minimizing the number o f cells required to meet power demands thus further
reducing stack size and weight.
The increased MEA size caused us to modify our membrane fabrication
procedure since we did not have the necessary facilities to prepare precursors o f this
size at the time. Therefore we established a working relationship with Elf Atochem,
N .A and transferred our precursor fabrication techniques to them in anticipation of
precursors 10-11 mils thick to be delivered to our laboratory. These precursors were
189
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
to be impregnated with polystyrene up to our baseline PSSA uptake value of 15%,
characterized and fabricated into MEAs using our high temperature procedure
(method HI) and assembled into the 5-cell stack. Elf Atochem, N .A initially had
difficulties maintaining a uniform thickness distribution o f 10-11 mils throughout the
precursor samples but after repeated efforts several precursors were prepared with
reasonable thickness (11-12 mils) and were incorporated into our membrane
preparatory cycle. The precursor samples were impregnated with polystyrene
following our standard procedure including sulfonation, hydrolysis and screening
according to our normal methods. The data in table 3.7 provides the proton
conductivity values and methanol diffusion coefficients for the prepared membranes
as compared to MEA 98-35 and Nafion®-l 17.
Table 3.7 PSSA uptake and related properties o f potential stack membranes as
compared to the same properties in membrane 98-35 and Nafion®-! 17.
Membrane PSSA Uptake Proton Conductivity CH3 OH Diffusion Coefficient
N afion-117 N/A 60,65 mS/cm 6.40 E-08 cm2 /s
98-35 15% 45/50 mS/cm 3.59 E-07 cm2 /s
98-76 15.50% 62/65 mS/cm 1.51 E-07 cm2 /s
98-77 14.50% 37/46 mS/cm 1.41 E-07 cm2 /s
98-78 (Cell 1) 14% 42/43 mS/cm 1.16 E-07 cm2 /s
98-80 (Cell 2) 13.50% 38/40 mS/cm 6.14 E-08 cm2 /s
98-82 (Cell 3) 14% 45/46 mS/cm 1.03 E-07 cm2 /s
98-86 (Cell 4) 12% 30/35 mS/cm 9.27 E-08 cm2 /s
98-87 (Cell 5) 14% 50/55 mS/cm 1.06 E-07 cm2 /s
190
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Despite the favorable PSSA uptake related properties (table 3.7) exhibited by
the stack-membrane candidates several problems still existed. Inadequate
polymerization pressure caused moderate to severe membrane wrinkling, attributed
to macroscopic phase separation, in some o f the membrane samples. The increased
thickness o f the PVDF precursors also resulted in thicker MEAs after the membrane
preparatory and MEA fabrication procedures. The overall physical properties o f the
stack candidates, as compared to MEA 98-35, have been summarized in table 3.8.
Normally some of these membranes would be rejected yet the difficulties in
obtaining uniform 10 mil precursors with 15% PSSA uptake per impregnation
maintained the importance o f these membranes as stack candidates.
Table 3.8 Overall physical properties o f the potential stack membranes as
compared to membrane 98-35 and Nafion®-! 17.
Membrane Precursor
Thickness (Average)
PVDF-PSSA
Thickness
Morphology MEA
Thickness
Nafion 7 mils N/A Smooth 20 mils
98-35 — 10 mils — 13 mils Smooth 25 mils
98-76 11.5 mils 14.5 mils Wrinkled 27 m ils
98-77 11.4 mils 14.5 mils Slightly Wrinkled 31 mils
98-78 (Cel! 1) 10.2 mils 14 mils Slightly Wrinkled 29 mils
98-80 (Cell 2) 11.5 mils 14.5 mils Wrinkled 27 mils
98-82 (Cell 3) 11.5 mils 14 mils Extremely Wrinkled 30 mils
98-86 (C eil 4) 12 mils 15 mils Smooth 29 mils
98-87 (C ell 5) 11.5 mils 14 mils Extremely Wrinkled 28 mils
191
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
3.4.1 Electrical Performance of PVDF-PSSA Membranes: 80 cm2 5-Cell DMFC
Stack
Despite the thickness and morphology problems associated with some
membrane samples the candidates were fabricated into MEAs and inserted into the
5-cell stack. One difficulty associated with the proprietary design of the 5-cell stack
hardware was the inability to pressurize the stack and thus condition the MEAs. We
attempted to overcome this by soaking the MEAs in a hot water bath prior to
assembly but realized additional conditioning as a result o f driving protonic current
through the membrane would have to be done utilizing ambient cathode flow rates.
Once the MEAs were assembled into the advanced stack hardware the cell
was heated to 90 °C and conditioned utilizing ambient oxygen at the cathode.
However, like MEA 98-35, cell performance at these conditions was poor and
characterized by high resistivity values. Therefore the majority of the initial stack
conditioning was performed at 75 °C with 1.0 M methanol utilizing ambient cathode
oxygen at 0.10 L/min per cell. Once the stack was thoroughly conditioned at 75 °C
the steady state resistivity values per cell were measured at 60 °C and reported in
table 3.9. The data indicates that the measured resistance values are varied per
membrane and, when adjusted for a 25 cm2 active area, are slightly, or extremely
higher than MEA 98-35.
192
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Table 3.9 Measured resistivity values for cells 1-5 in a direct methanol 5-cell
stack as compared to MEA 98-35 at 60 °C.
MEA PSSA Uptake
60 °C Cell Resistance
Nafion-117 N/A 9 mOhms
98-35 15% 16 m Ohm s
98-78 (Cell 1) 14% 18 m Ohm s
98-80 (Cell 2) 13.50% 28 m Ohm s
98-82 (Cell 3) 14% 32 m Ohm s
98-86 (Cell 4) 12% 30 m Ohm s
98-87 (Cell 5) 14% 40 m Ohm s
The data in table 3.9 indicates that the resistivity of cells 1 and 2 are the
closest to the value obtained for MEA 98-35 at 60 °C and exhibited the optimum cell
performance, as compared to cells 2 and 3, when tested at our proposed stack
operating conditions. The data in figure 3.40 indicates that the operational target o f
0.400 V was achieved at 60 mA/cm2 for cell 1 and 56 mA/cm2 for cell 2, far below
the 100 mA/cm2 baseline established by MEA 98-35. The performance of cells 3
and 4 was poor and can be partially explained by the high resistivity values reported
for these two MEAs. Incidentally cell number 5 was damaged during stack assembly
and was replaced by a Nafion®-117 blank.
193
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.40 Fuel cell performance o f cells 1-5 in a direct methanol 5-cell stack as
compared to MEA 98-35 at 55 °C with 0.50 M methanol utilizing ambient air at the
cathode.
Electrical Performance
use PVDF-PSSA Stack MEAs & 98-35
Stack * 80 cm2, 98-35 = 25 cm2
0.50 M M ethanol, 55 °C
0.10 L/min Ambient Air p er MEA
0.9
0.8
0.7
• MEA 98-35
0.6
>
V
Cell I Voltage (V)
Cell 2 Voltage (V)
Cell 3 Voltage (V)
Cdl 4 Voltage (V)
o i
m
0.5
o
>
U 0.4
0.3
0.2
210 60 90 120 180 30 150 0
C urrent Density (mA/cm*)
The high cell resistivity values and poor electrical performance o f the stack
MEAs was discouraging considering the favorable PSSA-related characteristics
exhibited by the membrane candidates prior to MEA fabrication. We performed an
anode polarization analysis on the stack in order to determine the contributing issues
to reduced cell performance. The experiment was conducted at 60 °C with 1.0 M
methanol in order to directly compare anode performance with MEA 98-35 and, as
the data in figure 3.41 indicates, exhibited a wide range in anode polarization results
per cell. Cell number 1 exhibited optimum anode performance as compared to the
other stack MEAs and was slightly better than MEA 98-35 at low current densities.
194
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The rest o f the cells exhibited severe polarization losses further contributing to the
poor cell performance previously discussed in figure 3.40.
Figure 3.41 IR corrected anode performance o f cells 1-4 in a direct methanol 5-
cell stack as compared to MEA 98-35 at 55 °C with 0.50 M methanol.
Anode Perform ance
0.7
u s e PVDF-PSSA Stack MEAs &, 98-35
Stack = 80 cm2 , 98-35 = 25 cm2
1.00 M M ethanol, 60 °C
0.6
0.5
0.4
Cell 4
e 0.3
Ceil 3
Cell 2
0.2
Cell 1
MEA 98-35
0.1
100 1000 1 0 1
Current Density (mA/cm2 )
The disparity in anode polarization per MEA was difficult to explain since
our previous MEAs had all exhibited fairly reproducible and favorable anode
performance. However, upon post-testing analysis o f the stack MEAs the source o f
many o f the performance problems became evident. The membranes that had
exhibited moderate too severe surface wrinkling (table 3.8) resulted in damage at
both anode and cathode electrodes. The wrinkled areas had preferentially swollen
and expanded thus shredding the electrode backing, reducing catalyst utilization and
contributing to the high resistivity values for those particular cells. It is important to
195
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
note that MEA number 1 exhibited preferred membrane morphology and resistivity
values similar to MEA 98-35 suggesting incomplete conditioning may have
contributed to the reduced performance.
The initial development of the 5-cell stack resulted in disappointing cell
performance as a result o f a number of issues. The inability to adequately condition
the MEAs coupled with the increased thickness o f most of the samples resulted in
high resistivity values as compared to MEA 98-35. Also the morphology
irregularities in several o f the MEAs resulted in moderate to severe electrode damage
inhibiting performance. However, before stack testing was complete we had already
prepared a second set o f alternate membranes with normal, preferred surface
morphology as a result o f maintaining polymerization pressure throughout styrene
impregnation. Unfortunately the difficulties in obtaining a large quantity o f 10-11
mil precursors of uniform thickness, coupled with the irregularity in PSSA uptake
per impregnation persisted and continued to prove troublesome.
3.5 Modified PVDF-PSSA Membrane Development
A majority o f the rejected precursors intended for incorporation into our
membrane preparatory cycle were actually 13 mils in thickness, and highly uniform
throughout. We decided to investigate the impact o f increasing the PSSA uptake
beyond that of the 15% baseline utilizing the thicker precursor samples in hopes o f
reducing cell resistance without significantly contributing to methanol crossover.
Based on our early styrene impregnation experiments it was determined that 10%
PSSA uptake could be polymerized per impregnation consistently and reproducibly
196
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
from sample to sample. Therefore we established a second, arbitrary baseline value
o f 20% PSSA, that was easier to obtain than the 15% PSSA uptake required per
impregnation for MEA 98-35 and the subsequent stack MEAs. The 20% PSSA
uptake membranes were screened according to our normal methods and compared
