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Electrochemical studies on selected oxides for intermediate temperature-solid oxide fuel cells

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Electrochemical studies on selected oxides for intermediate temperature-solid oxide fuel cells
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Jaiswal, Abhishek
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Anodes ( jstor )
Bismuth ( jstor )
Cathodes ( jstor )
Conductivity ( jstor )
Electrodes ( jstor )
Electrolytes ( jstor )
Ions ( jstor )
Oxides ( jstor )
Oxygen ( jstor )
Sintering ( jstor )

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ELECTROCHEMICAL STUDIES ON SELECTED OXIDES FOR INTERMEDIATE
TEMPERATURE SOLID OXIDE FUEL CELLS














By

ABHISHEK JAISWAL


A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY

UNIVERSITY OF FLORIDA


2004
































Copyright 2004

by

Abhishek Jaiswal
































Dedicated to my parents.















ACKNOWLEDGMENTS

I would like to thank my advisor, Dr. Eric Wachsman, for his guidance,

encouragement, and support through the course of my graduate study. It has been a

highly enriching experience working with him. I would also like to thank him for

allowing me to visit various conferences to gain exposure in the field.

I would like to thank my committee members: Dr. Daryl Butt, Dr. David Norton,

Dr. Wolfgang Sigmund and Dr. Mark Orazem. Dr. Butt and Dr. Sigmund allowed me to

use their laboratory facilities for my research work.

I would like to acknowledge my group members, especially Jamie Rhodes and

Guojing Zhang, who made the time at UF a very memorable experience. I would like to

thank Dr. Hee-Sung Yoon, Dr. Keith Duncan and Dr. Suman Chatterjee for their advice

and help. I would also like to thank Wayne Acre for helping me in using the various

characterization facilities at MAIC.

Lastly, I would thank my parents and my brothers. They have been very kind to

bear long separations just to let me achieve my dreams. My mother has not been in good

health for the last three years and it has been a tough time for the whole family. I pray

that I can justify their sacrifices.















TABLE OF CONTENTS

page

ACKNOW LEDGM ENTS ............................................................................................ iv

LIST OF FIGURES .......................................................................................................... vii

ABSTRACT ..................................................................................................................... xiv

CHAPTER

I FUEL CELLS: A GREEN SOLUTION ....................................................................... 1

1.1 Introduction ......................................................................................................... 1
1.2 Operating Principle of SOFCs .......................................................................... 4
1.3 Designing SOFCs ............................................................................................... 6
1.4 SOFC Structures ................................................................................................. 7
1.4.1 Tubular SOFCs ........................................................................................ 8
1.4.2 Planar SOFCs ........................................................................................ 10
1.4.3 New Designs: Metal Supported and Microtubular SOFCs ................... 11
1.5 Conclusion ........................................................................................................ 13


2 IONIC CONDUCTION IN SOLID ELECTROLYTES ........................................ 19

2.1 Introduction ...................................................................................................... 19
2.2 Fluorite Type Oxides ........................................................................................ 20
2.2.1 Structural Aspects .................................................................................. 20
2.2.2 Electrical Aspects ................................................................................. 20
2.2.3 Oxygen Ion Conductivity as a Function of Dopant Concentration and
Temperature ................................................................................................. 22
2.2.4 Grain Boundary Contribution to Total Conductivity ............................ 26
2.2.5 Oxygen Ion Conductivity as a Function of Time ................................. 27
2.3 Conclusion ........................................................................................................ 28


3 DIRECT CURRENT BIAS STUDIES ON (Bi2O3)os(Er2O3)0.2 ELECTROLYTE
AND Ag-(Bi203)o.8(Er2O3)0.2 CERMET ELECTRODE ........................................ 36

3.1 Introduction ...................................................................................................... 36
3.2 Experim ental .................................................................................................... 39








3.3 Results and Discussion ................................................................................... 40
3.3.1 Processing ............................................................................................. 40
3.3.2 DC Bias and Impedance Spectroscopy: ESB Electrolyte ...................... 41
3.3.3 DSC Studies: ESB Electrolyte ............................................................... 43
3.3.4 DC Bias and Impedance Spectroscopy: Ag-(Bi2O3)0.8(Er2O3)0.2 Electrode45
3.4 Conclusion ........................................................................................................ 48


4 BISMUTH RUTHENATE BASED CATHODES FOR IT-SOFC ....................... 66

4.1 Introduction ...................................................................................................... 66
4.1.1 Undoped Bismuth Ruthenate ................................................................. 66
4.1.2 Doped Bismuth Ruthenates ................................................................... 68
4.1.3 Bismuth Ruthenate and Stabilized Bismuth Oxide Composites ........... 70
4.2 Experimental W ork ........................................................................................... 71
4.3 Results and Discussion ................................................................................... 72
4.3.1 Processing ............................................................................................... 72
4.3.2 Impedance Spectroscopy Studies .......................................................... 73
4.3.2.1 Undoped Bi2Ru207 cathode ........................................................ 73
4.3.2.2 Doped Bi2Ru207 cathodes .......................................................... 75
4.3.2.3 Bi2Ru2O7.3-(Bi2O3)0.8(Er2O3)0.2 composite cathodes ................... 79
4.4 Conclusion ........................................................................................................ 81


5 ANODE SUPPORTED THICK FILM CERIA ELECTROLYTE UNIT CELLS FOR
IT-SOFC ................................................................................................................... 111

5.1 Introduction ......................................................................................................... 111
5.2 Experimental ....................................................................................................... 115
5.3 Results and Discussion ....................................................................................... 116
5.3.1 Initial Fabrication and Performance Results ............................................ 116
5.3.2 Improving the Performance of Anode Supported Ceria Unit Cells ......... 121
5.3.2.1 Reducing the co-sintering temperature of anode/electrolyte
bilayer ......................................................................................................... 123
5.3.2.2 Optimization of LSCF-GDC cathode sintering temperature .......... 124
5.3.2.3 Composite cathodes containing ESB as the electrolyte phase ....... 129
5.4 Conclusion .......................................................................................................... 131


6 CONCLUSION AND FUTURE W ORK ................................................................. 165

LIST OF REFERENCES ................................................................................................. 170

BIOGRAPHICAL SKETCH ........................................................................................... 177















LIST OF FIGURES


page

1-1 Fuel cell types with typical reactants ................................................................... 14

1-2 Schematic diagram of the current-voltage characteristics as well as the different
losses in a solid oxide fuel cell ............................................................................. 14

1-3 Placement of reactant energies relative of the edges of conduction and valence band
of the electrolyte in a thermodynamically stable electrochemical cell ............... 15

1-4 Siemens Westinghouse tubular SOFC design ...................................................... 15

1-5 Siemens Westinghouse flattened tubular SOFC design ...................................... 16

1-6 Theoretical and actual performance of Siemens Westinghouse tubular and flattened
tubular (HPD ) SOFC designs ............................................................................... 16

1-7 Planar SO FC design ............................................................................................ 17

1-8 Performance of a planar SOFC design at 800 'C ................................................. 17

1-9 Microstructure of an alloy supported SOFC based on YSZ electrolyte .............. 18

1-10 Microstructure of an alloy supported SOFC based on doped cerium oxide
electrolyte ................................................................................................................. 18

2-1 Oxygen ion conductivity of common solid electrolytes ...................................... 30

2-2 The fluorite structure .......................................................................................... 31

2-3 Isothermal conductivity of some ceria solid solution at temperature close to 200
.C .............................................................................................................................. 3 1

2-4 Schematic Arrhenius plots of lattice and grain boundary conductivity for clean
Ceo9Gdo.0O195, impure Ceo.9Gd0 101.95, and impure Ceo.8Gd.201.9 ................... 32

2-5 Conductivity of (Bi203)1-x,(Er2O3), in air ............................................................ 33

2-6 Decay in isothermal oxygen ion conductivity of 20 mol% ESB on annealing below
the transition tem perature ................................................................................... 34









2-7 Time constant as function of cation radii and lattice parameter of 25 mol% cation
doped bism uth oxide .......................................................................................... 35

2-8 Time constant as function of erbia concentration and lattice parameter of erbia
stabilized bism uth oxide ...................................................................................... 35

3-1 XRD patterns of calcined 20 mol% erbia stabilized bismuth oxide powders derived
from am orphous citrate route ............................................................................... 50

3-2 Cross-sectional micrographs of the ESB electrolyte and Ag-ESB electrode
interface .................................................................................................................... 5 1

3-3 Arrhenius plot of oxygen ion conductivity in ESB electrolyte ........................... 52

3-4 Percentage change in resistance of the ESB electrolyte under no bias at 300 'C, 500
'C and 625 'C .......................................................................................................... 53

3-5 Initial voltage vs. time plots at 500 'C under different bias currents .................. 54

3-6 Percentage change in resistance of the ESB electrolyte at 500 'C under different
bias currents ........................................................................................................ 55

3-7 Percentage change in resistance of the ESB electrolyte at 625 'C under different
bias currents ........................................................................................................ 56

3-8 Arrhenius plot of ASR (Klcm2) for Ag-ESB electrode ........................................ 58

3-9 Impedance plots of Ag-ESB electrode at 500 'C under different bias currents ....... 59

3-10 Ag-ESB electrode resistance vs bias current at 500 'C and 625 'C ................... 60

3-11 Impedance plots of Ag-ESB electrode at 625 'C under different bias currents ....... 61

3-12 Impedance plots of Ag-ESB electrode under 250 mA bias at 625 'C ................. 62

3-13 Cross-sectional microstructures of the electrolyte/electrode interface after annealing
experiments at 625 'C under 250 mA bias for -26 hours ................................... 63

3-14 Proposed electrode reaction mechanism for Ag-ESB electrodes ........................ 64

3-15 Phase diagram between Ag and Ag20 proposed by Assal et al .......................... 65

4-1 Phase diagram between Bi203 and Ru02 proposed by Prosychev et al .............. 83

4-2 Phase diagram between Bi203 and RuO2 proposed by Hrovat et al .................... 83

4-3 Cation sublattice for one-quarter unit cell of the pyrochlore structure ............... 84

4-4 XRD patterns for Bi2Ru2O7 after calcination at 900 'C for 10 hours .................. 85








4-5 XRD patterns for Bi2Ru207.3 after calcination at 775 C for 10 hours ................ 85

4-6 XRD pattern for Bi2Ru2O7 and GDC powder mixture after heat treatment at 850 C
for 10 hours .............................................................................................................. 86

4-7 XRD pattern for Bi2Ru2O73 and GDC powder mixture after heat treatment at 850
T for 10 hours ................................................................................................... 86

4-8 Surface and cross-section micrographs of Bi2Ru207 electrode on GDC
electrolyte ................................................................................................................. 87

4-9 Impedance plots of Bi2Ru207 electrode on GDC at 500 C ................................ 88

4-10 Impedance plots of Bi2Ru207 electrode on GDC at 700 C ................................ 89

4-11 Arrhenius plot of the Bi2Ru207 electrode ASR (Qcm2) in air ............................. 90

4-12 ln(ASR) vs. ln(po2) of Bi2Ru2O7 electrode at different temperatures with m in
parenthesis T(m ) ................................................................................................. 90

4-13 Impedance plots of Ca doped Bi2Ru2O7 electrodes on GDC at 500 C .............. 91

4-14 Impedance plots of Ca doped Bi2Ru2O7 electrodes on GDC at 700 C .............. 92

4-15 Arrhenius plot of Ca doped Bi2Ru207 electrode ASR (Kcm2) in air .................. 93

4-16 Impedance plots of Ag doped Bi2Ru2O7 electrodes on GDC at 500 C ............. 94

4-17 Impedance plots of Ag doped Bi2Ru207 electrodes on GDC at 700 C .............. 95

4-18 Arrhenius plot of Ag doped Bi2Ru2O7 electrode ASR (fQcm 2) in air .................. 96

4-19 Arrhenius plot of undoped, 5 mol% Ca and Ag doped Bi2Ru207 electrode ASR
(Q cm 2) in air ........................................................................................................ 97

4-20 Arrhenius plot of undoped, 10 mol% Ca, Ag and Sr doped Bi2Ru207 electrode ASR
(f cm 2) in air ........................................................................................................ 97

4-21 Impedance plots of 5 mol% Ca doped Bi2Ru2O7 electrodes on GDC at 500 C ...... 98

4-22 Impedance plots of 5 mol% Ca doped Bi2Ru2O7 electrodes on GDC at 700 'C ...... 99

4-23 Impedance plots of 5 mol% Ag doped Bi2Ru207 electrodes on GDC at 500 'C...100

4-24 Impedance plots of 5 mol% Ag doped Bi2Ru207 electrodes on GDC at 700 C. .. 101

4-25 ln(ASR) vs. ln(Po2) of 5 mol% Ca doped Bi2Ru2O7 electrode at different
temperatures with m in parenthesis T(m) ............................................................... 102








4-26 ln(ASR) vs. ln(P02) of 5 mol% Ag doped Bi2Ru207 electrode at different
temperatures with m in parenthesis T(m) ............................................................... 102

4-27 Impedance plots of 5 mol% Ca doped Bi2Ru207 electrodes on GDC at 500 C .... 103

4-28 Impedance plots of 5 mol% Ca doped Bi2Ru2O7 electrodes on GDC at 700 C .... 104

4-29 DSC plots of powder mixture of Bi2Ru207.3 with monoclinic a-Bi203 and fcc
fluorite (B i20 3)0 8(Er2O 3)02 .................................................................................... 105

4-30 SEM surface micrographs of bismuth ruthenate-bismuth oxide composite
electrode ................................................................................................................. 106

4-31 Impedance plots of Bi2Ru2O7.3-(Bi2O3)0.s(Er2O3)0.2 composite electrodes on GDC at
50 0 C ..................................................................................................................... 10 7

4-32 Impedance plots of Bi2Ru2O.3-(Bi2O3)0,8(Er2O3)0.2 composite electrodes on GDC at
500 C ..................................................................................................................... 10 8

4-33 ASR (Qcm2) of Bi2Ru2O7.3-(Bi203)0.8(Er2O3)02 composite electrodes in comparison
to B i2R u20 7 electrode ............................................................................................. 109

4-34 Arrhenius plot of Bi2Ru2O73-(Bi2O3)0.8(Er2O3)02 composite electrode ASR (Q2cm2)
in comparison with Bi2Ru207 electrode ................................................................. 109

4-35 Proposed electrode reaction mechanism for bismuth ruthenate electrodes ............ 110

5-1 Representative microstructures of the GDC film under different pre-sintering
conditions for a 1600 'C final sintering ................................................................. 133

5-2 Representative microstructures of the GDC film under different pre-sintering
conditions for a 1650 C final sintering ................................................................. 134

5-3 Porosity in the GDC film as a function of the pre-sintering and final sintering
tem perature ............................................................................................................. 135

5-4 Cross-sectional microstructures of the GDC film under different pre-sintering
conditions for a 1600 C final sintering ................................................................. 136

5-5 Surface microstructure of the GDC film after 850 'C pre-sintering and 1400 C
fi nal sintering .......................................................................................................... 136

5-6 Surface and cross-sectional microstructure of reduced Ni-GDC anode ................. 137

5-7 Cross-sectional microstructure of LSCF cathode sintered at 750 C & 9000C.....137

5-8 Cross-sectional microstructure of tested unit cell .................................................. 137








5-9 Open circuit potential of the cell as function of temperature ................................. 138

5-10 Average oxygen ion transference number of GDC electrolyte as function of
tem perature ............................................................................................................. 138

5-11 I-V characteristics of the cell with air at the cathode ............................................. 139

5-12 Power density of the cell with air at the cathode .................................................... 139

5-13 I-V characteristics of the cell with oxygen at the cathode ..................................... 140

5-14 Power density of the cell with oxygen at the cathode ............................................ 140

5-15 ln(A SR/T) vs l/T of the cell ................................................................................... 141

5-16 I-V characteristics of the cell with bilayer sintering temperature of 1350 C ........ 142

5-17 Power density of the cell with bilayer sintering temperature of 1350 C ............... 142

5-18 I-V characteristics at 600 C of the cell with bilayer sintering temperature of 1350
C before and after therm al cycling ........................................................................ 143

5-19 I-V characteristics at 650 C of the cell with bilayer sintering temperature of 1350
C before and after therm al cycling ........................................................................ 143

5-20 Cross-sectional and surface microstructures of the tested unit cell with bilayer
sintering tem perature of 1350 C ............................................................................ 144

5-21 I-V characteristics of the cell with bilayer sintering temperature of 1450 C ........ 145

5-22 Power density of the cell with bilayer sintering temperature of 1450 C ............... 145

5-23 I-V characteristics at 650 C of the cell with bilayer sintering temperature of 1450
C before and after therm al cycling ........................................................................ 146

5-24 Cross-sectional microstructure of the tested unit cell with bilayer sintering
tem perature of 1450 C ........................................................................................... 146

5-25 I-V characteristics of the cell with bilayer sintering temperature of 1550 C ........ 147

5-26 Power density of the cell with bilayer sintering temperature of 1550 C ............... 147

5-27 Cross-sectional microstructure of the tested unit cell with bilayer sintering
tem perature of 1550 C ........................................................................................... 148

5-28 Cross-sectional microstructures of the tested unit cell with cathode sintering
temperature between 850-1350 C, and surface microstructure showing cathode
layer sintered at 1250 C and current collecting layer sintered at 1000 C ............ 149








5-29 I-V characteristics of the cell with cathode sintering temperature of 850 C ......... 150

5-30 Power density of the cell with cathode sintering temperature of 850 C ............... 150

5-31 I-V characteristics of the cell with cathode sintering temperature of 1000 C ....... 151

5-32 Power density of the cell with cathode sintering temperature of 1000 C ............. 151

5-33 I-V characteristics at 650 C of the cell with cathode sintering temperature of 850
C before and after therm al cycling ........................................................................ 152

5-34 I-V characteristics at 750 C of the cell with cathode sintering temperature of 850
C before and after therm al cycling ........................................................................ 152

5-35 I-V characteristics at 650 C of the cell with cathode sintering temperature of 1000
C before and after thermal cycling ........................................................................ 153

5-36 I-V characteristics at 750 C of the cell with cathode sintering temperature of 1000
C before and after therm al cycling ........................................................................ 153

5-37 I-V characteristics of the cell with cathode sintering temperature of 1150 C ....... 154

5-38 Power density of the cell with cathode sintering temperature of 1150 C ............. 154

5-39 I-V characteristics of the cell with cathode sintering temperature of 1250 C ....... 155

5-40 Power density of the cell with cathode sintering temperature of 1250 C ............. 155

5-41 I-V characteristics of the cell with cathode sintering temperature of 1350 C ....... 156

5-42 Power density of the cell with cathode sintering temperature of 1350 C ............. 156

5-43 I-V characteristics at 650 0C of the cell with cathode sintering temperature of 1150
C before and after therm al cycling ........................................................................ 157

5-44 I-V characteristics at 650 C of the cell with cathode sintering temperature of 1250
C before and after therm al cycling ........................................................................ 157

5-45 Maximum power density as a function of cathode sintering temperature ............. 158

5-46 I-V characteristics at 500 C of the cell with cathode sintering temperature of 1350
C, with contributions from ohmic and non-ohmic polarization ............................ 159

5-47 I-V characteristics at 700 C of the cell with cathode sintering temperature of 1350
C, with contributions from ohmic and non-ohmic polarization ............................ 159

5-48 Non-ohmic polarization of the cells with cathode sintering temperature of 1250 0C
and 1350 T ........................................................................................................... 160









5-49 Ohmic polarization of the cells with cathode sintering temperature of 1250 C and
1350 0C ................................................................................................................... 160

5-50 Arrhenius plot of ASR of the ohmic polarization of the cells with cathode sintering
temperature of 1250 *C and 1350 'C, in comparison to -30 grm bulk GDC
electrolyte ............................................................................................................... 16 1

5-51 X RD patterns for LSCF-ESB ................................................................................. 162

5-52 XRD patterns for LSCuF-ESB ............................................................................... 162

5-53 I-V characteristics of the cell with LSCuF-ESB cathode sintered at 750 'C, in
comparison with the cell with LSCF-GDC cathode sintered at 850 "C ................. 163

5-54 Power density of the cell with LSCuF-ESB cathode sintered at 750 'C, in
comparison with the cell with LSCF-GDC cathode sintered at 850 C ................. 163

5-55 OCP of the cell with Ag-ESB cathode as a function of time ................................. 164

5-56 Cross-sectional microstructures of the tested unit cell with Ag-ESB cathode in
back-scatted electron m ode .................................................................................... 164














Abstract of Dissertation Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy

ELECTROCHEMICAL STUDIES ON SELECTED OXIDES FOR INTERMEDIATE
TEMPERATURE SOLID OXIDE FUEL CELLS

By

Abhishek Jaiswal

December 2004

Chair: Eric D. Wachsman
Major Department: Materials Science and Engineering

Fuel cell technology holds the promise to change the way power is generated,

transmitted, and utilized in our increasing demanding lifestyles. State of the art solid

oxide fuel cells (SOFCs) utilize an all ceramic design and operate at 750-1000 'C. Lower

operating temperatures will significantly improve the economics of power generation

using SOFCs. The aim of this dissertation was to evaluate and develop component

materials for SOFCs, which could work efficiently at temperatures between 500-750 'C.

Erbia stabilized bismuth oxide (ESB) shows one of the highest oxygen ion

conductivity among all solid electrolytes. However, due to positional and occupational

ordering the conductivity decays below the transition temperature (-600 'C). The effect

of direct current bias on the ordering phenomenon in ESB was studied using symmetrical

cells with Ag-ESB electrodes. At 500 'C, the endotherm, related to reverse transition, is

enhanced by the applied bias at short time but with negligible change in conductivity

decay. It is proposed that the conductivity decay with anneal time is related more to the









positional ordering than occupational ordering. Ag-ESB electrodes showed good

performance, though were unstable under high currents at 625 C due to Ag migration

with oxygen flux.

Novel bismuth ruthenate based cathodes were evaluated using impedance

spectroscopy with symmetric cells on gadolinium doped ceria (GDC) electrolytes.

Undoped bismuth ruthenate electrode showed area specific resistance (ASR) of 55.64

Qcm2 at 500 C and 1.45 Qcm2 at 700 C in air. Doping with similar size Ca2+, Ag+, or

Sr2+ on Bi3+ site did not improve the electrode performance significantly, while bismuth

ruthenate-ESB composites showed 3-4 times lower electrode ASR. Bismuth ruthenate-

ESB (62.5:37.5 wt%) composite showed the best performance of 18.4 f cm2 at 500 0C

and 0.32 Kcm2 at 700 C in air. Addition of the ESB phase is believed to reduce the rate

limiting surface diffusion in oxygen reduction reaction.

Anode supported thick film GDC electrolyte unit cells were developed for IT-

SOFCs. A colloidal deposition technique was used to fabricate dense, thick GDC

electrolyte films on porous Ni-GDC anode supports. Pre-sintering temperature of the

anode and final sintering temperature of the anode/electrolyte bilayer were found to be

the primary parameters determining the density of the film. The sintering temperature of

LSCF-GDC (70:30 wt%) composite cathode was optimized to 1250-1350 C, which

resulted in a maximum power density of 0.338 W/cm2 at 0.771 A/cm2 700 C. Current

interrupt showed that apart from the electrolyte layer, the ohmic polarization across the

cell has significant contributions from the electrodes.














CHAPTER 1
FUEL CELLS: A GREEN SOLUTION

1.1 Introduction

A fuel cell is an electrochemical device which converts the chemical energy of a

fuel and an oxidant into electrical energy. The process essentially involves an invariant

electrode-electrolyte system. In essence, fuel cells are similar to batteries although with

much longer lifetime; as in a fuel cell, the fuel is replenished and the product is

discharged continuously. The conversion is direct without the need of intermediate

conversion into heat and mechanical energy, as in the case of conventional

turbine/generator systems. The energy conversion in fuel cells is not limited by the

Carnot cycle and efficiency up to 50 % is achievable.' As fuel cells do not involve any

combustion process, there is no formation of pollutants including NOx, SOx,

hydrocarbons and particulates. Also, fuel cells do not consist of any moving parts and

hence do not generate any noise pollution and require low maintenance. Due to these

advantages over other energy generators, fuel cells seek application in stationary as well

as in tractionary applications.2

The concept of producing electrical energy from a simple electrochemical cell was

first shown by Grove in 1839.3 It took however another 120 years before Bacon4

assembled a fuel cell stack exhibiting useful power densities, which was later modified

by Pratt & Whitney for the onboard power sources for NASA Apollo space missions. At

present different types of fuel cells are under development, which are in general

identified by the type of electrolyte used: phosphoric acid fuel cells (PAFCs), proton






2


exchange membrane fuel cells (PEMFCs), molten carbonate fuel cells (MCFCs), solid

oxide fuel cells (SOFCs), alkaline fuel cells (AFCs). The different types of fuel cells with

typical reactants and operating temperatures are shown in figure 1-1. Low temperature

PAFCs, PEMFCs and AFCs require external reforming of the fuel, whereas high

temperature SOFCs and MCFCs can internally reform and work with hydrocarbon fuels

directly.5 Also, high temperature operation removes the need for expensive noble metals

to catalyze the electrode reactions. These factors result in lower cost, lesser complexity

and higher energy conversion efficiency of the high temperature systems. Compared to

molten carbonate fuel cells (MCFCs), solid oxide fuel cells (SOFCs) are advantageous as

the solid oxide electrolyte is usually very stable and does not show any migration

problems under the operating conditions.

Among all types of fuel cells, proton exchange membrane fuel cells (PEMFCs) and

solid oxide fuel cells (SOFCs) are considered to be the most advanced and closest to

wide-scale commercialization. However, there remain significant technological and

engineering challenges to be solved. When running on hydrocarbon fuels in addition to

external reforming, PEMFCs also require CO removal from the fuel feed as they are

susceptible to CO poisoning which results in lower conversion efficiencies. On the other

hand, SOFCs run at high temperatures and ideally can internally reform any hydrocarbon

fuel with high efficiencies without the need of expensive catalysts. In addition, the high-

quality waste heat from the SOFCs can be utilized in cogeneration to further improve the

overall efficiency of the system. The state of the art SOFCs operate at temperatures

between 750-1000 'C using yttria stabilized zirconia (YSZ) as the electrolyte material,

La(Sr)MnO3-YSZ composite as the cathode material, Ni-YSZ ceramic metal (cermet)









composite as the anode material and LaCrO3 as the interconnect material. Currently, there

are two configurations seeking commercialization: tubular and planar. Among the two,

the seal-less tubular design of Seimens Westinghouse is the most advanced with electrical

efficiency up to 46%.6 With combined heat and power cycle, efficiency up to 85% and in

pressurized SOFC-gas turbine hybrid system, efficiency greater than 55% are achieved.

However, the high operating temperatures of SOFCs put considerable limitations

on the choice of materials for the various components and also on the lifetime of the cell.

Over the years research in the industry and academia has resulted in the development of

exotic materials and fabrication techniques such that the cell components are able to

perform exceptionally well under extreme conditions, to withstand thermal mismatch, to

be microstructurally stable and to counter interactions between adjoining components.

However this has led to high material and fabrication costs, making SOFCs

uncompetitive with existing power generation technologies. Decreasing the operating

temperatures will considerably improve the economics of power generation using SOFCs

by enabling the use of cheap ferritic stainless steel alloys as the interconnect material

instead of expensive ceramics, cheaper balance of plant and insulation along with

increased lifetime.7 Lower operating temperatures will also result in faster startup which

is critical in certain applications. All these factors have led to a global drive towards

reducing the operating temperature of SOFCs from 750-1000 'C to intermediate

temperatures of 500-750 'C. However, efficient operation at intermediate temperatures

will require new electrolyte materials with higher conductivity and new electrode

materials with better catalytic activity at lower temperatures.









1.2 Operating Principle of SOFCs

The working of a fuel cell could be understood using a concentration cell model.

Let us assume that the cell has different oxygen partial pressure at the cathode and the

anode. The electrolyte is a gas impermeable solid with Pt electrodes attached to it. The

model can be represented as follows:

(anode) 02, Pt / Solid Electrolyte / Pt, 02 (cathode) 1-1

The cathodic and anodic reactions are as follows:

2e-(pt, c) + 1/2 02(g, c) = 02-(se, c) cathode 1-2

02-(se, a) = 2e-(pt, a) + 1/2 02(g, a) anode 1-3

where the subscripts represent the site of concerned species: a, c, g, se represent anode,

cathode, gas and solid electrolyte, respectively. The difference in the chemical potential

of 02 (P 2 = R TIn po2 ) at the anode and the cathode results in the generation of

electrical potential difference (A~p) across the cell.
1 it. (o,~d~o
A0 = I o2- T =I -f t o2- (Po2 ,T)RTd(lnpo2) 1-4



to2- (Po2 T) = 2- (Po2 ,T)
'r Z (po2,T)
1 1-5

Apart from the oxygen partial pressure difference, the electrolyte transport

properties in terms of the oxygen ion transference number (to2_) determines the electrical

potential difference generated. The oxygen ion transference number is given by the ratio

of oxygen ion conductivity to the total conductivity and at a constant temperature is a

function of oxygen partial pressure. The oxygen partial pressure varies along the









electrolyte thickness and the gradient (b Po2 / x) in turn depends on the transport

properties oi of the electrolyte.

In an actual fuel cell, fuel (H2, CO, hydrocarbons) at the anode develops a low

oxygen partial pressure and air/oxygen at the cathode develops a high oxygen partial

pressure. Electrical power is generated when an external load is attached to the cell, and

useful power is achieved by stacking cells through interconnects in series and parallel.

The efficiency of a cell is lower than theoretical because of the irreversible ohmic and

non-ohmic polarization losses across the cell, as shown in figure 1-2. The irreversible

losses appear as heat, and hence the thermal management of the fuel cell stacks is an

important design consideration.

V(I) E tlohmic (lact + qconanode (lact + ?lconcathode 1-6

The bulk of the ohmic polarization comes from the solid electrolyte. High oxygen

ion conductivity, to2 equal to I and small thickness are the desired characteristics of the

solid electrolyte. Non-ohmic polarizations from the electrodes consist of activation and

concentration polarizations. Activation polarization is related to the kinetics of chemical

and charge transfer reactions at the gas/solid interface. High catalytic activity for oxygen

reduction at the cathode and for fuel oxidation at the anode is required. Fast oxygen ion

transfer across the electrolyte/gas interface and electrolyte/electrode interface requires

good chemical and mechanical compatibility of the electrolyte/electrode interface.

Formation of blocking interface phases must be avoided. Concentration polarization is

related to the diffusion of gaseous species through porous electrodes to reach the reaction

sites. Microstructural engineering is required to avoid the diffusion limited regions in the

electrodes when operating at high currents.









1.3 Designing SOFCs

Ohmic and activation polarizations are thermally activated processes, resulting in

improved efficiency of the SOFCs with increasing temperatures. However, high

operating temperatures could result in chemical and mechanical compatibility issues

between adjoining components. Chemical and mechanical compatibility between

adjoining components are of paramount importance as the SOFC is expected to perform

for 40-100 thousand hours at elevated temperatures and for numerous thermal cycles

from the room temperature to the operating temperature. Mechanical compatibility

between neighboring components against thermal cycling requires close matching of the

thermal expansion coefficient (TEC). At times, SOFC designers have to compromise the

electrochemical performance in order to ensure a good thermal expansion match between

components. In general, it is achieved by introducing external dopants in the host

material to vary the TEC.

Processing of electrodes should ensure a good adhesion at the electrode/electrolyte

interface and a stable electrode microstructure over the period of operation. In general,

the processing temperatures are higher than the operating temperatures to ensure that the

electrode microstructure does not coarsen over the period of operation with resultant loss

in active sites for the electrode reaction and in consequent decay in cell performance.

However, sintering at high temperatures could lead to problems in certain electrode-

electrolyte combinations, as in the case of La(Sr)MnO3-YSZ where tertiary La2Zr207

phase appears at the interface. La2Zr207 is an oxygen ion blocking phase and its presence

at the interface degrades the cell performance considerably.

Chemical stability of all the components in their respective working environments

is also a requirement. The electrolyte faces both the oxidizing and the reducing









atmospheres and as described by Goodenough,8 its thermodynamic stability requires that

the bottom of the electrolyte conduction band is above the highest occupied molecular

orbital (HOMO) of the reductant and the top of the electrolyte valence band is below the

lowest unoccupied molecular orbital (LUMO) of the oxidant as shown in figure 1-3. The

Fermi energies of metallic electrodes should also lie in the energy gap Eg of the

electrolyte, with that of the anode rising to the HOMO of the reductant and that of the

cathode falling to the LUMO of the oxidant under operating conditions.

Although Ni as the metal component in the cermet anode satisfies the major

requirements for catalytic oxidation of H2 and CO fuels, its use for direct oxidation of

hydrocarbons encourages carbon deposition on the surface which with time reduces the

catalytic activity of anode. Cu has been shown to be more effective with hydrocarbon

fuels and it is likely that in future, Cu will replace Ni in SOFCs having internal reforming

capabilities.9 Due to the low melting point of Cu, there are still technical hurdles with Cu

based anodes in terms of fabrication and operating temperatures. Ni anodes can also be

poisoned in sulphur containing fuels due to formation of thermodynamically stable NiS.

Ceramic anodes based on doped SrTiO3 are under development and show significant

promise in terms of carbon and sulphur tolerance.'0

1.4 SOFC Structures

Based on the component that supports a unit cell, SOFCs can be divided into two

basic types: electrode supported and electrolyte supported. Chan et al.' have modeled the

sensitivity of the cell voltage to the thickness of each component. The model showed that

with the same thickness of each component, the cell voltage is most sensitive to the

thickness of the electrolyte followed by the cathode and then by the anode. Therefore,

anode or cathode supported cells should be preferred over electrolyte supported cells.









Compared to the cathode supported cells, the model showed that the anode supported

cells are superior in terms of operating current density range and available power density.

Apart from the fuel cell performance other factors like target application, fabrication and

operating costs, fuel choice and utilization, operating temperature, mechanical strength

and expected lifetime are also important considerations. Electrode supported design with

a thick film electrolyte perform better, but it also have issues in terms of strength, cost

and performance. Porous electrodes supports are mechanically weaker in comparison and

thermal cycling puts considerable amount of thermal and mechanical stresses on them.

Anode supported cells with Ni based cermet anode can face catastrophic failure during

loss of fuel supply or under high fuel utilizations, as oxidation of Ni to NiO results in

large volume expansion which generates cracks in the anode and in the thick electrolyte

film. Further, depositing dense (pin-hole free) thick electrolyte layers on porous supports

reliably, inexpensively, with good adhesion and without undue sintering of the porous

support over large areas still remains a technical challenge. Keeping these factors in

mind, different manufacturers are pursuing different SOFC design and manufacturing

routes.

1.4.1 Tubular SOFCs

Among all SOFC designs, Siemens Westinghouse's seal-less cathode supported

tubular design, shown in figure 1-4, is considered the most advanced. They are targeted

for the stationary power generation market and have been tested successfully for

hundreds of thousands of hours. The La(Sr)MnO3 cathode support tube (2.2 cm diameter)

is extruded and sintered, over which -40 jim thick, dense YSZ electrolyte is deposited

using electrochemical vapor deposition (EVD). To connect unit cells together, -0.9cm

wide and -85 gim thick strip of LaCrO3 based interconnect material is deposited using









plasma spraying. Finally, 100-150 urm thick Ni-YSZ cermet anode is deposited either by

Ni slurry coating followed by electrochemical vapor deposition of YSZ or by sintering of

a Ni-YSZ slurry coating. Single cells are then joined together using Ni felt in series and

parallel to achieve desired power ratings. Air flow is inside the cathode tube while fuel

flow is outside the cell.6

The advantage of the tubular geometry is in its simplicity. It avoids the use of seals

which at high temperatures are always an issue in terms of reactivity with other

components and failure during thermal cycling. The tubular structure is more rugged than

the planar structures which are discussed later. However, the cost of the tubular design is

still prohibitive for mass commercialization against competitive technologies. In the

tubular design, over 90% of the weight of the single cell comes from the cathode tubular

support and therefore cheaper raw materials for the cathode can bring in significant cost

reduction. Fabrication routes for the dense electrolyte and interconnect are also fairly

expensive and cheaper fabrication techniques like colloidal/electrophoretic deposition are

being evaluated for further cost reduction.6

Tubular single cells show a power density of -0.25 W/cm2 at 1000 OC (fuel: 89 %

H2 / 11 % H20, 85 % fuel utilization; oxidant: air), which is significantly lower than that

of the planar cells. Lower power density is primarily due to the long electronic current

path along the circumference of the cathode tube. Increasing the thickness of the tube

reduces the electronic resistance but also adds to the cathodic polarization as the oxygen

path from the gas phase is radially towards the electrode/electrolyte interface. In the

planar design, the current paths are much shorter being perpendicular to the thickness of

the layers and power density greater than 1 W/cm2 at 800 'C are in general achieved. A









new design combining the benefits of both the tubular and planar design is being

investigated by Siemens Westinghouse consisting of a flattened tube incorporating ribs in

the cathode structure as shown in figure 1-5. The new design lowers the internal

resistance of the cell with a shortened electronic current path along the ribs in the

cathode. Moreover, reduced resistance allows the use of thinner cathodes with lower

cathode polarization. The new flattened tubular geometry shows higher power density

compared to the tubular geometry as shown in figure 1-6.

1.4.2 Planar SOFCs

Generic design of a planar SOFC is shown in figure 1-7. The assembly consists of

thin flat components which are fabricated using low-cost ceramic techniques like tape

casting, colloidal deposition and screen printing. Different organizations are focusing on

different variations of the planar design and use different manufacturing processes to

fabricate planar fuel cell stacks.

At present electrolyte supported, cathode supported and anode supported designs

are under development. Electrolyte supported cells with 50-100 tm thick YSZ electrolyte

have high electrolyte resistance and therefore are suitable for operation at temperatures

-1000 C. The operating temperatures can be reduced by at least 200 'C by reducing the

thickness of the YSZ electrolyte to 5-20 p.m in electrode supported planar designs. Ni-

YSZ anode supported designs are preferred over cathode supported designs in terms of

better electrochemical performance, better thermal and electrical conductivity, and

minimal interaction with the YSZ electrolyte during the sintering process. Although as

mentioned before, there remains concern over Ni oxidation in the case of loss of fuel

supply. Also in terms of cost, it has been noted that the anode supported design has more

electrolyte content than the electrolyte supported design itself.