with MEA 98-35 and Nafion®-117, as shown in table 3.10. The results indicate that
the use o f a thicker precursor resulted in much thicker membranes and MEAs while
the increase in PSSA uptake resulted in high membrane water content and proton
/jSy
conductivity values surpassing even that o f Nafion -117. Unfortunately the
measured methanol diffusion coefficients had increased by a factor o f 3 as compared
to membrane 98-35 despite the dramatic increase in thickness.
Table 3.10 PSSA uptake and related properties o f membrane 98-94, 98-95 and
98-96 as compared to the same properties o f membrane 98-35 and Nafion®-! 17.
Membrane PSSA
Uptake
Thickness Proton
Conductivity
CHjOH
Diffusion Coefficieny
MEA
Thickness
Nafion-117 N/A 7 mils 60/65 mS/cm 3.60 E-07 cm 2/s 20 mils
98-35 15% 13 mils 45/50 mS/cm 6.40 E-08 cm 2/s 26 mils
98-94 20.60% 19.5 mils 77/80 mS/cm 2.30 E-07 cm2/s N/A
98-95 21% 18.5 mils 75/77 mS/cm 1.86 E-07 cm2/s 33 mils
98-96 21.50% 19 mils 80/76 mS/an 1.99 E-07 cm2/s N/A
3.5.1 Electrical Performance o f PVDF-PSSA Membranes: 2"X2" DMFCs
Despite the increase in the methanol diffusion coefficients, a 20% PSSA
uptake sample (98-95) was fabricated into an MEA using our high temperature
procedure (method ID) and evaluated for cell performance. Upon hydration at 75 °C
197
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
for 2 days the steady state resistivity value at 60 °C was 13.5 mOhms, achieved
without pressurized conditioning. The resistivity value was slightly lower than that
measured during the testing o f MEA 98-35 (15 mOhms) yet higher than a previous
MEA of identical PSSA uptake (MEA 98-40: 20% PSSA) that exhibited an OCV
resistance of 12 mOhms but was only 11.5 mils thick. We therefore reasoned that
the increased thickness of MEA 98-95 (17.5 mils) contributed to the difference in
resistivity as compared to MEA 98-40. The methanol crossover rates of MEA 98-95
measured at 60 °C with 1.0 M methanol were also slightly lower (53 mA/cm2) than
the value reported for MEA 98-40 (58 mA/cm2) further signifying the influence of
thickness on crossover and resistivity. Unfortunately the methanol crossover value
o f MEA 98-95 was almost 20 mA/cm2 higher than that measured for MEA 98-35 (33
mA/cm2) at 60 °C with 1.0 M methanol, a value 40% ofNafion®-l 17.
MEA 98-95 was inserted into cell hardware and evaluated for cell
performance at 60 °C with 1.0 M methanol utilizing ambient oxygen at the cathode.
The data in figure 3.42 indicates that a cell voltage o f 0.425 V at 160 mA/cm2 was
obtained utilizing 0.10 L/min o f ambient oxygen, surpassing the performance o f
MEAs 98-72 (figure 3.34) and 98-35 (figure 3.10) at the same operating conditions.
Also the cell performance o f MEA 98-95 was relatively unaffected by cathode flow
rates and only measured increased resistivity from cathode drying at flow rates
greater than 1.5 L/min. It is important to note that the cell performance was achieved
without utilizing high temperature pressurized conditioning, an important factor in
screening potential stack candidates.
198
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.42 Effect of ambient oxygen flow rates on the performance o f MEA 98-
95 in a direct methanol fuel cell at 60 °C with 1.0 M methanol.
Electrical Performance
u se PVDF-PSSA MEA 98-95
2"x2" MEA (25 cm2 )
1.0 M M ethanol, 60 °C, Ambient Oxygen
0.9
0.8
0.7
0.6
v
0.10 L/min Ambient Oxygen
v
S S .
O
>
0.50 L/min Ambient Oxygen 0.5
1.00 L/min Ambient Oxygen
£ 0.4
0.3
0.2
0.1
450 300 330 360 390 420 180 2 1 0 240 270 150 30 60 90 120 0
Current Density (mA/cm1 )
The improved cell performance o f MEA 98-95 utilizing ambient oxygen was
attributed to the reduced grain-boundary polarization losses correlated to the increase
in PSSA uptake. However, we conducted anode polarization analysis at both 60 and
90 °C in order to confirm polarization losses were not attributed to the anode. The
data in figure 3.43 indicates that the anode performance o f MEA 98-95 at 60 °C and
90 °C with 1.0 M methanol is consistent with and even slightly better than that
measured for MEA 98-35. Once again this validated our high temperature MEA
fabrication technique as a fairly reproducible process from sample to sample.
199
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.43 IR corrected anode performance o f MEA 98-95 in a direct methanol
fuel cell as compared to MEA 98-35 at 60 °C with 1.0 M methanol.
Anode Perform ance
0.7
u s e PVDF-PSSA MEA 98-95 & 98-35
2 "x2" MEA (25 cm2 )
1.0 M M ethanol, 60 °C
0.6
0.5
w
z
M
>
3
c
u
o
a.
MEA 98-35
MEA 98-95
0.2
0.1
1000 1 0 0 1 0 1
Current Density (mA/cmz)
Based on the cell and anode polarization o f MEA 98-95 the cathode
performance was calculated at 60 °C with 1.0 M methanol utilizing ambient oxygen
and plotted in figure 3.44. The data indicates that the cathode benefits from the
increased proton flux attributed to the lower grain-boundary resistivity thus
maintaining cathode water content and associated reduction kinetics despite the
significant increase in methanol crossover, as compared to MEA 98-35. The cathode
performance o f MEA 98-40 has also been included as a reference since that MEA
also contained 20% PSSA uptake. However the reduced thickness of MEA 98-40
resulted in an even lower grain-boundary resistivity inhibiting cathode performance
200
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
as a result o f flooding, exasperated by operating at such a low flow rate of 0.10
L/min ambient air.
Figure 3.44 IR corrected cathode performance o f MEA 98-95 in a direct methanol
fuel cell as compared to MEAs 98-35 and 98-40 at 60 °C with 1.0 M methanol
utilizing ambient oxygen.
Cathode Performance
0.9
0 .8
u s e PVDF-PSSA MEAs 98-95,98-35 & 98-40
2"x2" M EA (25 cm2)
1.0 M M ethanol, 60 °C
0.10 L/m in A m bient Oxygen
0.7
0 .6
>
MEA 98-95
MEA 98-35
U 0.4
MEA 98-40
0.3
0 .2
0 .1
1000 1 0 1 0 0 1
C u rren t Density (m A/cm 2 )
In our subsequent evaluation of the cell performance o f MEA 98-95 utilizing
ambient air we anticipated a reduction in performance, as compared to our baseline
value o f 0.400 V at 100 mA/cm2, as a result o f cathode flooding and related oxygen
mass transfer limitations. These issues were discussed previously in the cell testing
o f MEA 98-72 (figure 3.38) and were only alleviated by increasing the cathode flow
rate to effectively remove water thus maintaining oxygen access and diffusion.
However the data in figure 3.45 reports a cell voltage o f 0.422 V at 100 mA/cm2
201
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
surpassing the performance o f MEA 98-35 and meeting the target performance
required for the stack.
Figure 3.45 Performance o f MEA 98-95 in a direct methanol fuel cell as
compared to MEA 98-35 at 60 °C with 1.0 M methanol utilizing ambient air at the
cathode.
Electrical Performance
0.7
use PVDF-PSSA M EA 98-95 & 98-35
2"X 2" M EA (25 cm2 )
1.0 M M ethanol, 60 °C,
0.10 L/min A m bient A ir
0 .6
0.5
X T 0.4
MEA 98-95
MEA 98-35
0 .2
0 .1
150 2 0 0 250 1 0 0
C u rre n t Density (m A /cm ’)
0 5 0
The favorable cell performance o f MEA 98-95, as compared to MEA 98-35
at 60 °C utilizing ambient air, was surprising since our earlier work with high PSSA
uptake MEAs resulted in moderate to severe cathode polarization losses attributed to
poor oxygen solubility as a result o f cathode flooding. In order to understand this we
plotted the cathode performance o f MEA 98-95 at 60 °C with 1.0 M methanol
utilizing 0.10 L/min ambient air in figure 3.46 and, despite the increase in methanol
crossover, it is almost superimposable to that o f MEA 98-35. Therefore the
202
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
slight improvement in ambient air cell performance o f MEA 98-95 may be partly
attributed to the slight gains at the anode discussed previously in figure 3.43.
Figure 3.46 IR corrected cathode performance o f MEA 98-95 in a direct methanol
fuel cell as compared to MEA 98-35 at 60 °C with 1.0 M methanol utilizing ambient
air.
Cathode Perform ance
USC PVDF-PSSA MEA 98-95 & 98-35
2"x2" MEA (25 cm2 ) *=
1.0 M M ethanol, 60 °C
0.10 L/min Ambient A ir
0.9
0.8
07
Z
>
>
MEA 98-95 8
©
a.
MEA 98-35
0.3
0 .2
0.1
1 1 0 1 0 0 1 0 0 0
C urren t D a u b y (mA/cm*)
Since the cell performance o f MEA 98-95 exceeded that of MEA 98-35
utilizing both ambient oxygen and air at 60 °C with 1.0 M methanol it became an
immediate membrane candidate for stack development. Therefore we evaluated cell
performance at the proposed operating conditions o f the stack in hopes o f
maintaining our operation baseline o f 0.400 V at 100 mA/cm2. However, methanol
crossover rates were determined prior to ceil testing and measured 25 mA/cm2 at
55°C with 0.50 M methanol, 10 mA/cm2 higher than MEA 98-35. Despite the slight
203
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
increase in methanol crossover the data in figure 3.47 indicates a cell performance of
0.413 V at 100 mA/cm2, comparable to Nafion®-117 and superior to MEA 98-35,
utilizing ambient air at 0.10 L/min.
Figure 3.47 Performance o f MEA 98-95 in a direct methanol foel cell as
compared to MEA 98-35 and Nafion®-117 at 55 °C with 0.50 M methanol utilizing
ambient air at the cathode.
Electrical Performance
USC PVDF-PSSA MEAs 98-95,98-35 & N afion-117
2"X 2" M EA (25 cm2)
0.50 M M ethanol, 55 °C
0.10 L/min Am bient Air
0.9
0.8
0.7
0 .6
MEA 98-95
5 0-5
MEA 98-35
O 0.4
0.3
0 .2
0 .1
1 0 0 150 2 0 0 0 5 0 250
C a rre n t Density (mA/cm1)
3.5.1.1 Fuel and Fuel Cell Efficiency o f PVDF-PSSA Membranes: 2"X2" DMFCs
The consequence o f the increased methanol crossover o f MEA 98-95 as
compared to MEA 98-35 was a slight reduction in fuel utilization at the anode. The
data in figure 3.48 reports a fuel efficiency value for MEA 98-95 o f 79.3% at 100
mA/cm2 at 55 °C with 0.50 M methanol as compared to 85.4% for MEA 98-35
obtained at the same current density. Despite the slight decrease the fuel
204
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
efficiency o f MEA 98-95 is still superior to that o f Nafion®-117 that has a reported
fuel utilization o f 74%, also at an operating current density o f 100 mA/cm2.
Figure 3.48 Fuel efficiency o f MEA 98-95 in a direct methanol fuel cell as
compared to MEA 98-35 and Nafion®-! 17 at 55 °C with 0.50 M methanol.
Fuel Efficiency %
1 0 0
90
80
70
£
S'
a
ti
60
PVDF-PSSA MEAs 98-95,98-35 & Nafion-117
2"x2" M EA (25 cm2)
0.50 M M ethanol, 55 °C
u s e
u
e
u
50
3
40
30
MEA 98-35
2 0
MEA 98-95
1 0
150 250 2 0 0 1 0 0 5 0 0
C urrent Density (mA/cm2)
As previously mentioned in section 3.3.1.2, an improvement in the voltage
efficiency o f a cell can increase the overall fiiel cell efficiency compensating for the
reduction in fuel utilization associated with increased methanol crossover. The fuel
cell efficiency o f MEA 98-95 illustrated in figure 3.49 was 26.6% at 100 mA/cm2, a
slight decrease as compared to MEA 98-35 (27.8%) at the same operating current
density. This indicates that the improvement in the cell voltage o f MEA 98-95 as
compared to MEA 98-35 was not sufficient to account for the loss in fuel utilization
attributed to the increase in methanol crossover. However, MEAs 98-95 and 98-
205
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
35 still compared quite favorably to the fuel cell efficiency value o f Nafion®-117
(25.6%) at the same operating parameters.
Figure 3.49 Fuel cell efficiency o f MEA 98-95 in a direct methanol fuel cell as
compared to MEA 98-35 and Nafion®-117 at 55 °C with 0.50 M methanol utilizing
ambient air at the cathode.
Fuel Cell Efficiency %
35
u s e PVDF-PSSA M EAs 98-95,98-35 & Nafion-117
2"x2" M EA (25 cm2 )
0.50 M M ethanol, 55 °C
0.10 L/min Ambient A ir
30
25
2 0
u
s
MEA 98-95
1 0 -
Nafion-117
250 150 1 0 0 2 0 0 0 5 0
Cnrrent Demity (mA/cm1 )
3.6 Membrane Electrode Assembly: Fabrication Method IV
The ability o f MEA 98-95 to obtain the desired cell performance, despite the
slight increase in methanol crossover, without the need to aggressively condition the
membrane was a major accomplishment considering the stack design. However, in
our efforts to reproduce the cell performance in subsequent MEAs we have
encountered several issues still to be resolved. Our high temperature MEA
fabrication methodology has resulted in adequate inter fecial contact but dries the
membrane out completely during processing. Therefore upon membrane
206
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
hydration in the cell hardware the bonded electrode papers have lost their integrity
and shredded in some samples since the increased PSSA uptake (20%) o f the second-
generation membranes results in considerable swelling as compared to the reduced
PSSA content samples (15%) discussed earlier. In order to avoid MEA damage we
investigated and developed a non-bonded technique (method IV) for MEA
fabrication designed for high PSSA uptake samples.
The non-bonded methodology presented some concerns regarding adequate
contact at the membrane-catalyst and electrode-electrolyte interfeces. During the
fabrication o f MEAs utilizing the various methods (I,II,HI) previously discussed we
determined that temperature and pressure had a significant impact on establishing
adequate interfacial contact at these junctions. The feet that MEA 98-35 had a lower
resistance than MEA 97-33, despite a lower PSSA uptake value, validated the
benefits o f method HI (section 3.3) in establishing a low contact resistance.
Therefore we investigated methods to establish the same inter fecial contact using a
non-bonded methodology as that obtained using our bonded techniques. In
preparation o f our first non-bonded MEA we selected unused membrane portions
from MEA 98-95 in order to reduce any inconsistency between membrane samples.
We modified the MEA fabrication procedure used in method ID slightly in order to
maintain adequate contact at the membrane-catalyst interface while avoiding the hot-
press bonding o f the catalyzed-electrode papers. The anode and cathode catalyst ink
compositions were the same as described in method III resulting in a catalyst loading
o f 8-12 mg/cm2 at each electrode.
207
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
3.6.1 Electrical Performance of PVDF-PSSA MEAs: 2"x2" DMFCs
The initial membrane resistivity o f MEA 98-95B (non-bonded) was high as
compared to MEA 98-95 (bonded) and exhibited poor electrical performance.
However upon cell assembly we had included thick teflon gaskets in order to avoid
crushing the electrode papers due to the torque applied to the cell. We successively
removed the gaskets several mils in thickness at a time and the cell resistance
dropped accordingly. Ultimately MEA 98-95B exhibited an OCV resistance o f 13.5
mOhms at 60 °C and 11.0 mOhms at 90 °C, values almost identical to MEA 98-95
(14 mOhms at 60 °C and 12 mOhms at 90 °C). This suggested that our non-bonded
procedure (method IV) had resulted in adequate electrical contact as compared to our
bonded MEA. Incidentally the measured OCV methanol crossover rates for MEA
98-95B with 1.0 M methanol was 50 mA/cm2 at 60 °C, consistent with MEA 98-95.
Based on the favorable resistivity values exhibited by the non-bonded MEA
98-95B we evaluated the electrical performance as compared to the bonded MEA
98-95 at 60 °C with 1.0 M methanol utilizing 0.10 L/min ambient oxygen, as
illustrated in figure 3.50. The data indicates that despite the consistent resistivity
measurements between the two MEAs, MEA 98-95 exhibited superior performance
by 67 mV at an operating current density o f 160 mA/cm2. This suggested that
despite the adequate electrical contact established by the non-bonded MEA
technique, based on the favorable resistivity values, the poor performance may be
attributed to poor catalyst utilization at the electrode-electrolyte interface.
208
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.50 Performance o f MEA 98-95B prepared using a non-bonded MEA
fabrication technique in a direct methanol fuel cell as compared to MEA 98-95 at 60
°C with 1.00 M methanol utilizing ambient oxygen at the cathode.
Electrical Performance
0.9
use PVDF-PSSA M EAs 98-95 & 98-95B
2"x2" MEA (25 cm 2)
1.00 M M ethanol, 60 °C
0.10 L/min A m bient Oxygen
0.8
0.7
0.6
r 0 5
MEA 98-95: Bonded MEA
MEA 98-95B: Non-Bonded MEA
0.3
0.2
0.1
450 90 240 270 300 360 390 420 30 60 120 210 330 0 150 180
Current Density (mA/cm2)
In order to determine the source o f the discrepancy in cell performance
between the bonded and non-bonded MEAs we conducted an anode polarization
analysis on each MEA at 60 °C with 1.00 M methanol, as shown in figure 3.51. The
data indicates that the anode potential o f MEA 98-95B is 28 mV more polarized than
MEA 98-95 at 100 mA/cm2 and that disparity increases further with current density.
This suggests that the non-bonded MEA approaches methanol mass transfer limits at
much lower current densities than the bonded MEA. However, this does not
coincide with our results reported in section 3.3.1.1, obtained from MEA 98-35 that
indicate anode performance with 1.0 M methanol was sufficient to alleviate
209
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
methanol mass transfer limitations at current densities well beyond 100 mA/cm2.
Therefore the data suggests that the non-bonded anode catalyst utilization is
decreased as compared to the bonded MEA.
Figure 3.51 IR corrected anode performance o f MEA 98-95B in a direct methanol
fuel cell as compared to MEA 98-95 at 60 °C with 1.0 M methanol.
Anode Perform ance
0 .6
use PVDF-PSSA M EA 98-95 & 98-95B
2 "x 2 " M EA (25 cm2)
1.0 M M ethanol, 60 °C
0.5
0.4
0.3
98-95B: Non-Bonded MEA
98-95: Bonded MEA
0 .2
0 .1
0
100 1000 1 1 0
C u rren t Density (m A /cm 2)
Based on the non-bonded anode and cell performance o f MEA 98-95B we
calculated the IR corrected cathode performance at 60 °C with 1.0 M methanol
utilizing 0.10 L/min ambient oxygen, as illustrated in figure 3.52. The data indicates
that the cathode potential for the non-bonded MEA is 29 mV lower than that o f the
bonded MEA at 100 mA/cm2. Since MEA 98-95B exhibited comparable resistivity
to MEA 98-95 the polarization losses are not thought to be grain-boundary related.
Also, the similarity in the methanol crossover rates between the two MEAs would
210
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
suggest that the polarization losses of the non-bonded MEA are not attributed to any
additional methanol poisoning. Therefore this indirectly suggests that the catalyst
utilization is slightly reduced in the non-bonded MEA cathode as compared to the
bonded cathode o f MEA 98-95.
Figure 3.52 IR corrected cathode performance o f MEA 98-95B in a direct
methanol fuel cell as compared to MEA 98-95 at 60 °C with 1.00 M methanol
utilizing ambient oxygen.
Cathode Perform ance
1
0.9
0.8
use PVDF-PSSA M EAs 98-95 & 98-95B
2 ”x2" MEA (25 cm2)
1.0 M M ethanol, 60 °C
0.10 L/min Am bient Oxygen
0.7
Z 0.6
■ 2 0.5
98-95: Bonded MEA
98-95B: Non Bonded MEA - 0.4
0.3
0.2
0.1
0
1000 1 0 100 1
C urrent D ensity (m A/cm 2)
The reduction in the anode performance o f the non-bonded MEA (figure
3.51) presented a significant concern for cell testing utilizing ambient air since
cathode performance is already complicated. However, the data in figure 3.53
indicates that the non-bonded MEA exhibited a cell voltage of 0.393 V at 100
mA/cm2 at 60 °C with 0.50 M methanol utilizing 0.10 L/min ambient air. This
represented a 45 mV loss as compared to MEA 98-95 but was still acceptable in
211
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
regards to our baseline requirement o f 0.400 V at 100 mA/cm2. Therefore, despite
the reduced cell voltage as compared to MEA 98-95 and the inability to obtain the
0.400 V baseline at 55 °C the non-bonded MEA was still considered a viable stack
candidate.
Figure 3.53 Performance o f MEA 98-95B prepared using a non-bonded MEA
fabrication technique in a direct methanol fuel cell as compared to MEA 98-95 at 60
°C with 0.50 M methanol utilizing ambient air at the cathode.
Electrical Performance
0.9
use PVDF-PSSA MEAs 98-95 & 98-95B
2"x 2 " MEA (25 cm 1 )
0.50 M M ethanol, 60 °C
0.10 L/m in Ambient A ir
0.8
0.7
0.6
~ Z 0.5
ac
98-95: Bonded M EA
s 0.4
98-95B: Non-Bonded MEA
0.3
0.1
0
50 250 0 I C O 300 150 200
C u rre n t Density (mA/cm1)
The ability o f the non-bonded MEA 98-95B to exhibit reasonable ambient air
performance was surprising considering the relatively poor oxygen performance
discussed earlier in figure 3.50. This suggested that the use o f 0.50 M methanol in
our air tests benefited not only the cathode, by minimizing the impact o f crossover
rates on cathode potentials, but also at the anode. The data in figure 3.54 indicates
212
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
that the anode potential o f the non-bonded MEA is 19 mV more polarized than the
bonded MEA at 100 mA/cm2, much closer than the gap exhibited by the non-bonded
and bonded MEAs with 1.0 M methanol reported previously in figure 3.51.
Furthermore, the anode potential o f the non-bonded MEA with 1.00 M methanol
(Figure 3.51: 358 mV) is only 7 mV better than that measured with 0.5 M methanol
(Figure 3.54: 365 mV) at 100 mA/cm2. This suggests that the increase in
concentration does not necessarily alleviate any mass transfer limitations that may be