For the fabrication of the anode supported planar SOFC, the electrolyte layer is

deposited on a tape cast anode by using either screen printing, colloidal deposition or

electrophoretic deposition followed by the co-firing of the bilayer. Some manufactures

also use tape calendaring or lamination to form the anode/electrolyte bilayer. Finally, the

cathode is screen printed and sintered to complete the unit cell. Anode supported planar

design shows significantly higher performance compared to the tubular design and at 800

'C, planar single cells consisting ofO-10 jm YSZ layer have shown 1.8 W/cm2 (fuel: 97

% H2 / 3 % H20, low fuel utilization; oxidant: air), as shown in figure 1-8. Moreover,

reduced operating temperatures allow the possibility of using metal interconnects which

along with cheaper fabrication routes for the cell can bring down the cost significantly.

Successful development of low-cost and oxidation resistant metallic alloy interconnects

and long lasting seals for the separation of the fuel and the oxidant atmospheres will

enable the commercialization of planar SOFCs.

1.4.3 New Designs: Metal Supported and Microtubular SOFCs

Although the anode supported planar design show significantly higher performance

than the tubular design, they leave a lot to be desired in terms of mechanical strength,

start-up time and sealing. This has led to the development of metal supported and micro-

tubular SOFC designs.

In the metal supported SOFC design, a high temperature metal alloy supports the

electrode/electrolyte assembly as shown in figure 1-9 and 1-10. The assembly is

fabricated using the same low-cost techniques used to fabricate the anode supported

planar design. The metal support results in lower cost, higher mechanical strength and

higher thermal shock resistance of the structure which allows for rapid start-up.12,13 The

metal support also increases the electronic and thermal conductivity resulting in lower









losses in current collection and better thermal management of the stack. Moreover, metal

support allows welding, brazing and other joining techniques for sealing unit cells. 12,13

However, there are a number of technical hurdles with the metal support design. These

include the fabrication temperatures and atmosphere for the assembly which is limited by

the melting/oxidation temperatures of the metal alloy support. For example, if the metal

alloy support is made of stainless steel then the assembly can not be readily fired above

1000 0C, which might be inadequate for the sintering of some components resulting in

lower performance of the cell. Stability of the electrode/alloy interface is also a concern.

Alloys with Cr203 oxide scales in moist air atmosphere form volatile CrO2(OH)2 species,

which can deposit at the cathode surface to act as blocking species. On the anode side,

high sintering temperatures result in interdiffusion between Ni and Fe-Cr alloys with

resultant coarsening of the anode microstructure due to the formation of Cr203 oxide

scales. Further, elastic modulus mismatch between the alloy support and the ceramic

components can result in high stresses.

The microtubular design as the name suggests is similar to the Siemens

Westinghouse tubular design although with a much smaller diameter with either

electrolyte or anode supported geometry. The high length to diameter ratio provides for

its high strength and extraordinary thermal shock resistance. These tubes can withstand

temperatures gradient of 1000 C at one end to room temperature at the other end.

Therefore, the micro-tubular design can combine the benefits of a seal-less design with

high mechanical strength and rapid start-up. Low-cost ceramic techniques are being used

to fabricate in order to keep the costs down. However, the design is plagued with the

same performance issue as that of Siemens Westinghouse's tubular design: long current









collection paths. In fact the present microtubular design does not incorporate

interconnects and therefore the current collection path is actually longer.

Both metal supported and microtubular designs hold a lot of promise and the

solution of the above described technical hurdles will determine whether they remain

competitive with other designs.

1.5 Conclusion

SOFCs are efficient and environmental friendly energy generators. In this chapter,

the operating principle, the components and their requirements, and the various designs

of SOFCs were introduced. Tremendous progress has been made over the decades in the

development of materials, fabrication techniques and fuel cell design which enable a

stable performance over extended periods of time. However, the cost of SOFC system is

still prohibitive for mass commercialization against competitive technologies. New cost

effective approaches are being actively evaluated both in terms of fabrication techniques

and SOFC design. Decreasing the operating temperatures of SOFCs to intermediate

temperature range of 500-750 C will open a wide range of options for the different

components, which can help in further reducing the cost. As the performance of SOFC

deceases with temperature, the challenge is to develop new materials which can perform

well at intermediate temperatures.






















Internal
Reforming
H2, CO



External
Reforming
H2, CO2

External
Reforming -
H2, CO2
(CO removal)

H2


ANODE


ELECTROLYTE


CATHODE


I +


02(air)



02 (air)
4 C02



S02 (air)



4 02 (air)



4 02(air)
(C02 removal)


Figure 1-1. Fuel cell types with typical reactants (from ref. 5).


J (Amp/cm2)


Figure 1-2. Schematic diagram of the current-voltage characteristics as well as the
different losses in a solid oxide fuel cell.


f- SOFC (500-1000C)
H20
CO2 4 02-


MCFC (650C)
H20 Co32
CO2


PAFC (200C)
W 1 H20


PEMFC (80C)
H~ IH20


AFC (70C)
H20 4- Off













HOMO










Reductant


tEnergy
Conduction Band


I
Eq




ValnceBand/

SElectrolyte


X

Figure 1-3. Placement of reactant energies relative of the edges of conduction and
valence band of the electrolyte in a thermodynamically stable electrochemical
cell (from ref. 8).


Interconnection



Electrolyte

Air
Electrode





Flow =-Fuel Electrode


Figure 1-4. Siemens Westinghouse tubular SOFC design (from ref. 6).


Vo



LUMO



Oxidant















Air Ebctvde


Elctrolyte


Fuel ElectrcW*."


Itrconnction


44 f 4444


Figure 1-5. Siemens Westinghouse flattened tubular SOFC design (from ref. 6).


- -- I I **I~I~S~ I


Operatng Temperature: 10Iro
0,45 Oxant Air, 6.0 Sokhs
Fuel; 89% M2 11% I V. 85% Fuel Utzafio
0.40k


0.3SF


0 100 200 300 400 500


600 700 800 900 1000


Current Dwsty (mAcnr')

Figure 1-6. Theoretical and actual performance of Siemens Westinghouse tubular and
flattened tubular (HPD) SOFC designs (from ref. 6).


v I


w VYvv


BOB
e00
VO IPedmn e
v vHPD-SOFC
v22 cm OO Cyivkical

vHPD-OF
a 22cmODCyx rilic&
. I I I I I I =


rin


I I


| tJ






17



Anode

Interconnection
(Bipolar Plate)
Cathode



Electrolyte

Anode









Figure 1-7. Planar SOFC design (from ref. 14).



1,1" 2000
1.0 1800
0.9- 1600
0.8-
~1400
~0.7 I
1200
0.6-
0.5-1000
0-50.4
0.4- 80
0.3' CCU -60 .

0.2 40
0.1 200
0.0 0
0 1000 2000 3000 4000 50 60 7000
Current Density (n.icm2)


Figure 1-8. Performance of a planar SOFC design at 800 C (from ref. 15).



















Vt


YSZ
NiIYSZ


Fe/Cr alloy


Figure 1-9. Microstructure of an alloy supported SOFC based on YSZ electrolyte (from
ref. 12).


Contact Layer



Cathode


Electrolyte

Anode

Steel Substrate


Figure 1-10. Microstructure of an alloy supported SOFC based on doped cerium oxide
electrolyte (from ref. 13).













CHAPTER 2
IONIC CONDUCTION IN SOLID ELECTROLYTES

2.1 Introduction

Oxygen ion conducting solid electrolytes are the backbone of solid oxide fuel cells

(SOFCs), allowing selective transport of oxygen ions for electrochemical oxidation of

fuel to generate electrical power. The ionic conduction is possible due to the presence of

oxygen deficiency in the solid electrolyte. In general, the oxygen deficiency is introduced

by doping the host electrolyte with lower valent cations. The crystal structure of the host

material should be stable enough so that the introduction of the dopant ions does not

destroy the structure. Further, to make the activation energy for the migration of the

oxygen ions low, the electrolyte should have a relatively open crystal structure. The

saddle point critical radius in the lattice should be large enough for the passage of the

oxygen ions without creating lattice disturbances.

The electrolyte materials for SOFCs are mostly oxides having fluorite or perovskite

structure because of the inherent looseness in the structures and the ability to accept wide

range of dopants. Oxygen ion conductivity of common solid oxide electrolytes is shown

in figure 2-1. Better conductivity values are achieved with dopants having similar ionic

radii to the cation being substituted. The electrolyte in the fuel cell is exposed to both the

oxidant and the fuel atmospheres, and hence, it should be thermodynamically stable

under oxidizing and reducing conditions. To avoid the generation of electronic

conductivity, the electrolyte should have a large band gap and the dopants introduced into

the lattice should not exhibit multiple oxidation states.









As the oxygen ion conduction is a thermally activated process, the performance of

the electrolyte improves with temperature. However, there are limits in terms of operating

temperatures due to material compatibility issues with other components and operational

viability over extended periods of time. Reduced operating temperatures will open up

wide range of materials and significantly improve the economics of power generation

using SOFCs. However, operation of SOFCs at intermediate temperatures requires better

performance electrolyte and electrode materials.

2.2 Fluorite Type Oxides

2.2.1 Structural Aspects

The fluorite type unit cell is shown in figure 2-2. CeO2 and Th02 exhibit the cubic

fluorite structure from room temperature up to their melting points. The packing in the

fluorite structure is far from closed packed and there are cubic 8-coordinated interstices at

the center of the unit cell. This loose structure provides for the possibility of achieving

unusually wide range of the solid solution with alkaline earth and rare earth oxides such

as CaO and Y203.14 Generally, when the cation size of the host and the guest are almost

equal the solid solution is easily formed. In the case of ZrO2 and Bi203, the high

temperature cubic fluorite structure could be stabilized at lower temperatures by forming

a solid solution and hence, the terms stabilized zirconia and stabilized bismuth oxide.

2.2.2 Electrical Aspects

To maintain the electrical neutrality of the solid solution, ionic and/or electronic

defects are generated in the lattice. In the case of fluorite systems, anti-frenkel defect

pairs (oxygen vacancy Vo and oxygen interstitial O,") are the dominant ionic defects. On

doping with lower valent cations, oxygen vacancies with effective charge of +2 are

generated to counter the effective negative charge of the dopant species. For example in









the case of CaO doped Ce02, the defect and the electroneutrality equation can be written

as follows:

CaO ")Cace" + Oox + Vo 2-1

2[0,'7 + 2[Cace'] + le'] = 2[Vo I + [h 7 2-2

At a particular temperature, the electrolytic domain for the solid solution is the

regime of oxygen partial pressure where the ionic defects dominate i.e. in the present

example when

[Cace"] = [Vo] 2-3

As mentioned earlier for ZrO2, doping apart from the generation of oxygen vacancies also

helps in the stabilization of the high temperature fluorite phase at lower temperatures.

Electrical conductivity (a,, S/cm or ohmlcm') of a species is given by the product

of concentration of species (c,, no/cm3), charge of the species (Zie, coulomb), and

electrical mobility (,u, cm 2/sec-Volt). The total electrical conductivity of the system is the

sum of the contributions from all ionic and electronic species, while the partial

transference number (t,) of a species is given by the ratio of partial conductivity to the

total conductivity.

ur cZ e", 2-4

UT =U, = Xr c,'Zje'u 2-5

ti = ,/aT 2-6

Due to the loose packing in the fluorite structure, the oxygen ion mobility is

significantly higher than that of other crystal structures, but it is still orders of magnitude

lower than that of electronic defects (electron and electron hole). In order to make a

predominant ionic conductor, large amounts of doping is required to generate high









concentration of ionic defects. In fact, doping is a misnomer in that respect as levels

required are in the range of 5-30 mol%. In addition to keep the concentration of

electronic defects in the system low, both the host and the dopant cations should have a

fixed valence state in the temperature and oxygen partial pressure range of interest.

2.2.3 Oxygen Ion Conductivity as a Function of Dopant Concentration and
Temperature

As doping increases the oxygen vacancy concentration, oxygen ion conductivity

can be enhanced by increasing the dopant concentration but only up to a certain level. As

shown in figure 2-3 for solid solution based on ceria, beyond a certain dopant level the

ionic conductivity decreases. 15 This behavior has been explained in terms of association

of vacancies or formation of defect complexes in the dilute solution range and formation

of super-lattices or ordering of vacancies in concentrated solution range.

Introduction of the dopant cation changes the ionic potential of the lattice in the

vicinity of the dopant. This leads to trapping of the oxygen ion vacancy to the dopant

cation because of columbic attraction (between effectively -ve charged dopant and

effectively +ve charged oxygen ion vacancy) and/or relaxation of elastic strains

(generated by the dopant) in the lattice. Recent studies have shown that the elastic

relaxation has more to do with the trapping of oxygen ion vacancy with the dopant. As

shown in figure 2-3, dopants with the same charge do not show similar magnitudes for

oxygen ion conductivity and that the choice of the dopant for the optimum oxygen ion

conductivity in a particular system is critically related to the size of the dopant cation.

Dopants with similar size as that of the host show the best performance.

In the concentrated solution regime, long range ordering has been associated with

the conductivity decay. For example in the case of Ca stabilized ZrO2 (CSZ), ordered








small domains (-10 A) have been observed by electron diffraction studies after long

annealing times at elevated temperatures. 16 This observation has been explained in terms

of disproportationation of metastable CSZ with time into ZrO2 and CaZr205

(Zro.67Cao.3501.67). CaZr205 has a structure similar to that of C-type rare earth oxides

(with site specific oxygen vacancy) and thus with significantly lower conductivity.

Ordered sublattice essentially means that specific atoms have specific stable sites

with low energy, and therefore, for a conducting species ordering results in loss of mobile

species and in low conductivity. Y203 has C-type crystal structure which is similar to the

fluorite structure though with a site specific oxygen vacancy. This regularity results in

oxygen ion conductivity much smaller than that of CSZ, although the concentration of

oxygen vacancies is much larger. Similarly, Bi203 exhibits two metastable phases below

650 C: tetragonal -phase and body centered cubic y-phase. Both these phases have an

ordered oxygen ion sublattice with oxygen ion conductivity up to three orders of

magnitude lower than the high temperature 6-phase, which has a cubic fluorite structure

with disordered oxygen ion sublattice.17-19 The concentration of oxygen vacancy is

similar in all the three phases, though the concentration of mobile oxygen vacancy is

significantly smaller in 3-Bi203 and y-Bi203. It is also important to note that the crystal

structure also plays a key role in determining the oxygen ion conductivity in terms of

jump directions, jump paths and activation barrier for the oxygen ion motion as described

below.8

a = c'Ze-p = c- (Ze)2D/(k.T) 2-7

a = nvN. (Ze)2.D.exp(-AHrk.T)/(k.T) 2-8

D, fz.(1 n,)a2V.oexp(AS,1k)/6 2-9









Nemst-Einstein can use be used to relate the electrical mobility (a, cm2 /sec-Volt)

with diffusivity (D, cm 2/sec), while concentration (c) can be related to site fraction of

oxygen vacancy (n,) and concentration of oxygen ion sites per unit volume (N, no/cm3).

Diffusivity is related to jump distance (a, cm), jump frequency (v, Hz), structure factor (f

unit less on the order of 1), number of nearest neighbors (z), and number of available sites

for a vacancy jump (I n,). The jump frequency is related to the motional enthalpy (AH,)

for the oxygen ion jump to a neighboring vacant site through the common face bottle

neck. The motional enthalpy is related to the size of the mobile ion and distance between

the center of the common face and the peripheral ion. The Arrhenius relationship of

conductivity can be represented as

ln(a- T) = ZnA EI(k. T) 2-10

A = n,.(I-n,.)N (Ze)2fza2" voexp(ASm/k)/(6"k) 2-11

In the case of an ordered system, vacancies can be differentiated on the basis of the

site i.e. normal (occupied) site and interstitial (vacant) site at a higher energy level (AHg).

ZAHg is a function of temperature and vanishes at the transition temperature (Tt), above

which the system becomes disordered. The concentration of the vacancies on the normal

sites is thermally activated and increases with temperature. The fraction of vacant sites on

normal sites (n,) is related to -exp(-AHg'2kT) and the fraction of occupied sites on

ordered sites is related to m *n, where m is the ratio of number of normal sites with

respect to interstitial sites. This results in higher activation energy for oxygen ion

conductivity at temperature below T,.8

E, = AH,, + AHg/2 T < Tt 2-12

Ea = AH T> T, 2-13









In the case of doped system below the critical temperature T*, oxygen ion

vacancies get progressively trapped into ordered clusters due to coulombic attraction and

elastic relaxation, as described earlier. Goodenough8 has shown using solution

thermodynamics that ninv*, fraction of oxygen vacancies dissolved randomly on the

oxygen ion sublattice is given by

n/n,* = exp(-AHr (1 T/T*)/k.T) 2-14

where n, is the fraction of randomized vacant sites, n,* is the fraction of total vacant

(including both randomized and ordered) sites, and AH is trapping enthalpy.8 This results

in well know kink in the Arrhenius plot with a higher activation energy and a higher pre-

exponential term (A) at temperatures below T*.

Ea = AHm + AHt, A, = Aexp(AH/k'T*) T< Tt 2-15

E. = AHm A2 =A T> T 2-16

Different oxide systems show different magnitudes of T*, for example for yttria

stabilized zirconia (YSZ) -800 C, for gadolinium doped ceria (GDC) -400 'C and for

erbia stabilized bismuth oxide (ESB) -600 'C. Further it has been observed that on

increasing the doping concentration, high temperature Ea increases and eventually

becomes equal to the magnitude of low temperature E,, resulting in a single activation

energy. This is consistent with the idea that the concentration of trapped clusters

increases with doping and tends to dominate the activation energy for conduction. Or in

other words, the transition temperature T* increases with doping in order to provide the

necessary thermal energy to release the oxygen ion vacancies from the traps.

Hohnke20 has modeled the dependence of the pre-exponential factor and the

activation energy on the dopant level by using a pseudo-chemical equilibria and by









introducing terms for short range and long range interactions. For low dopant levels,

oxygen ion conductivity can be given by the following equations

a = (A/T)f(c).exp(-Eik T) 2-17

f(c) = c05; Ea = AH, + AHt T< T, and Yce' 2-18

f(c) = cO; Ea = H,, + m Ht T < T, and Cace" 2-19

f(c) = c; Ea = AHm T > Tt 2-20

indicating that at high temperatures and low dopant levels, conductivity increases linearly

with the oxygen vacancy concentration. At lower temperatures, there is an additional

terms in the activation energy due to the short range interaction. Further depending on the

dopant type, conductivity varies square root of concentration for a single charged dopant

and is independent of concentration for a double charged dopant. With higher dopant

levels, the activation energy at low temperatures is found to increase linearly with dopant

level. This phenomenon has been attributed to long range interaction and is included in

the following generalized conductivity expression with equilibrium constants for short

range (K,) and long range (K) interactions.20

Ea Hm + Ht+AHIc 2-21

a = a'1 + Kj (1 + K)]' 2-22

2.2.4 Grain Boundary Contribution to Total Conductivity

The kink in the Arrhenius plot of the conductivity can also be introduced by the

grain boundary contribution to the total conductivity. In doped ceria, the grain boundary

has a higher activation energy and a lower conductivity than the bulk either due to the

presence of impurities at the grain boundaries or due to higher dopant level than the bulk.

Due to the higher activation energy, the effect of the grain boundary on the total

conductivity is more prominent at low temperatures and it vanishes at higher









temperatures. The presence of the grain boundary contribution to the total conductivity

could result in suboptimum choice of dopant level as observed by Steele.21 As shown in

figure 2-4, lattice conductivity for 10 mol% gadolinium doped ceria (GDC) shows a

transition temperature of -400 'C with a high temperature activation energy of -0.64 eV

and a low temperature activation energy of -0.77 eV. By increasing the doping level to

20 mol%, the lattice conductivity decreases with a constant activation energy of -0.78

eV. The presence of impurities like SiO2 at the grain boundaries in 10 mol% GDC result

in a low grain boundary conductivity with activation energy of -0.8-1.0 eV. The effect of

grain boundaries on the total conductivity can be seen up to 900-1000 'C. In comparison,

10 mol% GDC without impurities at grain boundaries shows little influence of the grain

boundary contribution to the total conductivity above 500 'C. For impure 20 mol% GDC

the grain boundary conductivity is higher and contributes up to 700 'C, which results in a

higher total conductivity than impure 10 mol% GDC. Thus, the presence of impure grain

boundaries can lead to suboptimum choice of the higher dopant level in terms of the total

conductivity, when clearly the lower dopant level shows better bulk conductivity.

2.2.5 Oxygen Ion Conductivity as a Function of Time

ESB in the cubic fluorite structure is one of the highest known oxygen ion

conductors and like YSZ and GDC, also shows the kink in the Arrhenius plot of

conductivity for low dopant levels as shown in figure 2-5. The transition temperature is

-600 'C, with a high temperature activation energy of -0.68 eV and a low temperature
22
activation energy of-i .28 eV. Further, it has been observed by Wachsman et al. that on

isothermal annealing below the transition temperature the conductivity decays with time,

as shown in figure 2-6. The conducitivity decay has been attributed to continuous

trapping of oxygen ion vacancies in the growing ordered clusters resulting in loss in









mobile oxygen vacancies. Conductivity decay also occurs in CSZ on annealing at

elevated temperature and the presence of ordered small domains (-10 A') have been

observed by electron diffraction studies on annealed CSZ.16 Similarly, formation of

superstructure and oxygen ion displacement have been observed in ESB using electron

and neutron diffraction studies, respectively.2227 The conductivity decay in stabilized

bismuth oxide is comparatively much faster than in YSZ and can be represented by the

following empirical equation

a(t) = e(co) + [a(O) a()].exp[(-t/r)f] 2-23

where u(t) is the conductivity at time t, r is the pertinent time constant, and # is a

dimensionless parameter. The larger is the time constant r, the more stable is the system.

The time constant increases with the dopant cation radius and with the dopant

concentration as shown in figure 2-7 and 2-8, respectively.28 As in most practical

applications, solid electrolytes will be used in an electrochemical potential gradient with

resultant oxygen ion flux. It is plausible that the oxygen ion flux can affect the trapping

of oxygen ion vacancies in the condensed ordered clusters and hence the conductivity

decay below the transition temperature. Part of chapter 3 deals with studying this

phenomenon in erbia stabilized bismuth oxide solid electrolyte.

2.3 Conclusion

This chapter introduces the basics of oxygen ion conduction in solid oxide

electrolytes. Models from the literature for oxygen ion conductivity as a function of

dopant level, under different temperature regimes and as a function of time are described.

The primary reason for the deviant behavior from ideal oxygen ion conductivity is due to

the elastic relaxation on introducing the dopant in the host leading to formation of defect

complexes in the case of low dopant levels and formation of ordered clusters in the case






29


of high dopant levels. This results in well known kink in the Arrhenius plot of

conductivity with a high activation energy at low temperatures and a low activation

energy at high temperatures. Further, the ordering phenomenon could also lead to decay

in conductivity on isothermal annealing below the transition temperature, which is

believed to occur due to trapping of mobile vacancies in the condensed ordered clusters.




















Temperature (C)
900 800 700 600 500 400


0.6 1.0 1.2 1.4
1000JT (K)


1.6 1.8 2.0


Figure 2-1. Oxygen ion conductivity of common solid electrolytes (from ref. 5).


300


0



b
-2
0



.3

-4































Figure 2-2. The fluorite structure (Blue spheres Anions; Red spheres Cations) (from
ref. 17).


0 10 20 30
x (expressed as %) in Ce.M,02-,2

Figure 2-3. Isothermal conductivity of some ceria solid solution at temperature close to
200 'C (from ref. 17).


Gd





































CLEAN
CGOO 1)


1064eV







I I


Ib)


low
MSI ~


Uwnuu
0G(16)

Unatice
0-04cV


VMPURE 0 73eV
CGOOM 1700C




0 9-1 t


Figure 2-4. Schematic Arrhenius plots of lattice and grain boundary conductivity for (a)
clean Ceo 9Gdo 10 1.95 (b) impure Ceo0gGdo. 10 195 (c) impure Ceo.sGdo.2019 (from ref. 23).

















89 T( K) 60






lxx












U
xx






10-








IIII H III I i S I I I I I I I1 I
10 12 I4 16 18
1cOIT (K-)


Figure 2-5. Conductivity of (Bi203)1-x(Er2O3)x in air; (o) x = 0.2, (X) x = 0.25, (A) x =
0.3, (V) x = 0.35, (o) x = 0.4, (e) x = 0.45, (A) x = 0.5, (V) x = 0.6; broken
line represents the conductivity of pure Bi203 (from ref. 21).








19









!.0

4750C

---- seoC



6 0.6



0.4



~ 0.2



0.0

0 100 200
Time (kr)
Figure 2-6. Decay in isothermal oxygen ion conductivity of 20 mol% ESB on annealing
below the transition temperature (from ref. 31).










5.470
100 c--r--


0



Yb/


F


1.00 1.01


LATTICE PARAMETER (A)
5.480 5.490


o Cation Radii
Lottiec Poromctcr


1.02 1.03 1.04
CATION RADII (A)


Figure 2-7. Time constant as function of cation radii and lattice parameter of 25 mol%
cation doped bismuth oxide (from ref. 30).


LATTICE IAHAMk: I I:H (A)
5.480 5.500 5.520


1


VV




10


9


5.540


10 15 20 25
Er.O. CONCENTRATION (Mole%)


Figure 2-8. Time constant as function of erbia concentration and lattice parameter of
erbia stabilized bismuth oxide (from ref. 30).


1.05


Er.U3/ Y,


AA














CHAPTER 3
DIRECT CURRENT BIAS STUDIES ON (Bi203)o.8(Er2O3)0.2 ELECTROLYTE AND
Ag-(Bi203)o.8(Er2O3)o.2 CERMET ELECTRODE

3.1 Introduction

Bismuth oxide based electrolytes show one of the highest oxygen ion conductivity

among all solid electrolytes and are of considerable interest for application in solid oxide

fuel cells and oxygen sensors. 17-19 The high temperature 6-Bi203 has a cubic fluorite

structure with 25 % inherently vacant oxygen sites. The high concentration of disordered

oxygen ion vacancies in 6-Bi203 along with the high polarizability of Bi3+ cation results

in oxygen ion conductivity which is one to two orders of magnitude higher than that of

yttria stabilized zirconia at comparable temperatures. However, on cooling to lower

temperatures the high temperature 6-phase transforms to the stable monoclinic a-phase at

729 'C. Two intermediate metastable phases (tetragonal 1-phase and body centered cubic

y-phase) are also known to exist below 650 'C. The 03 and )-phase have an ordered

oxygen ion sublattice with oxygen ion conductivity up to three orders of magnitude lower

than the 6-phase, while the a-phase is a p-type electronic conductor. Further, the

transformation from 6-phase to P3-phase is accompanied by a sudden large volume change

which destroys the mechanical integrity of the material during thermal cycling.

Therefore, it is necessary to stabilize the high temperature and high oxygen ion

conductivity 6-phase at lower temperatures in order to utilize Bi203 in practical

applications. 7-'9 Takahashi et al.29'3 were first to show that by doping Bi203 with

isovalent rare earth oxides the high temperature 6-phase can be retained at lower









temperatures. Unlike stabilized zirconia, doping does not result in increase in the oxygen

ion vacancy concentration but only serves to stabilize the high temperature phase at lower

temperatures resulting in enhanced oxygen ion conductivity.

Arrhenius plots of oxygen ion conductivity of these stabilized bismuth oxides show

a characteristic kink at transition temperature (T*) -600 'C with a lower activation

energy above T* and a higher activation energy below T*. Moreover, the conductivity

shows a continuous decay with time on isothermal annealing in air below this transition
temperature.3 32 The conductivity decay has traits similar to those of nucleation

and growth. The rate of decay is fastest at -500 'C and is comparatively slower both

above and below 500 C.23,31,32 The conductivity decay is reversible and can be achieved

by heating above the transition temperature.23'2428

Transmission electron microscopy studies on stabilized bismuth oxides show that

the isothermal anneal below the transition temperature results in the formation of

superstructure with a lattice parameter twice that of the parent structure. 23-2 Neutron

diffraction studies have shown that after annealing the oxygen ions are also displaced

from the center of the cation tetrahedra (8c-sites) towards the faces of the cation

tetrahedra (32f-sites) along the <1 l1> directions.22'2326'27 Calorimetric studies on

annealed samples show the presence of an endotherm at about the transition temperature;

the magnitude of the endotherm increases with the annealing time.23,28,3133 The cooling

cycle does not show an exotherm indicating that the reverse process is comparatively

much slower.

The above described conductivity and structural changes in stabilized bismuth

oxides on annealing below the transition temperature have been attributed by Wachsman








and coworkers22,27 to occupational and positional ordering of the oxygen ion sub-lattice:

ordering of oxygen vacancies in directions resulting in the formation of a 2X2

super-lattice and displacement of oxygen ions from 8c to 32f site also along the <11 1>

directions. On isothermal annealing below the transition temperature, the mobile oxygen

vacancies are consumed in the growing ordered clusters and become unavailable for

conduction resulting in the conductivity decay with time. This process is reversible and

upon heating above the transition temperature the lattice disorders with resultant regain in

conductivity. The endotherm at the transition temperature is due to the heat energy

required for disordering the ordered lattice. The magnitude of the endotherm increases

with the annealing time indicating that the extent of ordering increases with the annealing

time.

In most applications, solid electrolytes will be used in an electrochemical potential

gradient with resultant oxygen ion flux, and therefore, it is important to understand the

effect of oxygen ion flux on the conductivity decay and the ordering kinetics of bismuth

oxide solid electrolytes. Among lanthanide stabilized bismuth oxides, 20 mol% erbia

stabilized bismuth oxide (ESB) shows the highest oxygen ion conductivity and is used in

the present study. To generate the oxygen ion flux, direct current bias was applied across

a symmetric cell consisting of ESB electrolyte and Ag-ESB cermet electrodes, and

electrochemical impedance spectroscopy was used to measure the oxygen ion

conductivity of the electrolyte as a function of temperature, time, and current bias.

Further, the heat of enthalpy of the order-disorder transition was measured using

differential scanning calorimetry (DSC) for samples annealed at different temperatures,

time periods, and current bias to understand the effect of oxygen ion flux on the ordering








kinetics. Ag-stabilized bismuth oxide cermets were used as electrodes in this study as

they have been reported to have high performance as cathodes for IT-SOFCs by Xia et

al.34 It was hoped that these electrodes would prevent the gas diffusion limited

decomposition of bismuth oxide based electrolyte to Bi metal at high currents as

mentioned by Takahashi et al.35 and Doshi et al.36 However, it was found during the

course of this study that there are some long term stability issues with these electrodes.

3.2 Experimental

Standard solid state method was used to fabricate (Bi203)0 8(Er2O3)02 electrolyte

discs. Bi203 (99.999 %, Alfa Aesar) and Er203 (99.99 %, Alfa Aesar) powders in desired

weight ratios were ball milled in acetone for 24 hours with zirconia ball media and

calcined at 800 'C for 10 hours. The resulting powder was again ball milled, sieved,

uniaxially pressed, cold isostatically pressed, and finally sintered at 890 'C for 10 hours.

For the electrochemical experiments, sintered discs were polished to a final thickness of

-0.76 mm. Diameter of the discs after sintering was -10.9 mm.

For the cermet electrodes, (Bi203)0.8(Er2O3)02 powder was prepared using an

amorphous citrate route. Bi(NO3)3.5H20 (99.999 %, Alfa Aesar) and Er(N03)3.5H20

(99.9 %, Alfa Aesar) in desired weight ratios were first dissolved in dilute nitric acid

solution, and then citric acid was added in a metal cation:citric acid molar ratio of 1:1.5.

The solution was gelled and foamed at 80-100 'C. The precursor was then calcined at

temperatures between 400 'C and 700 'C. Ag-(Bi203)0.8(Er2O3)02 electrode paste was

made by ball milling Ag20 (Alfa Aaesar) and (Bi203)0.8(Er2O3)02 in 60:40 wt.% ratio

with Heraeus V006 binder for 24 hours. For the symmetric cells, the electrode paste was

screen-printed on both sides of the electrolyte to which Ag lead wires were attached, and

the assembly was sintered at 750 C for 1 hour.









For the experimental setup, a frequency response analyzer (Solartron 1260) along

with an electrochemical interface (Solartron 1287) was used. A constant oxygen ion flux

was generated inside the electrolyte by using the electrochemical interface in the

galvanostatic mode. Impedance measurements were done every 30 minutes to separate

the electrolyte resistance from the electrode response. The cell was given a pretreatment

step in order to avoid the overloading of the current range during impedance

measurements. The samples were placed in a quartz experimental setup with Au lead

wires which was then put into a horizontal tube furnace, and the experiments were

conducted in air. Different samples were used for each temperature, time, and current

bias experiment. All samples were equilibrated at 650 'C and then cooled down to the test

temperature. The experiments were terminated at different time periods depending on

current bias and temperature in order to avoid the electrochemical decomposition of the

electrolyte.

Samples were characterized using XRD (APD-3720) and SEM (JEOL-6400).

Electron probe micro analysis (EPMA) was done on the sample cross-sections to

characterize the concentration profile of elements as a function of thickness. DSC (TA-

1290) studies were done to measure the heat of transformation of samples under different

conditions. The electrodes were polished off, and bulk samples of-9 mg mass were

heated from room temperature to 800 'C at 20 C/min in 100 sccm of argon.

3.3 Results and Discussion

3.3.1 Processing

Calcination of the oxide mixture at 800 C resulted in single phase powders and

sintering at 890 'C resulted in discs with more than 95 % density. XRD patterns for the

powders derived from the citrate route after calcination at different temperatures are









shown in figure 3-1. The citrate route did not significantly decrease the calcination

temperature to achieve the cubic fluorite structure; the primary phase up to 600 'C was

tetragonal P-Bi203. However for the cathode paste, powder calcined at 500 'C for 10

hours was used which eventually would form the cubic phase on sintering at 750 'C.

SEM micrographs of the cross-section of the electrode/electrolyte interface in the

secondary and back-scattered electron mode are shown in figure 3-2 (a) and (b),

respectively. The thickness of the electrode is -17 lim. After sintering, the electrode

seems to have low porosity possibly due to the low melting points of both Ag and

bismuth oxide. This microstructure is similar to that obtained by Xia et al.34 with Ag-

yttria stabilized bismuth oxide electrodes.

3.3.2 DC Bias and Impedance Spectroscopy: ESB Electrolyte

Figure 3-3 is a plot of log(oT) vs l/T for the ESB electrolyte in air. The graph

shows the well-known kink in the slope at T* -600 'C with activation energy equal to

0.68 eV above the transition temperature and 1.28 eV below the transition

temperature. 1923 Below the transition temperature, ordering of the oxygen ion sublattice

results in higher activation energy. Figure 3-4 is a plot of percentage change in resistance

of ESB as a function of time on isothermal annealing at 300 'C, 500 'C, and 625 'C. The

rate of increase is fastest at 500 'C, much slower at 300 C, and negligible at 625 'C. For

example after 40 hours, the resistance increase at 625 'C was -0.4 %, at 500 'C was

-898.4 %, and at 300 C was -12.2%. This suggests that the ordering has characteristics

typical to those of nucleation and growth, where the transformation rate depends on the

drive force as well as on the diffusion rate. At 625 'C, the driving force is small, and at

300 'C, the diffusion rate is slow.









Most of the dc bias experiments were carried out at 500 'C, where the ordering of

oxygen ion sublattice is fastest. On applying the bias at 500 'C, the initial voltage drop vs

time plots showed a valley; depth of the valley increased with the bias current as shown

in figure 3-5. This gives an impression that the bias is affecting the ordering process.

However, impedance spectroscopy done after the first 30 minutes of application of the

bias separated the electrolyte resistance from the electrode response and showed that the

valley is most likely related to the activation polarization at the electrode as described in

a following section. On the other hand, the electrolyte resistance increased with time at

500 'C without much affect of the dc bias as shown in figure 3-6. It was observed that at

high currents the increase in resistance was slightly faster though within the error limits

of the experiment. It should also be noted that the experiments at higher current bias were

much shorter, in comparison to lower currents, in order to avoid the electrochemical

decomposition of the electrolyte with voltage drops greater than 0.65 V across the

sample. For example, at 72 mA bias the experiment was terminated after -4 hours and in

the case of 12 mA bias the experiment was terminated after -42 hours. Another set of

samples were studied at 500 'C to confirm the results from 0-36 mA, and the electrolyte

resistance after 10 hours were essentially the same in all the samples.

At 625 'C i.e. above the transition temperature, the percentage change in resistance

as a function of current bias is shown in figure 3-7. The data has considerable noise,

which is basically due to the electrode's microstructural instability caused by the

electromigration of Ag as discussed in a following section. At 250 mA bias, the migration

was severe, and over the period of the experiment, the migration led to the formation of a

dense Ag layer on the counter electrode. This resulted in the eventual reduction of the









electrolyte to Bi metal up to -0.2 mm in depth as found by EPMA on the electrolyte

cross-section. However towards the end of the experiment, for samples with no bias and

for samples under bias on removal of the bias, the electrolyte resistance was the same as

that of the initial value. At 300 'C, dc bias experiments were not conducted due to high

resistance of the electrolyte.

3.3.3 DSC Studies: ESB Electrolyte

Stabilized bismuth oxides on annealing below the transition temperature show an

endotherm at about the transition temperature during subsequent heating.23283133 This

endotherm has been attributed to the heat energy required to disorder the ordered oxygen

ion sublattice. The ordering process is comparatively much slower, and hence, no

exothermic heat is observed during the cooling cycle and even after -10 hours of anneal

at 500 'C as found in this study. However, on annealing for -74 hours an endotherm

appears at 614.5 'C with a magnitude of 4.12 kJ/mol as given in Table 3-1. The ordering

process continues beyond 100 hours of anneal as observed by Wachsman et al.32 Samples

annealed under bias showed similar kinetic behavior, though the endotherm appeared at

shorter annealing times. After -10 hours anneal, the sample under no bias did not show

the endotherm, while the samples under 12-36 mA bias showed the endotherm as given

in Table 3-1. Moreover, it appears that for a constant annealing time the magnitude of the

endotherm increases with the bias current.

As expected, above the transition temperature at 625 'C there was no endothermic

peak observed even after annealing for 137 hours. Samples under 100 mA and 250 mA

bias at 625 'C also did not shown any endothermic peak. On annealing at a much lower

temperature of 300 0C for -49 hours, an endotherm of magnitude of 3.08 kJ/mol was

observed at 606.6 'C.









The DSC results on samples annealed at 500 'C do not correlate well with the

resistance measurements. After -10 hours, without bias the electrolyte showed more than

200 % increase in the resistance though without any trace in the DSC, and under bias the

sample showed similar increase in resistance but with an endotherm in the DSC profile.

This difference could be simply due to the detection limits and characteristics of the

technique employed or due to difference in the nucleation and growth characteristics of

the ordered domains under bias and without bias. The magnitude of the endotherm is

primarily dependent on the volumetric amount of the ordered phase present, while the

resistance depends on the parallel and series paths available for conduction. For example,

resistance will be more sensitive than DSC to minute amounts of the ordered domains

present at the grain boundaries.

As mentioned earlier, ordering of the oxygen ion sub-lattice comprise of two parts:

positional displacement of oxygen ions from 8c to 32f sites along <1I1> and

occupational ordering of oxygen vacancies along <11 1>.22,27 Positional displacement is

expected to be the faster of the two as the magnitude of the displacement is less than the

oxygen ion radii, and in fact, it has been observed to be partially present in un-annealed

samples.27 In contrast, occupational ordering involves rearrangement of the oxygen lattice

over larger distances on the order of the lattice parameter, and thus, it is expected to have

a higher time constant. The endotherm is expected to be primarily related to the

occupational ordering, and hence, it appears only after long anneal time under no bias

conditions. Since the endotherm at short time is enhanced by the applied bias with

negligible change in conductivity decay, it is possible that the decay in conductivity with

anneal time is more related to positional ordering than occupational ordering.









Kruidhof et al.37 have noticed that erbia stabilized bismuth oxides are non-

stoichiometric and slowly uptake oxygen on cooling below 550 'C in an oxygen

atmosphere. The ordering process, as evinced by the endotherm, is also a slow process

which occurs below -600 C. The matching temperature-time profile and species suggest

that the oxygen uptake and oxygen ion ordering (possibly occupational) could be related

processes. DSC results under bias conditions could also be explained using the

assumption that the current bias is assisting in the oxygen uptake and hence, in the

ordering process. These and other possibilities require further extensive studies, though it

is clear from the conductivity and DSC studies that the applied current bias is not

effecting the conductivity decay in erbia stabilized bismuth oxide below the transition

temperature but could be assisting in the occupational ordering.

3.3.4 DC Bias and Impedance Spectroscopy: Ag-(Bi2O3)o(Er2O3)o.2 Electrode

Area specific resistance (ASR) of the Ag-(Bi203)o8(Er2O3)02 electrode as a

function of temperature in air is shown in figure 3-8, with values of 3.08 92cm2 at 500 C

and 0.16 Qcm2 at 625 "C. The electrode ASR was calculated by multiplying the electrode

resistance by the electrode area (-0.93 cm 2) and dividing by 2 to account for the

symmetric cell. The performance of the cermet electrode is comparable to that reported

by Xia et al.34 and is encouraging for application as cathode in IT-SOFC.

Impedance plots of the electrode at 500 'C under 0-72 mA current bias are shown

in figure 3-9. The electrode impedance decreased with bias current as shown in figure 3-

10, which suggests that within the current range studied the electrode impedance is

dominated by activation polarization. As shown in figure 3-9 (c), the electrode impedance

consisted of at least two arcs: a low frequency arc with a characteristic frequency at -0.2

Hz and a high frequency arc with a characteristic frequency at -630 Hz. The plot of the









imaginary part of impedance (Z") vs. log frequency (f) provides a better representation of

the frequency distribution in the impedance data. On application of the bias current, the

dominant low frequency arc divided into two arcs with significantly smaller impedance,

while the impedance of the high frequency arc remained unaffected. The characteristic

frequency of the middle arc increased with the bias current, while on the other hand the

characteristic frequency of the low and high frequency arcs appear to be independent of

the applied current bias. The electrode arcs represent the different reaction steps in the

electrode reaction, and it is suggested that the low and middle frequency arcs are related

to charge transfer step, while the high frequency arc is related to the bulk transport in the

electrode.

Impedance plots of the electrode at 625 C under different current bias are shown in

figure 3-11. As expected, the activation polarization at 625 'C is much smaller in

comparison to 500 'C as shown in figure 3-10. The electrode impedance consisted of at

least two arcs: a low frequency arc with a characteristic frequency at -20 Hz and a high

frequency arc with a characteristic frequency at -400 Hz, as shown in figure 3-11 (b). As

in the case of 500 'C, impedance of the dominant low frequency arc decreased with the

bias current, while the impedance of the high frequency arc remained unaffected.

However, at 625 C the electrode performance was not stable under 250 mA bias, and the

electrode resistance increased by six times after -23 hours as shown in figure 3-12.

Magnitude of the high frequency electrode arc increased with time and eventually

became the dominant arc, which indicates that it might be related to some sort of

concentration polarization. Interestingly, magnitude of the high frequency arc reduced on









removing the bias at the end of the experiment, and once again, the low frequency arc

became the dominant arc.

After the experiment, the top surface of the working electrode looked silvery which

was initially light brown in color, and the counter electrode peeled off very easily with a

silver layer at the electrode/electrolyte interface. SEM micrographs of the

electrode/electrolyte interface at the working and the counter electrode after -26 hours at

625 C under 250 mA bias are shown in figure 3-13. Significant electromigration of Ag

occurred along with the oxygen flux to the electrode/electrolyte interface at the counter

electrode and to the top surface at the working electrode. A schematic of the proposed

electrode reaction mechanism is shown in figure 3-14. It is suggested that the low

frequency arc is related to the charge transfer step, while the high frequency arc is related

to the bulk transport of oxygen in the silver particles. The electrode response under no

bias is limited by the charge transfer step, and on application of the bias the charge

transfer resistance decreases, though along with Ag electromigration in the electrode. The

migration with time results in the formation of a dense Ag layer at the

electrode/electrolyte interface and in the bulk transport contribution in the electrode

impedance. As mentioned earlier, the formation of the dense Ag layer resulted in the

eventual reduction of the electrolyte at the counter electrode; however, it was not detected

in the impedance plots.

Oxygen diffusivity (Do) and solubility (C) in Ag layer for the case of finite

Warburg diffusion could be calculated using the following equations

62"(omax/2Do) = 4/1 3-1

ZD((O-+O) = R" T/(n2.F2.Co.Do) 3-2








where COmax is the frequency at the maximum value of the imaginary component, 6 is the

diffusion length, and ZD(o---O) is the real intercept.38 For Ag-ESB electrode at 625 'C,

250 mA after -26 hours, the calculated values for Do and C0, assuming 1 grm diffusion

length, are 6.1 9X 10-6 cm2/sec and 2.91 X1 0-6 moles/cm3, respectively. The calculated

values match well with experimental values for Do (7.42X10-6 cm2/sec) and C0

(1.94X1 0-6 moles/cm3) reported in the literature,39 which supports the argument that the

high frequency arc in the electrode impedance is related to the oxygen diffusion in Ag.

Phase diagram between Ag and Ag20, as proposed by Assal et al.,40 is shown in

figure 3-15. Ag forms a eutectic with Ag2O at 530 'C and 519 atm. 02. The liquidus

temperature in an oxygen free atmosphere is 962 C, which is lowered to 951 'C and 939

'C in air and 1 atm. 02, respectively. The high performance of the Ag electrodes for

oxygen reduction is because of its oxygen solubility which also results in the formation of

the eutectic, and as found in this study will consequently lead to electromigration of Ag

along with the oxygen flux above the eutectic temperature. Therefore, electrodes

containing Ag in sufficient proportions should not be expected to be microstructurally

stable for long periods of time even at intermediate temperatures. Ag-Pt and Ag-Pd alloys

with higher melting points can be better than Ag, though they will be much costlier

alternatives.

3.4 Conclusion

The effect of direct current bias on the ordering phenomenon of ESB electrolytes

and on the performance of Ag-ESB electrodes was studied. Impedance studies on ESB

electrolyte showed that on isothermal annealing below the transition temperature at 500

'C, the current bias did not have a significant effect on the rate of increase in resistance.

On the other hand, DSC studies showed that the endotherm related with the reverse









transition appeared at shorter annealing time periods on application of the current bias.

The above results are explained in terms of the relative effect of applied current bias on

the kinetics of occupational and positional ordering on the oxygen ion sub-lattice. It is

suggested that the rate of occupational ordering is enhanced under the bias current

resulting in the presence of endotherm at shorter time periods but without much change in

rate of increase in resistance. Ag-ESB cermet electrodes showed good electrochemical

performance. Impedance studies under current bias showed that the electrode response

consisted of charge transfer and bulk transport contributions. The electrode

microstructure was unstable at high bias currents at 625 'C, primarily due to the

electromigration of Ag along with the oxygen flux. The high oxygen solubility in Ag

results in good performance of the electrode, but it also results in the lowering of the

melting point of Ag and its consequent electromigration. Ag-Pt and Ag-Pd alloys with

higher melting points are suggested for replacing Ag, though they would increase the

cost.












8000


700C


00
4000
L, k 500jC



20000c
S0o0400o C

100 c-

20 30 40 50 60 70 80
20


Figure 3-1. XRD patterns of calcined 20 mol% erbia stabilized bismuth oxide powders
derived from amorphous citrate route.


































____=


ESB


Figure 3-2. Cross-sectional micrographs of the ESB electrolyte and Ag-ESB electrode
interface in (a) secondary electron and (b) back-scattered electron mode.



















0.68 eV *
4-


1 1.1 1.2


0


1.28 eV


1
1.3 1.4 1.5 1.6


1000/T (K-)


Figure 3-3. Arrhenius plot of oxygen ion conductivity in ESB electrolyte.


-2
0.9


1.7
















1000


800


0

0


600-


* OmA 300 oC
OmA 500 C
OmA 625 C


400:-


200


0
10 100 1000 10~


Time (min)

Figure 3-4. Percentage change in resistance of the ESB electrolyte under no bias at 300
T, 500 C, and 625 C.











0.525


0.475 72mA


0.425

0.375


,-0.325

'0.275

0.225

0.175


I


0.125

0.075


0 5 10 15 20 25 30 35

Time (min)


Figure 3-5. Initial voltage vs. time plots at 500 C under different bias currents.


36 mA


30 mA

24 mA


12 mA


--T


















OmA
12mA
* 24mA
A 30mA
v 36mA
0 72mA


500 C


100


1000


Time (min)


Figure 3-6. Percentage change in resistance of the ESB electrolyte at 500 'C under
different bias currents (Ro = 4.4 Q).


1000


800


0
0


600


400


200


0


104
















*OmA
I 1OOmA
250mA


1O


625 C


Time (min)


Figure 3-7. Percentage change in resistance of the ESB electrolyte at 625 'C under
different bias currents (Ro = 0.7 n).


50


40:

30

20


0

0


0

-10


1000


2000


3000

















Temp (C) Bias (mA) Time (hr) AHt (kJ/mol) Tt (C)

300 0 49.5 3.08 606.6

500 0 10 -

500 0 74 4.12 614.5

500 12 10 3.40 607.1

500 12 42 5.32 599.7

500 30 11.5 3.38 603.8

500 36 7 2.63 594.1

500 36 10 4.02 612.3

500 72 4 -

500 100 3.5 -

625 0 137 -

625 100 24 -

625 250 26 -


Table 3-1. Enthalpy for order-disorder transition and transition temperature of ESB
electrolyte on isothermal annealing at different temperatures, bias currents,
and time periods.















10.


Cl
S
0
2


*- 1.41 eV


0.1


0.01
1.05


1.1


1.15 1.2
1000/T (K-')


Figure 3-8. Arrhenius plot of ASR (Qcm2) for Ag-ESB electrode.


1.25


1.3















500 C


OmA


500 C


-1.5


12mA
-1 U

24mA
-0.5 630 Hz


0 2mA


0.5 '


4 4.5 5 5.5
z' (0)
(a)


6 6.5


0.2 Hz


OmA


72mA


MnOmmmmnmm 12mA
24mA
rf 30mA
36mA


4 5 6 7
z, (02)


8 9 10


S. -2.5


-0.6 24mA**0

-0.53m *
30m.A-


-0.4


-0.2


-0.1 so 72mA


1 10
f(Hz)


100 1000


04
0.1 1 10 100 1000 10
f(Hz)


Figure 3-9. Impedance plots of Ag-ESB electrode at 500 'C under different bias currents
(a) and (b) imaginary vs. real impedance (c) and (d) imaginary impedance vs.
frequency.


500 C


500 C


-1.5


0
0.1


OmA
* 12mA
* 24mA
30mA
36mA
* 72mA


6.5





60












7
6

05

"" 4
500C
S 3



625 C
0 -0
0 100 200 300 400
Bias Current (mA)

Figure 3-10. Ag-ESB electrode resistance vs bias current at 500 C and 625 C.










-0.3


625 C


-0.2


400
-0. 1 \



0



0.1 '
0.7


20 Hz


0 mA


'%
100 mA
mA


0.9
Z' (Q)


-0.14


-0.105


-0.07


-0.035


U
U
U *
** *,,.I m *
* +.. -.v


1 10 100
f (Hz)


1000 104


Figure 3-11. Impedance plots of Ag-ESB electrode at 625 C under different bias
currents (a) imaginary vs. real impedance (b) imaginary impedance vs.
frequency


625 C


0rnA
100 mA
250 mA
350 mA















625 C, 250mA


400 Hz


30 min
* 510 min
* 570 min
660 min
v 720 min
* 1290 min


0.6 0.8 1 1.2 1.4 1.6 1.8 2 2.2
Z' (Q)


-0.5


30 min
510 min
570 min
660 min
720 min
1290 min


V
ej V
OV V 0
V V
V.
"-T


625 C, 250 mA


1 10 1
f(Hz)


00 1000 104


Figure 3-12. Impedance plots of Ag-ESB electrode under 250 mA bias at 625 C (a)
imaginary vs. real impedance (b) imaginary impedance vs. frequency.


-1.2


-0.8


-0.4


0


-0.4 U

-0.3

-0.2

-0.1


0.1
0.1



















































Figure 3-13. Cross-sectional microstructures of the electrolyte/electrode interface after
annealing experiments at 625 C under 250 mA bias for -26 hours (a) working
electrode (b) electrolyte interface after counter electrode peeled off.














O2

Roads













(a)




02


Roads








(b)


Figure 3-14 Proposed electrode reaction mechanism for Ag-ESB electrodes.
















1300 I I :
;-1 0 1 2 .3 Iog(p z, bar)
1235
1200-1


1100
;! I,

1000

900- (Ag) liquid

800- 0


700- (v,

SI 88a I
0 0.05 0.10 0.15 0.20 0.25 0.30 0.35
Ag Mole fraction 0

Figure 3-15 Phase diagram between Ag and Ag20 proposed by Assal et al. (from ref.
42).













CHAPTER 4
BISMUTH RUTHENATE BASED CATHODES FOR IT-SOFC

4.1 Introduction

For efficient operation of SOFCs at intermediate temperatures, better performance

electrolytes and electrodes are required. Gadolinium doped ceria (GDC) shows

significantly higher ionic conductivity compared to YSZ at the same temperature and is

considered a promising electrolyte material for IT-SOFCs.41 For cathodes, perovskites

based on lanthanum manganite, lanthanum cobaltite, and lanthanum ferrite have been the

material of choice.42 Recently, pyrochlores based on bismuth ruthenate, lead ruthenate,

and yttrium ruthenate have been studied for application as cathodes in SOFCs. 4345

Ruthenium oxide is known to be catalytic active towards oxygen reduction and has been

studied as a cathode material.46 It can be expected that solid solutions containing

ruthenium oxide can also be beneficial as cathodes. Pyrochlore ruthenates are electrically

conductive and thus, satisfy the other requirement for the cathodes. In this work, cathodes

based on bismuth ruthenate were studied for IT-SOFCs based on ceria electrolytes.

4.1.1 Undoped Bismuth Ruthenate

Abraham et al.47 first showed that bismuth ruthenate apart from the fcc pyrochlore

structure Bi2Ru207 also exists in the oxygen enriched cubic KSbO3 structure as

Bi2Ru2073 (or Bi3Ru3O1 1). At 975 'C, Bi2Ru2073 transforms irreversibly to Bi2Ru207 in

air. Bi2Ru2O7 is phase stable at 800 'C and did not transform back even after a month.47

Bi203-RuO2 phase diagrams reported by Prosychev et al.48 and Hrovat et al.49 are shown

in figure 4-1 and 4-2, respectively. There appears some confusion regarding the transition









temperature from the KSbO3 structure to the fcc pyrochlore structure (975 C47 965 0C,48

950 C49). Hrovat et al.49 noted that the reaction of precursors with alumina crucible could

have affected the transition temperature.

Linquette-Mailley et al.50'5 have used both Bi2Ru2O7.3 and Bi2Ru2O7 as electrodes

to reduce the response time of YSZ based oxygen sensors at low temperatures. They

found that the oxygen content in Bi2Ru207 is slightly over-stoichiometric in the pressure

range of 3X103 to 105 Pa at 800 C, while Bi2Ru207.3 reduces irreversibly to Bi2Ru207 at

pressures less than 7.4X102 Pa at 800 C. Further, from redox potentiometric

measurements they concluded that Bi2Ru2O7 reduces and becomes a mixed ionic

electronic conductor at cathode polarizations higher than 700 mV/air at 373 'C; therefore,

the length of triple phase boundaries (TPBs) in the electrode structure is very important at

low polarizations.

Bae et al.43 studied Bi2Ru2073, Pb2Ru2065 and Y2Ru207 as cathode materials for

IT-SOFCs based on ceria electrolytes. They found that Bi2Ru207.3 and Pb2Ru2065 react

with ceria at the processing temperatures, while Y2Ru207 was stable with ceria at 900 'C

and was studied in detail using impedance spectroscopy. Y2Ru207 electrode showed an

area specific resistance (ASR) of 4000 Qcm2 at 627 'C, and on doping with 5 mol% Sr,

the electrode ASR reduced to 47 f cm 2. This improved performance of the electrode was

explained in terms of the enhanced ionic conductivity of Y2Ru2O7 with Sr doping.

Takeda et al.4 studied pyrochlores Bi2Ru207, Pb2Ru2O6.5 and perovskites

CaRuO3, SrRuO3 as cathode materials for YSZ based SOFCs. Bi2Ru207 and Pb2Ru206.5

were found to be stable with YSZ at 900 'C. However, the presence of sillenite type-

impurity phase in Bi2Ru2O7 could result in the formation of monoclinic zirconia. They









observed metallic behavior for pyrochlore ruthenates with almost temperature

independent conductivity (102_103 S/cm from room temperature up to 900 'C), which is

comparable to the best conventional cathode materials. The thermal expansion coefficient

of Bi2Ru207 and Pb2Ru2O65 between 700-900 'C was measured to be 0.99-1.00 X 10-5 K
'and 1.10-1.21 X l05 K- respectively. Pyrochlore ruthenates showed better electrode

performance than the perovskite ruthenates, and their performance was comparable to

lanthanum manganate but inferior to that of lanthanum cobaltite based cathodes.44

4.1.2 Doped Bismuth Ruthenates

A2B207 pyrochlore structure is essentially derived from an oxygen deficient cubic

fluorite structure with both ordered cation and anion sub-lattice, as shown in figure 4-3. It

exhibits Fd3m space group with eight formula units within a cubic unit cell.8 The cation

sublattice consists of bigger A3+ and smaller B4+ which order into alternate (110) rows in

every other (001) plane and in alternate( 10) rows in the other (001) planes. This cation

ordering provides three distinguishable tetrahedral sites for the oxygen ions: 8a-sites

surrounded by 4 A3+ cations, 8b-sites surrounded by 4 B4+ cations, and 48f-sites

surrounded by 2 A3+ and 2B4+ cations. 8a and 48f-sites are occupied, while 8b-sites are

vacant resulting in an ordered oxygen ion sub-lattice. The formula unit of the pyrochlore

can also be written as A2B2060' to distinguish between the oxygen ions occupying the

48f-site as 0 and those occupying the 8a-sites as 0'.

Although bismuth ruthenate pyrochlore has sufficient electronic conductivity to

perform as a good cathode, it will be beneficial to introduce oxygen ion conductivity into

the structure. Presence of ionic conductivity in the structure will not limit the oxygen

reduction reaction to the TPBs, leading to enhanced electrode performance. The oxygen









ion vacancies in the pyrochlore structure are ordered but still result in an ionic

conductivity, which is larger than that of undoped fluorites. It is expected that the

pyrochlore structure may exhibit a transition temperature above which the oxygen ion

sub-lattice will disorder, leading to enhanced oxygen ion conductivity.8 As the anion

ordering is related to the cation ordering, the transition temperature could be lowered by

manipulating the ion radii of the two cations. The smaller is rA/rB, the lower is the

expected transition temperature. As in the case of fluorites, oxygen ion conductivity can

also be introduced into the pyrochlore structure by generating additional oxygen ion

vacancies by doping with lower valent cations. These additional oxygen ion vacancies are

expected to primarily occupy 8a and 48f-sites and hence, could contribute to oxygen ion

conduction. However, one has to remember that apart from the concentration of mobile

oxygen vacancies, the oxygen ion conductivity also depends on structure factors such as

jump directions, jump paths, and activation barrier for the oxygen ion motion and on

dopant-vacancy interactions.

Fortunately, bismuth ruthenate forms solid solution with large solubility limits with

a number of dopants on the A-site, and the doping does not significantly affect the

electronic conductivity at least at room temperature.52 In the present work, Ca2+ (r = 1.12

A), Sr2 (r = 1.26 A), and Ag (r = 1.28 A) with comparable ionic radii with the host Bi3

(r = 1.17 A) were studied as dopants on A-site. This strategy was used keeping in mind

that, possibly as in fluorite systems, the dopant with better matched ionic radii with the

host would generate lesser elastic strains in the lattice and would show higher oxygen ion

conductivity.









4.1.3 Bismuth Ruthenate and Stabilized Bismuth Oxide Composites

It is well known that a composite cathode, consisting of an electrocatalyst and an

oxygen ion conductive phase, enhances the performance significantly by effectively

extending the reaction zone from the electrode/electrolyte interface into the electrode.

Chemical compatibility between the two phases is critical in order to avoid the formation

of resistive tertiary phases. Composite cathodes based on La(Sr)MnO3-6-YSZ and

La(Sr)MnO3-&-GDC for YSZ electrolytes and La(Sr)Co(Fe)O3.8-GDC for GDC

electrolytes have been very effective in reducing the polarization resistance compared to

single phase cathodes.355 The performance of the composite cathode depends on the

relative ratio, particle size distribution, and spacial distribution of the two phases, so as to

achieve high concentration of TPBs and percolation for both the phases. In addition, the

electrode microstructure should also be porous to provide for the gas diffusion. TPBs in

the electrode are the boundaries, between the electrocatalyst, the electrolyte, and the gas

phase, where the charge transfer reaction is believed to take place. High concentration of

TPBs (large number of reaction sites) and the percolation of the two phases (continuous

pathway for electron and oxygen ion transport) results in low electrode polarization.

In this part of the work, (Bi2Ru207.3)-(Bi203)o.(Er2O3)0.2 composite cathodes were

studied. As shown in the phase diagrams in figure 4-1 and 4-2, the low temperature

bismuth ruthenate (Bi2Ru207.3) is the thermodynamically stable phase with Bi203 in the

temperature range of interest and hence, was chosen as the electrocatalyst phase in the

composite electrode. The electrolyte phase in the composite electrode was 20 mol% erbia

stabilized bismuth oxide (ESB), one of the highest known oxygen ion conductors.

Bi2Ru207.3 forms a eutectic with Bi203 at relatively low temperatures, though there is

some discrepancy between the two proposed diagrams regarding the eutectic temperature









and composition. Prosychev et al.48 proposed that the eutectic is between Bi2Ru2O7.3 and

a-Bi203 at -37 mol% RuO2. The eutectic is at 730 'C, which incidentally is also the

transformation temperature from a to 8-Bi203. Hrovat et al.49 instead proposed that the

eutectic is between Bi2Ru2073 and 8-Bi203 at -20 mol% RuO2. The eutectic is at 745 'C,

above a to 6-Bi203 transformation temperature. The presence of the eutectic could

significantly undermine the stability of the electrode microstructure during processing

and/or under operation. Differential scanning calorimetry (DSC) studies were done to

study the eutectic temperature and composition.

4.2 Experimental Work

Standard solid state synthesis was used to fabricate the bismuth ruthenate powders.

Bi203 (99.9995%, Alfa Aesar), RuO2.XH20 (99.99%, Alfa Aesar), CaCO3 (99%, Fisher

Scientific), Sr(N03)2 (99.97%, Alfa Aesar) and Ag (99.9%, Alfa Aesar) powders were

mixed in stoichiometric amounts in an agate mortar and pestle, and calcined at 900 'C for

10 hours to achieve undoped and doped Bi2Ru207 pyrochlore phase. To achieve the low

temperature Bi2Ru2O73 phase, mixed powders were calcined at 775 0C for 10 hours. The

calcined powder was leached with dilute HNO3 to remove the sillenite type-impurity

phase. For the composite electrodes, (Bi203)0.s(Er2O3)0.2 powders were prepared using

amorphous citrate route as described in the experimental section of chapter 3. For the

reactivity tests of Bi2Ru2O7 and Bi2Ru2O73 with the electrolyte, 11 mol% gadolinium

doped ceria powder (GDC, Rhodia) was mixed with the respective powders and heat

treated at 850 0C for 10 hours. DSC (TA-1290) studies were done to study the eutectic

between bismuth ruthenate and bismuth oxide. Powder mixtures of -13 mg mass were

heated from room temperature to 800 'C at 20 C/min in 100 sccm of air.









Symmetrical cells for impedance spectroscopy studies were fabricated by brush

painting the electrode paste on GDC electrolyte pellets (uniaxially pressed and sintered at

1450 'C for 6 hours). The electrode paste was made by mixing the desired weight ratios

of the powders with Heraeus V006 binder. The electrodes were dried at 150 'C and

sintered at temperatures between 750 and 850 C for 2 hours with Pt lead wires. The

samples were characterized using XRD (APD 3720) and SEM (JEOL 6400).

Impedance spectroscopy was done using a Solartron 1260 in the frequency range of 0.1

Hz -32 MHz at temperatures between 350-700 C. Oxygen partial pressure was varied

between 0.04-1 atm by using 02/air-N2 gas mixtures.

4.3 Results and Discussion

4.3.1 Processing

XRD patterns of the calcined Bi2Ru2O7 and Bi2Ru2O7.3 powders before and after

leaching are shown in figure 4-4 and 4-5, respectively. Leaching with dilute HN03 is

effective in removing the sillenite type-impurity phase (Bii2RuO20), which results in

predominant single phase powders.

XRD patterns of the powder mixtures of GDC with Bi2Ru2O7 and Bi2Ru2O7.3 after

heat treatment at 850 'C for 10 hours are shown in figure 4-6 and 4-7, respectively. There

are no new peaks identifiable in the patterns, which indicate that there is no reaction

between GDC and either of the bismuth ruthenate phases. Similarly, Hrovat et al.56

studied the subsolidus phase equilibria in RuO2-Bi2O3-CeO2 system and no ternary

compound was found at 800 'C. The tie line is between Bi2Ru2O7 and CelxBixO2-x/2 (0 <

x < 0.33) solid solution. These results are in contrast with those of Bae et al.,43 who found

unknown reaction products between Bi2Ru2073 and GDC at 800 'C. It is possible that the

presence of sillenite type impurities in bismuth ruthenate could lead to a new product, as









Takeda et al.44 found between Bi2Ru207 and YSZ where the presence of sillenite type

impurities resulted in the transformation to monoclinic zirconia.

Surface and cross-section micrographs of Bi2Ru2O7 electrode on GDC electrolyte

are shown in figure 4-8 (a) and (b), respectively. Solid state synthesis of Bi2Ru2O7

powders at 900 'C has resulted in a coarse particle size (-3 gim). After sintering at 850 0C,

the electrode is porous with a thickness of -100 gtm. Initial sintering experiments showed

that lower sintering temperatures resulted in poor adhesion of the electrode film with the

substrate, while at higher temperatures there was evaporation of the electrode material.

4.3.2 Impedance Spectroscopy Studies

4.3.2.1 Undoped Bi2Ru207 cathode

Impedance plots of Bi2Ru2O7 electrode on GDC at 500 'C and 700 'C are shown in

figure 4-9 and 4-10, respectively. High frequency intercept corresponds to the bulk

conductivity of the GDC electrolyte, which is comparable to values reported in the

literature. The depressed semicircle at lower frequencies is due to the Bi2Ru207 electrode

with resistance of 158.29 92 and 4.13 Q in air at 500 'C and 700 'C, respectively. The

electrode ASR was calculated by multiplying the electrode resistance by the electrode

area (-0.7 cm2) and dividing by 2 to account for the symmetric cell. The Arrhenius plot

of the electrode ASR (fcm2) in air is shown in figure 4-11. ASR values are significantly

smaller than those reported by Linquette-Mailley et al.51 for Bi2Ru207 electrode on YSZ

electrolyte. For example, they reported ASR of -160 fcm2 compared to 55.64 f~cm2 in

this study on GDC at 500 0C and -20 k'cm2 compared to 3.03 Qcm2 in this study at 650

'C. Also, they reported a kink in the Arrhenius plot with activation energy of -i.3 eV

below 567 'C and -1.0 eV above 567 'C, and suggested that there are two different rate

limiting steps in the electrode reaction in the temperature range of study. However, in this









study a single activation energy of -1.26 eV in the temperature range of 450-700 'C was

observed.

To understand the mechanism of oxygen reduction at the electrode, impedance

measurements were done as a function of oxygen partial pressure. In general, ASR of the

electrode varies with the oxygen partial pressure according to the following equation

[ASR] = [ASR]o (Po2)-m 4-1

where the magnitude of m provides an insight into the rate limiting step in the oxygen

reduction reaction at the electrode. With metal and metal oxide electrodes on solid

electrolytes: m = 0.25 has been associated with the charge transfer reaction at the TPBs,

m = 0.5 with the surface diffusion of the dissociatively adsorbed oxygen at the electrode

to the TPBs, and m = 1 with the gaseous diffusion of oxygen molecules in the electrode

structure. In addition to the electrode-electrolyte combination, the microstructure also

plays a key role in determining the rate limiting step.38,51,57-59

Figure 4-12 is a graph of In [ASR] vs. In P02. The values of m range primarily

between 0.5 and 0.6 (with the exception of 700 'C), which suggests that the rate limiting

step in the present case is surface diffusion of the dissociatively adsorbed oxygen at the

electrode surface to the TPBs. In contrast, Linquette-Mailley et al.51 reported values of m

-0.25 at temperatures less than 567 'C and continuously increasing values for

temperatures greater than 567 'C for Bi2Ru207 electrode on YSZ electrolyte. The rate

limiting step at lower temperatures was attributed to the charge transfer step and at higher

temperatures a transition to surface diffusion was suggested.

The better performance of the Bi2Ru207 electrode on GDC electrolyte compared to

that on YSZ electrolyte along with a different rate limiting step again demonstrates the









role played by the electrolyte in the electrode polarization. For an electrode with

negligible ionic conductivity such as Bi2Ru207 at low polarizations, the

electrode/electrolyte interfaces are the active TPBs for the electrode reaction to take

place. Hence, it can be expected that the bulk and surface properties of the electrolyte can

have a significant effect on the overall performance of the electrode. Higher ionic

conductivity of GDC compared to that of YSZ and the capability of cerium to exist in

multiple valence states to provide sufficient electronic charge carriers could result in

improved charge transfer and better electrode performance.60

4.3.2.2 Doped Bi2Ru207 cathodes

To further improve the performance of Bi2Ru207 cathodes, doping with aliovalent

cations was studied on Bi3+ site to introduce oxygen ion vacancies and hence, oxygen ion

conductivity in the structure. The presence of mixed ionic and electronic conduction

would activate the oxygen reduction reaction on the complete electrode surface and

would just not be limited to the TPBs at the electrode/electrolyte interface. Although,

Bi2Ru207 forms solid solutions with a number of dopants, the choice was short listed to
cations (Ca2+ Sr'+, Ag) with fixed valence and with close ionic radii match with Bi3

host.

Impedance plots of Ca doped Bi2Ru207 [(Bi1-xCax)2Ru207-6 BCRx, x = 5-30

mol%; Bi2Ru207 BRO7] at 500 'C and 700 'C are shown in figure 4-13 and 4-14,

respectively. At 500 'C, all dopant levels led to increase in the electrode polarization,

while at 700 'C, 5 and 10 mol% doping resulted in decrease in the electrode polarization

compared to undoped bismuth ruthenate. With x > 20 mol%, the electrode polarization is

an order of magnitude higher at each temperature along with additional electrode arcs at

lower frequencies, which indicates that with high dopant levels there is drastic change in









the electrode properties leading to multiple rate limiting steps. Arrhenius plot of the

electrode ASR in air is shown in figure 4-15. Doping results in the increase in activation

energy from -1.26 eV for undoped Bi2Ru207 to -1.37 eV for Ca doped Bi2Ru207, though

5 and 10 mol% Ca doped Bi2Ru207 show better performance at higher temperatures.

Impedance plots of Ag doped Bi2Ru207 [(Bi I-Agx)2Ru207-8 a BARx between x

5-20 mol%] at 500 C and 700 0C are shown in figure 4-16 and 4-17, respectively. The

performance was similar to that of Ca doped systems, with higher activation energy

compared to undoped Bi2Ru2O7 and better performance at higher temperatures with 5 and

10 mol% Ag doped Bi2Ru2O7 as shown in figure 4-18. Performance of 5 mol% Ca and

Ag doped systems are compared in figure 4-19. Both show better performance than

undoped Bi2Ru2O7 above -550 'C with similar activation energy of- .37 eV. Possibly,

Ca is a better choice of dopant as it shows slightly better performance. 10 mol% Sr doped

Bi2Ru207 (BSR10) showed inferior performance than undoped, 10 mol% Ca and Ag

doped Bi2Ru2O7, as shown in figure 4-20 and hence was not studied further.

Electrode polarization of 5 mol% Ca and Ag doped Bi2Ru2O7 was studied as a

function of P02 and dc bias in order to understand the rate limiting steps with these

electrodes. Impedance plots for 5 mol% Ca doped system are shown in figure 4-21 and 4-

22, while for 5 mol% Ag doped system are shown in figure 4-23 and 4-24. Plots of

ln(ASR) vs. ln(Po2) at temperatures between 400-700 'C for 5 mol% Ca and Ag doped

systems are shown in figure 4-25 and 4-26, respectively. The value of m for 5 mol % Ca

doped Bi2Ru207 ranged between 0.6 and 0.8 at 0.11 < P02 (atm.) < 1 and at a lower P02 of

0.04 atm., the value of m increased indicating a change in the rate limiting step. On the

other hand for 5 mol% Ag doped Bi2Ru207, the value of m ranged consistently between








0.5 and 0.6 in the complete P02 range similar to that of undoped Bi2Ru2O7, suggesting

that the rate limiting step is the surface diffusion of dissociatively adsorbed oxygen at the

electrode surface to the TPBs. Initial impedance measurements under direct current bias

on 5 mol% Ca and Ag doped Bi2Ru2O7 also support the argument at least at high

temperatures. Impedance plots for 5 mol% Ca doped system as a function of current bias

at 500 'C and 700 'C are shown in figure 4-27 and 4-28, respectively. At 500 'C, on the

application of current bias the electrode polarization decreased which indicated that the

rate limiting step is charge transfer. The electrode polarization behavior was more

complex at 700 'C; electrode polarization increased with current bias up to 15 mA,

decreased slightly between 15-25 mA, and then increased beyond 25 mA along with

additional electrode arcs. At higher temperatures, the charge transfer is fast, and hence, a

comparatively slower step in the electrode reaction could be rate limiting, which in the

present case is diffusion related.

Additional information of the rate limiting steps in the electrode reaction is given

by the frequency distribution of the imaginary part of the impedance. At a particular

temperature, each reaction step is associated with a characteristic frequency and therefore

the dominant reaction steps can be identified if their characteristic frequencies could be

resolved in the frequency distribution. Sometimes it is adequate as in the case of 20 and

30 mol% Ca doped Bi2Ru2O7, and 5 mol% Ca doped Bi2Ru2O7 under different bias

currents, where additional humps in the electrode impedance appear indicating the

presence of additional steps in the electrode reaction. However, sometimes the frequency

distribution information of imaginary part is inadequate to resolve if the characteristic

frequencies of two steps are too close.









As shown in figure 4-9, 4-10, 4-21 4-24, the characteristic frequency of the

electrode impedance for undoped, 5 mol% Ca and Ag doped Bi2Ru2O7 in general

increases with P02 indicating a change in dominant rate limiting step. Possibly at lower

P02, multiple reaction steps are present which determine the larger total impedance. This

phenomenon is more clear in the case of 5 mol% Ca doped Bi2Ru2O7, where there is

increase in the value of m at lower P02. This in effect again reflects on the inadequacies

of using the magnitude of m and frequency distribution of imaginary part of impedance to

determine the rate limiting steps present in the electrode reaction. In general it is not

reliably possible to deconvolute multiple humps present in the electrode response and

hence the magnitude of ASR and m are average of the different reaction steps taking

place over the P02 range of study, whereas the frequency distribution of imaginary part

may also not provide adequate resolution and identification of individual steps. New

methodologies are being adopted to extract more information from the frequency

distribution of the impedance. In one strategy adopted by Ivers-Tiffee and co-

workers,61'62 the fourier transformation of the imaginary part of the impedance have been

shown to provide much better resolution of the reaction steps present.

Doped Bi2Ru207 pyrochlore did not improve the electrode performance as

significantly as observed by Bae et al.43 with their studies on 5 mol% Sr doped Y2Ru207.

It is difficult to explain, without speculating, the observed results due to lack of

understanding of the basic defect chemistry; the effect of doping and P02 on the structure

and electrical properties of pyrochlore ruthenates. Studies concerning the same are

currently underway in our lab and it is hoped that they will shed more light on the

performance of these electrodes. However, one thing is clear that doping with aliovalent









cations on the Bi3+ site is not a very effective strategy to improve the performance of

Bi2Ru2O7 pyrochlore cathodes.

4.3.2.3 Bi2Ru207_-(Bi2O3)o.8(Er2O3)0.2 composite cathodes

From the studies done on undoped and doped Bi2Ru207, it is clear that the surface

diffusion of dissociatively adsorbed oxygen at the electrode surface to the TPBs is one of

the rate limiting steps involved in the oxygen reduction reaction at the electrode. In this

part of the work, composite cathodes consisting of Bi2Ru207.3 and (Bi203)o.8(Er2O3)0.2

were evaluated.

DSC plots in the heating cycle for powder mixtures of Bi2Ru2O7.3 with monoclinic

t-Bi203 and fcc fluorite (Bi203)0.s(Er2O3)0.2 are shown in figure 4-29. The behavior of the

two powder mixtures is apparently different; Bi2Ru207.3 and at-Bi203 showed an

endotherm at -735 *C, while Bi2Ru2O7.3 and ESB did not show any endotherm. One

possible explanation for this behavior could be that the eutectic composition is between

Bi2Ru2O7.3 and a-Bi203, as proposed by Proshychev et al.as However, this is not

conclusive as the heating rates used in the present study were high, which may not allow

the system to equilibrate.