present at that operating current density (100 mA/cm2). This further supports our
contention that the non-bonded anode performance is limited by poor catalyst
utilization, a phenomenon only aggravated by increasing methanol concentration.
Figure 3.54 IR corrected anode performance o f MEA 98-95B in a direct methanol
fuel cell as compared to MEA 98-95 at 60 °C with 0.50 M methanol.
Anode Performance
0.7
use PVDF-PSSA M EA 98-95 & 98-95B
2"x2" MEA (25 cm 2)
0.50 M M ethanol, 60 °C
0.6
0.5
0.4
0.3
0.2
MEA 98-95: Bonded
MEA 98-95B: Non Bonded MEA
0.1
0
1 1000 10 100
C u rre n t Density (mA/cm2)
213
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
We then calculated the cathode performance o f MEAs 98-95B and 98-95 at
60 °C with 0.50 M methanol utilizing 0.10 L/min ambient air. The data in figure
3.55 indicates that there is a 26 mV reduction in the cathode potential o f MEA 98-
95B, as compared to MEA 98-95 at a current density o f 100 mA/cm2. Incidentally
the discrepancy between the non-bonded and bonded MEA cathode potentials at 100
mA/cm2 utilizing air was similar to the difference measured utilizing oxygen (figure
3.52: 29 mV). Therefore since the resistivity and methanol crossover rates o f the
two MEAs was comparable we reasoned that the slight reduction in performance was
again attributed to decreased catalyst utilization.
Figure 3.55 IR corrected cathode performance o f MEA 98-95B in a direct
methanol fuel cell as compared to MEA 98-95 at 60 °C with 0.50 M utilizing
ambient air.
Cathode Performance
0.9
08
use PVDF-PSSA MEAs 98-95 & 98-95B
2"x2" M EA (25 cm2 )
0.50 M M ethanol, 60 °C
0.10 L/min Ambient Air
0 7
S d
Z 0.6
9 t
■ 3 0 5
I
o
e.
= 04
t>
W
9S-95. Bonded MEA
98-95B: Non Bonded MEA
0.3
02
0 1
10 1 100 1000
Current Density (mA/an2 )
214
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Our non-bonded MEA fabrication technique resulted in a 20% PSSA uptake
MEA that exhibited comparable resistivity and crossover rates to our bonded MEA
and achieved the desired performance requirement for stack operation. However,
both anode and cathode performance was sensitive to methanol concentration and
polarization losses were increased as compared to our bonded MEAs. The impact of
methanol molarity on the cathode potentials was to be expected since crossover rates
increase with concentration. However, the significant anode polarization losses
evident at relatively low current densities was surprising and ultimately indicated
that reduced catalyst utilization is a consequence o f this non-bonded technique.
We are currently investigating methods to improve the catalyst utilization
using this non-bonded technique and have developed a relationship with Giner, Inc.
to institute a standard MEA fabrication technique to deliver MEAs for the 5-cell
stack and subsequent system. Incidentally Giner, Inc. has independently developed a
MEA fabrication technique for 20% PSSA uptake membranes and has prepared and
evaluated several 25 cm2 MEAs that we have provided. The data in figure 3.56
indicates that each sample has achieved and even exceeded the minimum potential o f
0.400 V at 100 mA/cm2 with 0.50 M methanol utilizing 0.10 L/min ambient air.
215
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.56 Performance o f MEAs Giner DE-6, 10 and 12 in a direct methanol
fuel cell at 60 °C with 0.50 M methanol utilizing 0.10 L/min ambient air.
Electrical Performance
0.900
use PVDF-PSSA MEAs
2"x2" MEA (25 cm2 )
0.50 M M ethanol, 60 °C
0.10 L/min Am bient A ir
G iner, inc.
0.800
0.700
0.600
T 0 .500
Giner DE-6
Giner DE-10
= 0.400
Gmcr DE-12
0 300
0.200
0.100
0.000
180 160 60 80 120 140 20 100 0 40
C u rren t Density (mA/cm2)
3.7 Generation II 5-Cell Stack Development
The ability o f MEA 98-95B to meet the required performance baseline was a
significant achievement in our attempt to deliver a 5-cell stack and subsequent 65
cell operating system using a PVDF-PSSA membrane. From a membrane
impregnation standpoint the 20% PSSA uptake baseline was readily obtained and
reproducible in subsequent precursor samples. It was also at this time that precursor
development at USC resulted in the ability to fabricate 13 mil precursors o f uniform
thickness and adequate size for development o f 80 cm2 MEAs. We incorporated the
USC precursors into a batch containing samples supplied by Elf Atochem, N.A. and
prepared several second-generation membrane candidates. The important
216
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
physical characteristics and PSSA related properties were determined according to
our usual screening methods and are reported in table 3.11.
Table 3.11 PSSA uptake and related properties for potential stack membranes as
compared to the same properties in membrane 98-95 and Nafion®-! 17.
M em brane PSSA
Uptake
Thickness W ater
C ontent
Proton
Conductivity
C H 3OH
Diffusion Coefficient
Nafion-117 N/A 7 mils 33% 60/65 mS/cm 3.60 E-07 cm2/s
98-95 21% 18.5 mils 37% 75/77 mS/cm 1.86 E-07 cm2/s
99-08 (Elf Atochcm.N.A.) 21.50% 18 mils 34% 76/78 mS/cm 1.99 E-07 cm2/s
99-14 (USC) 19% 17.5 mils 38% 75/80 mS/cm 1.29 E-07 cm2/s
99-15 (USC) 20.60% 18 mils 37% 73/79 mS/cm N/A
99-16 (USC) 19.80% 18.5 mils 36% 76/81 mS/cm N/A
99-17 (USC) 21% 18 mils 38% 75/76 mS/cm N/A
99-17a (USC) 21% 18 mils 38% 75/76 mS/cm N/A
99-18 (USC) 21% 18.5 mils 37% 74/79 mS/cm N/A
3.7.1 Electrical Performance o f PVDF-PSSA MEAs: 80 cm2 Single Cell DMFC
Based on our experience with developing a non-bonded MEA fabrication
procedure (method IV) and the success o f the MEAs tested at Giner, Inc. we
prepared a single cell 80 cm2 MEA to validate our membrane and MEA
methodology on a scaled up MEA. The initial resistivity of the MEA was high and
the performance was poor. However, upon post-testing analysis it was determined
that the anode paper was crushed due to excessive torque applied to the cell. We
then prepared a second sample, MEA 99-18 and continued cell testing.
217
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
MEA 99-18 exhibited adequate resistivity at 60 °C (16 mOhms) as compared
to MEAs 98-95 (14 mOhms) and 98-95B (13.5 mOhms) when adjusted for active
area size. MEA 99-18 was then tested at 60 °C with 0.50 M methanol utilizing
ambient air at the cathode, as shown in figure 3.57. The cell exhibited a cell voltage
o f 0.408 V at 100 mA/cm2 inline with our desired performance baseline. This was
particularly exciting since the performance was achieved utilizing only ambient
conditioning at temperatures ranging from 60-70 °C and requiring only 3-4 scans to
meet the performance. Incidentally the ambient air flow rate measured was 0.70
L/min and was attributed to difficulties we had sealing the inlet to the cell. We have
since modified the inlet for both the single cell test hardware and 5-cell stack in order
to avoid this issue.
Figure 3.57 Performance o f MEA 99-18 (80 cm2) in a direct methanol fuel cell as
compared to MEAs 98-95 and 98-95B at 60 °C with 0.50 M methanol utilizing
ambient air at the cathode.
Electrical Performance
0.9
USC PVDF-PSSA MEAs 98-95, 98-95B & 99-18
Bonded & Noil-Bonded
0.50 M M ethanol, 60 °C, 0.10 L/min Am bient A ir
0.8
0.7
MEA 98-95: Bonded - 25 cm2
0. 6
MEA 9 9 -18: Non-Bonded - 80 cm2
u
9
MEA 98-95B: Non-Bonded - 25 cm2
0.3
0.1
0 50 100 250
150 300 200
Current Density (mA/cm2 )
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The ability o f MEA 99-18 to achieve the desired performance baseline o f
0.400 V at 100 mA/cm2 validated the combined MEA fabrication methodology
designed by Giner, Inc. and our JPL team. More importantly though it confirmed
our ability to scale-up our membranes and MEA to an 80 cm2 active area while
m aintain in g the desired electrical performance. We have currently prepared seven
advanced MEAs and assembled them into the 5-cell stack. Preliminary results
indicate that the resistivity values for each MEA is ideal and we are currently
modifying the air inlet(s) to avoid leaks and anticipate promising electrical
performance in the near future.
3.8 Experimental Section
The initial PVDF precursors used in the preparation o f PVDF-PS
interpenetrating polymer networks and subsequent PVDF-PSSA polymer electrolytes
were prepared by compression molding Kynar 460 powder according to a technique
developed by our research team. A measured amount o f PVDF powder was spread
between two Kapton sheets then inserted into a hot press and held at 220 °C under
pressure for 15-30 minutes followed by immediate cooling. This process resulted in
sheets o f fairly uniform thickness ranging from 7-13 mils depending on the applied
pressure. In the experiments that used Kynar 740 precursors, Westlake Plastics and
Elf Atochem, N.A. supplied the films directly. These commercial sheets were then
incorporated into our IPN methodology without any modifications.
219
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
The PVDF-PSSA membranes were prepared according to the method
described by Hodgdon and Boyak in which polystyrene is polymerized and
crosslinked in an inert PVDF matrix and then sulfonated.1 8 4 The precursor PVDF
films were first swollen in an acetone bath at 35-40 °C for 24 hours then submerged
in a second bath containing styrene monomer, divinyl benzene and initiator (AIBN).
The swollen precursor was allowed to equilibrate in the bath for 1-2 hours before
being sandwiched between two sheets o f stiff aluminum foil and inserted into a
stainless steel die. The die was then inserted into a hot press where it was heated to
60 °C for one hour then 90 °C for a second hour under 200 PSI of pressure and
subsequently cooled to room temperature. The PVDF-PS membranes were then
placed in a vacuum oven at 100 °C for 24 hours to remove any unpolymerized
monomer and then weighed to determine the polystyrene uptake, the amount of
which was dependent upon the initiator concentration and number of impregnations,
as shown in table 3.12.
Table 3.12 Impact o f AIBN initiator concentration on the polystyrene uptake
value(s) incorporated during precursor impregnation.
Initiator C oncentration Polystyrene-DVB Uptake
Per Im pregnation
0.43 wt % 14-16%
0.27 wt % 9-11%
0.11 w t% 5-6%
220
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Once the desired polystyrene uptake value had been achieved the membranes
were sulfonated by immersion in a CHCI3 solution o f 30% CISO3 H for 3-5 days
resulting in almost 1 sulfonic acid group per aromatic ring, as shown in figure 3.58.
Upon removal from the acid bath the membranes were hydrolyzed and washed in a
deionized water bath at 80 °C for an additional 3-4 days. During the sulfonation