Sintering temperatures were limited to 750 'C in order to avoid the transformation

of Bi2Ru2O.3 to the high temperature pyrochlore phase. With this limitation, the

electrode posed significant problems in terms of sinterability with the GDC electrolyte. A

number of trials ended in failure to prepare the single phase electrode with good adhesion

with the electrolyte. Composite cathodes were however prepared successfully, possibly

due to the high sinterability of the bismuth oxide phase. SEM surface micrographs of one

such composite cathode are shown in figure 4-30. Linquette-Mailley et al.5' were able to

deposit single phase Bi2Ru2O7.3 electrodes using a pyrosol deposition technique, though








with poor adhesion with the substrate. The performance of the Bi2Ru2073 electrode was

similar to that of Bi2Ru2O7 pyrochlore phase with comparable activation energy and

frequency distribution.51 Hence, it was felt appropriate in this study to compare the

performance of the composite electrodes with that of the pyrochlore single phase

electrode.

Impedance plots of the composite cathodes [Bi2Ru2O7.3:(Bi2O3)0 s(Er2O3)0.2 = X:Y

wt. ratio =- X:Y] at 500 'C and 700 'C are shown in figure 4-31 and 4-32, respectively.

Performance of the Bi2Ru207 pyrochlore phase is also shown for comparison.

Introduction of the ESB phase in the electrode resulted in the reduction of the electrode

impedance at each temperature along with a lower characteristic frequency, which

indicated a change in the rate limiting step for the oxygen reduction reaction in the

composite electrodes compared to that of single phase pyrochlore electrode. As shown in

figure 4-33, the three composite compositions showed better performance than the

pyrochlore cathode at each temperature, with best performance of 18.4 Qcm2 at 500 0C

and 0.32 ncm2 at 700 'C with a composite electrode containing 37.5 wt% ESB.

Arrhenius plots of the composite electrodes and single phase pyrochlore electrode is

shown in figure 4-34. Among the three compositions of the composite electrode, the

activation energy of the electrode reaction decreases with the addition of the ionic phase.

This trend is similar to those observed for other composite electrodes,53-55 where the

addition of ionic phase leads to improved ionic conductivity of the composite electrode

and decreases the activation energy. The proposed electrode reaction mechanism for the

bismuth ruthenate electrodes is shown in figure 4-35. The presence of the ESB in the

electrode will not only increase the concentration of TPBs and ionic conductivity of the








electrode, but possible also shortens the rate limiting surface diffusion path to the TPBs

and hence, improves the electrode performance.

The performance of the composite electrode is highly promising and compares well

with established cathode materials. Further optimization of the electrode morphology in

terms of relative ratio, particle size and spatial distribution of the two phases will realize

high performance cathodes for IT-SOFCs.

4.4 Conclusion

Bismuth ruthenate based electrodes were evaluated as a prospective cathode

material for IT-SOFCs based on ceria electrolytes. No reaction product was found

between bismuth ruthenate and GDC electrolyte after heat treating at 850 'C for 10 hours.

Bismuth ruthenate pyrochlore electrodes showed ASR values of 55.64 fcm2 at 500 C

and 1.45 Qcm2 at 700 *C in air. The surface diffusion of dissociatively adsorbed oxygen

at the electrode to TPBs is believed to be the rate limiting step. Doping with lower valent

cations on Bi3 site, to improve the ionic conductivity, was not found to be very effective

the improving the electrode performance. 5 mol% Ca and Ag doping were found to only

slightly improve the electrode performance over undoped bismuth ruthenate pyrochlores.

Rate limiting step for the oxygen reduction reaction in 5 mol% Ag doped bismuth

ruthenate electrode was found to be the surface diffusion of dissociatively adsorbed

oxygen, while for 5 mol% Ca doped system multiple rate limiting steps were observed in

the oxygen partial pressure range of study. Composite electrodes consisting of Bi2Ru207.3

and (Bi203)0.8(Er2O3)0.2 showed 3-4 times better performance than Bi2Ru207 pyrochlore

electrodes. Best performance of 18.4 ncm2 at 500 'C and 0.32 f~cm2 at 700 'C in air was

observed with a composition containing 37.5 wt% (Bi203)0 s(Er2O3)0.2. The performance

is very promising, and it is expected that further optimization in electrode composition






82


and microstructure will result in high performance cathodes for IT-SOFCs. Lastly, a need

was felt to develop new methodologies which could more effectively distinguish the

different rate limiting steps in the electrode impedance.































Figure 4-1. Phase diagram between Bi203 and RuO2 proposed by Prosychev et al. (from
ref. 50)



RuO2' Bi2Ru2O? \
T ,Bi2Ru207 Liquid
Scl
WO -
IL quid 7f '



*Ru02Liquid 203Luid


u00


I I I I I I I I I
2 Ii3Ru3011



Figure 4-2. Phase diagram between Bi203 and RuO2 proposed by Hrovat et al. (from ref.
51)
















....... a r(E) A (ORIGIN)
a I OA






88(a)
0 B

8(0) -









Figure 4-3. Cation sublattice for one-quarter unit cell of the pyrochlore structure. Two of
the 48f-anions and one of the 8b-sites are also shown for 34a (from ref. 8).










2000 Bi RUOo


1500


100



0
500 1 I

*!


20 30 40 50 60 70
20
Figure 4-4. XRD patterns for Bi2Ru2O7 after calcination at 900 C for 10 hours (a) before
and (b) after leaching.


1400
*Bi RuO
1200 12 20:

1000

800

0600 I ,
U I
400
[t *
200 *

0
20 30 40 50 60 70
20

Figure 4-5. XRD patterns for Bi2Ru2073 after calcination at 775 C for 10 hours (a)
before and (b) after leaching.




Full Text
ELECTROCHEMICAL STUDIES ON SELECTED OXIDES FOR INTERMEDIATE
TEMPERATURE - SOLID OXIDE FUEL CELLS
By
ABHISHEK JAISWAL
A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY
UNIVERSITY OF FLORIDA
2004

Copyright 2004
by
Abhishek Jaiswal

Dedicated to my parents.

ACKNOWLEDGMENTS
I would like to thank my advisor, Dr. Eric Wachsman, for his guidance,
encouragement, and support through the course of my graduate study. It has been a
highly enriching experience working with him. I would also like to thank him for
allowing me to visit various conferences to gain exposure in the field.
I would like to thank my committee members: Dr. Daryl Butt, Dr. David Norton,
Dr. Wolfgang Sigmund and Dr. Mark Orazem. Dr. Butt and Dr. Sigmund allowed me to
use their laboratory facilities for my research work.
I would like to acknowledge my group members, especially Jamie Rhodes and
Guojing Zhang, who made the time at UF a very memorable experience. I would like to
thank Dr. Hee-Sung Yoon, Dr. Keith Duncan and Dr. Suman Chatteijee for their advice
and help. I would also like to thank Wayne Acre for helping me in using the various
characterization facilities at MAIC.
Lastly, I would thank my parents and my brothers. They have been very kind to
bear long separations just to let me achieve my dreams. My mother has not been in good
health for the last three years and it has been a tough time for the whole family. I pray
that I can justify their sacrifices.
IV

TABLE OF CONTENTS
page
ACKNOWLEDGMENTS iv
LIST OF FIGURES vii
ABSTRACT xiv
CHAPTER
1 FUEL CELLS: A GREEN SOLUTION 1
1.1 Introduction 1
1.2 Operating Principle of SOFCs 4
1.3 Designing SOFCs 6
1.4 SOFC Structures 7
1.4.1 Tubular SOFCs 8
1.4.2 Planar SOFCs 10
1.4.3 New Designs: Metal Supported and Microtubular SOFCs 11
1.5 Conclusion 13
2 IONIC CONDUCTION IN SOLID ELECTROLYTES 19
2.1 Introduction 19
2.2 Fluorite Type Oxides 20
2.2.1 Structural Aspects 20
2.2.2 Electrical Aspects 20
2.2.3 Oxygen Ion Conductivity as a Function of Dopant Concentration and
Temperature 22
2.2.4 Grain Boundary Contribution to Total Conductivity 26
2.2.5 Oxygen Ion Conductivity as a Function of Time 27
2.3 Conclusion 28
3 DIRECT CURRENT BIAS STUDIES ON (Bi2O3)0.8(Er2O3)0.2 ELECTROLYTE
AND Ag-(Bi2O3)0 8(Er2O3)0 2 CERMET ELECTRODE 36
3.1 Introduction 36
3.2 Experimental 39
v

3.3 Results and Discussion 40
3.3.1 Processing 40
3.3.2 DC Bias and Impedance Spectroscopy: ESB Electrolyte 41
3.3.3 DSC Studies: ESB Electrolyte 43
3.3.4 DC Bias and Impedance Spectroscopy: Ag-(BÍ203)o.8(Er203)o.2Electrode45
3.4 Conclusion 48
4 BISMUTH RUTHENATE BASED CATHODES FOR IT-SOFC 66
4.1 Introduction 66
4.1.1 Undoped Bismuth Ruthenate 66
4.1.2 Doped Bismuth Ruthenates 68
4.1.3 Bismuth Ruthenate and Stabilized Bismuth Oxide Composites 70
4.2 Experimental Work 71
4.3 Results and Discussion 72
4.3.1 Processing 72
4.3.2 Impedance Spectroscopy Studies 73
4.3.2.1 Undoped BÍ2RU2O7 cathode 73
4.3.2.2 Doped BÍ2RU2O7 cathodes 75
4.3.2.3 Bi2Ru207.3-(BÍ203)o.8(Er203)o.2 composite cathodes 79
4.4 Conclusion 81
5 ANODE SUPPORTED THICK FILM CERIA ELECTROLYTE UNIT CELLS FOR
IT-SOFC Ill
5.1 Introduction Ill
5.2 Experimental 115
5.3 Results and Discussion 116
5.3.1 Initial Fabrication and Performance Results 116
5.3.2 Improving the Performance of Anode Supported Ceria Unit Cells 121
5.3.2.1 Reducing the co-sintering temperature of anode/electrolyte
bilayer 123
5.3.2.2 Optimization of LSCF-GDC cathode sintering temperature 124
5.3.2.3 Composite cathodes containing ESB as the electrolyte phase 129
5.4 Conclusion 131
6 CONCLUSION AND FUTURE WORK 165
LIST OF REFERENCES 170
BIOGRAPHICAL SKETCH 177
vi

LIST OF FIGURES
page
1-1 Fuel cell types with typical reactants 14
1-2 Schematic diagram of the current-voltage characteristics as well as the different
losses in a solid oxide fuel cell 14
1 -3 Placement of reactant energies relative of the edges of conduction and valence band
of the electrolyte in a thermodynamically stable electrochemical cell 15
1-4 Siemens Westinghouse tubular SOFC design 15
1 -5 Siemens Westinghouse flattened tubular SOFC design 16
1-6 Theoretical and actual performance of Siemens Westinghouse tubular and flattened
tubular (HPD) SOFC designs 16
1-7 Planar SOFC design 17
1-8 Performance of a planar SOFC design at 800 °C 17
1-9 Microstructure of an alloy supported SOFC based on YSZ electrolyte 18
1-10 Microstructure of an alloy supported SOFC based on doped cerium oxide
electrolyte 18
2-1 Oxygen ion conductivity of common solid electrolytes 30
2-2 The fluorite structure 31
2-3 Isothermal conductivity of some ceria solid solution at temperature close to 200
°C 31
2-4 Schematic Arrhenius plots of lattice and grain boundary conductivity for clean
Ceo.9Gdo.1O195, impure Ceo.9Gdo.1O1.95, and impure Ceo.gGdo^Oi.g 32
2-5 Conductivity of (BÍ203)i_x(Er203)x in air 33
2-6 Decay in isothermal oxygen ion conductivity of 20 mol% ESB on annealing below
the transition temperature 34
Vll

2-7 Time constant as function of cation radii and lattice parameter of 25 mol% cation
doped bismuth oxide 35
2-8 Time constant as function of erbia concentration and lattice parameter of erbia
stabilized bismuth oxide 35
3-1 XRD patterns of calcined 20 mol% erbia stabilized bismuth oxide powders derived
from amorphous citrate route 50
3-2 Cross-sectional micrographs of the ESB electrolyte and Ag-ESB electrode
interface 51
3-3 Arrhenius plot of oxygen ion conductivity in ESB electrolyte 52
3-4 Percentage change in resistance of the ESB electrolyte under no bias at 300 °C, 500
°C, and 625 °C 53
3-5 Initial voltage vs. time plots at 500 °C under different bias currents 54
3-6 Percentage change in resistance of the ESB electrolyte at 500 °C under different
bias currents 55
3-7 Percentage change in resistance of the ESB electrolyte at 625 °C under different
bias currents 56
3-8 Arrhenius plot of ASR (Qcm2) for Ag-ESB electrode 58
3-9 Impedance plots of Ag-ESB electrode at 500 °C under different bias currents 59
3-10 Ag-ESB electrode resistance vs bias current at 500 °C and 625 °C 60
3-11 Impedance plots of Ag-ESB electrode at 625 °C under different bias currents 61
3-12 Impedance plots of Ag-ESB electrode under 250 mA bias at 625 °C 62
3-13 Cross-sectional microstructures of the electrolyte/electrode interface after annealing
experiments at 625 °C under 250 mA bias for ~26 hours 63
3-14 Proposed electrode reaction mechanism for Ag-ESB electrodes 64
3-15 Phase diagram between Ag and Ag2Ü proposed by Assal et al 65
4-1 Phase diagram between BÍ2O3 and RuCft proposed by Prosychev et al 83
4-2 Phase diagram between BÍ2O3 and RuCb proposed by Hrovat et al 83
4-3 Cation sublattice for one-quarter unit cell of the pyrochlore structure 84
4-4 XRD patterns for BÍ2RU2O7 after calcination at 900 °C for 10 hours 85
viii

4-5 XRD patterns for BÍ2R.U2O7.3 after calcination at 775 °C for 10 hours 85
4-6 XRD pattern for BÍ2RU2O7 and GDC powder mixture after heat treatment at 850 °C
for 10 hours 86
4-7 XRD pattern for BÍ2RU2O7 3 and GDC powder mixture after heat treatment at 850
°C for 10 hours 86
4-8 Surface and cross-section micrographs of BÍ2RU2O7 electrode on GDC
electrolyte 87
4-9 Impedance plots of BÍ2RU2O7 electrode on GDC at 500 °C 88
4-10 Impedance plots of BÍ2RU2O7 electrode on GDC at 700 °C 89
4-11 Arrhenius plot of the BÍ2RU2O7 electrode ASR (Qcm2) in air 90
4-12 ln(ASR) vs. ln(po2) of BÍ2RU2O7 electrode at different temperatures with m in
parenthesis T(m) 90
4-13 Impedance plots of Ca doped BÍ2RU2O7 electrodes on GDC at 500 °C 91
4-14 Impedance plots of Ca doped BÍ2RU2O7 electrodes on GDC at 700 °C 92
4-15 Arrhenius plot of Ca doped BÍ2RU2O7 electrode ASR (Gem2) in air 93
4-16 Impedance plots of Ag doped BÍ2RU2O7 electrodes on GDC at 500 °C 94
4-17 Impedance plots of Ag doped BÍ2RU2O7 electrodes on GDC at 700 °C 95
4-18 Arrhenius plot of Ag doped BÍ2RU2O7 electrode ASR (f2cm ) in air 96
4-19 Arrhenius plot of undoped, 5 mol% Ca and Ag doped BÍ2RU2O7 electrode ASR
(Gem2) in air 97
4-20 Arrhenius plot of undoped, 10 mol% Ca, Ag and Sr doped BÍ2RU2O7 electrode ASR
(Qcm2) in air 97
4-21 Impedance plots of 5 mol% Ca doped BÍ2RU2O7 electrodes on GDC at 500 °C 98
4-22 Impedance plots of 5 mol% Ca doped Bi2Ru207 electrodes on GDC at 700 °C 99
4-23 Impedance plots of 5 mol% Ag doped BÍ2RU2O7 electrodes on GDC at 500 °C. ..100
4-24 Impedance plots of 5 mol% Ag doped BÍ2RU2O7 electrodes on GDC at 700 °C. ..101
4-25 ln(ASR) vs. ln(po2) of 5 mol% Ca doped BÍ2RU2O7 electrode at different
temperatures with m in parenthesis T(m) 102
IX

4-26 ln(ASR) vs. ln(po2) of 5 mol% Ag doped BÍ2R112O7 electrode at different
temperatures with m in parenthesis T(m) 102
4-27 Impedance plots of 5 mol% Ca doped BÍ2RU2O7 electrodes on GDC at 500 °C....103
4-28 Impedance plots of 5 mol% Ca doped BÍ2RU2O7 electrodes on GDC at 700 °C....104
4-29 DSC plots of powder mixture of BÍ2RU2O7 3 with monoclinic 01-BÍ2O3 and fee
fluorite (Bi203)o 8(Er2O3)0.2 105
4-30 SEM surface micrographs of bismuth ruthenate-bismuth oxide composite
electrode 106
4-31 Impedance plots of BÍ2Ru207.3-(BÍ203)o.8(Er203)o.2 composite electrodes on GDC at
500 °C 107
4-32 Impedance plots of BÍ2Ru207.3-(BÍ203)o.8(Er203)o.2 composite electrodes on GDC at
500 °C 108
4-33 ASR (Gem ) of BÍ2RU2O7 3-(BÍ203)o.8(Er203)o.2 composite electrodes in comparison
to BÍ2RU2O7 electrode 109
4-34 Arrhenius plot of BÍ2RU2O7 3-(BÍ203)o.8(Er203)o.2 composite electrode ASR (Qcm2)
in comparison with BÍ2RU2O7 electrode 109
4-35 Proposed electrode reaction mechanism for bismuth ruthenate electrodes 110
5-1 Representative microstructures of the GDC film under different pre-sintering
conditions for a 1600 °C final sintering 133
5-2 Representative microstructures of the GDC film under different pre-sintering
conditions for a 1650 °C final sintering 134
5-3 Porosity in the GDC film as a function of the pre-sintering and final sintering
temperature 135
5-4 Cross-sectional microstructures of the GDC film under different pre-sintering
conditions for a 1600 °C final sintering 136
5-5 Surface microstructure of the GDC film after 850 °C pre-sintering and 1400 °C
final sintering 136
5-6 Surface and cross-sectional microstructure of reduced Ni-GDC anode 137
5-7 Cross-sectional microstructure of LSCF cathode sintered at 750 °C & 900°C 137
5-8 Cross-sectional microstructure of tested unit cell 137
x

5-9 Open circuit potential of the cell as function of temperature 138
5-10 Average oxygen ion transference number of GDC electrolyte as function of
temperature 138
5-11 I-V characteristics of the cell with air at the cathode 139
5-12 Power density of the cell with air at the cathode 139
5-13 I-V characteristics of the cell with oxygen at the cathode 140
5-14 Power density of the cell with oxygen at the cathode 140
5-15 ln(ASR/T) vs 1/T of the cell 141
5-16 I-V characteristics of the cell with bilayer sintering temperature of 1350 °C 142
5-17 Power density of the cell with bilayer sintering temperature of 1350 °C 142
5-18 I-V characteristics at 600 °C of the cell with bilayer sintering temperature of 1350
°C before and after thermal cycling 143
5-19 I-V characteristics at 650 °C of the cell with bilayer sintering temperature of 1350
°C before and after thermal cycling 143
5-20 Cross-sectional and surface microstructures of the tested unit cell with bilayer
sintering temperature of 1350 °C 144
5-21 I-V characteristics of the cell with bilayer sintering temperature of 1450 °C 145
5-22 Power density of the cell with bilayer sintering temperature of 1450 °C 145
5-23 I-V characteristics at 650 °C of the cell with bilayer sintering temperature of 1450
°C before and after thermal cycling 146
5-24 Cross-sectional microstructure of the tested unit cell with bilayer sintering
temperature of 1450 °C 146
5-25 I-V characteristics of the cell with bilayer sintering temperature of 1550 °C 147
5-26 Power density of the cell with bilayer sintering temperature of 1550 °C 147
5-27 Cross-sectional microstructure of the tested unit cell with bilayer sintering
temperature of 1550 °C 148
5-28 Cross-sectional microstructures of the tested unit cell with cathode sintering
temperature between 850-1350 °C, and surface microstructure showing cathode
layer sintered at 1250 °C and current collecting layer sintered at 1000 °C 149
xi

5-29 I-V characteristics of the cell with cathode sintering temperature of 850 °C 150
5-30 Power density of the cell with cathode sintering temperature of 850 °C 150
5-31 I-V characteristics of the cell with cathode sintering temperature of 1000 °C 151
5-32 Power density of the cell with cathode sintering temperature of 1000 °C 151
5-33 I-V characteristics at 650 °C of the cell with cathode sintering temperature of 850
°C before and after thermal cycling 152
5-34 I-V characteristics at 750 °C of the cell with cathode sintering temperature of 850
°C before and after thermal cycling 152
5-35 I-V characteristics at 650 °C of the cell with cathode sintering temperature of 1000
°C before and after thermal cycling 153
5-36 I-V characteristics at 750 °C of the cell with cathode sintering temperature of 1000
°C before and after thermal cycling 153
5-37 I-V characteristics of the cell with cathode sintering temperature of 1150 °C 154
5-38 Power density of the cell with cathode sintering temperature of 1150 °C 154
5-39 I-V characteristics of the cell with cathode sintering temperature of 1250 °C 155
5-40 Power density of the cell with cathode sintering temperature of 1250 °C 155
5-41 I-V characteristics of the cell with cathode sintering temperature of 1350 °C 156
5-42 Power density of the cell with cathode sintering temperature of 1350 °C 156
5-43 I-V characteristics at 650 °C of the cell with cathode sintering temperature of 1150
°C before and after thermal cycling 157
5-44 I-V characteristics at 650 °C of the cell with cathode sintering temperature of 1250
°C before and after thermal cycling 157
5-45 Maximum power density as a function of cathode sintering temperature 158
5-46 I-V characteristics at 500 °C of the cell with cathode sintering temperature of 1350
°C, with contributions from ohmic and non-ohmic polarization 159
5-47 I-V characteristics at 700 °C of the cell with cathode sintering temperature of 1350
°C, with contributions from ohmic and non-ohmic polarization 159
5-48 Non-ohmic polarization of the cells with cathode sintering temperature of 1250 °C
and 1350 °C 160
Xll

5-49 Ohmic polarization of the cells with cathode sintering temperature of 1250 °C and
1350 °C 160
5-50 Arrhenius plot of ASR of the ohmic polarization of the cells with cathode sintering
temperature of 1250 °C and 1350 °C, in comparison to ~30 pm bulk GDC
electrolyte 161
5-51 XRD patterns for LSCF-ESB 162
5-52 XRD patterns for LSCuF-ESB 162
5-53 I-V characteristics of the cell with LSCuF-ESB cathode sintered at 750 °C, in
comparison with the cell with LSCF-GDC cathode sintered at 850 °C 163
5-54 Power density of the cell with LSCuF-ESB cathode sintered at 750 °C, in
comparison with the cell with LSCF-GDC cathode sintered at 850 °C 163
5-55 OCP of the cell with Ag-ESB cathode as a function of time 164
5-56 Cross-sectional microstructures of the tested unit cell with Ag-ESB cathode in
back-scatted electron mode 164
Xlll

Abstract of Dissertation Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy
ELECTROCHEMICAL STUDIES ON SELECTED OXIDES FOR INTERMEDIATE
TEMPERATURE - SOLID OXIDE FUEL CELLS
By
Abhishek Jaiswal
December 2004
Chair: Eric D. Wachsman
Major Department: Materials Science and Engineering
Fuel cell technology holds the promise to change the way power is generated,
transmitted, and utilized in our increasing demanding lifestyles. State of the art solid
oxide fuel cells (SOFCs) utilize an all ceramic design and operate at 750-1000 °C. Lower
operating temperatures will significantly improve the economics of power generation
using SOFCs. The aim of this dissertation was to evaluate and develop component
materials for SOFCs, which could work efficiently at temperatures between 500-750 °C.
Erbia stabilized bismuth oxide (ESB) shows one of the highest oxygen ion
conductivity among all solid electrolytes. However, due to positional and occupational
ordering the conductivity decays below the transition temperature (-600 °C). The effect
of direct current bias on the ordering phenomenon in ESB was studied using symmetrical
cells with Ag-ESB electrodes. At 500 °C, the endotherm, related to reverse transition, is
enhanced by the applied bias at short time but with negligible change in conductivity
decay. It is proposed that the conductivity decay with anneal time is related more to the
xiv

positional ordering than occupational ordering. Ag-ESB electrodes showed good
performance, though were unstable under high currents at 625 °C due to Ag migration
with oxygen flux.
Novel bismuth ruthenate based cathodes were evaluated using impedance
spectroscopy with symmetric cells on gadolinium doped ceria (GDC) electrolytes.
Undoped bismuth ruthenate electrode showed area specific resistance (ASR) of 55.64
Gem2 at 500 °C and 1.45 Gem2 at 700 °C in air. Doping with similar size Ca2+, Ag+, or
Sr2+ on Bi ,+ site did not improve the electrode performance significantly, while bismuth
ruthenate-ESB composites showed 3-4 times lower electrode ASR. Bismuth ruthenate-
ESB (62.5:37.5 wt%) composite showed the best performance of 18.4 Gem2 at 500 °C
and 0.32 Gem2 at 700 °C in air. Addition of the ESB phase is believed to reduce the rate
limiting surface diffusion in oxygen reduction reaction.
Anode supported thick film GDC electrolyte unit cells were developed for IT-
SOFCs. A colloidal deposition technique was used to fabricate dense, thick GDC
electrolyte films on porous Ni-GDC anode supports. Pre-sintering temperature of the
anode and final sintering temperature of the anode/electrolyte bilayer were found to be
the primary parameters determining the density of the film. The sintering temperature of
LSCF-GDC (70:30 wt%) composite cathode was optimized to 1250-1350 °C, which
resulted in a maximum power density of 0.338 W/cm2 at 0.771 A/cm2, 700 °C. Current
interrupt showed that apart from the electrolyte layer, the ohmic polarization across the
cell has significant contributions from the electrodes.
xv

CHAPTER 1
FUEL CELLS: A GREEN SOLUTION
1.1 Introduction
A fuel cell is an electrochemical device which converts the chemical energy of a
fuel and an oxidant into electrical energy. The process essentially involves an invariant
electrode-electrolyte system. In essence, fuel cells are similar to batteries although with
much longer lifetime; as in a fuel cell, the fuel is replenished and the product is
discharged continuously. The conversion is direct without the need of intermediate
conversion into heat and mechanical energy, as in the case of conventional
turbine/generator systems. The energy conversion in fuel cells is not limited by the
Carnot cycle and efficiency up to 50 % is achievable.1 As fuel cells do not involve any
combustion process, there is no formation of pollutants including NOx, SOx,
hydrocarbons and particulates. Also, fuel cells do not consist of any moving parts and
hence do not generate any noise pollution and require low maintenance. Due to these
advantages over other energy generators, fuel cells seek application in stationary as well
as in tractionary applications.
The concept of producing electrical energy from a simple electrochemical cell was
first shown by Grove in 1839. ' It took however another 120 years before Bacon4
assembled a fuel cell stack exhibiting useful power densities, which was later modified
by Pratt & Whitney for the onboard power sources for NASA Apollo space missions. At
present different types of fuel cells are under development, which are in general
identified by the type of electrolyte used: phosphoric acid fuel cells (PAFCs), proton
1

2
exchange membrane fuel cells (PEMFCs), molten carbonate fuel cells (MCFCs), solid
oxide fuel cells (SOFCs), alkaline fuel cells (AFCs). The different types of fuel cells with
typical reactants and operating temperatures are shown in figure 1-1. Low temperature
PAFCs, PEMFCs and AFCs require external reforming of the fuel, whereas high
temperature SOFCs and MCFCs can internally reform and work with hydrocarbon fuels
directly.’ Also, high temperature operation removes the need for expensive noble metals
to catalyze the electrode reactions. These factors result in lower cost, lesser complexity
and higher energy conversion efficiency of the high temperature systems. Compared to
molten carbonate fuel cells (MCFCs), solid oxide fuel cells (SOFCs) are advantageous as
the solid oxide electrolyte is usually very stable and does not show any migration
problems under the operating conditions.
Among all types of fuel cells, proton exchange membrane fuel cells (PEMFCs) and
solid oxide fuel cells (SOFCs) are considered to be the most advanced and closest to
wide-scale commercialization. However, there remain significant technological and
engineering challenges to be solved. When running on hydrocarbon fuels in addition to
external reforming, PEMFCs also require CO removal from the fuel feed as they are
susceptible to CO poisoning which results in lower conversion efficiencies. On the other
hand, SOFCs run at high temperatures and ideally can internally reform any hydrocarbon
fuel with high efficiencies without the need of expensive catalysts. In addition, the high-
quality waste heat from the SOFCs can be utilized in cogeneration to further improve the
overall efficiency of the system. The state of the art SOFCs operate at temperatures
between 750-1000 °C using yttria stabilized zirconia (YSZ) as the electrolyte material,
La(Sr)Mn03-YSZ composite as the cathode material, Ni-YSZ ceramic metal (cermet)

3
composite as the anode material and LaCr03 as the interconnect material. Currently, there
are two configurations seeking commercialization: tubular and planar. Among the two,
the seal-less tubular design of Seimens Westinghouse is the most advanced with electrical
efficiency up to 46%.6 With combined heat and power cycle, efficiency up to 85% and in
pressurized SOFC-gas turbine hybrid system, efficiency greater than 55% are achieved.
However, the high operating temperatures of SOFCs put considerable limitations
on the choice of materials for the various components and also on the lifetime of the cell.
Over the years research in the industry and academia has resulted in the development of
exotic materials and fabrication techniques such that the cell components are able to
perform exceptionally well under extreme conditions, to withstand thermal mismatch, to
be microstructurally stable and to counter interactions between adjoining components.
However this has led to high material and fabrication costs, making SOFCs
uncompetitive with existing power generation technologies. Decreasing the operating
temperatures will considerably improve the economics of power generation using SOFCs
by enabling the use of cheap ferritic stainless steel alloys as the interconnect material
instead of expensive ceramics, cheaper balance of plant and insulation along with
increased lifetime.7 Lower operating temperatures will also result in faster startup which
is critical in certain applications. All these factors have led to a global drive towards
reducing the operating temperature of SOFCs from 750-1000 °C to intermediate
temperatures of 500-750 °C. However, efficient operation at intermediate temperatures
will require new electrolyte materials with higher conductivity and new electrode
materials with better catalytic activity at lower temperatures.

4
1.2 Operating Principle of SOFCs
The working of a fuel cell could be understood using a concentration cell model.
Let us assume that the cell has different oxygen partial pressure at the cathode and the
anode. The electrolyte is a gas impermeable solid with Pt electrodes attached to it. The
model can be represented as follows:
(anode) O2, Pt / Solid Electrolyte / Pt, O2 (cathode) 1-1
The cathodic and anodic reactions are as follows:
2e (pt, cj + 1/2 (P(g, c) PP (se, c)
cathode
1-2
PP (se, a) ~ 2e (pt_ a) 1/2 02(g, a)
anode
1-3
where the subscripts represent the site of concerned species: a, c, g, se represent anode,
cathode, gas and solid electrolyte, respectively. The difference in the chemical potential
of O2 (p0 = RTln p()i) at the anode and the cathode results in the generation of
electrical potential difference (A(p) across the cell.
A = i^itoiJdPo2 =^\t0i-(Po2J)RTd(lnp02) 1-4
°,(Po2’T)
1-5
Apart from the oxygen partial pressure difference, the electrolyte transport
properties in terms of the oxygen ion transference number (t 2-) determines the electrical
potential difference generated. The oxygen ion transference number is given by the ratio
of oxygen ion conductivity to the total conductivity and at a constant temperature is a
function of oxygen partial pressure. The oxygen partial pressure varies along the

5
electrolyte thickness and the gradient (<5 pQ^ / 6x) in turn depends on the transport
properties o¡ of the electrolyte.
In an actual fuel cell, fuel (H2, CO, hydrocarbons) at the anode develops a low
oxygen partial pressure and air/oxygen at the cathode develops a high oxygen partial
pressure. Electrical power is generated when an external load is attached to the cell, and
useful power is achieved by stacking cells through interconnects in series and parallel.
The efficiency of a cell is lower than theoretical because of the irreversible ohmic and
non-ohmic polarization losses across the cell, as shown in figure 1-2. The irreversible
losses appear as heat, and hence the thermal management of the fuel cell stacks is an
important design consideration.
V(I) = E°- r/ohmic - (rj act VJcone)anode ~ (tfact cone)cathode \~6
The bulk of the ohmic polarization comes from the solid electrolyte. High oxygen
ion conductivity, tQ2- equal to 1 and small thickness are the desired characteristics of the
solid electrolyte. Non-ohmic polarizations from the electrodes consist of activation and
concentration polarizations. Activation polarization is related to the kinetics of chemical
and charge transfer reactions at the gas/solid interface. High catalytic activity for oxygen
reduction at the cathode and for fuel oxidation at the anode is required. Fast oxygen ion
transfer across the electrolyte/gas interface and electrolyte/electrode interface requires
good chemical and mechanical compatibility of the electrolyte/electrode interface.
Formation of blocking interface phases must be avoided. Concentration polarization is
related to the diffusion of gaseous species through porous electrodes to reach the reaction
sites. Microstructural engineering is required to avoid the diffusion limited regions in the
electrodes when operating at high currents.

6
1.3 Designing SOFCs
Ohmic and activation polarizations are thermally activated processes, resulting in
improved efficiency of the SOFCs with increasing temperatures. However, high
operating temperatures could result in chemical and mechanical compatibility issues
between adjoining components. Chemical and mechanical compatibility between
adjoining components are of paramount importance as the SOFC is expected to perform
for 40-100 thousand hours at elevated temperatures and for numerous thermal cycles
from the room temperature to the operating temperature. Mechanical compatibility
between neighboring components against thermal cycling requires close matching of the
thermal expansion coefficient (TEC). At times, SOFC designers have to compromise the
electrochemical performance in order to ensure a good thermal expansion match between
components. In general, it is achieved by introducing external dopants in the host
material to vary the TEC.
Processing of electrodes should ensure a good adhesion at the electrode/electrolyte
interface and a stable electrode microstructure over the period of operation. In general,
the processing temperatures are higher than the operating temperatures to ensure that the
electrode microstructure does not coarsen over the period of operation with resultant loss
in active sites for the electrode reaction and in consequent decay in cell performance.
However, sintering at high temperatures could lead to problems in certain electrode¬
electrolyte combinations, as in the case of La(Sr)Mn03-YSZ where tertiary La2Zr207
phase appears at the interface. La2Zr207 is an oxygen ion blocking phase and its presence
at the interface degrades the cell performance considerably.
Chemical stability of all the components in their respective working environments
is also a requirement. The electrolyte faces both the oxidizing and the reducing

7
atmospheres and as described by Goodenough,8 its thermodynamic stability requires that
the bottom of the electrolyte conduction band is above the highest occupied molecular
orbital (HOMO) of the reductant and the top of the electrolyte valence band is below the
lowest unoccupied molecular orbital (LUMO) of the oxidant as shown in figure 1-3. The
Fermi energies of metallic electrodes should also lie in the energy gap Eg of the
electrolyte, with that of the anode rising to the HOMO of the reductant and that of the
cathode falling to the LUMO of the oxidant under operating conditions.
Although Ni as the metal component in the cermet anode satisfies the major
requirements for catalytic oxidation of H2 and CO fuels, its use for direct oxidation of
hydrocarbons encourages carbon deposition on the surface which with time reduces the
catalytic activity of anode. Cu has been shown to be more effective with hydrocarbon
fuels and it is likely that in future, Cu will replace Ni in SOFCs having internal reforming
capabilities.9 Due to the low melting point of Cu, there are still technical hurdles with Cu
based anodes in terms of fabrication and operating temperatures. Ni anodes can also be
poisoned in sulphur containing fuels due to formation of thermodynamically stable NiS.
Ceramic anodes based on doped SrTiC>3 are under development and show significant
promise in terms of carbon and sulphur tolerance.10
1.4 SOFC Structures
Based on the component that supports a unit cell, SOFCs can be divided into two
basic types: electrode supported and electrolyte supported. Chan et al.11 have modeled the
sensitivity of the cell voltage to the thickness of each component. The model showed that
with the same thickness of each component, the cell voltage is most sensitive to the
thickness of the electrolyte followed by the cathode and then by the anode. Therefore,
anode or cathode supported cells should be preferred over electrolyte supported cells.

8
Compared to the cathode supported cells, the model showed that the anode supported
cells are superior in terms of operating current density range and available power density.
Apart from the fuel cell performance other factors like target application, fabrication and
operating costs, fuel choice and utilization, operating temperature, mechanical strength
and expected lifetime are also important considerations. Electrode supported design with
a thick film electrolyte perform better, but it also have issues in terms of strength, cost
and performance. Porous electrodes supports are mechanically weaker in comparison and
thermal cycling puts considerable amount of thermal and mechanical stresses on them.
Anode supported cells with Ni based cermet anode can face catastrophic failure during
loss of fuel supply or under high fuel utilizations, as oxidation of Ni to NiO results in
large volume expansion which generates cracks in the anode and in the thick electrolyte
film. Further, depositing dense (pin-hole free) thick electrolyte layers on porous supports
reliably, inexpensively, with good adhesion and without undue sintering of the porous
support over large areas still remains a technical challenge. Keeping these factors in
mind, different manufacturers are pursuing different SOFC design and manufacturing
routes.
1.4.1 Tubular SOFCs
Among all SOFC designs, Siemens Westinghouse’s seal-less cathode supported
tubular design, shown in figure 1-4, is considered the most advanced. They are targeted
for the stationary power generation market and have been tested successfully for
hundreds of thousands of hours. The La(Sr)Mn03 cathode support tube (2.2 cm diameter)
is extruded and sintered, over which ~40 pm thick, dense YSZ electrolyte is deposited
using electrochemical vapor deposition (EVD). To connect unit cells together, ~0.9cm
wide and ~85 pm thick strip of LaCr03 based interconnect material is deposited using

9
plasma spraying. Finally, 100-150 pm thick Ni-YSZ cermet anode is deposited either by
Ni slurry coating followed by electrochemical vapor deposition of YSZ or by sintering of
a Ni-YSZ slurry coating. Single cells are then joined together using Ni felt in series and
parallel to achieve desired power ratings. Air flow is inside the cathode tube while fuel
flow is outside the cell.6
The advantage of the tubular geometry is in its simplicity. It avoids the use of seals
which at high temperatures are always an issue in terms of reactivity with other
components and failure during thermal cycling. The tubular structure is more rugged than
the planar structures which are discussed later. However, the cost of the tubular design is
still prohibitive for mass commercialization against competitive technologies. In the
tubular design, over 90% of the weight of the single cell comes from the cathode tubular
support and therefore cheaper raw materials for the cathode can bring in significant cost
reduction. Fabrication routes for the dense electrolyte and interconnect are also fairly
expensive and cheaper fabrication techniques like colloidal/electrophoretic deposition are
being evaluated for further cost reduction.6
Tubular single cells show a power density of-0.25 W/cm2 at 1000 °C (fuel: 89 %
H2 /11 % H2O, 85 % fuel utilization; oxidant: air), which is significantly lower than that
of the planar cells. Lower power density is primarily due to the long electronic current
path along the circumference of the cathode tube. Increasing the thickness of the tube
reduces the electronic resistance but also adds to the cathodic polarization as the oxygen
path from the gas phase is radially towards the electrode/electrolyte interface. In the
planar design, the current paths are much shorter being perpendicular to the thickness of
the layers and power density greater than 1 W/cm2 at 800 °C are in general achieved. A

10
new design combining the benefits of both the tubular and planar design is being
investigated by Siemens Westinghouse consisting of a flattened tube incorporating ribs in
the cathode structure as shown in figure 1-5. The new design lowers the internal
resistance of the cell with a shortened electronic current path along the ribs in the
cathode. Moreover, reduced resistance allows the use of thinner cathodes with lower
cathode polarization. The new flattened tubular geometry shows higher power density
compared to the tubular geometry as shown in figure 1 -6.
1.4.2 Planar SOFCs
Generic design of a planar SOFC is shown in figure 1-7. The assembly consists of
thin flat components which are fabricated using low-cost ceramic techniques like tape
casting, colloidal deposition and screen printing. Different organizations are focusing on
different variations of the planar design and use different manufacturing processes to
fabricate planar fuel cell stacks.
At present electrolyte supported, cathode supported and anode supported designs
are under development. Electrolyte supported cells with 50-100 pm thick YSZ electrolyte
have high electrolyte resistance and therefore are suitable for operation at temperatures
-1000 °C. The operating temperatures can be reduced by at least 200 °C by reducing the
thickness of the YSZ electrolyte to 5-20 pm in electrode supported planar designs. Ni-
YSZ anode supported designs are preferred over cathode supported designs in terms of
better electrochemical performance, better thermal and electrical conductivity, and
minimal interaction with the YSZ electrolyte during the sintering process. Although as
mentioned before, there remains concern over Ni oxidation in the case of loss of fuel
supply. Also in terms of cost, it has been noted that the anode supported design has more
electrolyte content than the electrolyte supported design itself.

11
For the fabrication of the anode supported planar SOFC, the electrolyte layer is
deposited on a tape cast anode by using either screen printing, colloidal deposition or
electrophoretic deposition followed by the co-firing of the bilayer. Some manufactures
also use tape calendaring or lamination to form the anode/electrolyte bilayer. Finally, the
cathode is screen printed and sintered to complete the unit cell. Anode supported planar
design shows significantly higher performance compared to the tubular design and at 800
°C, planar single cells consisting of ~10 pm YSZ layer have shown 1.8 W/cm2 (fuel: 97
% H2 / 3 % H2O, low fuel utilization; oxidant: air), as shown in figure 1-8. Moreover,
reduced operating temperatures allow the possibility of using metal interconnects which
along with cheaper fabrication routes for the cell can bring down the cost significantly.
Successful development of low-cost and oxidation resistant metallic alloy interconnects
and long lasting seals for the separation of the fuel and the oxidant atmospheres will
enable the commercialization of planar SOFCs.
1.4.3 New Designs: Metal Supported and Microtubular SOFCs
Although the anode supported planar design show significantly higher performance
than the tubular design, they leave a lot to be desired in terms of mechanical strength,
start-up time and sealing. This has led to the development of metal supported and micro¬
tubular SOFC designs.
In the metal supported SOFC design, a high temperature metal alloy supports the
electrode/electrolyte assembly as shown in figure 1-9 and 1-10. The assembly is
fabricated using the same low-cost techniques used to fabricate the anode supported
planar design. The metal support results in lower cost, higher mechanical strength and
higher thermal shock resistance of the structure which allows for rapid start-up.1213 The
metal support also increases the electronic and thermal conductivity resulting in lower

12
losses in current collection and better thermal management of the stack. Moreover, metal
12 13
support allows welding, brazing and other joining techniques for sealing unit cells. ’
However, there are a number of technical hurdles with the metal support design. These
include the fabrication temperatures and atmosphere for the assembly which is limited by
the melting/oxidation temperatures of the metal alloy support. For example, if the metal
alloy support is made of stainless steel then the assembly can not be readily fired above
1000 °C, which might be inadequate for the sintering of some components resulting in
lower performance of the cell. Stability of the electrode/alloy interface is also a concern.
Alloys with Cr2Ü3 oxide scales in moist air atmosphere form volatile Cr02(0H)2 species,
which can deposit at the cathode surface to act as blocking species. On the anode side,
high sintering temperatures result in interdiffusion between Ni and Fe-Cr alloys with
resultant coarsening of the anode microstructure due to the formation of CT2O3 oxide
scales. Further, elastic modulus mismatch between the alloy support and the ceramic
components can result in high stresses.
The microtubular design as the name suggests is similar to the Siemens
Westinghouse tubular design although with a much smaller diameter with either
electrolyte or anode supported geometry. The high length to diameter ratio provides for
its high strength and extraordinary thermal shock resistance. These tubes can withstand
temperatures gradient of 1000 °C at one end to room temperature at the other end.
Therefore, the micro-tubular design can combine the benefits of a seal-less design with
high mechanical strength and rapid start-up. Low-cost ceramic techniques are being used
to fabricate in order to keep the costs down. However, the design is plagued with the
same performance issue as that of Siemens Westinghouse’s tubular design: long current

13
collection paths. In fact the present microtubular design does not incorporate
interconnects and therefore the current collection path is actually longer.
Both metal supported and microtubular designs hold a lot of promise and the
solution of the above described technical hurdles will determine whether they remain
competitive with other designs.
1.5 Conclusion
SOFCs are efficient and environmental friendly energy generators. In this chapter,
the operating principle, the components and their requirements, and the various designs
of SOFCs were introduced. Tremendous progress has been made over the decades in the
development of materials, fabrication techniques and fuel cell design which enable a
stable performance over extended periods of time. However, the cost of SOFC system is
still prohibitive for mass commercialization against competitive technologies. New cost
effective approaches are being actively evaluated both in terms of fabrication techniques
and SOFC design. Decreasing the operating temperatures of SOFCs to intermediate
temperature range of 500-750 °C will open a wide range of options for the different
components, which can help in further reducing the cost. As the performance of SOFC
deceases with temperature, the challenge is to develop new materials which can perform
well at intermediate temperatures.

14
Internal
Reforming
H2, CO
External
Reforming
Hj, C02
External
Reforming
h2, co2
(CO removal)
>
>
h2 â–º
ANODE
CATHODE
1
L
ELECTROLYTE
J
L
h2o
SOFC (500-1OOOC)
co2
< o2
f~~
h2o
co2
MCFC (650C)
< co32
PAFC (200C)
m —►
h2o
PEMFC (80C)
H* â–º
h2o
h2o
AFC (70C)
< OH-
02 (air)
02(air)
CO2
02 (air)
02 (air)
02 (air)
(C02 removal)
Figure 1-1. Fuel cell types with typical reactants (from ref. 5).
J (Amp/cm2)
'H Ohmic
'H Cathode
Anode
Figure 1-2. Schematic diagram of the current-voltage characteristics as well as the
different losses in a solid oxide fuel cell.

15
Ec
t
V
oc
LUMO
Oxidant
Figure 1-3. Placement of reactant energies relative of the edges of conduction and
valence band of the electrolyte in a thermodynamically stable electrochemical
cell (from ref. 8).
Figure 1-4. Siemens Westinghouse tubular SOFC design (from ref. 6).

16
Figure 1-5. Siemens Westinghouse flattened tubular SOFC design (from ref. 6).
n
E
o
5
«
E
3
O
>
2
S2
i
o
Q.
0.50
* Operating Temperature: 1000X
“I—'—i—«—i—•—I—»—r
0.45
- Oxidant: Air, 6.0 Stoichs
. Fuel: 89% H2 + 11 % H20. 85% Fuel Utilization T â–¼ t
0.40
â–¼
-
T
*
0.35
â–¼
“
0.30
â–¼
VVVVV
“
0.25
â–  â–  a
-
0.20
• _ ^ 1
V 8°°°
â– 
â– 
V O n
—
v _0«°
Theoretical Performance
0.15
m
a
â– 
>
K>
t HPD-SOFC
_
V o
V *
â–  22 cm OD Cylindrical
â– 
0.10
-
Actual Performance
-
v o
v HPD-SOFC
â– 
0.05
o 22 cm OD Cylindrical
0 00
' ■ > 1 1 1 1—4 1 1—
_1 i 1 . 1 â–  1 i_A
0 100 200 300 400 500 600 700 800 900 1000
Current Density (mA/cm*)
Figure 1-6. Theoretical and actual performance of Siemens Westinghouse tubular and
flattened tubular (HPD) SOFC designs (from ref. 6).

17
Anode
Interconnection
(Bipolar Plate)
Cathode
Electrolyte
Anode
Figure 1-7. Planar SOFC design (from ref. 14).
Current Density (mA cm2)
Figure 1-8. Performance of a planar SOFC design at 800 °C (from ref. 15).
Power Density (mW/cm2)

18
YS2
Ni/YSZ
Fe/Cr alloy
Figure 1-9. Microstructure of an alloy supported SOFC based on YSZ electrolyte (from
ref. 121.
Contact Layer
Cathode
Electrolyte
Anode
Steel Substrate
Figure 1-10. Microstructure of an alloy supported SOFC based on doped cerium oxide
electrolyte (from ref. 13).

CHAPTER 2
IONIC CONDUCTION IN SOLID ELECTROLYTES
2.1 Introduction
Oxygen ion conducting solid electrolytes are the backbone of solid oxide fuel cells
(SOFCs), allowing selective transport of oxygen ions for electrochemical oxidation of
fuel to generate electrical power. The ionic conduction is possible due to the presence of
oxygen deficiency in the solid electrolyte. In general, the oxygen deficiency is introduced
by doping the host electrolyte with lower valent cations. The crystal structure of the host
material should be stable enough so that the introduction of the dopant ions does not
destroy the structure. Further, to make the activation energy for the migration of the
oxygen ions low, the electrolyte should have a relatively open crystal structure. The
saddle point critical radius in the lattice should be large enough for the passage of the
oxygen ions without creating lattice disturbances.
The electrolyte materials for SOFCs are mostly oxides having fluorite or perovskite
structure because of the inherent looseness in the structures and the ability to accept wide
range of dopants. Oxygen ion conductivity of common solid oxide electrolytes is shown
in figure 2-1. Better conductivity values are achieved with dopants having similar ionic
radii to the cation being substituted. The electrolyte in the fuel cell is exposed to both the
oxidant and the fuel atmospheres, and hence, it should be thermodynamically stable
under oxidizing and reducing conditions. To avoid the generation of electronic
conductivity, the electrolyte should have a large band gap and the dopants introduced into
the lattice should not exhibit multiple oxidation states.
19

20
As the oxygen ion conduction is a thermally activated process, the performance of
the electrolyte improves with temperature. However, there are limits in terms of operating
temperatures due to material compatibility issues with other components and operational
viability over extended periods of time. Reduced operating temperatures will open up
wide range of materials and significantly improve the economics of power generation
using SOFCs. However, operation of SOFCs at intermediate temperatures requires better
performance electrolyte and electrode materials.
2.2 Fluorite Type Oxides
2.2.1 Structural Aspects
The fluorite type unit cell is shown in figure 2-2. CeÜ2 and ThÜ2 exhibit the cubic
fluorite structure from room temperature up to their melting points. The packing in the
fluorite structure is far from closed packed and there are cubic 8-coordinated interstices at
the center of the unit cell. This loose structure provides for the possibility of achieving
unusually wide range of the solid solution with alkaline earth and rare earth oxides such
as CaO and Y2O3.14 Generally, when the cation size of the host and the guest are almost
equal the solid solution is easily formed. In the case of ZrC>2 and BÍ2O3, the high
temperature cubic fluorite structure could be stabilized at lower temperatures by forming
a solid solution and hence, the terms stabilized zirconia and stabilized bismuth oxide.
2.2.2 Electrical Aspects
To maintain the electrical neutrality of the solid solution, ionic and/or electronic
defects are generated in the lattice. In the case of fluorite systems, anti-frenkel defect
pairs (oxygen vacancy Vo " and oxygen interstitial Ot,r) are the dominant ionic defects. On
doping with lower valent cations, oxygen vacancies with effective charge of +2 are
generated to counter the effective negative charge of the dopant species. For example in

21
the case of CaO doped Ce02, the defect and the electroneutrality equation can be written
as follows:
CaO Cace" + Oo^ + Vo 2-1
2[0,"J + 2[CaCe"J + [e] = 2[V0 ] + [h ] 2-2
At a particular temperature, the electrolytic domain for the solid solution is the
regime of oxygen partial pressure where the ionic defects dominate i.e. in the present
example when
[Cace"] = [Vo] 2-3
As mentioned earlier for ZrC>2, doping apart from the generation of oxygen vacancies also
helps in the stabilization of the high temperature fluorite phase at lower temperatures.
Electrical conductivity (07, S/cm or ohm''em'1) of a species is given by the product
of concentration of species (c„ no/cm3), charge of the species (Z,e, coulomb), and
electrical mobility (ju¡, cm /sec-Volt). The total electrical conductivity of the system is the
sum of the contributions from all ionic and electronic species, while the partial
transference number (/,) of a species is given by the ratio of partial conductivity to the
total conductivity.
a, = CiZfiHi
2-4
oj = ¿07 = ¿ c,-Z,en,
2-5
t, = Oi/ oT
2-6
Due to the loose packing in the fluorite structure, the oxygen ion mobility is
significantly higher than that of other crystal structures, but it is still orders of magnitude
lower than that of electronic defects (electron and electron hole). In order to make a
predominant ionic conductor, large amounts of doping is required to generate high

22
concentration of ionic defects. In fact, doping is a misnomer in that respect as levels
required are in the range of 5-30 mol%. In addition to keep the concentration of
electronic defects in the system low, both the host and the dopant cations should have a
fixed valence state in the temperature and oxygen partial pressure range of interest.
2.2.3 Oxygen Ion Conductivity as a Function of Dopant Concentration and
Temperature
As doping increases the oxygen vacancy concentration, oxygen ion conductivity
can be enhanced by increasing the dopant concentration but only up to a certain level. As
shown in figure 2-3 for solid solution based on ceria, beyond a certain dopant level the
ionic conductivity decreases.15 This behavior has been explained in terms of association
of vacancies or formation of defect complexes in the dilute solution range and formation
of super-lattices or ordering of vacancies in concentrated solution range.
Introduction of the dopant cation changes the ionic potential of the lattice in the
vicinity of the dopant. This leads to trapping of the oxygen ion vacancy to the dopant
cation because of columbic attraction (between effectively -ve charged dopant and
effectively +ve charged oxygen ion vacancy) and/or relaxation of elastic strains
(generated by the dopant) in the lattice. Recent studies have shown that the elastic
relaxation has more to do with the trapping of oxygen ion vacancy with the dopant. As
shown in figure 2-3, dopants with the same charge do not show similar magnitudes for
oxygen ion conductivity and that the choice of the dopant for the optimum oxygen ion
conductivity in a particular system is critically related to the size of the dopant cation.
Dopants with similar size as that of the host show the best performance.
In the concentrated solution regime, long range ordering has been associated with
the conductivity decay. For example in the case of Ca stabilized ZrC>2 (CSZ), ordered

23
small domains (~10 Á) have been observed by electron diffraction studies after long
annealing times at elevated temperatures.16 This observation has been explained in terms
of disproportationation of metastable CSZ with time into ZrC>2 and CaZ^Os
(Zro.67Cao.3501.67). CaZr205 has a structure similar to that of C-type rare earth oxides
(with site specific oxygen vacancy) and thus with significantly lower conductivity.
Ordered sublattice essentially means that specific atoms have specific stable sites
with low energy, and therefore, for a conducting species ordering results in loss of mobile
species and in low conductivity. Y2O3 has C-type crystal structure which is similar to the
fluorite structure though with a site specific oxygen vacancy. This regularity results in
oxygen ion conductivity much smaller than that of CSZ, although the concentration of
oxygen vacancies is much larger. Similarly, BÍ2O3 exhibits two metastable phases below
650 °C: tetragonal P-phase and body centered cubic y-phase. Both these phases have an
ordered oxygen ion sublattice with oxygen ion conductivity up to three orders of
magnitude lower than the high temperature 5-phase, which has a cubic fluorite structure
with disordered oxygen ion sublattice.17’19 The concentration of oxygen vacancy is
similar in all the three phases, though the concentration of mobile oxygen vacancy is
significantly smaller in P-BÍ2O3 and Y-BÍ2O3. It is also important to note that the crystal
structure also plays a key role in determining the oxygen ion conductivity in terms of
jump directions, jump paths and activation barrier for the oxygen ion motion as described
below.8
a = c-Ze-p = câ–  (Ze)2D/(kT) 2-7
2-8
a = nv-Nâ–  (Ze):D0-exp (-A HJk- T)/(hT)
D0 =fz-(l - nv)-a2-Vo-exp(AS„/k)/6
2-9

24
Nemst-Einstein can use be used to relate the electrical mobility (p, cm /sec-Volt)
with diffusivity (D, cm2/sec), while concentration (c) can be related to site fraction of
oxygen vacancy (nv) and concentration of oxygen ion sites per unit volume (N, no/cm ).
Diffusivity is related to jump distance (a, cm), jump frequency (v, Hz), structure factor (f,
unit less on the order of 1), number of nearest neighbors (z), and number of available sites
for a vacancy jump {1 - nv). The jump frequency is related to the motional enthalpy {AHm)
for the oxygen ion jump to a neighboring vacant site through the common face bottle
neck. The motional enthalpy is related to the size of the mobile ion and distance between
the center of the common face and the peripheral ion. The Arrhenius relationship of
conductivity can be represented as
In(c-T) = InA - Ea/(kT)
2-10
A = nv-(l-nv)-N-(Ze)2-fz-a2-v0-exp(ASm/k)/(6-k)
2-11
In the case of an ordered system, vacancies can be differentiated on the basis of the
site i.e. normal (occupied) site and interstitial (vacant) site at a higher energy level (AHg).
AHg is a function of temperature and vanishes at the transition temperature (7j), above
which the system becomes disordered. The concentration of the vacancies on the normal
sites is thermally activated and increases with temperature. The fraction of vacant sites on
normal sites (nv) is related to ~exp(-AH¡/2kT) and the fraction of occupied sites on
ordered sites is related to m *nv, where m is the ratio of number of normal sites with
respect to interstitial sites. This results in higher activation energy for oxygen ion
conductivity at temperature below 7).8
Ea = AHm + AH¿2
T 2-12
Ea = AHm
T>T,
2-13

25
In the case of doped system below the critical temperature T*, oxygen ion
vacancies get progressively trapped into ordered clusters due to coulombic attraction and
elastic relaxation, as described earlier. Goodenough8 has shown using solution
thermodynamics that fraction of oxygen vacancies dissolved randomly on the
oxygen ion sublattice is given by
njn, * = exp(-AH( (.I - T/T*)/k- T) 2-14
where nv is the fraction of randomized vacant sites, nv* is the fraction of total vacant
(including both randomized and ordered) sites, and AH, is trapping enthalpy. This results
in well know kink in the Arrhenius plot with a higher activation energy and a higher pre¬
exponential term (A) at temperatures below T*.
Ea = AHm + AH,,
A, = A-exp(AH/k-T*) T < T,
2-15
Ea = AHm,
a2 = a t>t,
2-16
Different oxide systems show different magnitudes of T*, for example for yttria
stabilized zirconia (YSZ) -800 °C, for gadolinium doped ceria (GDC) -400 °C and for
erbia stabilized bismuth oxide (ESB) -600 °C. Further it has been observed that on
increasing the doping concentration, high temperature Ea increases and eventually
becomes equal to the magnitude of low temperature Ea resulting in a single activation
energy. This is consistent with the idea that the concentration of trapped clusters
increases with doping and tends to dominate the activation energy for conduction. Or in
other words, the transition temperature T* increases with doping in order to provide the
necessary thermal energy to release the oxygen ion vacancies from the traps.
Hohnke20 has modeled the dependence of the pre-exponential factor and the
activation energy on the dopant level by using a pseudo-chemical equilibria and by

26
introducing terms for short range and long range interactions. For low dopant levels,
oxygen ion conductivity can be given by the following equations
a = (AJT)f(c)-exp(-EJk-T)
2-17
O
II
Ea = AHm + AH,
T< TtandYce
2-18
II
Ea = AHm + AH,
T < T, and Cace"
2-19
f(c) = c;
Ea = AHm
T> T,
2-20
indicating that at high temperatures and low dopant levels, conductivity increases linearly
with the oxygen vacancy concentration. At lower temperatures, there is an additional
terms in the activation energy due to the short range interaction. Further depending on the
dopant type, conductivity varies square root of concentration for a single charged dopant
and is independent of concentration for a double charged dopant. With higher dopant
levels, the activation energy at low temperatures is found to increase linearly with dopant
level. This phenomenon has been attributed to long range interaction and is included in
the following generalized conductivity expression with equilibrium constants for short
range (K,) and long range (K¡) interactions.
Ea = AHm + AH, + AHfC
2-21
a = a0 [l + Kr (1 + KJ]-1
2-22
2.2.4 Grain Boundary Contribution to Total Conductivity
The kink in the Arrhenius plot of the conductivity can also be introduced by the
grain boundary contribution to the total conductivity. In doped ceria, the grain boundary
has a higher activation energy and a lower conductivity than the bulk either due to the
presence of impurities at the grain boundaries or due to higher dopant level than the bulk.
Due to the higher activation energy, the effect of the grain boundary on the total
conductivity is more prominent at low temperatures and it vanishes at higher

27
temperatures. The presence of the grain boundary contribution to the total conductivity
could result in suboptimum choice of dopant level as observed by Steele. As shown in
figure 2-4, lattice conductivity for 10 mol% gadolinium doped ceria (GDC) shows a
transition temperature of -400 °C with a high temperature activation energy of -0.64 eV
and a low temperature activation energy of -0.77 eV. By increasing the doping level to
20 mol%, the lattice conductivity decreases with a constant activation energy of-0.78
eV. The presence of impurities like SÍO2 at the grain boundaries in 10 mol% GDC result
in a low grain boundary conductivity with activation energy of -0.8-1.0 eV. The effect of
grain boundaries on the total conductivity can be seen up to 900-1000 °C. In comparison,
10 mol% GDC without impurities at grain boundaries shows little influence of the grain
boundary contribution to the total conductivity above 500 °C. For impure 20 mol% GDC
the grain boundary conductivity is higher and contributes up to 700 °C, which results in a
higher total conductivity than impure 10 mol% GDC. Thus, the presence of impure grain
boundaries can lead to suboptimum choice of the higher dopant level in terms of the total
conductivity, when clearly the lower dopant level shows better bulk conductivity.
2.2,5 Oxygen Ion Conductivity as a Function of Time
ESB in the cubic fluorite structure is one of the highest known oxygen ion
conductors and like YSZ and GDC, also shows the kink in the Arrhenius plot of
conductivity for low dopant levels as shown in figure 2-5. The transition temperature is
-600 °C, with a high temperature activation energy of -0.68 eV and a low temperature
activation energy of-1.28 eV. Further, it has been observed by Wachsman et al.22 that on
isothermal annealing below the transition temperature the conductivity decays with time,
as shown in figure 2-6. The conducitivity decay has been attributed to continuous
trapping of oxygen ion vacancies in the growing ordered clusters resulting in loss in

28
mobile oxygen vacancies. Conductivity decay also occurs in CSZ on annealing at
elevated temperature and the presence of ordered small domains (~10 A0) have been
observed by electron diffraction studies on annealed CSZ.16 Similarly, formation of
superstructure and oxygen ion displacement have been observed in ESB using electron
and neutron diffraction studies, respectively.22"27 The conductivity decay in stabilized
bismuth oxide is comparatively much faster than in YSZ and can be represented by the
following empirical equation
a(t) = erf co j + [o(0) - a(co)]exp[\(-t/xf] 2-23
where a(t) is the conductivity at time t, r is the pertinent time constant, and /? is a
dimensionless parameter. The larger is the time constant r, the more stable is the system.
The time constant increases with the dopant cation radius and with the dopant
concentration as shown in figure 2-7 and 2-8, respectively. As in most practical
applications, solid electrolytes will be used in an electrochemical potential gradient with
resultant oxygen ion flux. It is plausible that the oxygen ion flux can affect the trapping
of oxygen ion vacancies in the condensed ordered clusters and hence the conductivity
decay below the transition temperature. Part of chapter 3 deals with studying this
phenomenon in erbia stabilized bismuth oxide solid electrolyte.
2.3 Conclusion
This chapter introduces the basics of oxygen ion conduction in solid oxide
electrolytes. Models from the literature for oxygen ion conductivity as a function of
dopant level, under different temperature regimes and as a function of time are described.
The primary reason for the deviant behavior from ideal oxygen ion conductivity is due to
the elastic relaxation on introducing the dopant in the host leading to formation of defect
complexes in the case of low dopant levels and formation of ordered clusters in the case

29
of high dopant levels. This results in well known kink in the Arrhenius plot of
conductivity with a high activation energy at low temperatures and a low activation
energy at high temperatures. Further, the ordering phenomenon could also lead to decay
in conductivity on isothermal annealing below the transition temperature, which is
believed to occur due to trapping of mobile vacancies in the condensed ordered clusters.

30
Temperature (°C)
900 800 700 600 500 400 500
1000/T (K-1)
Figure 2-1. Oxygen ion conductivity of common solid electrolytes (from ref. 5).

31
Figure 2-2. The fluorite structure (Blue spheres - Anions; Red spheres - Cations) (from
ref. 17).
Figure 2-3. Isothermal conductivity of some ceria solid solution at temperature close to
200 °C (from ref. 17).

32
Figure 2-4. Schematic Arrhenius plots of lattice and grain boundary conductivity for (a)
clean Ceo.9Gdo.101.95 (b) impure Ceo.9Gdo.1O195 (c) impure Ceo.gGdo 20i.9 (from ref. 23).

33
KXXD/KK-')
Figure 2-5. Conductivity of (Bi203)i.x(Er203)x in air; (â–¡) x = 0.2, (X) x = 0.25, ( A) x =
0.3, (V) x = 0.35, (o) x = 0.4, (•) x = 0.45, (A) x = 0.5, (▼) x = 0.6; broken
line represents the conductivity of pure BÍ2O3 (from ref. 21).

Relative Conductivity a(tVoiO)
34
19
Figure 2-6. Decay in isothermal oxygen ion conductivity of 20 mol% ESB on annealing
below the transition temperature (from ref. 31).

35
LATTICE PARAMETER (A)
1.00 1.01 1.02 1.03 1.04 1.05
CATION RADII (A)
Figure 2-7. Time constant as function of cation radii and lattice parameter of 25 mol%
cation doped bismuth oxide (from ref. 30).
LATTICE HAHAMbltH (A)
5.480 5.500 5.520 5.540
Figure 2-8. Time constant as function of erbia concentration and lattice parameter of
erbia stabilized bismuth oxide (from ref. 30).

CHAPTER 3
DIRECT CURRENT BIAS STUDIES ON (Bi2O3)0 8(Er203)o.2 ELECTROLYTE AND
Ag-(Bi2O3)0 g(Er2O3)0 2 CERMET ELECTRODE
3.1 Introduction
Bismuth oxide based electrolytes show one of the highest oxygen ion conductivity
among all solid electrolytes and are of considerable interest for application in solid oxide
fuel cells and oxygen sensors.17'14 The high temperature 8-BÍ203 has a cubic fluorite
structure with 25 % inherently vacant oxygen sites. The high concentration of disordered
oxygen ion vacancies in 8-BÍ203 along with the high polarizability of Bi cation results
in oxygen ion conductivity which is one to two orders of magnitude higher than that of
yttria stabilized zirconia at comparable temperatures. However, on cooling to lower
temperatures the high temperature 8-phase transforms to the stable monoclinic a-phase at
729 °C. Two intermediate metastable phases (tetragonal P-phase and body centered cubic
y-phase) are also known to exist below 650 °C. The P and y-phase have an ordered
oxygen ion sublattice with oxygen ion conductivity up to three orders of magnitude lower
than the 8-phase, while the a-phase is a p-type electronic conductor. Further, the
transformation from 8-phase to P-phase is accompanied by a sudden large volume change
which destroys the mechanical integrity of the material during thermal cycling.
Therefore, it is necessary to stabilize the high temperature and high oxygen ion
conductivity 8-phase at lower temperatures in order to utilize Bi203 in practical
applications.17"19 Takahashi et al.29j0 were first to show that by doping Bi203 with
isovalent rare earth oxides the high temperature 8-phase can be retained at lower
36

37
temperatures. Unlike stabilized zirconia, doping does not result in increase in the oxygen
ion vacancy concentration but only serves to stabilize the high temperature phase at lower
temperatures resulting in enhanced oxygen ion conductivity.
Arrhenius plots of oxygen ion conductivity of these stabilized bismuth oxides show
a characteristic kink at transition temperature (T*) -600 °C with a lower activation
energy above T* and a higher activation energy below T*. Moreover, the conductivity
shows a continuous decay with time on isothermal annealing in air below this transition
temperature.23 24 28 3132 The conductivity decay has traits similar to those of nucleation
and growth. The rate of decay is fastest at -500 °C and is comparatively slower both
above and below 500 °C.23,31,32 The conductivity decay is reversible and can be achieved
by heating above the transition temperature. ’ ’
Transmission electron microscopy studies on stabilized bismuth oxides show that
the isothermal anneal below the transition temperature results in the formation of
superstructure with a lattice parameter twice that of the parent structure. ' Neutron
diffraction studies have shown that after annealing the oxygen ions are also displaced
from the center of the cation tetrahedra (8c-sites) towards the faces of the cation
tetrahedra (32f-sites) along the <111> directions.22 23 26,27 Calorimetric studies on
annealed samples show the presence of an endotherm at about the transition temperature;
the magnitude of the endotherm increases with the annealing time." ’ ’ ‘ The cooling
cycle does not show an exotherm indicating that the reverse process is comparatively
much slower.
The above described conductivity and structural changes in stabilized bismuth
oxides on annealing below the transition temperature have been attributed by Wachsman

38
and coworkers ’ to occupational and positional ordering of the oxygen ion sub-lattice:
ordering of oxygen vacancies in <111> directions resulting in the formation of a 2X2
super-lattice and displacement of oxygen ions from 8c to 32f site also along the <111>
directions. On isothermal annealing below the transition temperature, the mobile oxygen
vacancies are consumed in the growing ordered clusters and become unavailable for
conduction resulting in the conductivity decay with time. This process is reversible and
upon heating above the transition temperature the lattice disorders with resultant regain in
conductivity. The endotherm at the transition temperature is due to the heat energy
required for disordering the ordered lattice. The magnitude of the endotherm increases
with the annealing time indicating that the extent of ordering increases with the annealing
time.
In most applications, solid electrolytes will be used in an electrochemical potential
gradient with resultant oxygen ion flux, and therefore, it is important to understand the
effect of oxygen ion flux on the conductivity decay and the ordering kinetics of bismuth
oxide solid electrolytes. Among lanthanide stabilized bismuth oxides, 20 mol% erbia
stabilized bismuth oxide (ESB) shows the highest oxygen ion conductivity and is used in
the present study. To generate the oxygen ion flux, direct current bias was applied across
a symmetric cell consisting of ESB electrolyte and Ag-ESB cermet electrodes, and
electrochemical impedance spectroscopy was used to measure the oxygen ion
conductivity of the electrolyte as a function of temperature, time, and current bias.
Further, the heat of enthalpy of the order-disorder transition was measured using
differential scanning calorimetry (DSC) for samples annealed at different temperatures,
time periods, and current bias to understand the effect of oxygen ion flux on the ordering

39
kinetics. Ag-stabilized bismuth oxide cermets were used as electrodes in this study as
they have been reported to have high performance as cathodes for IT-SOFCs by Xia et
al.34 It was hoped that these electrodes would prevent the gas diffusion limited
decomposition of bismuth oxide based electrolyte to Bi metal at high currents as
mentioned by Takahashi et al.3' and Doshi et al.36 However, it was found during the
course of this study that there are some long term stability issues with these electrodes.
3.2 Experimental
Standard solid state method was used to fabricate (BÍ203)o,8(Er203)o.2 electrolyte
discs. BÍ2O3 (99.999 %, Alfa Aesar) and E^Cb (99.99 %, Alfa Aesar) powders in desired
weight ratios were ball milled in acetone for 24 hours with zirconia ball media and
calcined at 800 °C for 10 hours. The resulting powder was again ball milled, sieved,
uniaxially pressed, cold isostatically pressed, and finally sintered at 890 °C for 10 hours.
For the electrochemical experiments, sintered discs were polished to a final thickness of
-0.76 mm. Diameter of the discs after sintering was ~10.9 mm.
For the cermet electrodes, (BÍ203)o.8(Er203)o.2 powder was prepared using an
amorphous citrate route. Bi(N03)3.5H20 (99.999 %, Alfa Aesar) and Er(N03)3.5H20
(99.9 %, Alfa Aesar) in desired weight ratios were first dissolved in dilute nitric acid
solution, and then citric acid was added in a metal cationxitric acid molar ratio of 1:1.5.
The solution was gelled and foamed at 80-100 °C. The precursor was then calcined at
temperatures between 400 °C and 700 °C. Ag-(BÍ203)o.s(Er203)o.2 electrode paste was
made by ball milling Ag2Ü (Alfa Aaesar) and (Bi203)o.8(Er203)o.2 in 60:40 wt.% ratio
with Heraeus V006 binder for 24 hours. For the symmetric cells, the electrode paste was
screen-printed on both sides of the electrolyte to which Ag lead wires were attached, and
the assembly was sintered at 750 °C for 1 hour.

40
For the experimental setup, a frequency response analyzer (Solartron 1260) along
with an electrochemical interface (Solartron 1287) was used. A constant oxygen ion flux
was generated inside the electrolyte by using the electrochemical interface in the
galvanostatic mode. Impedance measurements were done every 30 minutes to separate
the electrolyte resistance from the electrode response. The cell was given a pretreatment
step in order to avoid the overloading of the current range during impedance
measurements. The samples were placed in a quartz experimental setup with Au lead
wires which was then put into a horizontal tube furnace, and the experiments were
conducted in air. Different samples were used for each temperature, time, and current
bias experiment. All samples were equilibrated at 650 °C and then cooled down to the test
temperature. The experiments were terminated at different time periods depending on
current bias and temperature in order to avoid the electrochemical decomposition of the
electrolyte.
Samples were characterized using XRD (APD-3720) and SEM (JEOL-6400).
Electron probe micro analysis (EPMA) was done on the sample cross-sections to
characterize the concentration profile of elements as a function of thickness. DSC (TA-
1290) studies were done to measure the heat of transformation of samples under different
conditions. The electrodes were polished off, and bulk samples of ~9 mg mass were
heated from room temperature to 800 °C at 20 °C/min in 100 seem of argon.
3.3 Results and Discussion
3.3.1 Processing
Calcination of the oxide mixture at 800 °C resulted in single phase powders and
sintering at 890 °C resulted in discs with more than 95 % density. XRD patterns for the
powders derived from the citrate route after calcination at different temperatures are

41
shown in figure 3-1. The citrate route did not significantly decrease the calcination
temperature to achieve the cubic fluorite structure; the primary phase up to 600 °C was
tetragonal (TBÍ2O3. However for the cathode paste, powder calcined at 500 °C for 10
hours was used which eventually would form the cubic phase on sintering at 750 °C.
SEM micrographs of the cross-section of the electrode/electrolyte interface in the
secondary and back-scattered electron mode are shown in figure 3-2 (a) and (b),
respectively. The thickness of the electrode is ~17 pm. After sintering, the electrode
seems to have low porosity possibly due to the low melting points of both Ag and
bismuth oxide. This microstructure is similar to that obtained by Xia et al.34 with Ag-
yttria stabilized bismuth oxide electrodes.
3.3.2 DC Bias and Impedance Spectroscopy: ESB Electrolyte
Figure 3-3 is a plot of log(aT) vs 1/T for the ESB electrolyte in air. The graph
shows the well-known kink in the slope at T* -600 °C with activation energy equal to
0.68 eV above the transition temperature and 1.28 eV below the transition
temperature.19'23 Below the transition temperature, ordering of the oxygen ion sublattice
results in higher activation energy. Figure 3-4 is a plot of percentage change in resistance
of ESB as a function of time on isothermal annealing at 300 °C, 500 °C, and 625 °C. The
rate of increase is fastest at 500 °C, much slower at 300 °C, and negligible at 625 °C. For
example after 40 hours, the resistance increase at 625 °C was ~0.4 %, at 500 °C was
-898.4 %, and at 300 °C was -12.2%. This suggests that the ordering has characteristics
typical to those of nucleation and growth, where the transformation rate depends on the
drive force as well as on the diffusion rate. At 625 °C, the driving force is small, and at
300 °C, the diffusion rate is slow.