procedure the membranes took on a dark black color that turned to yellow-orange
upon hydrolysis. The resulting membranes were flexible and exhibited excellent
physical properties.
Figure 3.58 Sulfonated polystyrene product from the 30% v/v chlorosulfonic
acid/chloroform bath.
CHC1
The thermal analysis o f the PVDF precursors, PVDF-PS membranes and
PVDF-PSSA polymer electrolytes were conducted by the Analytical Chemistry
Laboratory at the Jet Propulsion Laboratory in Pasadena, CA. Membrane samples
were provided to operator Gary Plett who conducted DSC analysis using a
Differential Scanning Calorimeter (Model 2910) made by T.A. Instruments, Inc.
The TGA experiments were performed using a Thermo-Mechanical Analyzer
221
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
(Model 2940) also made by T.A. Instruments, Inc. All thermal analysis experiments
were conducted on PVDF-PSSA membrane samples o f 20% PSSA uptake. The
Energy Dispersive X-Ray Analysis (EDAX) used to determine sulfonic acid content
and distribution within membrane samples was conducted by Ron Ruiz at the Failure
Analysis Laboratory located at JPL.
The membrane water content was determined according to the following
equation:
W eight (g) H20 - W eight (g) Dry
---------------------- X 100%
W eight (g) Dry
Membrane strips approximately 1 cm X 4 cm were dried in a vacuum oven at 90 °C
for 24 hours and then equilibrated in a DI water bath at 80-90 °C. The membrane
samples were removed every hour, padded dry to remove surface water, weighed and
then returned to the water bath. This process was repeated until the membranes
exhibited reproducible uptake values usually seen at a minimum bath time o f 4
hours.
The proton conductivity o f the PVDF-PSSA and Nafion membrane samples
was determined using a "four-probe" apparatus built in our laboratory that was
similar to a technique described by Cahan and Wainright.1 8 5 The membrane sample
is placed on a Pl'FE surface containing two platinum strip electrodes separated by a
fixed geometry (1.026 cm) and connected to a potentiostat, as shown in figure 3.59.
222
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Figure 3.59 Schematic of the apparatus used in the "four probe" technique to
measure the proton conductivity o f membrane samples (Reprinted from Reference
173).
Platinum
Electrodes
\
Membrane Sample
Platinum Foil
The IR drop measured across the membrane area exposed between the two platinum
electrodes is used to calculate the conductivity according to the following equation:
The 2-probe resistivity measurements used to assess bulk conductivity along
the transverse axis involved clamping a membrane sample between two graphite
electrodes using an O-ring joint. The bulk resistivity was then measured using a
Hewlit-Packard Milliohmeter and converted to conductivity values using the
aforementioned conductivity equation.
Our laboratory team also devised an apparatus used to measure the methanol
permeability characteristics o f the PVDF-PSSA and Nafion membrane samples. The
schematic diagram shown in figure 3.60 consists o f two vessels joined together by a
Rg = Membrane Resistance
L = Membrane Thickness
A = Membrane Area
223
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
ground down O-ring joint with a membrane sample clamped in between. Each
vessel has a 250 ml volume and the inner diameter o f the joint is 2.4 cm (area = 1.44
cm2).
Figure 3.60 Diagram o f the apparatus used to evaluate methanol permeability
(Reprinted from Reference 75).
3.0 M
CHjOH
Solution
Deion.
H,0
A B
The methanol permeability experiment involves adding 250 ml o f a 3.0 M
methanol solution to vessel A and 250 ml o f deionized water to vessel B. During the
experiment both vessels were stirred adequately to prevent any concentration build
up. Over specific periods o f time a 1 ml aliquot was removed from vessel B and
analyzed for methanol content using a Varian 3300 Gas Chromatograph equipped
with a car bo wax column. An internal standard, such as iso propanol, was used as a
reference for the determination o f concentration for each GC injection sample. The
calculated methanol concentrations were then plotted vs. time in order to determine
the mass transfer coefficient from the slope o f the plot. This value was then used to
determine the methanol diffusion coefficient according to the following equation:
224
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
m
gamma
C*
D A
t V
m = slope
C* = Concentration o f Vessel A
gamma = Mass Transfer Coefficient
D = Methanol Diffusion Coefficient
A = Area o f O-Ring Joint
t = membrane thickness
V = volume o f vessel B
The electrical testing o f the PVDF-PSSA MEAs using a 25 cm2 active area
was performed using a single cell apparatus purchased from Electrochem. The cell
consisted o f two serpentine flow fields for both the anode and cathode
compartments. The methanol concentrations tested at the anode ranged from 0.5 -
prevent evaporation but allow CO2 rejection from the system. No external heaters
were used to further heat the cell block and methanol was delivered to the anode at a
rate of 1.0-1.5 L/min. The cathode flow rates were measured prior to entering the
cathode compartment and pressurized conditions were maintained by the use o f a
valve located on the exit stream.
A CO 2 analyzer made by Horiba was used to determine the methanol
crossover rates by measuring the CO2 volume percent in the cathode exit stream.
Prior to cell testing the analyzer was calibrated using gases o f known CO2 content.
The CO2 volume percent generated from methanol crossover was then converted to
current density according to the following equations:
2.0 M and were heated at the fuel reservoir, which was equipped with a condensor to
225
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
(% Vol) (L/min) = L/min C02 = V
100
P = 1.00 atm = 0 PSIG
R = 0.0821 L-°K-Atm
PV = nRT T = 298 K
P V
n = -------- = moles/min
RT
V (1.0 atm) , , 0. --Q 4.
n (moles/sec) = - V (6.8 IE )
(0.0821) (298 K) (60 sec)
(96,487 C) (6 e-)
(n) ----------------------- = (23,156) (n) = A/cm
25 cm2
2
The electrical testing of the PVDF-PSSA ME As using an 80 cm2 active area
was performed using a 5-cell stack designed by our JPL team and constructed by
Giner, Inc. We also had an 80 cm2 single cell machined that was based on the same
stack design for use in the evaluation o f MEAs prior to stack assembly. The stack
design is currently proprietary but the methanol concentrations tested at the anode
and the method with which fuel was heated and delivered is the same as our testing
with the 25 cm2 single cell.
226
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
3.9 Conclusion
The use o f PVDF-PSSA se/w-sequential interpenetrating polymer networks
for application in direct methanol fuel cells that exhibit low crossover and high
performance has been demonstrated. In preparation o f membrane samples we
identified the factors affecting the degree of phase separation between polymers
during IPN synthesis. We determined that the degree and distribution o f crystalline
domains within the PVDF precursor dictated the distribution o f crosslinked-
polystyrene and ultimately polystyrene-sulfonic acid (upon sulfonation) during the
in-situ polymerization. Thermally extruded Kynar 740 precursors prepared from low
molecular weight polymer pellets had highly ordered crystalline domains of narrow
molecular weight distribution. Membrane samples prepared from Kynar 740
precursors resulted in polymer electrolytes with poor surface morphology attributed
to macroscopic phase separation o f PVDF and polystyrene. Compression molded
Kynar 460 precursors o f reduced crystallinity were prepared from high molecular
weight polymer powders and resulted in membranes with preferred surface
morphology suitable for membrane electrode fabrication and subsequent electrical
testing.
A number o f analytical techniques were used to characterize the PVDF-PS
and PVDF-PSSA membranes. Thermal analysis was used to characterize the PVDF
precursors), determine the physical interactions between polymer phases and
identify the thermal stability o f the polymer electrolytes. Differential Scanning
Calorimetry analysis was specifically used to identify the degree o f crystallinity
227
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
within the PVDF precursors) and to probe the degree o f miscibility between the
PVDF-PS and PVDF-PSSA phases. Thermogravimetric Analysis was used to
determine the thermal stability o f the polymer electrolytes, information critical to the
development o f MEA fabrication procedures. We also used Energy Dispersive X-
Ray Analysis (EDAX) to monitor the distribution o f PSSA within the PVDF matrix;
results used to modify our se/m'-sequential IPN methodology and screen MEA
candidates.
The water management properties of the PVDF-PSSA polymer electrolyte
membranes were also characterized and correlated strongly with membrane PSSA
uptake. We measured membrane water content and proton conductivity for various
membrane samples and determined that both properties increased linearly with PSSA
uptake and exhibited values comparable to Nafion®-117 above specific uptake
values. We also measured the methanol permeability rates for the PVDF precursor
and PVDF-PSSA polymer electrolytes and determined that permeability was
completely dependent upon the PSSA phase o f the membrane and increased with
membrane water content. However, even membrane samples that had comparable
membrane water content and proton conductivity to Nafion®-117 exhibited reduced
methanol diffusion coefficients by 75%. We also submitted PVDF-PSSA membrane
samples for independent analysis using 'H NMR techniques that reported reduced
water and methanol diffusion coefficients as compared to Nafion®-l 17.
Membrane candidates exhibiting preferred surface morphology and favorable
water management properties were fabricated into membrane electrode
228
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
assemblies with an active area o f 25 cm2. Based on MEA resistivity and electrical
performance we determined that the temperature and pressure used in the MEA
fabrication procedure strongly influenced interfacial contact at the catalyst-
membrane and electrode-electrolyte interface(s). Optimization o f membrane
parameters such as PSSA content and thickness coupled with modifications to our
MEA fabrication methodology resulted in low crossover, high performance PVDF-
PSSA MEAs.
Based on the promising electrical performance o f MEA 98-35 our JPL team
conducted parametric studies of stack and system requirements that established a
minimum baseline performance required by a MEA for an operating system. We