42
Most of the dc bias experiments were carried out at 500 °C, where the ordering of
oxygen ion sublattice is fastest. On applying the bias at 500 °C, the initial voltage drop vs
time plots showed a valley; depth of the valley increased with the bias current as shown
in figure 3-5. This gives an impression that the bias is affecting the ordering process.
However, impedance spectroscopy done after the first 30 minutes of application of the
bias separated the electrolyte resistance from the electrode response and showed that the
valley is most likely related to the activation polarization at the electrode as described in
a following section. On the other hand, the electrolyte resistance increased with time at
500 °C without much affect of the dc bias as shown in figure 3-6. It was observed that at
high currents the increase in resistance was slightly faster though within the error limits
of the experiment. It should also be noted that the experiments at higher current bias were
much shorter, in comparison to lower currents, in order to avoid the electrochemical
decomposition of the electrolyte with voltage drops greater than 0.65 V across the
sample. For example, at 72 mA bias the experiment was terminated after ~4 hours and in
the case of 12 mA bias the experiment was terminated after ~42 hours. Another set of
samples were studied at 500 °C to confirm the results from 0-36 mA, and the electrolyte
resistance after 10 hours were essentially the same in all the samples.
At 625 °C i.e. above the transition temperature, the percentage change in resistance
as a function of current bias is shown in figure 3-7. The data has considerable noise,
which is basically due to the electrode’s microstructural instability caused by the
electromigration of Ag as discussed in a following section. At 250 mA bias, the migration
was severe, and over the period of the experiment, the migration led to the formation of a
dense Ag layer on the counter electrode. This resulted in the eventual reduction of the

43
electrolyte to Bi metal up to ~0.2 mm in depth as found by EPMA on the electrolyte
cross-section. However towards the end of the experiment, for samples with no bias and
for samples under bias on removal of the bias, the electrolyte resistance was the same as
that of the initial value. At 300 °C, dc bias experiments were not conducted due to high
resistance of the electrolyte.
3.3.3 DSC Studies: ESB Electrolyte
Stabilized bismuth oxides on annealing below the transition temperature show an
endotherm at about the transition temperature during subsequent heating. ’ ’ " This
endotherm has been attributed to the heat energy required to disorder the ordered oxygen
ion sublattice. The ordering process is comparatively much slower, and hence, no
exothermic heat is observed during the cooling cycle and even after -10 hours of anneal
at 500 °C as found in this study. However, on annealing for -74 hours an endotherm
appears at 614.5 °C with a magnitude of 4.12 kJ/mol as given in Table 3-1. The ordering
process continues beyond 100 hours of anneal as observed by Wachsman et al.32 Samples
annealed under bias showed similar kinetic behavior, though the endotherm appeared at
shorter annealing times. After -10 hours anneal, the sample under no bias did not show
the endotherm, while the samples under 12-36 mA bias showed the endotherm as given
in Table 3-1. Moreover, it appears that for a constant annealing time the magnitude of the
endotherm increases with the bias current.
As expected, above the transition temperature at 625 °C there was no endothermic
peak observed even after annealing for -137 hours. Samples under 100 mA and 250 mA
bias at 625 °C also did not shown any endothermic peak. On annealing at a much lower
temperature of 300 °C for -49 hours, an endotherm of magnitude of 3.08 kJ/mol was
observed at 606.6 °C.

44
The DSC results on samples annealed at 500 °C do not correlate well with the
resistance measurements. After ~10 hours, without bias the electrolyte showed more than
200 % increase in the resistance though without any trace in the DSC, and under bias the
sample showed similar increase in resistance but with an endotherm in the DSC profile.
This difference could be simply due to the detection limits and characteristics of the
technique employed or due to difference in the nucleation and growth characteristics of
the ordered domains under bias and without bias. The magnitude of the endotherm is
primarily dependent on the volumetric amount of the ordered phase present, while the
resistance depends on the parallel and series paths available for conduction. For example,
resistance will be more sensitive than DSC to minute amounts of the ordered domains
present at the grain boundaries.
As mentioned earlier, ordering of the oxygen ion sub-lattice comprise of two parts:
positional displacement of oxygen ions from 8c to 32f sites along <111> and
occupational ordering of oxygen vacancies along <111>. ’ Positional displacement is
expected to be the faster of the two as the magnitude of the displacement is less than the
oxygen ion radii, and in fact, it has been observed to be partially present in un-annealed
samples. In contrast, occupational ordering involves rearrangement of the oxygen lattice
over larger distances on the order of the lattice parameter, and thus, it is expected to have
a higher time constant. The endotherm is expected to be primarily related to the
occupational ordering, and hence, it appears only after long anneal time under no bias
conditions. Since the endotherm at short time is enhanced by the applied bias with
negligible change in conductivity decay, it is possible that the decay in conductivity with
anneal time is more related to positional ordering than occupational ordering.

45
Kruidhof et al.37 have noticed that erbia stabilized bismuth oxides are non-
stoichiometric and slowly uptake oxygen on cooling below 550 °C in an oxygen
atmosphere. The ordering process, as evinced by the endotherm, is also a slow process
which occurs below -600 °C. The matching temperature-time profile and species suggest
that the oxygen uptake and oxygen ion ordering (possibly occupational) could be related
processes. DSC results under bias conditions could also be explained using the
assumption that the current bias is assisting in the oxygen uptake and hence, in the
ordering process. These and other possibilities require further extensive studies, though it
is clear from the conductivity and DSC studies that the applied current bias is not
effecting the conductivity decay in erbia stabilized bismuth oxide below the transition
temperature but could be assisting in the occupational ordering.
3.3.4 DC Bias and Impedance Spectroscopy: Ag-(BÍ203)o.8(Er203)o.2 Electrode
Area specific resistance (ASR) of the Ag-(BÍ203)o.s(Er203)o.2 electrode as a
function of temperature in air is shown in figure 3-8, with values of 3.08 Qcm at 500 °C
and 0.16 Qcm2 at 625 °C. The electrode ASR was calculated by multiplying the electrode
resistance by the electrode area (-0.93 cm2) and dividing by 2 to account for the
symmetric cell. The performance of the cermet electrode is comparable to that reported
by Xia et al.34 and is encouraging for application as cathode in IT-SOFC.
Impedance plots of the electrode at 500 °C under 0-72 mA current bias are shown
in figure 3-9. The electrode impedance decreased with bias current as shown in figure 3-
10, which suggests that within the current range studied the electrode impedance is
dominated by activation polarization. As shown in figure 3-9 (c), the electrode impedance
consisted of at least two arcs: a low frequency arc with a characteristic frequency at -0.2
Hz and a high frequency arc with a characteristic frequency at -630 Hz. The plot of the

46
imaginary part of impedance (Z") vs. log frequency (f) provides a better representation of
the frequency distribution in the impedance data. On application of the bias current, the
dominant low frequency arc divided into two arcs with significantly smaller impedance,
while the impedance of the high frequency arc remained unaffected. The characteristic
frequency of the middle arc increased with the bias current, while on the other hand the
characteristic frequency of the low and high frequency arcs appear to be independent of
the applied current bias. The electrode arcs represent the different reaction steps in the
electrode reaction, and it is suggested that the low and middle frequency arcs are related
to charge transfer step, while the high frequency arc is related to the bulk transport in the
electrode.
Impedance plots of the electrode at 625 °C under different current bias are shown in
figure 3-11. As expected, the activation polarization at 625 °C is much smaller in
comparison to 500 °C as shown in figure 3-10. The electrode impedance consisted of at
least two arcs: a low frequency arc with a characteristic frequency at ~20 Hz and a high
frequency arc with a characteristic frequency at -400 Hz, as shown in figure 3-11 (b). As
in the case of 500 °C, impedance of the dominant low frequency arc decreased with the
bias current, while the impedance of the high frequency arc remained unaffected.
However, at 625 °C the electrode performance was not stable under 250 mA bias, and the
electrode resistance increased by six times after ~23 hours as shown in figure 3-12.
Magnitude of the high frequency electrode arc increased with time and eventually
became the dominant arc, which indicates that it might be related to some sort of
concentration polarization. Interestingly, magnitude of the high frequency arc reduced on

47
removing the bias at the end of the experiment, and once again, the low frequency arc
became the dominant arc.
After the experiment, the top surface of the working electrode looked silvery which
was initially light brown in color, and the counter electrode peeled off very easily with a
silver layer at the electrode/electrolyte interface. SEM micrographs of the
electrode/electrolyte interface at the working and the counter electrode after ~26 hours at
625 °C under 250 mA bias are shown in figure 3-13. Significant electromigration of Ag
occurred along with the oxygen flux to the electrode/electrolyte interface at the counter
electrode and to the top surface at the working electrode. A schematic of the proposed
electrode reaction mechanism is shown in figure 3-14. It is suggested that the low
frequency arc is related to the charge transfer step, while the high frequency arc is related
to the bulk transport of oxygen in the silver particles. The electrode response under no
bias is limited by the charge transfer step, and on application of the bias the charge
transfer resistance decreases, though along with Ag electromigration in the electrode. The
migration with time results in the formation of a dense Ag layer at the
electrode/electrolyte interface and in the bulk transport contribution in the electrode
impedance. As mentioned earlier, the formation of the dense Ag layer resulted in the
eventual reduction of the electrolyte at the counter electrode; however, it was not detected
in the impedance plots.
Oxygen diffusivity (D0) and solubility (C0) in Ag layer for the case of finite
Warburg diffusion could be calculated using the following equations
32-(comax/2D0) = 4/n 3-1
ZD(oo-^0) = RT/(n2F*C0D0)
3-2

48
where comax is the frequency at the maximum value of the imaginary component, S is the
diffusion length, and Zd(co—*0) is the real intercept.38 For Ag-ESB electrode at 625 °C,
250 mA after -26 hours, the calculated values for D„ and C0, assuming 1 pm diffusion
length, are 6.19X10' cm /sec and 2.91X10' moles/cm , respectively. The calculated
ft 9
values match well with experimental values for D0 (7.42X10' cm /sec) and C0
(1.94X10' moles/cm ) reported in the literature, which supports the argument that the
high frequency arc in the electrode impedance is related to the oxygen diffusion in Ag.
Phase diagram between Ag and Ag2Ü, as proposed by Assal et al.,40 is shown in
figure 3-15. Ag forms a eutectic with Ag20 at 530 °C and 519 atm. O2. The liquidus
temperature in an oxygen free atmosphere is 962 °C, which is lowered to 951 °C and 939
°C in air and 1 atm. O2, respectively. The high performance of the Ag electrodes for
oxygen reduction is because of its oxygen solubility which also results in the formation of
the eutectic, and as found in this study will consequently lead to electromigration of Ag
along with the oxygen flux above the eutectic temperature. Therefore, electrodes
containing Ag in sufficient proportions should not be expected to be microstructurally
stable for long periods of time even at intermediate temperatures. Ag-Pt and Ag-Pd alloys
with higher melting points can be better than Ag, though they will be much costlier
alternatives.
3.4 Conclusion
The effect of direct current bias on the ordering phenomenon of ESB electrolytes
and on the performance of Ag-ESB electrodes was studied. Impedance studies on ESB
electrolyte showed that on isothermal annealing below the transition temperature at 500
°C, the current bias did not have a significant effect on the rate of increase in resistance.
On the other hand, DSC studies showed that the endotherm related with the reverse

49
transition appeared at shorter annealing time periods on application of the current bias.
The above results are explained in terms of the relative effect of applied current bias on
the kinetics of occupational and positional ordering on the oxygen ion sub-lattice. It is
suggested that the rate of occupational ordering is enhanced under the bias current
resulting in the presence of endotherm at shorter time periods but without much change in
rate of increase in resistance. Ag-ESB cermet electrodes showed good electrochemical
performance. Impedance studies under current bias showed that the electrode response
consisted of charge transfer and bulk transport contributions. The electrode
microstructure was unstable at high bias currents at 625 °C, primarily due to the
electromigration of Ag along with the oxygen flux. The high oxygen solubility in Ag
results in good performance of the electrode, but it also results in the lowering of the
melting point of Ag and its consequent electromigration. Ag-Pt and Ag-Pd alloys with
higher melting points are suggested for replacing Ag, though they would increase the
cost.

COUNTS
50
Figure 3-
20
. XRD patterns of calcined 20 mol% erbia stabilized bismuth oxide powders
derived from amorphous citrate route.

51
Figure 3-2. Cross-sectional micrographs of the ESB electrolyte and Ag-ESB electrode
interface in (a) secondary electron and (b) back-scattered electron mode.

52
1000/T (K'1)
Figure 3-3. Arrhenius plot of oxygen ion conductivity in ESB electrolyte.

53
10 100 1000 10
Time (min)
Figure 3-4. Percentage change in resistance of the ESB electrolyte under no bias at 300
°C, 500 °C, and 625 °C.

A(|> (Volt)
54
0.525
0
36 mA
10 15 20 25 30
Time (min)
35
Figure 3-5. Initial voltage vs. time plots at 500 °C under different bias currents.

55
Figure 3-6. Percentage change in resistance of the ESB electrolyte at 500 °C under
different bias currents (R<, = 4.4 Í2).

56
50
1 1 1
40
? • 0mA
^ 100mA
N®
"o 30
<¡> o 250mA
O:
O.
20
AÁ:
°r
0°:
4-*
oí
f>
^ 10
J 1? ! _ fl
\ fflSp 625 °c
0
O 1000 2000 3000
Time (min)
Figure 3-7. Percentage change in resistance of the ESB electrolyte at 625 °C under
different bias currents (Ro = 0.7 Q).

57
Temp (°C)
Bias (mA)
Time (hr)
AH, (kJ/mol)
T,(°C)
300
0
49.5
3.08
606.6
500
0
10
_
500
0
74
4.12
614.5
500
12
10
3.40
607.1
500
12
42
5.32
599.7
500
30
11.5
3.38
603.8
500
36
7
2.63
594.1
500
36
10
4.02
612.3
500
72
4
_
_
500
100
3.5
_
_
625
0
137
_
_
625
100
24
_
_
625
250
26
_
_
Table 3-1. Enthalpy for order-disorder transition and transition temperature of ESB
electrolyte on isothermal annealing at different temperatures, bias currents,
and time periods.

ASR
58
Figure 3-8. Arrhenius plot of ASR (Qcm2) for Ag-ESB electrode.

59
(a)
-5
-4
-3
-2
-1
O
1
4 5 6 7 8 9 10
Z' (Q)
(b)
~i 1 r
500 °C
0.2 Hz
0mA
-0.6 24mA4o,M>«^
500 °C
! ,1TTTi r—^
500 °C
0mA
â–  12mA
♦ 24mA
30mA
T 36mA
• 72mA
10 100 1000
f (Hz)
-0.4
^ _0-3 36mA
-0.2
•• 72mA
-u.i
10 100 1000
f (Hz)
-2.5
Figure 3-9. Impedance plots of Ag-ESB electrode at 500 °C under different bias currents
(a) and (b) imaginary vs. real impedance (c) and (d) imaginary impedance vs.
frequency.

60
G
o
c
cd
-*—>
C/2
o
O
_CD
DP
7
■Í'
6 1
5 4
4
500 °C
3
2 \
1 625 °C
o" * â–  - -
0 100 200 300
Bias Current (mA)
400
Figure 3-10. Ag-ESB electrode resistance vs bias current at 500 °C and 625 °C.

61
Figure 3-
(a)
-0.14
-0.105
a -°-07
-0.035
0
0.1 1 10 100 1000 104
f(Hz)
(b)
1. Impedance plots of Ag-ESB electrode at 625 °C under different bias
currents (a) imaginary vs. real impedance (b) imaginary impedance vs.
frequency
625 °C
0 mA
100 mA

62
-1.2
-0.8
a
í -0.4
0
0.6 0.8 1 1.2 1.4 1.6 1.8 2 2.2
Z' (Q)
(a)
-0.5
•%
30 min * •
-0.4 ■ 510 min * •
♦ 570 min
-0.3 A 660 min •
0.1 I—■— ■■■- — I4
0.1 1 10 100 1000 104
f (Hz)
(b)
625 °C, 250mA
400 Hz
\
• • • • •
30 min
â–  510 min
♦ 570 min
660 min
T 720 min
• 1290 min
. 20 Hz
r-
XV.
Figure 3-12. Impedance plots of Ag-ESB electrode under 250 mA bias at 625 °C (a)
imaginary vs. real impedance (b) imaginary impedance vs. frequency.

63
(a)
(b)
Figure 3-13. Cross-sectional microstructures of the electrolyte/electrode interface after
annealing experiments at 625 °C under 250 mA bias for ~26 hours (a) working
electrode (b) electrolyte interface after counter electrode peeled off.

64
o2
Oads
e')
\ij° Ag
\
O2- ESB
t
(b)
Figure 3-14 Proposed electrode reaction mechanism for Ag-ESB electrodes.

65
Ag Mole fraction O
Figure 3-15 Phase diagram between Ag and Ag2Ü proposed by Assal et al. (from ref.
42).

CHAPTER 4
BISMUTH RUTHENATE BASED CATHODES FOR IT-SOFC
4.1 Introduction
For efficient operation of SOFCs at intermediate temperatures, better performance
electrolytes and electrodes are required. Gadolinium doped ceria (GDC) shows
significantly higher ionic conductivity compared to YSZ at the same temperature and is
considered a promising electrolyte material for IT-SOFCs.41 For cathodes, perovskites
based on lanthanum manganite, lanthanum cobaltite, and lanthanum ferrite have been the
material of choice.42 Recently, pyrochlores based on bismuth ruthenate, lead ruthenate,
and yttrium ruthenate have been studied for application as cathodes in SOFCs.4345
Ruthenium oxide is known to be catalytic active towards oxygen reduction and has been
studied as a cathode material.46 It can be expected that solid solutions containing
ruthenium oxide can also be beneficial as cathodes. Pyrochlore ruthenates are electrically
conductive and thus, satisfy the other requirement for the cathodes. In this work, cathodes
based on bismuth ruthenate were studied for IT-SOFCs based on ceria electrolytes.
4.1.1 Undoped Bismuth Ruthenate
Abraham et al.47 first showed that bismuth ruthenate apart from the fee pyrochlore
structure BÍ2RU2O7 also exists in the oxygen enriched cubic KSbOj structure as
BÍ2RU2O7 3 (or BÍ3RU3O11). At 975 °C, BÍ2RU2O7 3 transforms irreversibly to BÍ2RU2O7 in
air. BÍ2RU2O7 is phase stable at 800 °C and did not transform back even after a month.47
BÍ2O3-RUO2 phase diagrams reported by Prosychev et al 48 and Hrovat et al.49 are shown
in figure 4-1 and 4-2, respectively. There appears some confusion regarding the transition
66

67
temperature from the KSbCT structure to the fee pyrochlore structure (975 °C,47 9 65 °C,48
950 °C49). Hrovat et al.49 noted that the reaction of precursors with alumina crucible could
have affected the transition temperature.
Linquette-Mailley et al.50,51 have used both BÍ2RU2O7.3 and BÍ2R.U2O7 as electrodes
to reduce the response time of YSZ based oxygen sensors at low temperatures. They
found that the oxygen content in BÍ2R.U2O7 is slightly over-stoichiometric in the pressure
range of 3X103 to 10' Pa at 800 °C, while BÍ2R.U2O7 3 reduces irreversibly to BÍ2R.U2O7 at
pressures less than 7.4X102 Pa at 800 °C. Further, from redox potentiometric
measurements they concluded that BÍ2R.U2O7 reduces and becomes a mixed ionic
electronic conductor at cathode polarizations higher than 700 mV/air at 373 °C; therefore,
the length of triple phase boundaries (TPBs) in the electrode structure is very important at
low polarizations.
Bae et al.43 studied BÍ2R.U2O7 3, Pb2Ru2Ü6 5 and Y2RU2O7 as cathode materials for
IT-SOFCs based on ceria electrolytes. They found that BÍ2RU2O7 3 and Pb2Ru2Ü6 5 react
with ceria at the processing temperatures, while Y2RU2O7 was stable with ceria at 900 °C
and was studied in detail using impedance spectroscopy. Y2RU2O7 electrode showed an
area specific resistance (ASR) of 4000 Qcm2 at 627 °C, and on doping with 5 mol% Sr,
the electrode ASR reduced to 47 if cm2. This improved performance of the electrode was
explained in terms of the enhanced ionic conductivity of Y2RU2O7 with Sr doping.
Takeda et al.44,45 studied pyrochlores BÍ2RU2O7, Pb2Ru2C>6.5 and perovskites
CaRu03, SrRuC>3 as cathode materials for YSZ based SOFCs. Bi2Ru207 and Pb2Ru206 5
were found to be stable with YSZ at 900 °C. However, the presence of sillenite type-
impurity phase in BÍ2RU2O7 could result in the formation of monoclinic zirconia. They

68
observed metallic behavior for pyrochlore ruthenates with almost temperature
independent conductivity (102-103 S/cm from room temperature up to 900 °C), which is
comparable to the best conventional cathode materials. The thermal expansion coefficient
ofBi2Ru207 and Pb2Ru206 5 between 700-900 °C was measured to be 0.99-1.00 X 10'^ K'
1 and 1.10-1.21 X 10° K"1, respectively. Pyrochlore ruthenates showed better electrode
performance than the perovskite ruthenates, and their performance was comparable to
lanthanum manganate but inferior to that of lanthanum cobaltite based cathodes.44
4.1.2 Doped Bismuth Ruthenates
A2B207 pyrochlore structure is essentially derived from an oxygen deficient cubic
fluorite structure with both ordered cation and anion sub-lattice, as shown in figure 4-3. It
exhibits Fd 3 m space group with eight formula units within a cubic unit cell. The cation
sublattice consists of bigger A3+ and smaller B4+ which order into alternate (110) rows in
every other (001) plane and in alternate (110) rows in the other (001) planes. This cation
ordering provides three distinguishable tetrahedral sites for the oxygen ions: 8a-sites
surrounded by 4 AJ~ cations, 8b-sites surrounded by 4 B4+ cations, and 48f-sites
surrounded by 2 A3+ and 2B4+ cations. 8a and 48f-sites are occupied, while 8b-sites are
vacant resulting in an ordered oxygen ion sub-lattice. The formula unit of the pyrochlore
can also be written as A^BiC^O’ to distinguish between the oxygen ions occupying the
48f-site as O and those occupying the 8a-sites as O’.
Although bismuth ruthenate pyrochlore has sufficient electronic conductivity to
perform as a good cathode, it will be beneficial to introduce oxygen ion conductivity into
the structure. Presence of ionic conductivity in the structure will not limit the oxygen
reduction reaction to the TPBs, leading to enhanced electrode performance. The oxygen

69
ion vacancies in the pyrochlore structure are ordered but still result in an ionic
conductivity, which is larger than that of undoped fluorites. It is expected that the
pyrochlore structure may exhibit a transition temperature above which the oxygen ion
# Q #
sub-lattice will disorder, leading to enhanced oxygen ion conductivity. As the anion
ordering is related to the cation ordering, the transition temperature could be lowered by
manipulating the ion radii of the two cations. The smaller is rA/re, the lower is the
expected transition temperature. As in the case of fluorites, oxygen ion conductivity can
also be introduced into the pyrochlore structure by generating additional oxygen ion
vacancies by doping with lower valent cations. These additional oxygen ion vacancies are
expected to primarily occupy 8a and 48f-sites and hence, could contribute to oxygen ion
conduction. However, one has to remember that apart from the concentration of mobile
oxygen vacancies, the oxygen ion conductivity also depends on structure factors such as
jump directions, jump paths, and activation barrier for the oxygen ion motion and on
dopant-vacancy interactions.
Fortunately, bismuth ruthenate forms solid solution with large solubility limits with
a number of dopants on the A-site, and the doping does not significantly affect the
electronic conductivity at least at room temperature?2 In the present work, Ca2+ (r = 1.12
Á), Sr2+ (r = 1.26 Á), and Ag+ (r = 1.28 Á) with comparable ionic radii with the host Bi3+
(r = 1.17 A) were studied as dopants on A-site. This strategy was used keeping in mind
that, possibly as in fluorite systems, the dopant with better matched ionic radii with the
host would generate lesser elastic strains in the lattice and would show higher oxygen ion
conductivity.

70
4.1.3 Bismuth Ruthenate and Stabilized Bismuth Oxide Composites
It is well known that a composite cathode, consisting of an electrocatalyst and an
oxygen ion conductive phase, enhances the performance significantly by effectively
extending the reaction zone from the electrode/electrolyte interface into the electrode.
Chemical compatibility between the two phases is critical in order to avoid the formation
of resistive tertiary phases. Composite cathodes based on La(Sr)Mn03.§-YSZ and
La(Sr)Mn03-s-GDC for YSZ electrolytes and La(Sr)Co(Fe)03-5-GDC for GDC
electrolytes have been very effective in reducing the polarization resistance compared to
single phase cathodes.53"20 The performance of the composite cathode depends on the
relative ratio, particle size distribution, and spacial distribution of the two phases, so as to
achieve high concentration of TPBs and percolation for both the phases. In addition, the
electrode microstructure should also be porous to provide for the gas diffusion. TPBs in
the electrode are the boundaries, between the electrocatalyst, the electrolyte, and the gas
phase, where the charge transfer reaction is believed to take place. High concentration of
TPBs (large number of reaction sites) and the percolation of the two phases (continuous
pathway for electron and oxygen ion transport) results in low electrode polarization.
In this part of the work, (BÍ2Ru207.3)-(BÍ203)o.8(Er203)o.2 composite cathodes were
studied. As shown in the phase diagrams in figure 4-1 and 4-2, the low temperature
bismuth ruthenate (BÍ2R.U2O7 3) is the thermodynamically stable phase with BÍ2O3 in the
temperature range of interest and hence, was chosen as the electrocatalyst phase in the
composite electrode. The electrolyte phase in the composite electrode was 20 mol% erbia
stabilized bismuth oxide (ESB), one of the highest known oxygen ion conductors.
BÍ2RU2O7 3 forms a eutectic with BÍ2O3 at relatively low temperatures, though there is
some discrepancy between the two proposed diagrams regarding the eutectic temperature

71
and composition. Prosychev et al.48 proposed that the eutectic is between BÍ2RU2O7 3 and
(X-BÍ2O3 at ~37 mol% Ru02. The eutectic is at 730 °C, which incidentally is also the
transformation temperature from a to 8-BÍ2O3. Hrovat et al.49 instead proposed that the
eutectic is between BÍ2RU2O7 3 and 5-BÍ2O3 at ~20 mol% RuÜ2. The eutectic is at 745 °C,
above a to 8-BÍ2O3 transformation temperature. The presence of the eutectic could
significantly undermine the stability of the electrode microstructure during processing
and/or under operation. Differential scanning calorimetry (DSC) studies were done to
study the eutectic temperature and composition.
4.2 Experimental Work
Standard solid state synthesis was used to fabricate the bismuth ruthenate powders.
Bi203 (99.9995%, Alfa Aesar), Ru02.XH20 (99.99%, Alfa Aesar), CaC03 (99%, Fisher
Scientific), Sr(N03)2 (99.97%, Alfa Aesar) and Ag (99.9%, Alfa Aesar) powders were
mixed in stoichiometric amounts in an agate mortar and pestle, and calcined at 900 °C for
10 hours to achieve undoped and doped BÍ2RU2O7 pyrochlore phase. To achieve the low
temperature BÍ2RU2O73 phase, mixed powders were calcined at 775 °C for 10 hours. The
calcined powder was leached with dilute HNO3 to remove the sillenite type-impurity
phase. For the composite electrodes, (BÍ203)o.8(Er203)o.2 powders were prepared using
amorphous citrate route as described in the experimental section of chapter 3. For the
reactivity tests of BÍ2RU2O7 and BÍ2RU2O7.3 with the electrolyte, 11 mol% gadolinium
doped ceria powder (GDC, Rhodia) was mixed with the respective powders and heat
treated at 850 °C for 10 hours. DSC (TA-1290) studies were done to study the eutectic
between bismuth ruthenate and bismuth oxide. Powder mixtures of ~13 mg mass were
heated from room temperature to 800 °C at 20 °C/min in 100 seem of air.

72
Symmetrical cells for impedance spectroscopy studies were fabricated by brush
painting the electrode paste on GDC electrolyte pellets (uniaxially pressed and sintered at
1450 °C for 6 hours). The electrode paste was made by mixing the desired weight ratios
of the powders with Heraeus V006 binder. The electrodes were dried at 150 °C and
sintered at temperatures between 750 and 850 °C for 2 hours with Pt lead wires. The
samples were characterized using XRD (APD - 3720) and SEM (JEOL - 6400).
Impedance spectroscopy was done using a Solartron 1260 in the frequency range of 0.1
Hz -32 MHz at temperatures between 350-700 °C. Oxygen partial pressure was varied
between 0.04-1 atm by using Ü2/air-N2 gas mixtures.
4.3 Results and Discussion
4.3.1 Processing
XRD patterns of the calcined BÍ2RU2O7 and BÍ2RU2O7 3 powders before and after
leaching are shown in figure 4-4 and 4-5, respectively. Leaching with dilute HNO3 is
effective in removing the sillenite type-impurity phase (BÍ12RUO20), which results in
predominant single phase powders.
XRD patterns of the powder mixtures of GDC with BÍ2RU2O7 and BÍ2RU2O7 3 after
heat treatment at 850 °C for 10 hours are shown in figure 4-6 and 4-7, respectively. There
are no new peaks identifiable in the patterns, which indicate that there is no reaction
between GDC and either of the bismuth ruthenate phases. Similarly, Hrovat et al.56
studied the subsolidus phase equilibria in Ru02-BÍ203-Ce02 system and no ternary
compound was found at 800 °C. The tie line is between Bi2Ru207 and Cei-xBix02-^/2 (0<
x < 0.33) solid solution. These results are in contrast with those of Bae et al.,43 who found
unknown reaction products between Bi2Ru207 3 and GDC at 800 °C. It is possible that the
presence of sillenite type impurities in bismuth ruthenate could lead to a new product, as

73
Takeda et al.44 found between BÍ2R.U2O7 and YSZ where the presence of sillenite type
impurities resulted in the transformation to monoclinic zirconia.
Surface and cross-section micrographs of BÍ2R.U2O7 electrode on GDC electrolyte
are shown in figure 4-8 (a) and (b), respectively. Solid state synthesis of BÍ2R.U2O7
powders at 900 °C has resulted in a coarse particle size (~3 pm). After sintering at 850 °C,
the electrode is porous with a thickness of-100 pm. Initial sintering experiments showed
that lower sintering temperatures resulted in poor adhesion of the electrode film with the
substrate, while at higher temperatures there was evaporation of the electrode material.
4.3.2 Impedance Spectroscopy Studies
4.3.2.1 Undoped BÍ2RU2O7 cathode
Impedance plots of BÍ2RU2O7 electrode on GDC at 500 °C and 700 °C are shown in
figure 4-9 and 4-10, respectively. High frequency intercept corresponds to the bulk
conductivity of the GDC electrolyte, which is comparable to values reported in the
literature. The depressed semicircle at lower frequencies is due to the BÍ2R.U2O7 electrode
with resistance of 158.29 Q and 4.13 Q in air at 500 °C and 700 °C, respectively. The
electrode ASR was calculated by multiplying the electrode resistance by the electrode
area (-0.7 cm ) and dividing by 2 to account for the symmetric cell. The Arrhenius plot
of the electrode ASR (Dcm2) in air is shown in figure 4-11. ASR values are significantly
smaller than those reported by Linquette-Mailley et al.51 for Bi2Ru207 electrode on YSZ
electrolyte. For example, they reported ASR of-160 Dcm2 compared to 55.64 Dcm2 in
this study on GDC at 500 °C and —20 Qcm2 compared to 3.03 Qcm2 in this study at 650
°C. Also, they reported a kink in the Arrhenius plot with activation energy of-1.3 eV
below 567 °C and -1.0 eV above 567 °C, and suggested that there are two different rate
limiting steps in the electrode reaction in the temperature range of study. However, in this

74
study a single activation energy of ~1.26 eV in the temperature range of 450-700 °C was
observed.
To understand the mechanism of oxygen reduction at the electrode, impedance
measurements were done as a function of oxygen partial pressure. In general, ASR of the
electrode varies with the oxygen partial pressure according to the following equation
[ASR] = [ASR]0 (p02)m 4-1
where the magnitude of m provides an insight into the rate limiting step in the oxygen
reduction reaction at the electrode. With metal and metal oxide electrodes on solid
electrolytes: m = 0.25 has been associated with the charge transfer reaction at the TPBs,
m = 0.5 with the surface diffusion of the dissociatively adsorbed oxygen at the electrode
to the TPBs, and m = 1 with the gaseous diffusion of oxygen molecules in the electrode
structure. In addition to the electrode-electrolyte combination, the microstructure also
plays a key role in determining the rate limiting step.38,51'57'''9
Figure 4-12 is a graph of In [ASR] vs. In po2- The values of m range primarily
between 0.5 and 0.6 (with the exception of 700 °C), which suggests that the rate limiting
step in the present case is surface diffusion of the dissociatively adsorbed oxygen at the
electrode surface to the TPBs. In contrast, Linquette-Mailley et al.51 reported values of m
-0.25 at temperatures less than 567 °C and continuously increasing values for
temperatures greater than 567 °C for Bi2Ru207 electrode on YSZ electrolyte. The rate
limiting step at lower temperatures was attributed to the charge transfer step and at higher
temperatures a transition to surface diffusion was suggested.
The better performance of the Bi2Ru207 electrode on GDC electrolyte compared to
that on YSZ electrolyte along with a different rate limiting step again demonstrates the

75
role played by the electrolyte in the electrode polarization. For an electrode with
negligible ionic conductivity such as BÍ2R.U2O7 at low polarizations, the
electrode/electrolyte interfaces are the active TPBs for the electrode reaction to take
place. Hence, it can be expected that the bulk and surface properties of the electrolyte can
have a significant effect on the overall performance of the electrode. Higher ionic
conductivity of GDC compared to that of YSZ and the capability of cerium to exist in
multiple valence states to provide sufficient electronic charge carriers could result in
improved charge transfer and better electrode performance.60
4.3.2.2 Doped BÍ2R112O7 cathodes
To further improve the performance of BÍ2R.U2O7 cathodes, doping with aliovalent
cations was studied on Bi3+ site to introduce oxygen ion vacancies and hence, oxygen ion
conductivity in the structure. The presence of mixed ionic and electronic conduction
would activate the oxygen reduction reaction on the complete electrode surface and
would just not be limited to the TPBs at the electrode/electrolyte interface. Although,
BÍ2R.U2O7 forms solid solutions with a number of dopants, the choice was short listed to
cations (Ca2+, Sr2+, Ag+) with fixed valence and with close ionic radii match with Bi3+
host.
Impedance plots of Ca doped BÍ2R.U2O7 [(Bi].xCax)2Ru207.5 = BCRx, x = 5-30
mol%; BÍ2R.U2O7 = BR07] at 500 °C and 700 °C are shown in figure 4-13 and 4-14,
respectively. At 500 °C, all dopant levels led to increase in the electrode polarization,
while at 700 °C, 5 and 10 mol% doping resulted in decrease in the electrode polarization
compared to undoped bismuth ruthenate. With x > 20 mol%, the electrode polarization is
an order of magnitude higher at each temperature along with additional electrode arcs at
lower frequencies, which indicates that with high dopant levels there is drastic change in

76
the electrode properties leading to multiple rate limiting steps. Arrhenius plot of the
electrode ASR in air is shown in figure 4-15. Doping results in the increase in activation
energy from ~1.26 eV for undoped BÍ2RU2O7 to ~1.37 eV for Ca doped BÍ2RU2O7, though
5 and 10 mol% Ca doped BÍ2RU2O7 show better performance at higher temperatures.
Impedance plots of Ag doped BÍ2RU2O7 [(Bii.xAgx)2Ru207^ = BARx between x =
5-20 mol%] at 500 °C and 700 °C are shown in figure 4-16 and 4-17, respectively. The
performance was similar to that of Ca doped systems, with higher activation energy
compared to undoped BÍ2RU2O7 and better performance at higher temperatures with 5 and
10 mol% Ag doped BÍ2RU2O7 as shown in figure 4-18. Performance of 5 mol% Ca and
Ag doped systems are compared in figure 4-19. Both show better performance than
undoped BÍ2RU2O7 above -550 °C with similar activation energy of-1.37 eV. Possibly,
Ca is a better choice of dopant as it shows slightly better performance. 10 mol% Sr doped
BÍ2RU2O7 (BSR10) showed inferior performance than undoped, 10 mol% Ca and Ag
doped BÍ2RU2O7, as shown in figure 4-20 and hence was not studied further.
Electrode polarization of 5 mol% Ca and Ag doped BÍ2RU2O7 was studied as a
function of P02 and dc bias in order to understand the rate limiting steps with these
electrodes. Impedance plots for 5 mol% Ca doped system are shown in figure 4-21 and 4-
22, while for 5 mol% Ag doped system are shown in figure 4-23 and 4-24. Plots of
ln(ASR) vs. ln(p02) at temperatures between 400-700 °C for 5 mol% Ca and Ag doped
systems are shown in figure 4-25 and 4-26, respectively. The value of m for 5 mol % Ca
doped BÍ2RU2O7 ranged between 0.6 and 0.8 at 0.11 < P02 (atm.) < 1 and at a lower P02 of
0.04 atm., the value of m increased indicating a change in the rate limiting step. On the
other hand for 5 mol% Ag doped BÍ2RU2O7, the value of m ranged consistently between

77
0.5 and 0.6 in the complete p02 range similar to that of undoped Bi2Ru207, suggesting
that the rate limiting step is the surface diffusion of dissociatively adsorbed oxygen at the
electrode surface to the TPBs. Initial impedance measurements under direct current bias
on 5 mol% Ca and Ag doped Bi2Ru207 also support the argument at least at high
temperatures. Impedance plots for 5 mol% Ca doped system as a function of current bias
at 500 °C and 700 °C are shown in figure 4-27 and 4-28, respectively. At 500 °C, on the
application of current bias the electrode polarization decreased which indicated that the
rate limiting step is charge transfer. The electrode polarization behavior was more
complex at 700 °C; electrode polarization increased with current bias up to 15 mA,
decreased slightly between 15-25 mA, and then increased beyond 25 mA along with
additional electrode arcs. At higher temperatures, the charge transfer is fast, and hence, a
comparatively slower step in the electrode reaction could be rate limiting, which in the
present case is diffusion related.
Additional information of the rate limiting steps in the electrode reaction is given
by the frequency distribution of the imaginary part of the impedance. At a particular
temperature, each reaction step is associated with a characteristic frequency and therefore
the dominant reaction steps can be identified if their characteristic frequencies could be
resolved in the frequency distribution. Sometimes it is adequate as in the case of 20 and
30 mol% Ca doped Bi2Ru2C>7, and 5 mol% Ca doped Bi2Ru2C>7 under different bias
currents, where additional humps in the electrode impedance appear indicating the
presence of additional steps in the electrode reaction. However, sometimes the frequency
distribution information of imaginary part is inadequate to resolve if the characteristic
frequencies of two steps are too close.