prepared one such 25 cm2 (active area) MEA (98-35) that demonstrated 0.401 V at
100 mA/cm2 with 0.50 M methanol at 55 °C utilizing 0.10 L/min ambient air at the
cathode; performance comparable to Nafion®-l 17. The methanol crossover rate was
25% that o f Nafion®-l 17 and translated into superior fuel utilization and a fuel cell
efficiency o f 28%. We delivered a 5-cell stack based on this baseline MEA that
exhibited poor electrical performance due to our inability to effectively scale-up the
MEA active area to 80 cm2 while maintaining necessary membrane properties such
as uniform thickness, preferred morphology and reproducible PSSA uptake.
In order to accommodate the 5-cell stack we adjusted our precursor
fabrication methodology to prepare membranes o f scaled-up active area size and
modified membrane properties in order to maintain important PSSA uptake-related
characteristics, thus establishing a second baseline membrane. We prepared one
229
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
such 25 cm2 (active area) MEA (98-95) that demonstrated 0.413 V at 100 mA/cm2
with 0.50 M methanol at 55 °C utilizing 0.10 L/min ambient air at the cathode. The
performance was comparable to our previous MEA (98-35) baseline but featured
methanol crossover 50% that o f Nafion®-l 17 realizing a slightly reduced fuel cell
efficiency o f 27%. However, the ability to effectively reproduce the PVDF
precursors) and PSSA uptake value o f the membrane(s), and relative ease with
which the MEA was conditioned, provided the opportunity to maintain stack and
system development.
In our efforts to prepare subsequent MEAs based on the modified PSSA
uptake baseline we were forced to redesign our MEA fabrication methodology to
accommodate increased membrane swelling. We developed one such method based
on a non-bonded MEA fabrication technique that resulted in comparable
performance (25 cm2 active area) as compared to MEAs prepared using our
previously established MEA fabrication methodology. This new method was also
used to prepare a scaled-up 80 cm2 MEA (99-18) that exhibited desirable electrical
performance consistent with the established stack requirements. We have prepared a
number o f membrane samples based on this second baseline that have exhibited
reproducible membrane properties and are dedicated to stack development and
testing.
230
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
4.0 References
177. Watkins, D., Dircks, K., Epp, D., Harkness, A., in "Proceedings o f 32n d
Power Sources Conference", Cherry Hill, NJ, June 9-12, 1986, The
Electrochemical Society, Inc., p. 590.
178. Raistrick, I.D., in "Diaphragms, Separators, and Ion Exchange Membranes",
Van Zee, J.W., White, R.E., Kinoshita, K., Burney, H.S., Editors, p. 172, The
Electochemical Society Softbound Proceeding Series, Pennington, New
Jersey, 1986.
179. Wilson, M.S., Valerio, J.A., Gottesfeld, S., Electrochimica Acta, 40, 3
(1995), 355-363.
180. Wilson, M.S., Gottesfeld, S., J. Appl. Electrochem. 22, 1 (1992).
181. Wilson, M.S., Gottesfeld, S., J. Electrochem. Soc. 139, L28 (1992).
182. Ren, X., Wilson, M.S., Gottesfeld, S., J. Electrochem. Soc. 143, L12 (1996).
183. Smart, M., "Chemical and Electrochemical Oxidation o f Small Organic
Molecules", UMI Ann Arbor, MI, 1998.
184. Narayanan, S.R., Smart, M., Prakash, G.K.S., Atti, A.R., Olah, G.A., Novel
Technology Report, RF 182, A0849.
185. Prakash, G.K.S., Smart, M., Atti, A., Olah, G.A., Narayanan, S.R., Valdez,
T., Surampudi, S., Abstract #409, The Electrochemical Society Meeting, 99-
2, October 17-22, Honolulu, HI, 1999.
186. Hodgdon, R.B., Boyack, J.R., J. Polymer Sci.: Part A 3, 1463 (1965).
187. Cahan, B., Wainright, J., J. Electrochem. Soc. 140, L185 (1993).
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
5.0 Bibliography
Allen, G., Bowden, M.J., Blundell, D.J., Hutchinson, F.G., Jeffs, G.M., Vyvoda, J.,
Polymer 14, 597 (1973).
Anastasiadis, S.H., Gancarz, I., Koberstein, J.T., Macromolecules 22, 1449-1453
(1989).
Appleby, A.J., Foulkes, F.R., "Fuel Cell Handbook", Krieger Publishing Co.,
Malabar, Florida, 1993.
Appleby, A.J., Yeager, E.B., Energy 11, 137 (1986).
Bachmann, M.A., Gordon, W., Koenig, J.L., Lando, J.B., J. Appl. Phys. 50, 6106
(1979).
Barensen, W.M., Heikens, D., Piet, P., Polymer 15, 119 (1974).
Barlow, J.W., Paul, D.R., Polym. Eng. Sci. 27, 1482 (1987).
Barlow, J.W., Paul, D.R., Polym. Eng. Sci.24, 525 (1984).
Benedetti, E., Catanorchi, S., D'Alessio, A., Moggi, G., Vergamini, P., Pracella, M.,
Ciardelli, F., Polymer International 41, 35-41 (1996).
Bentvelzen, J.H., Kimura-Yeh, F., Hopfenberg, H.B., Stannet, V., J. Appl.Polym.Sci.
17 (1973) 809.
Bouchet, R., Siebert, E., Solid State Ionics 118, 287-299 (1999).
Buchi, F.N., Gupta, B., Haas, O., Scherer, G.G., Electrochim. Acta 40, 345 (1995).
Buchi, F.N., Gupta, B., Haas, O., Scherer, G.G., J. Electrochem. Soc. 142, 3044
(1995).
Buchi, F.N., Gupta, B., Haas, O., Scherer, G.G., J. Electrochem. Soc. 142, 3044
(1995).
Cahan, B., Wainright, J., J. Electrochem. Soc. 140, L I85 (1993).
Ceynowa, J., Polymer 19, 73 (1978).
Chapman, L.E., Proc. 7th Intersoc. Energy Conv. Eng. Conf.', 466 (1972).
232
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Cowie, in "Integration o f Fundamental Polymer Science and Technology",
(Eds. Kleintjens, L.A., Lemstra, P.J.), Elsevier, London, 1986, p. 138.
Del Rio, C., Acosta, J.L., Eur. Polym. J. 32, 7 (1996) 913-917.
Discussions with E lf Atochem, N.A.
Donatelli, A.A., Sperling, L.H., Thomas, D.A., Macromolecules 9, 671 (1976).
Eisenberg, A., King, M., "Ion Containing Polymers" Academic, New York, 1977.
Eisenberg, A., Yeager, H.L., (eds.), "Perfuorinated Ionomer Membranes" ACS
Symposium Series 180, American Chemical Society, Washington, D.C.,
1982.
Encyclopedia o f Polymer Science and Engineering, 17, 2n d Ed., p. 532, John Wiley
& Sons, Inc., 1989.
Falk, M., Can. J. Chem. 58, 1495 (1980).
Falk, M., in Perfluorinated Ionomer Membranes, ACS Series, 1982, Ch.8, p. 139.
Fayt, R., Jerome, R., Teyssie, Ph., J. Polym. Sci. Polym. Lett. Ed. 24, 25 (1986).
Ferguson, R.C., Baume, E.G., J. Phys. Chem. 83, 1379 (1979).
Flint, S.D., Slade, R.C.T., Solid State Ionics 97, 299-307 (1997).
Frisch, K.C., Klempner, D., Mukheijee, S.K., Frisch, H.L., J. Appl. Polym. Sci. 18,
689(1974).
Fuller, T.F., Newman, J., J. Electrochem. Soc. 140, 1218 (1993).
Gailard, P., Ossenbach-Sauter, M., Riess, G., Makromol. Chem., Rapid Commun. 1,
771 (1980).
Gierke, T.D., Munn, G.E., Wilson, F.C., J. Polym. Sci., Polym. Phys. Ed. 19, 1687
(1981).
Gorlitz, M., Minke, R., Trautvetter, W., Weisberger, G., Angew. Makromol. Chem.
29/30, 137(1973).
Greenbaum, S.G., Work in Progress.
233
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Gregorio, R., Cestari, M., Journal o f Polymer Science: Part B: Polymer Physics, 32,
859-870 (1994).
Grotthuss, C.J.D., Ann. Chim. 58, 54 (1806).
Grove, W.R., Phil. Mag. 14, 127 (1839).
Grubb, W.T., J. Electrochem. Soc. 106, 275-278 (1959).
Grubb, W.T., Niedrach, L.W., J. Electrochem. Soc. 107, 131-134(1960).
Guerra, G., Karasz, F.E., MacKnight, W.J., Macromolecules 19, 1935 (1986).
Guo, Q., Pintauro, P.N., Tang, H., O'Connor, S., Journal o f Membrane Science 154,
175-181 (1999).
Harrison, J.W., (General Electric Company), the Potential of Fuel Cells for
Generation o f Electricity from Hydrogen, G.E. Direct Energy Conversion o f
Electricity from Hydrogen, G.E. Direct Energy Conversion Programs, Rep.
(undated).
Hietala, S., Koel, M., Skou, E., Elomaa, M., Sundholm, F., J. Mater. Chem. 8(5),
1127-1132(1998).
Hinatsu, J.T., Mizuhata, M., Tokenaka, H., J. Electrochem. Soc. 141 (1994) 1493.
Hodgdon, R.B., Boyack, J.R., J. Polymer Sci.: Part A 3, 1463 (1965).
Hodgdon, R.B., Enos, J.F., Aiken, E.J., (U.S.P. 3,341,366; September 12, 1967;
assigned to General Electric Company).
Holiday, L., (ed.), "Ionic Polymers", Halstead, New York, 1975.
Hourston, D.J., Huson, M.G., Journal o f Applied Polymer Science 46, 973-979
(1992).
Hudson, S.D., Davis, D.D., Lovinger, A.J., Macromolecules 25, 1759 (1992).
Huelk, V., Thomas, D.A., Sperling, L.H., Macromolecules 5, 348 (1972).
Jasso, C.F., Martinez, J.J., Mendizabal, E., Laguna, O., Journal of Applied Polymer
Science, 58, 2207 (1995).
K. Kordesch, G. Simader, "Fuel Cells and Their Applications", VCH
Verlagsgesellschaft mbH, Weinheim, Germany, 1996.
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Keith, H.D., Padden, F.J., J. Appl. Phys. 34, 2409 (1963).
Keith, H.D., Padden, F.J., J. Appl. Phys. 35, 1270 (1964).
Kim, B.S., Lee, D.S., Kim, S.C., Macromolecules 19, 2589-2593 (1986).
Kim, K J., Cho, Y.J., Kim, Y.H., Vibrational Spectroscopy, 9, 147-159 (1995).
Kim, S.C., Klempner, D., Frisch, K.C., Radigan, W., Frisch, H.L., Macromolecules
2, 258 (1976).
Kobayashi, T., Rikukawa, M., Sanui, K., Ogata, N., Solid State Ionics 106, 219-225
(1998).
Kreuer, K.D., Chem. Mater. 8, 610-641 (1996).
Kreuer, K.D., Weppner, W., Rabenau, A., Angew. Chem., Int. Ed. Engl. 21, 208
(1982).
LaConti, A.B., (General Electric Company), Introduction to SPE Coal Technology,
G.E. Direct Energy Conversion Programs, Rep. (undated).
Lee, D.S., Kim, S.C., Macromolecules 17, 268-272 (1984).
Lee, D.S., Kim, S.C., Macromolecules 18, 2173-2179 (1985).
Lee, J.H., Kim, S.C., Macromolecules 19, 644-648 (1986).
Litt, M.H., Lando, J.B., J. Polym. Sci. Polym. Phys. Ed. 20, 535-552 (1982).
Liu, Z., Marechal, P., Jerome, R., Polymer 38, 19, 4925-4929 (1997).
Ludwig, B.W., Urban, M.W., Polymer 39, 24 (1998).
Maricle, D.L., Murach, B.L., Van Dine, L.L., Abstract 35, The Electrochemical
Society Extended Abstracts, Vol. 94-1, San Francisco, CA, meeting, May 22-
27, 1994.
Marie, P., Selb, J., Rameau, A., Duplessix, R., Gallot, Y., in "Polymer Blends and
Mixtures", Walsh, , D.J., Higgins, J.S., Maconnachie, A., Eds.; NATO
Advanced Science Institute Series; Mortinus Nijhoff: Dordrecht, The
Netherlands, 1985; p. 449.
Marvel, C.S., Denon, C.E., J. Am. Chem. Soc., 60, 280 (1938).
235
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Mattem, D.E., Fu-Tyan, L., Hercules, D.M., Anal. Chem. 56, 2762-2769 (1984).
McElroy, J.F., (General Electric Company), Status o f Solid Polymer Electrolyte Fuel
Cell Technology, G.E. Direct Energy Conversion Program, Rep. (undated).
Michaels, A., Baddour, R.F., Bixter, H.J., Choo, C.Y.,
Ind. Eng. Chem. Process. Des. Dev. 1 (1962) 14.
Mijovic, J., Luo, H.L., Han, C.D., Polym. Eng. Sci. 22(4), 234 (1982).
Millar, J.R., J. Chem. Soc. 1311 (1960).
Moore, R.B., Martin, C.R., Macromolecules 21, 1334-1339 (1998).
Morel, G., Jozefonvicz, J., ApteL, P., Journal o f Applied Polymer Science 23, 2397-
2407 (1979).
Morra, B.S., Stein, R.S., J. Polym. Sci., Polym. Phys. Ed. 20, 2261 (1982).
Nandan, D., Pushpa, K., Kartha, P.K., Wahi, P.K., Iyer, R.M., Indian J. Chem. 33A,
395 (1994).
Nandon, D., Mohan, H., Iyer, R.M., J. Membrane Sci. 71, 69 (1992).
Narayanan, S.R., Frank, H., Jeffries-Nakamura, B., Smart, M., Chun W., Halpert,
G., Kosek, J., Cropley, C., "Studies on Methanol Crossover in Liquid-Feed
Direct Methanol PEM Fuel Cells", Electrochemical Proceedings Volume 95-
23, 278 (1995).
Narayanan, S.R., Kindler, A., Jeffries-Nakamura, B., Chun. W., Frank, H., Smart,
M., Surampudi, S., Halpert, G., "Performance o f PEM Liquid Feed Direct
Methanol-Air Fuel Cells", Electrochemical Society Proceedings Volume 95-
23, 261 (1995).
Narayanan, S.R., Kindler, A., Jeffries-Nakamura, B., Chun W., Frank, H., Smart,
M., Valdez, T., Surampudi, S., Halpert, G., Kosek, J., Cropley, C., "Recent
Advances in PEM Liquid-Feed Direct Methanol Fuel Cells", Proceedings o f
the IEEE 11th Annual Battery Conference on Applications and Advances,
p.l 13, Long Beach, CA (Jan 1996).
Narayanan, S.R., Kindler, A., Jeffries-Nakamura, B., Chun. W., Frank, H., Smart,
M., Surampudi, S., Halpert, G., "Performance o f PEM Liquid Feed Direct
Methanol-Air Fuel Cells", Electrochemical Society Proceedings Volume 95-
23, 261 (1995).
236
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Narayanan, S.R., Proceedings o f the 11th Annual Battery Conference on Applications
and Advances, Long Beach, CA., Jan 9-12, 1996. IEEE pg. 113-122.
Narayanan, S.R., Smart, M., Prakash, G.K.S., Atti, A.R., Olah, G.A., Novel
Technology Report, RF 182, A0849.
Narayanan, S.R., Vamos, E., Surampudi, S., Frank, H., Halpert, G., Abstract 73,
p. 126, The Electrochemical Society Extended Abstracts, Vol 93-2, New
Orleans, LA., Meeting, October, 1993.
Natansohn, A., Murali, R., Eisenberg, A., Chemtech 6,1990.
Nguyen, T.V., White, R.E., J.Electrochem.Soc. 140, 2178 (1993).
Nohe, R.. Ledjeff, K., Bauer, M., Mulhaupt, R., Journal o f Membrane Science 83,
211-220(1993).
Odian, G., "Principles o f Polymerization", 3rd Edition, Wiley-Interscience, 1991.
Olabisi, O., Robeson, L.M., Shaw, M.T., in "Polymer-Polymer Miscibility", Chpt. 3,
Academic, New York (1979).
Ostrovskii, D., Brodin, A., ToreU, L.M., Solid State Ionics 85, 323 (1996).
Ostrovskii, D.I., Torell, L.M., Paronen, M., Hietala, S., Sundholm, F., Solid State
Ionics 97, (1997) 315-321.
Ouhadi, T., Fayt, R., Jerome, R., Teyssie, Ph., J. Appl. Polym. Sci. 32, 5647 (1986).
Ouhadi, T., Fayt, R., Jerome, R., Teyssie, Ph., J. Polym. Sci., Polym. Phys. Ed 24,
973 (1986).
Ouhadi, T., Fayt, R., Jerome, R., Teyssie, Ph., Polym. Commun. 27, 212 (1986).
Papavoine, C.H.M., Maas, W.E.J.R., Veeman, W.S., Buning, G.H.W., Vankan,
J.M.J., Macromolecules 26, 6611-6616 (1993).
Parthasarathy, A., Dave, B., Srinivasan, S., Appleby, A.J., Martin, C.R., J.
Electrochem. Soc. 139, 1634-1640 (1992).
Paul, D.R., Barlow, J.W., J. Macromol. Sci. Rev. Macromol. Chem. C18, 109 (1980).
237
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Paul, D.R., Barlow, J.W., Keskkula, H., in "Polymer Blends", pp.399-461 in
"Encyclopedia o f Polymer Science and Engineering," Vol. 12, Mark, H.F.,
Bikales, N.M., Overberger, C.G., Menges, G., Eds., Wiley-Interscience, New
York, 1988.
Paul, D.R., in "Polymer Blends", Paul, D.R., Newman, S., Eds., Academic Press,
New York, 1978, 2, p.35.
Paul, D.R., Newman, S., in "Polymer Blends", Chpt. 5, Academic, New York
(1978).
Petermann, J., Gohill, R.M., J. Mater. Sci. 14, 2260, (1979).
Pourcelly, G., Oikonomou, A., Hurwitz, H.D., Gavach, C., J. Electroanal. Chem. 43,
287(1990).
Prakash, G.K.S., Smart, M., Atti, A., Olah, G.A., Narayanan, S.R., Valdez, T.,
Surampudi, S., Abstract #409, The Electrochemical Society Meeting, 99-2,
October 17-22, Honolulu, HI, 1999.
Prakash, G.K.S., Smart, M.C., Wang, Q.J., Atti, A., Olah, G.A., Narayanan, S.R.,
Chun, W., Valdez, T., Surampudi, S., J. Electrochem. Soc. Paper to be
submitted.
Przyluski, J., Dabrowski, A , Stys, S., Wieczorek, W., Solid State Ionics 60, 141-146
(1993).
Pu, C., Huang, W., Ley, K.L., Smotkin, E.S., J. Electrochem. Soc. 142, 7 (1995).
Raistrick, I.D., in "Diaphragms, Separators, and Ion Exchange Membranes", Van
Zee, J.W., White, R.E., Kinoshita, K., Burney, H.S., Editors, p. 172, The
Electochemical Society Softbound Proceeding Series, Pennington, New
Jersey, 1986.
Ramos, A.R., Cohen, R.E., Polym. Eng. Sci. 17, 639 (1977).
Ravikumar, M.K., Shukla, A.K., J. Electrochem. Soc. 143, 8 (1996).
Ren, X., Henderson, W., Gottesfeld, S., J. Electrochem. Soc. 144, 9 (1997).
Ren, X., Wilson, M.S., Gottesfeld, S., J. Electrochem. Soc. 143, L13 (1996).
Ren X., Wilson, M.S., Gottesfeld, S., Electrochem. Soc. Proc., PV 95-23 (1995) The
Electrochemical Society, Pennington, NJ.
238
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Ren, X., Wilson, M.S., Gottesfeld, S., J. Electrochem. Soc. 143, L12 (1996).
Ren, X., Zawodzinski, T.A., Uribe, F., Dai. H., Gottesfeld, S., in Proc.First
Int.Symp.Proton Conducting Membrane Fuel Cells I, Gottesfeld, S., Halpert,
G., Landgrebe, A., (Eds.) The Electrochemical Society, 1995.
Riess, G., Jolivet, Y., in "Copolymers, Polyblends, and Composites", Platzer, N.A.J.,
Ed.; Advances in Chemistry Series 142; American Chemical Society:
Washington, D.C., 1975; p. 243.
Rosenberg, Y., Sigmann, A., Narkis, M., Shkolnik, S., J. Appl. Polym. Sci. 43, 535
(1991).
Rouilly, M.V., Kotz, E.R., Haas, O., Schere, G.G., Chapiro, A., J. Membrane Sci. 81,
89-95 (1993).
Saito, H., Stuhn, B., Macromolecules 27, 216-218 (1994).
Savinell, R., Yeager, E., Tryk, D., Landau, U., Wainright, J., Weng, D., Lux, K., Litt,
M., Rogers, C., J. Electrochem. Soc. 141, 4 (1994).
Scherer, G.G, Solid State Ionics 94, 249-257 (1997).
Sergeeva, L.M., Gorbach, L.A., Russian Chemical Reviews 65, (4) 345-354 (1996).
Shibayama, K., Kobunshi Kaga/cu 19, 219 (1962).
Shibayama, K., Kobunshi Kagaku 20, 221 (1963).
Shibayama, K., Suzuki, Y., Kobunshi Kagaku 23, 24 (1966).
Shibayama, K., Zairyo 12, 362 (1963).
Shukla, A.K., Christensen, P.A., Hamnett, M., Hogarth, M., J. Power Sources 55, 87
(1997).
Siquiera, D.F., Galembeck, F., Nunes, S.P., Polymer 32, 6 (1991).
Skou, E., Kauranen, P., Hentschel, J., Solid State Ionics 97 333-337 (1997).
Smart, M., "Chemical and Electrochemical Oxidation o f Small Organic Molecules”,
UMI Ann Arbor, MI, 1998.
Sondheimer, S.J., Bunce, N.J., Fyfe, C.A., JMS-Rev. Macromol. Chem. Phys.
C26(3), 353-413 (1986).
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Sperling, L.H., "Interpenetrating Polymer Networks and Related Materials", Plenum
Press, New York, 1981.
Sperling, L.H., Chiu, T.W., Hartman, C.P., Thoma, D.A., Int. J. Polym. Mater. 1,
331 (1972).
Sperling, L.H., Chiu, T.W., Thomas, D.A., J. Appl.Poly.Sci. 17, 2443 (1973).
Sperling, L.H., Donatelli, A.A., Thomas, D.A., in "Recent Advances in Polymer
Blends, Grafts and Blocks", Plenum Press, New York, 1974, p. 375.
Sperling, L.H., Friedman, D.W., J. Polym. Sci., Part A-2 7, 425 (1969).
Sperling, L.H., George, H.F., Huelck, V., Thomas, D.A., J. Appl. Sci. 14, 2815
(1970).
Sperling, L.H., Mihalakis, E.N., J. Appl.Polym.Sci. 17, 3811 (1973).
Sperling, L.H., Sarge, H.D., J. Appl.Polym.Sci. 16, 3041 (1972).
Sperling, L.H., Taylor, D.W., Kirkpatrick, M.L., George, H.F., Bardman, D.R., J.
Appl. Polym. Sci. 14, 73 (1970).
Sperling, L.H., Thomas, D.A., Covitch, M.J., Curtius, A J., Polym. Eng. Sci. 12, 101
(1972).
Sperling, L.H., Thomas, D.A., Huelk, V., Macromolecules 5, 340 (1972).
Springer, T., Zawodzinski, T.A., Gottesfeld, S., J. Electrochem. Soc. 140, 1981
(1993).
Stein, R.S., Khambatta, F.B., Warner, F.P., Russell, T., Escala, A., Balizer, E., J.
Polym. Sci., Polym. Symp. 63, 313 (1978).
Sugiyama, S., Tsuneda, S., Saito, K., Furusaki, S., Sugo, T., Makuuchi, K., React.
Polym. 21, 187-191 (1993).
Surampudi, S., Narayanan, S.R., Vamos, E., Frank, H., Halpert, G., LaConti, A.,
Kosek, J., Prakash, G.K.S., Olah, G.A., J. Power Sources 47, 377 (1994).
Tazaki, M., Wada, R., Okabe, M., Homma, T., Kobunshi Ronbunshu, 50, 533 (1993).
Tialdo, G.C., Canepa, P., Munari, S., Journal o f Membrane Science 9, (1981) 191-
196.
240
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Uribe, F.A., Springer, T.E., Gottesfeld, S., J. Electrochem. Soc. 139 765 (1992).
Valdez, TJ., Narayanan, S.R., Rohatgi, N., "Development o f a 150-W Direct
Methanol Fuel Cell System” Proceedings of the 15th Annual Battery
Conference on Applications and Advances, Long Beach, CA., Jan 11-14,
2000. IEEE pg. 37-40, Eds.
Das, R.S.L., Frank, H., Valdez, T J., Proceedings o f Proton Conducting Membrane
Fuel Cells II, ECS Proceeding Series, Pennington, NJ, 1998 pg. 380-387.
Verbrugge, M.R., J. Electrochem. Soc. 136, 2 (1989).
Verbrugge, M.R., J. Electrochem. Soc. 136, 417 (1989).
Walsh, D.J., in "Integration o f Fundamental Polymer Science and Technology",
(Eds. Kleintjens, L.A., Lemstra, P.J.), Elsevier, London, 1986, p. 119.
Wang, J.T., Wasmus, S., Savinell, R.F., J. Electrochem. Soc. 143,4 (1996).
Warner, F.P., Macknight, W.J., Stein, R.S., J. Polym. Sci. Polym. Phys. Ed. 15, 2113
(1977).
Watkins, D., Dircks, K., Epp, D., Harkness, A., in "Proceedings o f 32n d Power
Sources Conference", Cherry Hill, NJ, June 9-12, 1986, The Electrochemical
Society, Inc., p. 590.
Weinhold, S., Litt, M.H., Lando, J.B., J. Polym. Sci. Polym. Lett. Ed. 17, 585 (1979).
Widmaier, J., Macromol. Symp. 93, 179-186 (1995).
Widmaier, J.M., Tabka, M.T., Eur. Polym. J. 28, 5 (1992).
Wilson, M.S., Gottesfeld, S., J. Appl. Electrochem. 22, 1 (1992).
Wilson, M.S., Gottesfeld, S., J. Electrochem. Soc. 139, L28 (1992).
Wilson, M.S., Valerio, J.A., Gottesfeld, S., Electrochimica Acta, 40, 3 (1995), 355-
363.
Yang, D., Thomas, E.L., J. Mater. Sci. Lett. 3, 929 (1984).
Yeager, H.L., Chap. 4 in Reference 22.
Yeo, R.S., Chan, S.F., Lee, 5., Journal o f Membrane Science 9 273-283 (1981).
241
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.
Yeo, R.S., J. Electrochem. Soc. 130, 533 (1983).
Zawodzinski, T.A., Derouin, C., Radzinski, S., Sherman, R.J., Smith, V.T., Springer,
T.E., Gottesfeld, S., J. Electrochem. Soc. 140, 1041 (1993).
Zawodzinski, T.A., Dorouin, S., Radzinski, S., Sherman, R.J., Smith, V.T., Springer,
T.E., Gottesfeld, J. Electrochem. Soc., 140, (1993) 1041.
Zawodzinski, T.A., Gottesfeld, S., Shoichet, S., McCarthy, J., J. Appl. Electrochem.
23 (1993) 86.
Zawodzinski, T.A., Springer, T., Davey, J., J. Electrochem. Soc. 140, 1981 (1993).
Zhou, P., Frisch, H.L., Journal o f Polymer Science: Part A: Polymer Chemistry 30,
887-893 (1992).
Reproduced with permission of the copyright owner. Further reproduction prohibited without permission.