78
As shown in figure 4-9, 4-10, 4-21 - 4-24, the characteristic frequency of the
electrode impedance for undoped, 5 mol% Ca and Ag doped BÍ2R.U2O7 in general
increases with P02 indicating a change in dominant rate limiting step. Possibly at lower
P02, multiple reaction steps are present which determine the larger total impedance. This
phenomenon is more clear in the case of 5 mol% Ca doped BÍ2R.U2O7, where there is
increase in the value of m at lower po2- This in effect again reflects on the inadequacies
of using the magnitude of m and frequency distribution of imaginary part of impedance to
determine the rate limiting steps present in the electrode reaction. In general it is not
reliably possible to deconvolute multiple humps present in the electrode response and
hence the magnitude of ASR and m are average of the different reaction steps taking
place over the P02 range of study, whereas the frequency distribution of imaginary part
may also not provide adequate resolution and identification of individual steps. New
methodologies are being adopted to extract more information from the frequency
distribution of the impedance. In one strategy adopted by Ivers-Tiffee and co-
workers, ’ the fourier transformation of the imaginary part of the impedance have been
shown to provide much better resolution of the reaction steps present.
Doped BÍ2RU2O7 pyrochlore did not improve the electrode performance as
significantly as observed by Bae et al.43 with their studies on 5 mol% Sr doped Y2RU2O7.
It is difficult to explain, without speculating, the observed results due to lack of
understanding of the basic defect chemistry; the effect of doping and P02 on the structure
and electrical properties of pyrochlore ruthenates. Studies concerning the same are
currently underway in our lab and it is hoped that they will shed more light on the
performance of these electrodes. However, one thing is clear that doping with aliovalent

79
cations on the Bi3+ site is not a very effective strategy to improve the performance of
BÍ2R.U2O7 pyrochlore cathodes.
4.3.2.3 BÍ2Ru207.3-(BÍ203)o.8(Er203)o.2 composite cathodes
From the studies done on undoped and doped BÍ2R.U2O7, it is clear that the surface
diffusion of dissociatively adsorbed oxygen at the electrode surface to the TPBs is one of
the rate limiting steps involved in the oxygen reduction reaction at the electrode. In this
part of the work, composite cathodes consisting of BÍ2R.U2O7 3 and (BÍ203)o.8(Er203)o.2
were evaluated.
DSC plots in the heating cycle for powder mixtures of BÍ2R.U2O7 3 with monoclinic
a-BÍ203 and fee fluorite (Bi203)o 8^203)0.2 are shown in figure 4-29. The behavior of the
two powder mixtures is apparently different; BÍ2R.U2O7 3 and 01-BÍ2O3 showed an
endotherm at -735 °C, while BÍ2R.U2O7 3 and ESB did not show any endotherm. One
possible explanation for this behavior could be that the eutectic composition is between
BÍ2R.U2O7.3 and (X-BÍ2O3, as proposed by Proshychev et al. However, this is not
conclusive as the heating rates used in the present study were high, which may not allow
the system to equilibrate.
Sintering temperatures were limited to 750 °C in order to avoid the transformation
of BÍ2R.U2O7.3 to the high temperature pyrochlore phase. With this limitation, the
electrode posed significant problems in terms of sinterability with the GDC electrolyte. A
number of trials ended in failure to prepare the single phase electrode with good adhesion
with the electrolyte. Composite cathodes were however prepared successfully, possibly
due to the high sinterability of the bismuth oxide phase. SEM surface micrographs of one
such composite cathode are shown in figure 4-30. Linquette-Mailley et al.51 were able to
deposit single phase BÍ2R.U2O7.3 electrodes using a pyrosol deposition technique, though

80
with poor adhesion with the substrate. The performance of the BÍ2R.U2O7.3 electrode was
similar to that of BÍ2R.U2O7 pyrochlore phase with comparable activation energy and
frequency distribution/1 Hence, it was felt appropriate in this study to compare the
performance of the composite electrodes with that of the pyrochlore single phase
electrode.
Impedance plots of the composite cathodes [BÍ2R.U2O7 3:(BÍ203)o.g(Er203)o.2 = X:Y
wt. ratio = X:Y] at 500 °C and 700 °C are shown in figure 4-31 and 4-32, respectively.
Performance of the BÍ2R.U2O7 pyrochlore phase is also shown for comparison.
Introduction of the ESB phase in the electrode resulted in the reduction of the electrode
impedance at each temperature along with a lower characteristic frequency, which
indicated a change in the rate limiting step for the oxygen reduction reaction in the
composite electrodes compared to that of single phase pyrochlore electrode. As shown in
figure 4-33, the three composite compositions showed better performance than the
pyrochlore cathode at each temperature, with best performance of 18.4 Qcm at 500 °C
and 0.32 Qcm2 at 700 °C with a composite electrode containing 37.5 wt% ESB.
Arrhenius plots of the composite electrodes and single phase pyrochlore electrode is
shown in figure 4-34. Among the three compositions of the composite electrode, the
activation energy of the electrode reaction decreases with the addition of the ionic phase.
This trend is similar to those observed for other composite electrodes,53'55 where the
addition of ionic phase leads to improved ionic conductivity of the composite electrode
and decreases the activation energy. The proposed electrode reaction mechanism for the
bismuth ruthenate electrodes is shown in figure 4-35. The presence of the ESB in the
electrode will not only increase the concentration of TPBs and ionic conductivity of the

81
electrode, but possible also shortens the rate limiting surface diffusion path to the TPBs
and hence, improves the electrode performance.
The performance of the composite electrode is highly promising and compares well
with established cathode materials. Further optimization of the electrode morphology in
terms of relative ratio, particle size and spatial distribution of the two phases will realize
high performance cathodes for IT-SOFCs.
4.4 Conclusion
Bismuth ruthenate based electrodes were evaluated as a prospective cathode
material for IT-SOFCs based on ceria electrolytes. No reaction product was found
between bismuth ruthenate and GDC electrolyte after heat treating at 850 °C for 10 hours.
Bismuth ruthenate pyrochlore electrodes showed ASR values of 55.64 Gem at 500 °C
and 1.45 Gem at 700 °C in air. The surface diffusion of dissociatively adsorbed oxygen
at the electrode to TPBs is believed to be the rate limiting step. Doping with lower valent
cations on Bi3+ site, to improve the ionic conductivity, was not found to be very effective
the improving the electrode performance. 5 mol% Ca and Ag doping were found to only
slightly improve the electrode performance over undoped bismuth ruthenate pyrochlores.
Rate limiting step for the oxygen reduction reaction in 5 mol% Ag doped bismuth
ruthenate electrode was found to be the surface diffusion of dissociatively adsorbed
oxygen, while for 5 mol% Ca doped system multiple rate limiting steps were observed in
the oxygen partial pressure range of study. Composite electrodes consisting of BÍ2RU2O7 3
and (Bi203)o.8(Er203)o.2 showed 3-4 times better performance than BÍ2RU2O7 pyrochlore
electrodes. Best performance of 18.4 Gem2 at 500 °C and 0.32 Gem2 at 700 °C in air was
observed with a composition containing 37.5 wt% (Bi2O3)0 8(Er203)o.2- The performance
is very promising, and it is expected that further optimization in electrode composition

82
and microstructure will result in high performance cathodes for IT-SOFCs. Lastly, a need
was felt to develop new methodologies which could more effectively distinguish the
different rate limiting steps in the electrode impedance.

83
Figure 4-1. Phase diagram between BÍ2O3 and RuC>2 proposed by Prosychev et al. (from
ref. 50)
Figure 4-2. Phase diagram between BÍ2O3 and R.UO2 proposed by Hrovat et al. (from ref.
51)

84
48(f)
Figure 4-3. Cation sublattice for one-quarter unit cell of the pyrochlore structure. Two of
the 48i-anions and one of the 8b-sites are also shown for 34a (from ref. 8).

85
2000
1500
00
H
Z1000
o
u
500
0
20 30 40 50 60 70
20
Figure 4-4. XRD patterns for BÍ2RU2O7 after calcination at 900 °C for 10 hours (a) before
and (b) after leaching.
1400
1200
1000
H 800
z
o 600
U
400
200
0
20 30 40 50 60 70
20
Figure 4-5. XRD patterns for BÍ2RU2O7 3 after calcination at 775 °C for 10 hours (a)
before and (b) after leaching.
*
* - Bi RuO
12 20
i ^ id w v IA) Jwvw' !Ui w; h1W * ‘' A/V
id
* *
* V ^i yl»'W W‘J■'r>'lAtfw
â–  !
*
-Bi
12
RuO
20
v'|»Amw.
' ; 'I' I “ ’ Sf!
<**"“J*‘*rf W wVv'w*/

86
üO
H
O
U
1500
1000
500
0
20 30 40 50 60 70
20
Figure 4-6. XRD pattern for BÍ2R.U2O7 and GDC powder mixture after heat treatment at
850 °C for 10 hours.
700
600
500
H 400
O 300
u
200
100
0
20 30 40 50 60 70
20
Figure 4-7. XRD pattern for BÍ2RU2O7 3 and GDC powder mixture after heat treatment at
850 °C for 10 hours.
* Bi Ru O +GDC (*)
2 2 7.3 v '
r

87
(b)
Figure 4-8. Surface (a) and cross-section (b) micrographs of BÍ2R.U2O7 electrode on GDC
electrolyte.

88
-300
-250
500 °C
50 100 150 200 250 300 350 400
Z' (Q)
(a)
(b) (c)
Figure 4-9. Impedance plots of BÍ2R.U2O7 electrode on GDC at 500 °C (a) imaginary vs.
real impedance (b) and (c) imaginary impedance vs. frequency.

Z” (O) Z" (fi)
89
G
9 10 11 12 13 14 15 16
Z' (Q)
-20
-15
-10
-5
0
5
__t I . ..... -1_1 . . . 1 ]— 1 I L 1 1 1 ».. i J 1 . I I 1
5 10 15 20 25 30
Z’ (Q)
(a)
(b)
(c) (d)
Figure 4-10. Impedance plots of BÍ2R.U2O7 electrode on GDC at 700 °C (a) and (b)
imaginary vs. real impedance (c) and (d) imaginary impedance vs. frequency.

90
1000/T (K )
Figure 4-11. Arrhenius plot of the BÍ2RU2O7 electrode ASR (ilcm2) in air
8 400 °C (0.56)
' n
6 450 °C (0.^3)
r -
p
^ ' r
500 °C (0.50)
•
on
4 550 °C (0.50)
â–¼
, *•
c
'c
600 °C (0.5*5)
A
A T T
2 650 °C (0.61)
A
A
700 °C (0.79)
0
r1
‘!>v ' ^
T(m)
-O
?
1 1
-5 -4 -3
-2
-1 0 1
ln(p )
vi02y
Figure 4-12. ln(ASR) vs. ln(po2) of BÍ2RU2O7 electrode at different temperatures with m
in parenthesis T(m).

Z" (Q)
91
-250
-200
-150
-100
-50
0
500 °C
' ■, ■ 1 ! ' i ■' ■ '—r -> ■ - ■
BCR20 a
A
A '
A v G
BCR30
A v
A v
A
A v
A
A
Jr
BR07^J^BCR5
-3000 500 °C
-2000
50 100 150 200 250 300
Z' (Q)
(a)
â– 1500
-1000
4
G
-500
0.1
A
A
'A
â– A
VA
A
"’a
VA
v A
V A
VA
• BR07
â–  BCR5
♦ BCR10
BCR20
BCR30
(b)
• BR07
â–  BCR5
♦ BCR10
a BCR20
BCR30
10 100
f (Hz)
(c)
10 100 1000
f (Hz)
(d)
Figure 4-13. Impedance plots of Ca doped BÍ2RU2O7 electrodes on GDC at 500 °C (a)
and (b) imaginary vs. real impedance (c) and (d) imaginary impedance vs.
frequency.

Z" (Q) Z" (Q)
92
-5 • BR07
â–  BCR5
♦ BCR 10
a BCR20
BCR30
A
700 °C
A
-4
-3
-2
-1
0
A
A
A v
A v
A7
£
10 11
12 13
Z'(Q)
(a)
14
Z' (Q)
700 °C
-1.5
(b)
• BR07
â–  BCR5
♦ BCR10
BCR20
a BCR30
-0.5
10 100 1000 10
f (Hz)
10 100 1000 10 10J
f (Hz)
(c)
(d)
Figure 4-14. Impedance plots of Ca doped BÍ2RU2O7 electrodes on GDC at 700 °C (a)
and (b) imaginary vs. real impedance (c) and (d) imaginary impedance vs.
frequency.

93
1000
£ 100
o
a
üO
<
10
t
/y
1.1
ft
t
l
♦
-A-
E
BR07 (1.26eV)
BCR5 (1.37eV)
BCR10 (1.36eV)
BCR20 (1.37eV)
BCR30 (1.40eV)
1.2
1.3
1.4
1000/T (K’1)
Figure 4-15. Arrhenius plot of Ca doped BÍ2RU2O7 electrode ASR (Í2cm2) in air.

94
-100
L
-80
~ -60
a
^ -40
-20
Z' (Q)
(a)
500 °C/
A
A
A.
BR07
BAR5
BARIO
* BAR20
-ffR
â–²
â–²
"tfc,.
â–¡ X
ck
rP
H
jy
*
A
<

)r,V>J ‘ ■
0
0.1 1
â–²
% i
ft
10 100 1000 10H
f (Hz)
(b)
Figure 4-16. Impedance plots of Ag doped Bi2Ru207 electrodes on GDC at 500 °C (a)
imaginary vs. real impedance (b) imaginary impedance vs. frequency.

95
a
-6
-4
-2
0
2
700 °C
BR07
BAR5
BARIO
* BAR20
10
-3
12 14
Z' (O)
(a)
16
(b)
Figure 4-17. Impedance plots of Ag doped BÍ2R.U2O7 electrodes on GDC at 700 °C (a)
imaginary vs. real impedance (b) imaginary impedance vs. frequency.

96
100
£
o
a
c¿
C/3
<
10
A
A
A
//
'//
¿/
¿I • BR07 (1.26eV)
â– n- BAR5 (1.38eV)
♦ BAR10(1.32eV)
BAR20(1.40 eV)
1.1 1.2 1.3 1.4
1000/T (K’1)
Figure 4-18. Arrhenius plot of Ag doped BÍ2RU2O7 electrode ASR (Qcm2) in air.

97
100
E
o
q
cC
oo
<
10
• BR07 (1.26eV)
BCR5 (1.37eV)
♦ BAR5 (1.38eV)
//
y
'/
//
Ss
//
1
/,
/ v
_L_i_j_i j. .1. i., i.. .—i—
1 1.1 1.2 1.3 1.4
1000/T (K'1)
Figure 4-19. Arrhenius plot of undoped, 5 mol% Ca and Ag doped BÍ2R.U2O7 electrode
ASR (Qcm2) in air.
A
100
♦ BR07 (1.26eV)
BCR10 (1.36eV)
♦ BARIO (1.32eV)
A,
E
u
q
00
<
BSR10 (1.20eV)
f a' /
y
* //
A'
10
A
w
//
,/
A'
1
1 1.05 1.1 1.15 1.2 1.25 1.3 1.35 1.4
1000/T (K'1)
Figure 4-20. Arrhenius plot of undoped, 10 mol% Ca, Ag and Sr doped BÍ2RU2O7
electrode ASR (ficm2) in air.

98
-150
-100
0.01 atm.
a
50
100
150
Z'(Q)
(a)
500 °C
-2500
-2000
-1500
0.21 atm.
\
%
200
250
f ' ' ’ ' l
500 °C
0.01 atm.
0 500 1000 1500 2000 2500 3000
Z' (Q)
(b)
(c) (d)
Figure 4-21. Impedance plots of 5 mol% Ca doped BÍ2R.U2O7 electrodes on GDC at 500
°C (a) and (b) imaginary vs. real impedance (c) and (d) imaginary impedance
vs. frequency.

99
-120
-100
700 °C
20 40 60 80 100 120 140
Z’ (Q)
(b)
(c)
(d)
Figure 4-22. Impedance plots of 5 mol% Ca doped BÍ2R.U2O7 electrodes on GDC at 700
°C (a) and (b) imaginary vs. real impedance (c) and (d) imaginary impedance
vs. frequency.

100
-300
-250
-200
g-150
^ -100
-50
0
500 °C
V
8 Hz
0.01 atm.
1 atm.
0.61 atm.
0.21 atm.
\
50
50 100 150 200 250 300 350 400
Z' (O)
-400
-350
-300
-250
S -200
-150
-100
-50
(a)
T â–¼
â–¼
V
N
500 °C
1 atm.
0.61 atm.
0.21 atm.
A 0.01 atm.
T -10° atm.
10 100 1000 10
f (Hz)
(b)
Figure 4-23. Impedance plots of 5 mol% Ag doped Bi2Ru207 electrodes on GDC at 500
°C (a) imaginary vs. real impedance (b) imaginary impedance vs. frequency.

Z" (Q)
101
Z' (Q)
(a)
f (Hz)
(b)
-20
1 VV
T
â–¼
700 °C
-15
T~10"5 atm.
â–¼
â–¼
(c)
Figure 4-24. Impedance plots of 5 mol% Ag doped BÍ2R.U2O7 electrodes on GDC at 700
°C (a) imaginary vs. real impedance (b) and (c) imaginary impedance vs.
frequency.

102
-4-3-2-10 1 2 3
ln(p )
V102y
Figure 4-25. ln(ASR) vs. ln(po2)of 5 mol% Ca doped BÍ2R.U2O7 electrode at different
temperatures with m in parenthesis T(m).
8 400 °C (0.62)
“T r
1 • '
' • 1 ! 1 ■ !
r
E
450 °C (0.59)
6 •
n
w
-
500 °C (0.57)
•
•
E -
00
4 550 °C (0.54)
T
â–¼
•
600 °C (0.54)
â–²
â–²
â–¼
â–¼
2 650 °C (0.60)
700 °C (0.69)
- k
A
â–²
' J-L-
" O _
0
T (m)
1 â–  .
1 1 â–  â–  â– 
ln(p )
vi02y
Figure 4-26. ln(ASR) vs. ln(po2)of 5 mol% Ag doped BÍ2RU2O7 electrode at different
temperatures with m in parenthesis T(m).

103
a
-60
-50
-40
-30
-20
-10
0
10
aindnnnDD[
íT*S\
10 mA 7.5 mA
a5 mA
500 °C
0 mA
2.5 mA
60 70 80 90 100 110 120 130
Z’ (Q)
(a)
-20
I ,'3-ti
0 mA
500 °C
Figure 4-27. Impedance plots of 5 mol% Ca doped BÍ2R.U2O7 electrodes on GDC at 500
°C (a) imaginary vs. real impedance (b) imaginary impedance vs. frequency.

104
14 16
Z’ (Q)
(a)
0 mA
0 25 mA
’ □
5 mA
« 30 mA
-4
~ O
10 mA
a 35 mA
â–²
15 mA
40 mA
a
-3
â–¼
20 mA
<5%
O rftj,
0 tm
-2
0
0
i3 Ob
□“ Op
On
700 °C
0.1 1 10 100 1000 1
Z’ (Q)
(b)
Figure 4-28. Impedance plots of 5 mol% Ca doped BÍ2R.U2O7 electrodes on GDC at 700
°C (a) imaginary vs. real impedance (b) imaginary impedance vs. frequency.

Heat Flow (mW)
105
-10
-15
-20
-25
-30
-35
-40
400 500 600 700 800
Temperature (°C)
Figure 4-29. DSC plots of powder mixture of BÍ2R.U2O7 3 with monoclinic 01-BÍ2O3 and
fee fluorite (BÍ203)o.8(Er203)o.2-

106
Figure 4-30. SEM surface micrographs of bismuth ruthenate-bismuth oxide composite
electrode in (a) secondary electron mode and (b) back scattered mode.

107
60 100 140 180 220
Z' (O)
(a)
-60 , .
f (Hz)
(b)
Figure 4-31. Impedance plots of BÍ2RU2O7 3-(Bi203)o 8^203)0.2 composite electrodes on
GDC at 500 °C (a) imaginary vs. real impedance (b) imaginary impedance vs.
frequency.

108
-5
-4
^ -3
g
^ -2
-1
0
13 14 15 16 17 18
Z' (Q)
(a)
-1.5 I —-n
g
-0.
Figure 4-32. Impedance plots of BÍ2R.U2O7 3-(BÍ203)o.8(Er2C>3)o.2 composite electrodes on
GDC at 500 °C (a) imaginary vs. real impedance (b) imaginary impedance vs.
frequency.
A
A
A
A
A
A
A
A
A
A
A
A
700 °C
100 1000 10
f (Hz)
(b)
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110
(a)
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Figure 4-35 Proposed electrode reaction mechanism for bismuth ruthenate electrodes.

CHAPTER 5
ANODE SUPPORTED THICK FILM CERIA ELECTROLYTE UNIT CELLS FOR IT-
SOFC
5.1 Introduction
State of the art SOFCs operating at 750-1000 °C employ yttria stabilized zirconia
(YSZ) as the electrolyte material. High operation temperatures are required to overcome
the resistive loss across the YSZ electrolyte. There is considerable driving force to reduce
the operating temperatures of SOFCs to 500-750 °C; the advantages include use of cheap
ferritic stainless steel as the interconnect material, lower operating cost, and faster startup
• • AT AA
time for mobile applications.
Towards reducing the operating temperatures of SOFCs, two approaches have been
actively pursued. The first approach is towards finding better electrolyte (with high ionic
conductivity and high ionic transference number) and better electrode (with high catalytic
activity and high mixed conductivity) materials. At the same temperature, doped ceria
shows oxygen ion conductivity which is significantly higher than that of YSZ.41
However, doped ceria is thermodynamically unstable in the anode atmosphere of the fuel
cell due to the reduction from Ce4+ to Ce3 valence state. The polaronic charge transfer on
the mixed valent Ce4+/Ce3+ array results in electronic conduction inside the electrolyte,
which under fuel cell conditions results in an internal short circuit and a lower available
power than theoretical.8 It has been shown that this internal short circuit loss across
doped ceria decreases on lowering the operating temperatures, due to a wider electrolytic
111

112
domain, which makes doped ceria one of the most promising electrolyte materials for IT-
SOFCs.66
The second approach towards reducing the operating temperatures of SOFCs is by
reducing the thickness of the electrolyte and attaining higher powder densities at lower
temperatures with electrode supported electrolyte unit cells.67 For processing of thick
films on porous supports, the colloidal route has advantages over other film deposition
techniques like vapor deposition (CVD/EVD) and chemical routes (sol gel) in its
simplicity, cost effectiveness, flexibility (thickness ranging from 10-100 pm can be
deposited), and upscalibility.68 Also, exceptional performance has been reported for
SOFCs fabricated from colloidal deposition technique by de Souza et al.69,70 and Kim et
al.71 using a 10 pm thick YSZ electrolyte.
In this chapter, fabrication and performance results of anode supported thick film
ceria electrolyte unit cells are reported. A colloidal dip/spray coating method was used to
fabricate the ceria films. Dense, thick films of ceria on porous anodes was achieved by
varying the pre-sintering temperature of the anode substrate and final sintering
temperature of the anode/electrolyte bilayer, which results in an optimum condition
where the magnitude of shrinkage in the anode substrate is similar to that of the
electrolyte film.
Candidate materials for cathodes should be mixed ionic electronic conductors
(MIEC) with high catalytic activity for oxygen molecule dissociation and oxygen
reduction. The material should be compatible with other cell components with respect to
chemical reaction and thermal expansion coefficient (TEC). Oxides with perovskite
structure (ABO3) have been the material of choice for the cathodes; where A is a rare

113
earth element and B is a transition metal (Fe, Mn, Co). On doping the A site with alkaline
earth cations (Sr2+, Ca2 ' ), the charge compensation occurs by valence change of the
transition metal cations and under certain conditions by oxygen vacancy formation
resulting in an MIEC.72
La(Sr)Mn03 (LSM)-YSZ composites have been the material of choice for YSZ
based SOFCs. The oxygen reduction happens at the triple phase boundaries (TPBs)
between the electrocatalyst, the electrolyte and the gas phase. Furthermore, the TPBs are
active only if electrons, oxygen ions and oxygen gas can transport to or away from the
TPBs. Hence, the performance of the composite cathode depends critically on the relative
ratio, particle size, and spatial distribution of the two phases, so as to achieve high
concentration of TPBs and percolation for both the phases. It is often found that the high
polarization losses are primarily associated with inadequate ionic transport within the
electrode structure. Hence, at lower temperatures LSM (a poor ionic conductor) is
replaced by La(Sr)Co(Fe)03_5 (LSCF) or Sr(Sm)Co03 (SSC), which have considerable
ionic conductivities.73'74 YSZ is replaced by doped ceria, with higher oxygen ion
conductivity, as the electrolyte phase to develop better performance LSCF-GDC
cathodes.53'55
LSCF has been extensively studied as a candidate cathode material for SOFCs
based on ceria electrolytes.75,76 LSCF is a good electronic conductors (e.g.
Lao.6Sro.4Coo.2Feo 8O3 - 300S/cm at 750 °C) and also shows fast oxygen surface exchange
with higher oxygen ion conductivity than that of LSM (e.g. LaoeSro^Coo^FeosCb - 10'3
S/cm at 750 °C).55 They also appear to be chemically stable with ceria electrolytes as the
pyrochlore compound La2Ce2C>7 does not exist.60 Among the family, the Fe rich

114
compositions are more attractive than Co rich compositions, as they have lower TEC and
hence are better matched with ceria electrolytes.75 Lao.6Sro.4Coo.2Feo 803-Cei-xGdx02-s
composite cathodes have been studied by Dusastre et al. and Murray et al. and were
found to be very effective in improving the performance over single phase
Lao.6Sro.4Coo.2Feo 8O3 cathodes. Optimum composition was found to be between 30-50
wt% Ceo.9Gdo.i02-s with area specific resistance (ASR) values of 0.33-0.6 Qcm2 at 600
°C. In this work, both single phase Lao.6Sro.4Coo.2Feo.8O3 and composite
Lao.6Sro.4Coo.2Feo.803-Ceo.89Gdo 11O2-8 (70:30 wt. ratio) cathodes were studied on anode
supported ceria unit cells. One cathode composition based on Lao.6Sro.4Cuo.1Feo.9O3-
(Bi203)o.8(Er203)o.2 (70:30 wt. ratio) was also studied. Ag-(BÍ203)o.75(Er2s03)o.2 and Ag-
BIMEVOX cermet cathodes have been also reported for low temperature
applications.34,77,78 Ag-(BÍ203)o.8(Er203)o.2 cermets was studied as a prospective cathode
material, but instability issues were encountered as mentioned in chapter 3.
Ni based cermet composites have been adopted as the anode material by most of
the SOFC teams. Ni as the metal component satisfies the major requirements for the
electrochemical oxidation of H2 and CO fuels. Addition of the ceramic component makes
TEC of the composite comparable to that of the electrolyte.72 For lower operating
temperatures, YSZ has been replaced by doped ceria as the ceramic component of the
anode. Doped ceria is a MIEC under the reducing conditions at the anode, and the
reaction zone is not restricted to the electrocatalyst/electrolyte interface which results in
enhanced performance. Chemical composition and microstructure play important roles in
the performance of the anodes as shown by Ohara et al.79 They studied the performance
of Ni-samaria doped ceria cermet anodes as a function of Ni content and found that a

115
cermet with a Ni content of around 50 vol% showed the lowest anodic polarization (~30
mV at 300 mA/cm2, 800 °C). Thus, in this work NiO-Ceo 89Gdo,ii02-s anodes were
fabricated with Ni content of 50 vol%. Reduction of NiO to Ni was accomplished in situ
in the cell which would ideally generate ~26 % porosity in the structure. Initial cells were
fabricated using relatively coarse NiO powders from Alfa Aesar, and upon realization
that coarse NiO powders may result in incomplete reduction to Ni,80 finer NiO powders
from J. T. Baker were used in later studies.
5.2 Experimental
Ceo 89Gdo.ii02-s (GDC, d50= 0.64 pm, SA = 24 m2/gm, Rhodia) and NiO powders
(-325 #, Alfa-Aeaser or ~3 pm, J. T. Baker) were ball milled in ethanol for 24 hours with
zirconia ball media. The milled NiO-GDC powders were dried, sieved with 140 #
stainless steel mesh, and pressed uniaxially in 1.25” or 1.5” die. The green bodies were
then pre-sintered at temperatures between 500 °C and 1100 °C. For the colloidal
suspensions, 5 gm of GDC powder was dispersed in 100 ml of isoproponal using an
ultrasonic bath. The GDC electrolyte layer was deposited on the pre-sintered NiO-GDC
discs by dip/spray coating and then finally sintered at temperatures between 1350 °C and
1650 °C. In the dip-coating process, the anode substrate was placed in a sample holder
and dipped in the colloidal sol. Multiple coatings of the sol were done to get thicker
films. In the spray-coating process, the sol was fed to a commercial spray gun and
sprayed on the anode substrates with nitrogen as carrier gas. Some anode substrates were
reduced in H2 atmosphere at 800 °C for 5 hours to measure the anode density after
reduction, using the Archimede’s principle.

116
Lao 6Sro.4Coo.2Feo 8O3 (LSCF, d50= 0.7 pm, SA = 6 m2/gm, Praxiar Specialty
Chemicals), Lao 6Sro,4Cuo iFeooCb (LSCuF, SA = 5 m2/gm from NexTech Materials),
GDC powders (Rhodia), (BÍ203)o.8(Er203)o.2 (ESB, made in house using amorphous
citrate route as described in chapter 3), and Ag2Ü (Alfa Aesar) powders were ball milled
in desired weight ratios with Heraeus V006 binder for 24 hours to form the cathode paste.
The paste was brush painted on the GDC electrolyte film and sintered to form the cathode
layer. Finally, Pt/Ag mesh was attached to the anode with Pt paste and to the cathode
with the cathode paste, and sintered to act as current collectors.
The samples were characterized using XRD (APD-3720) and SEM (JEOL-6400) in
surface and cross-sectional view. The percentage porosity in the GDC films was
estimated by stereological point counting method using a 25 point grid.
Unit cells were sealed on an alumina tube with Aremco cement (516 & 571), as
described by de Souza et al.70 and gold lead wires were attached to the cell. The cells
were tested with air/oxygen at the cathode side and H2 with 3% H2O at the anode side at
flow rates of 10-30 seem between 450 °C and 750 °C. The flow rate at the anode and the
cathode was kept equal. I-V characteristics of unit cells were measured using a Solatron
electrochemical interface (SI 1287). Current interrupt technique was used to separate the
ohmic and non-ohmic polarizations of the cell.
5.3 Results and Discussion
5.3.1 Initial Fabrication and Performance Results
For application in SOFCs, the electrodes need to be sufficiently porous to allow for
the gaseous diffusion, and the electrolyte needs to be highly dense to act as a barrier for
the fuel and the oxidant at the two electrodes. Surface microstructures of the dip-coated
ceria electrolyte films on anode substrates (containing NiO fom Alfa Aesar) under

117
different pre-sintering conditions (800-1100 °C for 4 hours) and different final sintering
conditions (1600 °C for 6 hours and 1650 °C for 10 hours) are shown in figure 5-1 and 5-
2. Figure 5-3 shows the percentage porosity in the GDC film under these sintering
profiles. For each condition, the volume fraction of the porosity was estimated by the
mean fraction of the points which lie on the pores; the error bars correspond to the
standard deviation of the mean. Varying the pre-sintering temperature led to minima in
the porosity of the GDC film at 850 °C for both the final sintering temperatures of 1600
°C and 1650 °C with mean values of 1.97 % and 1.67 %, respectively. At higher pre¬
sintering temperature, the non-shrinking nature of the anode substrate leads to
constrained sintering of the film and to a highly porous microstructure. At -850 °C, the
shrinkage in the substrate is similar to that of the film which resulted in a comparatively
denser film. Final sintering at 1650 °C led to excessive warpage of the bilayer and to
coarsening of the microstructure as shown in figure 5-2. Both grain and pore size
increased, and became almost twice the size of that sintered at 1600 °C. Bigger pore size
in the thick film electrolyte microstructure will have greater chances of failure in
partitioning fuel and air atmospheres. Also thermodynamic studies on film stability have
shown that for the film to be stable, the film thickness should be larger than the grain size
of the film. Otherwise, the film would lower its energy by breaking into isolated islands
and exposing the substrate. The sampling done to measure the porosity in the film may
not be ideal as the surface porosity of the film might be different from that of the bulk.
However, it was assumed that the trends in the surface density of the film will be similar
to that of the bulk density. Cross-sectional micrographs of the films seem to support the
argument as can be seen in figure 5-4. Efforts to reduce the final sintering temperature to

118
avoid the co-sintering of the anode resulted in a highly porous film, as shown in figure 5-
5 for the film pre-sintered at 850 °C and finally sintered at 1400 °C. Hence, it was decided
to work with samples pre-sintered at 850 °C and finally sintered at 1600 °C for I-V
measurements in a fuel cell configuration.
The density of the anode substrate after sintering at 1600 °C was measured to be
6.72 gm/cc (~97 % of the theoretical density). On exposure to H2 atmosphere at 800 °C,
NiO in the anode reduced to Ni resulting in an open porosity of ~20 % and a total
porosity of ~29 % in the structure. Surface and cross-sectional micrographs of the anode
after reduction in back-scattered electron mode are shown in figure 5-6 (a) and (b),
respectively. Sintering at 1600 °C has led to a coarsened anode microstructure with low
concentration of TPBs and low porosity. The porosity at the surface is enclosed within Ni
grains, while in the cross-section it appears to be more uniformly distributed. There are
also concerns regarding the incomplete reduction of coarse NiO grains to Ni . The
reduced anode, however, was conductive at room temperature, and XRD patterns did not
show any peaks for NiO within the detection limits of the instrument. Further
improvement may be possible by reducing the final sintering temperature of the ceria
films to avoid the co-sintering of the anode and by addition of pore-formers in the anode
composition to improve the porosity in the structure.
The high surface area LSCF powders had high sinterability and densified very
easily at moderate temperatures; similar results were observed by Murray et al.55 Cross-
sectional microstructures of the LSCF cathode sintered at 750 °C and 900 °C for 1 hour
are shown in figure 5-7 (a) and (b), respectively. The cathodes have insufficient porosity,
so for the I-V measurements a cathode sintered at 750 °C was used. Fractured cross-

119
section of a tested unit cell with the reduced Ni-GDC anode (~1 mm), GDC electrolyte
(~15 jam) and LSCF cathode (~85 (am) is shown in figure 5-8. The active cathode and
anode areas were approximately 1.6 cm and 4.5 cm , respectively.
Open circuit potential (OCP) and current-voltage behavior of the cell were
measured as a function of temperature, cathode gas atmosphere and flow rate. Figure 5-9
and 5-10 are plots of OCP and average oxygen ion transference number vs. temperature
of the cell, respectively. The electrolytic domain of ceria reduces and its ionic
transference number decreases with increasing temperatures, which results in lower OCP.
With oxygen at the cathode, the cell showed higher OCP in comparison with air as
expected. However, the theoretical OCP increased even more and the average
transference number of the cell with oxygen was slightly lower than that of air. This
indicates that electrode polarization effects are present even under the open circuit
condition due to electronic leakage in the ceria electrolyte.
I-V characteristics of the cell with air at the cathode are shown in figure 5-11. The
performance has been normalized with respect to the cathode area. In the open circuit
condition, the cell potential decreased on increasing the temperature as shown in figure 5-
9. On drawing current from the cell, the cell potential decreased because of the various
polarization losses across the cell. At 550 °C, the curvature at low current densities
suggests that there is some activation polarization at the electrodes, which reduces on
going to higher temperatures. Figure 5-12 shows the power density of the cell with air at
the cathode which has been extrapolated to higher current densities. At 650 °C with a
flow rate of 30 seem, the cell showed a measured power density of 0.24 W/cm2 at a
current density of 0.4 A/cm2. At 650 °C, the extrapolated power density is about 0.27

120
W/cm2 at 0.62 A/cm2. The performance is comparable to power densities reported in the
literature for thick film ceria IT-SOFCs.82 83
The performance of the cell was compared with oxygen at the cathode, and similar
trends were observed in the I-V characteristics and power density of the cell as shown in
figure 5-13 and 5-14, respectively. The cell with oxygen showed a higher measured
power density (0.29 W/cm2 at 0.4 A/cm2, 650 °C) compared to air at cathode. At 650 °C,
the extrapolated power density is about 0.5 W/cm at 1.14 A/cm .
Figure 5-15 is a plot of ln(ASR/T) vs 1/T with air and oxygen at the cathode. Area
specific resistance (ASR) of the cell is the slope of I-V plots in the ohmic region. With air
at the cathode, there is a significant decrease in ASR on increasing the flow rate but with
similar activation energy (-0.58 eV). For example at 650 °C, ASR of the cell decreased
from 0.77 Qcm to 0.47 Qcm on increasing the air flow rate from 20 seem to 30 seem.
With oxygen, there is a further decrease in ASR but without much effect of flow rate and
with a higher activation energy (-0.82 eV) compared to that of air. At 650 °C, ASR of the
cell with oxygen was about 0.35 Qcm for both the flow rates. This suggests that there
are two different rate limiting steps, which determine the cell performance, with air and
oxygen at the cathode. With air, the effect of flow rate and a lower activation energy
suggests that the rate limiting step is the gas diffusion at the relatively dense, -85 pm
thick cathode, and with oxygen, charge transport across the electrolyte seems to be the
rate limiting step.
Due to the asymmetric size of the electrodes (cathode arearanode area = 1:2.8), the
measured power densities, normalized to the cathode area, could have been easily over¬
estimated; specifically in terms of the anode concentration polarization as a larger anode,

121
compared to cathode, would have a larger area for fuel transport. As pointed out by Jiang
et al.,84 the over-estimation in performance depends on the transport properties and the
thickness of the anode, and on the relative size of the cathode.
5.3.2 Improving the Performance of Anode Supported Ceria Unit Cells
After the initial demonstration of fabrication and testing of anode supported thick
film ceria unit cells with acceptable performance characteristics, a need was felt for
further development and improvement in a number of areas. NiO particle size in the
anode has to be reduced in order to ensure complete reduction under operating
conditions. Sintering temperature of the anode/electrolyte bilayer has to be also reduced
so as to avoid excessive sintering of the anode. As mentioned earlier, asymmetric size of
the electrodes may lead to over-estimation of the cell performance. To avoid this issue,
bigger diameter anode substrates were fabricated which allowed comparable anode and
cathode geometries. Dip-coating was replaced by spray-coating to deposit electrolyte
films. Different types of composite cathodes were evaluated to further improve the cell
performance. Lastly, the testing schedule and performance criterion of unit cells were
standardized to enable better comparison between cells. These and other modifications
are described in detail in the following paragraphs.
The co-sintering of the anode/electrolyte bilayer should result in highly dense
electrolyte though with a porous anode support. It helps that the in situ reduction of NiO
to Ni generates porosity in the anode structure. However, there are concerns with high co¬
sintering temperatures in terms of incomplete reduction of coarsened NiO grains,
presence of low porosity, and low concentration of TPBs in the anode structure. To avoid
these issues, finer NiO powders from J. T. Baker was evaluated. Finer particle size will