Linked assets

University of Southern California Dissertations and Theses

Conceptually similar

PDF

Electrochemical studies of membranes, catalysts, and fuels for direct oxidation fuel cells

PDF

Investigation of silylated onium ions

PDF

Investigations in selective fluorinations: Novel synthetic methodologies and material syntheses

PDF

Activation and functionalization of carbon-hydrogen bonds catalysed by oxygen ligated iridium metal complexes

PDF

Design and synthesis of novel lipid mediators

PDF

Design and synthesis of novel heterocycles and peptidomimetics from organoboronic acids, amines and carbonyl compounds

PDF

ipso substitutions of aryl boronic acids and aryltrifluoroborates

PDF

Association of perfluorocarbon containing water soluble polymers

PDF

Fluorinated carbocations and carboxonium ions

PDF

Synthesis And Characterization Of Functionalized Silicon Polymers

PDF

Crosslinked siloxanes: Preparation and properties

PDF

I. Metal nanoparticles on high surface area latex supports: Preparation and applications, and, II. Modified semiconductor quantum dots as luminescent label for in situ hybridizations

PDF

Design and synthesis of foscarnet -peptide conjugates, halogenated bisphosphonates and phosphonocarboxylates and biochemical sensing using film -bulk -acoustic -resonators

PDF

Chemical and electrochemical oxidation of small organic molecules

PDF

Alkylamino-substituted styrylpyrazine two -photon fluorophores: Structure-nonlinear optical property relationships and the effect of the chemical environment

PDF

Superelectrophilic Reactions Under Superacidic Catalysis

PDF

Stereo controlled synthesis of substituted cyclic ethers

PDF

Development of polymer electrolyte membranes for fuel cell applications

PDF

Design and syntheses of polymeric materials for visible and near -infrared emitting applications

PDF

Atropisomerism In Sterically Hindered Carbocations

Asset Metadata

Creator Atti, Anthony Richard (author)

Core Title Development of high-performance polymer electrolyte membranes for direct methanol fuel cells

School Graduate School

Degree Doctor of Philosophy

Degree Program Chemistry

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag chemistry, organic,chemistry, polymer,energy,engineering, materials science,OAI-PMH Harvest

Language English

Contributor Digitized by ProQuest (provenance)

Advisor Prakash, G.K. Surya (committee chair), Olah, George A. (committee member), Weber, William P. (committee member)

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c16-148856

Unique identifier UC11334730

Identifier 3042408.pdf (filename),usctheses-c16-148856 (legacy record id)

Legacy Identifier 3042408-0.pdf

Dmrecord 148856

Document Type Dissertation

Rights Atti, Anthony Richard

Type texts

Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the au...

Repository Name University of Southern California Digital Library

Repository Location USC Digital Library, University of Southern California, University Park Campus, Los Angeles, California 90089, USA