122
have better reduction properties and with higher sinterability may also ease in the co¬
sintering of the electrolyte film.
Anode supports were pre-sintered at 500 °C to burn off the binders and then spray
coated with GDC colloidal sol. Unlike dip-coating, spray-coating does not require
mechanically strong anode supports, and therefore, higher pre-sintering temperatures
were unnecessary. Further, a number of anode supports could be spray-coated together
which resulted in reduced sample to sample variability within a batch; however, this did
not solve problems with batch to batch variability. Spray-coated bilayers were finally
sintered at temperatures between 1350-1550 °C for 6 hours. Warpage in the samples
increased with the final sintering temperature. LSCF-GDC cathode paste was brush
painted on the electrolyte side and sintered at temperatures between 850-1350 °C for 2
hours to complete the unit cell. The active cathode and anode areas were approximately
3.0 cm and 4.5 cm , respectively.
After sealing the cell, H2 was flowed at 650 °C for ~24 hours to reduce NiO to Ni in
the anode. On starting H2 flow, OCP of the cell increased sharply and then stabilized with
time. The performance of the cell was first evaluated in terms of the magnitude of OCP,
which is indicative of effective sealing, dense electrolyte film and well performing
electrode/electrolyte interfaces. The magnitude of OCP of a good performing cell was
-0.75 V at 700 °C and -1.0 V at 500 °C. The performance of the cell was then evaluated
in terms of I-V characteristics and available power density at each temperature. The cells
were also evaluated in terms of stability under thermal cycling from 750 °C to 500 °C, at
least a couple of times over a time period of 5-7 days. Further, current interrupt technique
was used to separate ohmic and non-ohmic polarization losses of the cell. Ohmic

123
polarization comes primarily from the electrolyte and the non-ohmic polarization from
the electrodes. As in general the anode polarization is much smaller in comparison to the
cathode polarization, the non-ohmic polarization of the cell was used to compare
cathodes with different compositions and sintering conditions.
5.3.2.1 Reducing the co-sintering temperature of anode/electrolyte bilayer
In this part of the work, co-sintering temperature of the anode/electrolyte bilayer
was optimized using LSCF-GDC composite cathodes sintered at 1250 °C. I-V
characteristics and power density of the cell with bilayer sintering temperature of 1350 °C
is shown in figure 5-16 and 5-17, respectively. Numerals before the temperatures in the
graphs represent the day of the experiment. Performance of the cell is poor with a
maximum power density at 650 °C of only 0.114 W/cm2 at 0.244 A/cm2. At 700 °C and
750 °C, performance is much lower with a very low OCP and limiting current density.
Further, it was observed that the performance of the cell decayed with thermal cycling at
600 °C and 650 °C, both in terms of OCP and limiting current density, as shown in figure
5-18 and 5-19, respectively. Post-test micrographs in figure 5-20 show that the GDC
electrolyte film is cracked extensively. Due to the low sintering temperature, the
electrolyte film is not highly dense and possibly with interconnected porosity. Poor
performance of the cell can be explained in terms of air crossover through the porous
electrolyte which resulted in the eventual oxidation of the anode. As the Ni to NiO re¬
oxidation results in volume expansion, the electrolyte film cracked under the stress
generated at the interface. Indeed, anode at the interface with the electrolyte had a
greenish tinge of NiO indicating re-oxidation.
I-V characteristics and power density of the cell with final sintering temperature of
1450 °C is shown in figure 5-21 and 5-22, respectively. The cell showed better

124
9 9
performance with a maximum power density at 600 °C of 0.146 W/cm at 0.336 A/cm .
Maximum power density available at 600 °C, 650 °C, and 700 °C were comparable. OCP
of the cell at 500 °C was 0.97 V which indicated that the electrolyte film is dense.
However, the OCP at 750 °C was much lower than expected at 0.65 V, the reasons behind
which are unclear. As shown in figure 5-23, the cell performance at 650 °C improved
after thermal cycling from 750 °C which is believed to occur due to the slow kinetics of
NiO reduction. Post-test micrograph in figure 5-24 shows that sintering at 1450 °C has
resulted in a denser electrolyte film and in improved cell performance.
I-V characteristics and power density of the cell with final sintering temperature of
1550 °C is shown in figure 5-25 and 5-26, respectively. Performance of the cell was the
worst among the three sintering conditions with a maximum power density at 650 °C of
only 0.040 W/cm at 0.092 A/cm . OCP and the limiting current density of the cell were
low at each temperature. Post-test micrographs in figure 5-27 show that the electrolyte is
highly dense but the anode is not porous as expected. Further, the anode was not
conductive at room temperature and was green in color. As the electrolyte is dense, gas
cross over is ruled out to cause re-oxidation of the anode. The most likely reason could be
the incomplete initial reduction of NiO, probably due to the high density of the anode
after sintering at 1550 °C. Thus, 1450 °C appears to be the best sintering temperature for
the spray-coated anode/electrolyte bilayer with finer NiO powders. Moreover with
comparable anode and cathode geometries, the performance estimations are closer to real.
5.3.2.2 Optimization of LSCF-GDC cathode sintering temperature
Cathode contributes a significant part to the total polarization losses of the cell and
hence demands an optimization in both composition and microstructure. Dusastre et al.53
and Murray et al.55 studied LSCF-GDC type cathodes using impedance spectroscopy and

125
found the best performance with a composition in the range of 30-50 wt% GDC. In this
part of the work, an attempt was made to optimize the sintering temperature of LSCF-
GDC composite cathode with 30 wt % GDC. The anode/electrolyte bilayer for all
samples was sintered at 1450 °C.
Cross-sectional micrographs of unit cells with LSCF-GDC cathodes sintered at
temperatures between 850-1300 °C are shown in figure 5-28 (a)-(e). Two effects of the
sintering conditions on the cathode microstructure are clearly evident. Firstly, at lower
sintering temperatures a very fine cathode microstructure is present in terms of both
particle and pore size. On increasing the sintering temperatures from 850 °C to 1350 °C,
both the particle and pore size increase by about 6 times. Secondly, with higher sintering
temperatures there is better bonding and necking between cathode particles and between
cathode and electrolyte layer. The above described effects are more evident in figure 5-28
(f), which is a surface micrograph of the interface between the cathode layer and the
current collector layer. Both layers have the same composition but with different
sintering temperatures; the cathode layer is sintered at 1250 °C while the current collector
layer is sintered at 1000 °C.
I-V characteristics and power density of unit cells with LSCF-GDC cathodes
sintered at 850 and 1000 °C are shown in figure 5-29 - 5-32. The cells showed good
performance in terms of OCP (-1.0 V at 500 °C, -0.77 V at 750 °C), but the performance
in terms of limiting current density was poor. This resulted in a maximum power density
at 700°C of only 0.036 W/cm2 at 0.08 A/cm2 for the cathode sintered at 850 °C and of
0.047 W/cm" at 0.11 A/cm" for the cathode sintered at 1000°C. Further after thermal
cycling between 500-750 °C, it was observed that the performance of both of the cells

126
decayed as shown in figure 5-33 - 5-36. But interestingly, OCP of both cells remained
constant with thermal cycling which indicated that the component responsible for the
decay is other than the electrolyte. This is in contrast to the performance decay in the unit
cell with anode/electrolyte bilayer sintering temperature of 1350 °C, where both the OCP
value and the limiting current density decreased as a result of thermal cycling. Indeed
after the test it was found that the cathode was sintered poorly with the electrolyte, which
would have been aggravated during the thermal cycling and consequently resulted in the
performance decay.
I-V characteristics and power density of unit cells with higher cathode sintering
temperatures of 1150 °C, 1250 °C, and 1350 °C are shown in figure 5-37 - 5-42. The cells
showed expected OCP values and significantly higher limiting current densities, which
resulted in 4-6 times higher maximum power density compared to the cells with lower
cathode sintering temperatures. For example at 700 °C, cathode sintering temperature of
1150 °C resulted in a maximum power density of 0.201 W/cm2 at 0.393 W/cm2, cathode
sintering temperature of 1250 °C resulted in 0.244 W/cm2 at 0.422 A/cm2 (extrapolated to
0.272 W/cm at 0.591 A/cm ) and cathode sintering temperature of 1350 °C resulted in
0.282 W/cm2 at 0.465 A/cm2 (extrapolated to 0.338 W/cm2 at 0.771 A/cm2). Further,
after thermal cycling from 750 °C the performance of the cell improved, as shown in
figure 5-43 and 5-44, which is believed to occur due to the slow kinetics of NiO
reduction. The maximum power density in the temperature range of 500-750 °C for unit
cells with different cathode sintering temperatures is shown in figure 5-45. Evidently,
depending on the operating temperature the cathode sintered either at 1250 °C or 1350 °C
should be preferred; at temperatures below 650 °C, the cell with the cathode sintered at

127
1250 °C showed better performance, whereas at higher temperatures the cathode sintered
at 1350 °C showed better performance.
The current interrupt technique was used to further investigate the performance
characteristics of the two sintering conditions by separating out the ohmic and the non-
ohmic polarizations of the cell. In the current interrupt technique, the current is stopped
for a very short time on the order of micro-seconds and the voltage profile is observed
using a high speed oscilloscope. The ohmic polarization has a smaller time constant and
responds instantaneously to the current interrupt, while the non-ohmic polarizations
respond slowly which helps in separating the two polarizations. I-V characteristics with
separated ohmic and non-ohmic polarizations at 500 °C and 700 °C for the cell with
cathode sintering temperature of 1350 °C are shown in figure 5-46 and 5-47, respectively.
The ohmic and non-ohmic polarizations as a function of current density for the two
sintering conditions are shown in figure 5-48 and 5-49, respectively. As both the
polarizations are thermally activated, they decrease with increasing operating
temperatures.
ASR values of the ohmic polarization for the two cells are compared with that of
the bulk GDC electrolyte (extrapolated to 30 pm thickness) in figure 5-50. Activation
energy for the ionic conduction in the bulk GDC electrolyte is -0.66 eV, which is similar
to the values reported by Steele and Goodenough. Ohmic polarization of both cells
show comparable activation energies, which suggests that its origins are in the oxygen
ion conduction in GDC. However, the polarization magnitude for both cells is about 3
times larger in comparison to the 30pm GDC electrolyte. It is believed that apart from the
thick film electrolyte, ionic transport in the anode and cathode layers through the GDC

128
particles is also contributing to the total ohmic polarization. Therefore to decrease the
total ohmic polarization across the cell, one has to focus not only on the electrolyte
properties and thickness but also on the ionic conduction in the electrodes and interfaces.
At temperatures between 500-650 °C, as shown in figure 5-48 and 5-49, the non-
ohmic polarization comprises the larger fraction of the total polarization for both the
cells, especially at low current densities where it is limited by the activation polarization.
At higher temperatures the activation polarization is absent, and the non-ohmic and
ohmic polarizations are of comparable magnitudes. Both the samples showed similar
trends; though, it is clear that at lower temperatures the cathode sintered at 1250 °C
showed lower polarizations and hence better performance, while at higher temperatures
as both the samples showed similar polarizations it is not clear why the sample sintered at
1350 °C showed better performance.
In any case, the present study showed the importance of the cathode sintering
condition and the resultant microstructural effects on the performance of solid oxide fuel
cells. Higher sintering temperatures resulted in coarsening of the cathode microstructure
with larger particle and pore size, which reduced the concentration of TPBs but on the
other hand improve gas diffusion through the cathode. Higher sintering temperatures also
improve the bonding between electrode particles and between electrode and the
electrolyte layer, which reduces the interfacial resistance. Therefore, there are a number
of contradictory requirements in terms of the cathode microstructure towards lowering
the cathode polarization, and in this study with LSCF-GDC cathodes, the optimum
sintering temperature was found to be around 1250-1350 °C. This result is in
contradiction with the studies done on LSCF-GDC cathodes using impedance

129
ci cS
spectroscopy by Dusastre et al. and Murray et al., where they found that the optimum
sintering temperature is around 800-900 °C. This can be attributed to two reasons. Firstly,
the final microstructure depends not only on the sintering temperature but also on the
starting particle size distribution (which was similar in the three cases), slurry
composition and method of deposition. Secondly and more importantly, impedance
spectroscopy measurements under no bias currents are not ideal representatives of the
electrode ASR (for example at maximum power density conditions), and therefore, the
trends generated using impedance spectroscopy should be verified under actual fuel cell
conditions.
In cells with ceria electrolyte, the power density depends on the conflicting
parameters of OCP and ASR of the cell. On increasing the temperature, OCP decreases
along with ASR which results in a critical temperature (~700°C) at which the maximum
power density can be achieved. By depositing an electron blocking and high oxygen ion
conducting layer of stabilized bismuth oxide on the cathode side of the doped ceria
electrolyte, it has been shown that the open circuit potential can be improved without
increasing the ASR of the cell.8' Thus, this limitation could be overcome with
ceria/bismuth oxide bilayer electrolytes resulting in improved SOFC performance at
intermediate temperatures.
5.3.2.3 Composite cathodes containing ESB as the electrolyte phase
As found in the earlier section, the ohmic polarization across the cell has
contributions not only from the electrolyte layer but also from the ionic conduction in the
electrode structure. This confirms the trends observed in the literature where to improve
the performance of composite cathodes, LSM electrocatalyst (poor ionic conductor) has
been replaced by higher oxygen ion conductors like LSCF or SSC, and YSZ as the

130
electrolyte component has been replaced by GDC. ESB has significantly higher oxygen
ion conductivity than GDC and in this part has been investigated to replace GDC as the
electrolyte component in composite cathodes.
ESB was tested for reactivity with LSCF and LSCuF perovskites. XRD patterns for
powder mixtures after different heat treatments are shown in figure 5-51 and 5-52.
Additional peaks appear for each combination after the heat treatment which have been
identified as rhombohedral BÍ2O3, possibly forming by the dissolution of CoJ+ and/or Sr2^
into ESB. For the LSCF-ESB mixture, the rhombohedral phase appears to form at even
660 °C after 28 hours. The rhombohedral BÍ2O3 phase is highly undesirable as it is known
to have up to two orders of magnitude lower oxygen ion conductivity than the fee BÍ2O3
phase. LSCuF has been reported to have high sinterability and was evaluated as a
composite cathode with ESB, in order to avoid issues with tertiary phase formation at
higher sintering temperatures.
I-V characteristics and power density of the cell with LSCuF-ESB (70:30 wt. ratio)
cathode sintered at 750 °C is shown in figure 5-53 and 5-54, respectively. Performance of
the cell with LSCF-GDC cathode sintered at 850 °C is also shown for comparison. The
anode/electrolyte bilayers for both the cells were sintered at 1450 °C. OCP values of the
two cells are similar at 600 °C and 650 °C. At 650 °C, the cell with LSCuF-ESB cathode
is limited by the activation polarization at low current densities, but the ASR of the cell in
the ohmic region is comparable to that of LSCF-GDC which results in a maximum power
density of 0.021 W/cm2 at 0.052 A/cm2. However at 600 °C, there is a sudden increase in
ASR of the cell compared to that of LSCF-GDC resulting in maximum power density of

131
only 0.006 W/cm2 at 0.015 A/cm2. The sudden increase in ASR is believed to occur due
to the presence of resistive, tertiary rhombohedral phase in the cathode microstructure.
Although there have been reports in the literature of synthesizing Lai.xSrxCo03
electrode layers on stabilized bismuth oxide electrolytes by thermal decomposition
without secondary phase formation, ’ it is highly unlikely that interface between the
perovskite cathode and the BÍ2O3 electrolyte can deliver stable performance at elevated
temperatures over extended periods of time. To circumvent the issue of the tertiary
rhomohedral phase, cermet cathodes containing Ag as metal component were also
studied. With Ag containing cathodes, issues in terms of microstructural stability were
observed as also mentioned in chapter 3. OCP of the cell at 600 °C with Ag20-ESB
(60:40 wt. ratio) cathode was unstable as shown is shown in figure 5-55. Post-test
micrograph in figure 5-56 shows the presence of a thin layer of Ag at the
cathode/electrolyte interface which is believed to from due to Ag migration at high
temperatures resulting in unstable OCP.
5.4 Conclusion
This chapter focused on the development of Ni-GDC anode supported thick film
GDC electrolyte solid oxide fuel cells. Fabrication of dense GDC electrolyte films was
done by a colloidal deposition technique. Initial sintering studies showed that the density
of the film depends on the sintering schedule - the pre-sintering and the final sintering
conditions. Unit cells were fabricated and tested with LSCF-GDC composite cathodes to
optimize the sintering temperature of the anode/electrolyte bilayers. The optimum
temperature to obtain sufficiently dense electrolyte film without excessive sintering of the
anode was found to be 1450 °C. Effect of the LSCF-GDC cathode sintering temperature
was also studied and the best performance was observed at 1250-1350 °C with a

132
maximum power density of 0.338 W/cm" at 0.771 A/cm", 700 °C. Cathode microstructure
under different sintering conditions showed the importance of good bonding between the
electrode particles and between the electrode and the electrolyte layer for optimum
performance. Current interrupt technique was used to separate ohmic and non-ohmic
polarizations of the cell. Apart form the electrolyte, the ohmic polarization had significant
contributions from the ionic transport in the electrodes. ESB containing cathodes were
studied to improve the cathode performance but without much success. ESB reacted with
perovskite cathodes to form resistive rhomohedral phase which resulted in low cell
performance, while Ag-ESB cathodes have issues in terms of stability because of Ag
migration. BÍ2R.U2O7.3-ESB composite cathodes, developed in chapter 4, have not yet
been tested on anode supported ceria electrolyte unit cells. The performance of these
cathodes is very promising for operation at intermediate temperatures, and it is expected
that with further optimization in electrode composition and microstructure, high power
density IT-SOFCs could be realized.

133
(c) 900 °C
(d) 1000 °C
(e) 1100 °C
Figure 5-1. Representative microstructures of the GDC film under different pre-sintering
conditions for a 1600 °C final sintering.

134
(b) 850 °C
(a) 800 °C
(c) 900 °C
(d) 1000 °C
(e) 1100 °C
Figure 5-2. Representative microstructures of the GDC film under different pre-sintering
conditions for a 1650 °C final sintering.

135
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C
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24%
20%
16%
12%
8%
4%
0%
Final Sintering Temperature
□ 1600 °C
El 1650 °C
a
800
850
900
1000
1100
Pre-sintering Temperature (°C)
Figure 5-3. Porosity in the GDC film as a function of the pre-sintering and final sintering
temperature.

136
(c) 1100 °C
Figure 5-4. Cross-sectional microstructures of the GDC film under different pre-sintering
conditions for a 1600 °C final sintering.
Figure 5-5. Surface microstructure of the GDC film after 850 °C pre-sintering and 1400
°C final sintering.

137
Figure 5-6. Surface (a) and cross-sectional (b) microstructure of reduced Ni-GDC anode.
(a) (b)
Figure 5-7. Cross-sectional micro structure of LSCF cathode sintered at (a) 750 °C & (b)
900°C.
Figure 5-8. Cross-sectional microstructure of tested unit cell - Ni-GDC anode (~1
mm)/GDC electrolyte (~15 |am)/LSCF cathode (~85 jam).

138
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0.96
0.94
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0.86
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500 550 600 650 700 750
Temperature (°C)
Figure 5-9. Open circuit potential of the cell as function of temperature.
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Temperature (°C)
Figure 5-10. Average oxygen ion transference number of GDC electrolyte as function of
temperature.
30 seem O
30 seem air
30 seem O
30 seem air

139
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550 °C
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Current Density (A/cnT)
Figure 5-11. I-V characteristics of the cell with air at the cathode.
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Current Density (A/cm )
Figure 5-12 Power density of the cell with air at the cathode.

140
Current Density (A/cm )
Figure 5-13. I-V characteristics of the cell with oxygen at the cathode.
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Current Density (A/cm )
Figure 5-14. Power density of the cell with oxygen at the cathode.

141
-5
-5.5
^ -6
P
< -6.5
j:
-7
-7.5
-8
0.001 0.00105 0.0011 0.00115 0.0012 0.00125 0.0013
1/T (K'1)
Figure 5-15. ln(ASR/T) vs 1/T of the cell.
• 20 air
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142
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2
Current Density (A/cm")
Figure 5-16. I-V characteristics of the cell with bilayer sintering temperature of 1350 °C.
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Current Density (A/cm")
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Figure 5-17. Power density of the cell with bilayer sintering temperature of 1350 °C.

143
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0.7
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Current Density (A/cm )
n—t—i—’—* r
%
2-600 °C
.3-600 °C
Figure 5-18. I-V characteristics at 600 °C of the cell with bilayer sintering temperature of
1350 °C before and after thermal cycling.
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Figure 5-19. I-V characteristics at 650 °C of the cell with bilayer sintering temperature of
1350 °C before and after thermal cycling.

144
Figure 5-20. Cross-sectional (a) and surface (b) microstructures of the tested unit cell
with bilayer sintering temperature of 1350 °C.

145
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Figure 5-22. Power density of the cell with bilayer sintering temperature of 1450 °C.

146
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b 0.7
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Current Density (A/cirf)
Figure 5-23. I-V characteristics at 650 °C of the cell with bilayer sintering temperature of
1450 °C before and after thermal cycling.
I ' 1 ' 1 l
• 2-650 °C
" 3-650 °C
4 4-650°C
’ \
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Figure 5-24. Cross-sectional microstructure of the tested unit cell with bilayer sintering
temperature of 1450 °C.

147
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â– 
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Figure 5-25. I-V characteristics of the cell with bilayer sintering temperature of 1550 °C.
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Current Density (A/cm“)
Figure 5-26. Power density of the cell with bilayer sintering temperature of 1550 °C.

148
Figure 5-27. Cross-sectional microstructure of the tested unit cell with bilayer sintering
temperature of 1550 °C.

149
(e) (f)
Figure 5-28. Cross-sectional microstructures of the tested unit cell with cathode sintering
temperature of 850 °C (a), 1000 °C (b), 1150 °C (c), 1250 °C (d), 1350 °C (e).
(f) is the surface microstructure showing cathode layer sintered at 1250 °C and
current collecting layer sintered at 1000 °C.

150
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Figure 5-29. I-V characteristics of the cell with cathode sintering temperature of 850 °C.
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Figure 5-30. Power density of the cell with cathode sintering temperature of 850 °C.
•AAA
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Current Density (A/cm )

151
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Figure 5-31. I-V characteristics of the cell with cathode sintering temperature of 1000
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4-700 °C
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Figure 5-32. Power density of the cell with cathode sintering temperature of 1000 °C.

152
1
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0.2
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Current Density (A/cm )
Figure 5-33. I-V characteristics at 650 °C of the cell with cathode sintering temperature
of 850 °C before and after thermal cycling.
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2 0.4
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Current Density (A/cm“)
2-750 °C
6-750 °C
1â–  . . t
Figure 5-34. I-V characteristics at 750 °C of the cell with cathode sintering temperature
of 850 °C before and after thermal cycling.

153
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0.8
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0.6
V->
e

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13
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0.2
0
0 0.04 0.08 0.12 0.16
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Current Density (A/cm“)
1 ' I 1 1 r—T"
• 1-650C
â–  3-650C
3-650 °C *
1-650°C
I ... I - L.-
Figure 5-35. I-V characteristics at 650 °C of the cell with cathode sintering temperature
of 1000 °C before and after thermal cycling.
Figure 5-36. I-V characteristics at 750 °C of the cell with cathode sintering temperature
of 1000 °C before and after thermal cycling.

154
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0 0.1 0.2 0.3 0.4
0.5
2s
0.6
Current Density (A/cm )
Figure 5-37. I-V characteristics of the cell with cathode sintering temperature of 1150
°C.
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£ 0.05
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Current Density (A/cm )
Figure 5-38. Power density of the cell with cathode sintering temperature of 1150 °C.
2-700 °C
j*$W
3-750 °C
at,
,*• ' 1-650°C
4 aAAa‘aaAa
1-600 °C
2-550 °C
11-500 °C
15-450°C

155
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0.2
2-550 °C
7-500 °C
0
0.1 0.2 0.3 0.4 0.5
Current Density (A/cm")
0.6
Figure 5-39. I-V characteristics of the cell with cathode sintering temperature of 1250
°C.
0.3
^ 0.25
g
| 0.2
t °'15
Q
$3 0.1
£
o
Cu
0.05
0
0 0.1 0.2 0.3 0.4 0.5 ,0.6 0.7
Current Density (A/cm")
Figure 5-40. Power density of the cell with cathode sintering temperature of 1250 °C.
I ' : ' I
3-700 °C
4*
3-750°C
1-650 °C
2-600 °C
\
4
f*
A? 7-500 °C
3
\
2-550 °C

156
1
>
13
+->
c
a>
O
cx
u
0.93
0.8
0.7
0.6
0.5
0.4
0.3
0.2
0
%
Jjim *
a
♦
J- A 3-700 °C
Qk
♦ A J-750°C
* 4 2-650 °C
2-600 °C
2-550 °C
4-500 °C#
0.6
0.1 0.2 0.3 0.4 0.5
Current Density (A/cm")
Figure 5-41. I-V characteristics of the cell with cathode sintering temperature of 1350
°C.
0.4
0.35
0.3
I 0.25
I °-2
? 0.15
I 0.1
eu
0.05
0
0 0.2 0.4 0.6 2 0.8
Current Density (A/cm")
i—1—«—'—i—1—1 '—i ■i ’ ;
3-700 °C
3-750 °C
LT
BA*’'
W,
2-650 °C
/♦
jT 2-600 °C
aT 4-500oC 2-550°C
Figure 5-42. Power density of the cell with cathode sintering temperature of 1350 °C.

157
1
0.9
r 0,8
5
§ 0.7
o
CU
0.6
u
0.5
0.4
0 0.05 0.1 0.15 0.2 0.25 0.3^0.35 0.4
Current Density (A/cm")
Figure 5-43. I-V characteristics at 650 °C of the cell with cathode sintering temperature
of 1150 °C before and after thermal cycling.
# 1-650 °C
* 4-650 °C
♦ 10-650 °C
i .... i . . . l..
• ♦
1
0.9
>
^5 0.8
C
■4—>
o
^ °-7
"S
U
0.6
0.5
0 0.1 0.2 0.3 0.4 0.5
Current Density (A/cnC)
Figure 5-44. I-V characteristics at 650 °C of the cell with cathode sintering temperature
of 1250 °C before and after thermal cycling.

158
Temperature (°C)
Figure 5-45. Maximum power density as a function of cathode sintering temperature.

159
>
"5
c
u
â– *->
o
cu
13
u
1
0.9
0.8
0.7
0.6
0.5
0.4
0.3
0.2
r
• •
o
E°
j—' O
E - ri - ri
cathode anode
1
E° - n - ri - IR
cathode anode
500 °C
0.05
0.1
0.15
. ♦
0.2
Current Density (A/cm-)
Figure 5-46. I-V characteristics at 500 °C of the cell with cathode sintering temperature
of 1350 °C, with contributions from ohmic and non-ohmic polarization.
0.85
0.8
>
.3 0.75
+->
c
+->
o
^ °-7
13
U
0.65
0.6
0 0.1 0.2 0.3 0.4 0.5
Current Density (A/cm-)
Figure 5-47. I-V characteristics at 700 °C of the cell with cathode sintering temperature
of 1350 °C, with contributions from ohmic and non-ohmic polarization.
* •
4 E° - r| - r|
cathode anode
E° - r, - n - IR
cathode anode
700 °C

160
G
O
c3
N
‘C
o
a.
o
O
i
G
O
Current Density (A/cnT)
Figure 5-48. Non-ohmic polarization of the cells with cathode sintering temperature of
1250 °C (filled symbols) and 1350 °C (open symbols).
0.35
_ 0.3 f-
>
1 °-25
03
.a o.2
Vh
•
-2
a. 0.15
o
•
"g
M 0.1
• ■
O •
i ♦
0.05 •
♦*>
500 °C
550 °C
o
600 °C
650 °C
A
700 °C
At k.
: V 750 °C
V'ts
0 0.1 0.2 0.3 0.4 0.5 0.6
Current Density (A/cnf)
Figure 5-49. Ohmic polarization of the cells with cathode sintering temperature of 1250
°C (filled symbols) and 1350 °C (open symbols).

161
-6
-6.5
-7
-7.5
-8
-8.5
-9
-9.5
0.9 1 1.1 1.2 1.3 1.4
1000/T (K')
I . I 1 I 1 T 1 1 t I ! 1 I I ! ' 1 r- 7»
• 1250 °C (0.62 eV) •
■ 1350 °C (0.69 eV) ,/ ♦
♦ BULK (0.66 eV) V
♦
• ♦
â– 
t "
♦
I J 1 l. i. I I J . , 1 , . I .
Figure 5-50. Arrhenius plot of ASR of the ohmic polarization of the cells with cathode
sintering temperature of 1250 °C and 1350 °C, in comparison to ~30 pm bulk
GDC electrolyte.

162
20
Figure 5-51. XRD patterns for LSCF-ESB (a) no heat treatment (b) 660 °C / 28 hours (c)
740 °C / 28 hours (d) 800 °C / 4 hours.
20
Figure 5-52. XRD patterns for LSCuF-ESB (a) no heat treatment (b) 800 °C / 4 hours.

163
>
c
o
CU
a>
U
0.8 •
Ü3
13 0.6
• 2-600 °C
■ 2-650°C
3-600 °C (LSCF-GDC)
1-650 °C (LSCF-GDC)
â–¡
a
0.4
0.2
â–¡
â–¡
â–¡
m
â–¡
0
0
0.02 0.04 0.06 0.08 0.1 0.12 0.14
Current Density (A/cm“)
Figure 5-53. I-V characteristics of the cell with LSCuF-ESB cathode sintered at 750 °C,
in comparison with the cell with LSCF-GDC cathode sintered at 850 °C.
0.03
^0.025
|
| °-°2
•s 0.015
G
Q
fc 0.01
£
o
^0.005
0
0 0.02 0.04 0.06 0.08 0.1
2
Current Density (A/cm )
Figure 5-54. Power density of the cell with LSCuF-ESB cathode sintered at 750 °C, in
comparison with the cell with LSCF-GDC cathode sintered at 850 °C.
n'
â–¡
â– B-Q-
p c
’0< .
' -0
f
D'
a
$
¡i
j
•••
• 2-600 °C
■ 2-650 °C
â–¡
B
\
--B-
3-600 °C (LSCF-GDC)
1-650 °C (LSCF-GDC)

164
1
0.8
>
Tí 0.6
c
-t-J
o
^ °-4
U
0.2
0
0 0.5 1 1.5 2 2.5 3 3.5 4
Time (hour)
Figure 5-55. OCP of the cell with Ag-ESB cathode as a function of time.
Figure 5-56. Cross-sectional microstructures of the tested unit cell with Ag-ESB cathode
in back-scatted electron mode.

CHAPTER 6
CONCLUSION AND FUTURE WORK
Reduction in the operating temperatures will make solid oxide fuel cells (SOFCs)
competitive with other power generation technologies in both stationary and tractionary
markets. The aim of this dissertation was to evaluate and develop component materials
for SOFCs, which could work efficiently at intermediate temperatures. Effect of direct
current bias on the ordering phenomenon in erbia stabilized bismuth oxide (ESB), a
candidate electrolyte material, was studied using impedance spectroscopy and differential
scanning calorimetry (DSC). Ag-ESB cermets, undoped and doped bismuth ruthenates,
and bismuth ruthenate-ESB composites were evaluated as candidate cathode materials for
IT-SOFCs using impedance spectroscopy. Unit cells consisting of thick film gadolinium
doped ceria (GDC) electrolyte supported on Ni-GDC anode were developed for operation
at intermediate temperatures.
The effect of the DC bias on oxygen ion conductivity decay in erbia stabilized
bismuth oxide on isothermal annealing below the transition temperature (-600 °C) was
studied using impedance spectroscopy. Decay in oxygen ion conductivity on annealing
below the transition temperature occurs due to the trapping of mobile oxygen vacancies
in the condensed ordered clusters. The trapping of the oxygen vacancies could be affected
by the jump frequency which depends on the electrochemical potential gradient and the
resultant oxygen ion flux. Below the 600 °C transition temperature (at 500 °C), no
significant effect on the conductivity decay was observed for bias currents up to 72 mA.
However, the enthalpy change in annealed samples, which has previously been related to
165

166
the order to disorder transition, did not match the trends of the conductivity decay. After
~10 hours, without bias the electrolyte showed more than 200 % increase in the
resistance though without any observed endotherm using DSC, and under bias the sample
showed similar increase in resistance but with an endotherm in the DSC profile. As the
ordering of the oxygen ion sub-lattice comprises of two parts: positional and occupational
ordering, one possible reason could be that the kinetics of the conductivity decay and the
endotherm have origins to different type of ordering process, which progress at different
rate under the applied current bias. Positional displacement is the faster of the two, as the
displacement is on the order of less than an oxygen ion radii and is also observed to be
partially present in un-annealed samples. In contrast, occupational ordering involves
rearrangement of the oxygen lattice over larger distances on the order of lattice parameter
and thus would be expected to have a higher time constant. The endotherm is expected to
be primarily related to the occupational ordering, and hence, it appears only after long
anneal time under no bias conditions. Since the endotherm at short time is enhanced by
the applied bias with negligible change in conductivity decay, it is possible that the decay
in conductivity with anneal time is more related to positional ordering than occupational
ordering. This explanation has to be carefully verified using electron and neutron
diffraction studies, which could elucidate the effect of the oxygen ion flux on the kinetics
of positional and occupational ordering, and on the growth of ordered clusters.
Ag-ESB cermet cathodes showed good performance with area specific resistance
(ASR) values of 3.08 Qcm2 at 500 °C and 0.16 Dcm2 at 625 °C in air. Impedance studies
under current bias showed that the electrode response consisted of charge transfer and
bulk transport contributions. At 500 °C, the electrode impedance is dominated by the

167
charge transfer step and it decreased on application of current bias. At 625 °C, the
electrode performance at high current bias was not stable and the bulk transport
contribution increased with time. After the experiment, microstructural analysis revealed
significant migration of Ag along with the oxygen flux in the electrode. The Ag-Ag2Ü
phase diagram shows the presence of a eutectic at 530 °C and 519 atm. O2. The high
performance of the Ag cermet electrodes for oxygen reduction is due to oxygen solubility
of Ag, which also leads to formation of the eutectic, and therefore can result in
microstructural instabilities. Therefore, electrodes containing Ag in sufficient proportions
should not be expected to be microstructurally stable for long periods of time above the
eutectic temperature. Electrodes with higher melting point Ag-Pt or Ag-Pd alloy, or
electrodes with lower Ag fraction with other metal or metallic oxides to ensure electronic
percolation could provide better long term performance.
Bismuth ruthenates based cathodes were evaluated using impedance spectroscopy
in an attempt to develop high performance cathodes for IT-SOFCs. Bismuth ruthenate
electrodes showed ASR values of 55.64 Í2cm2 at 500 °C and 1.45 Qcm2 at 700 °C in air.
Impedance studies as a function of oxygen partial pressure revealed that the rate limiting
step in the oxygen reduction reaction is the surface diffusion of dissociatively adsorbed
oxygen at the electrode to the triple phase boundaries (TPBs). Two strategies were
employed to improve the electrode performance: doping on Bi3+ site with similar size
aliovalent cations in order to introduce oxygen ion conductivity and formation of
composite electrodes consisting of bismuth ruthenate and stabilized bismuth oxide. It was
hoped that the introduction of oxygen ion conductivity in bismuth ruthenate will activate
the electrode surface for the oxygen reduction reaction, and that by forming composite

168
cathodes, the concentration of TPBs could be increased along with improved ionic
conductivity in the electrode structure. Doping bismuth ruthenates with Ca , Ag , and
Sr2+ did not improve the performance significantly. On the other hand, composite
electrodes showed much better performance and the composition with 37.5 wt%
(BÍ203)o.8(Er203)o.2 showed ASR values of 18.4 Qcm2 at 500 °C and 0.32 Qcm2 at 700 °C
in air. The performance of the composite cathodes is very promising, and it is expected
that further optimization in electrode composition and microstructure will result in high
performance cathodes for IT-SOFCs. Lastly, it was felt that new methodologies are
required to extract more useful information out of the impedance data, which could help
in identifying the rate limiting steps present in the electrode reaction and guide
researchers towards solving relevant issues with intrinsic material properties and
electrode microstructure.
The inter-relationship between the processing parameters of anode supported ceria
electrolyte unit cells, the resultant microstructure, and the fuel cell performance was
studied. Thick film GDC electrolyte unit cells were fabricated on porous Ni-GDC anode
supports using a colloidal deposition technique. Sintering studies showed that the density
of the electrolyte film depends on the sintering schedule: pre-sintering temperature of
anode substrate and final sintering temperature of anode/electrolyte bilayer. Unit cells
were completed with LSCF-GDC composite cathodes and tested under fuel cell
configuration. Performance results showed that the optimum temperature for the sintering
of anode/electrolyte layer was about 1450 °C. The cathode sintering temperature also has
a major effect due to requirement of good bonding between cathode particles and of the
cathode with the electrolyte layer. Best performance was observed with a cathode

169
sintered at 1350 °C with a maximum power denisty of 0.338 W/cm: at 0.771 A/cm2, 700
°C. Current interrupt studies showed that apart form the electrolyte, the total ohmic
polarization across the unit cell had significant contributions from the ionic transport in
the electrode structure. To improve cell performance, composite cathodes containing ESB
as the electrolyte phase were studied but without much success. ESB reacted with the
perovskite cathodes to form resistive rhombohedral phase which resulted in low fuel cell
performance, while Ag-ESB cathodes were found to be microstructurally unstable due to
Ag migration at high temperatures. The performance of the cells developed in this work
is comparable with those reported in the literature. To further improve the cell
performance, effort needs to be put towards improving anode and cathode
microstructures, reducing thickness of the electrolyte without compromising its density,
and developing better cathode materials.

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BIOGRAPHICAL SKETCH
Abhishek Jaiswal was bom on October 10th, 1976 in the north Indian city of
Lucknow. He was an average student during his high school days and spent most of his
time playing and watching cricket. During his final years at high school, he became more
serious about his future, and after a few unsuccessful attempts, he made to the India’s
premier engineering school-Indian Institute of Technology (IIT). He decided to go to IIT-
Bombay for his B. Tech, in Metallurgical Engineering and Material Science. Like any
other IITian, he feels that the four years at IIT were the best of his life. During his senior
year project, he was introduced to the solid oxide fuel cell technology and found it
interesting enough to carry on working for his doctorate at University of Florida. The last
four years at UF and in United States have been highly enriching both academically and
personally to him.
177

I certify that I have read this study and that in my opinion it conforms to acceptable
standards of scholarly presentation and is fully adequate, in scope and quality, as a
dissertation for the degree of Doctor of Philosophy.
Eric D. Wachsman, Chairman
Professor of Materials Science and
Engineering
I certify that I have read this study and that in my opinion it conforms to acceptable
standards of scholarly presentation and is fully adequate, in scope and quality, as a
dissertation for the degree of Doctor of Philosoph)
Darryl
Associate Professor of Materials Science
and Engineering
I certify that I have read this study and that in my opinion it conforms to acceptable
standards of scholarly presentation and is fully adequate, in scope and quality, as a
dissertation for the degree of Doctor of Philosophy.
Wolfgang M. Sigmund
Associate Professor of Materials Science
and Engineering
I certify that I have read this study and that in my opinion it conforms to acceptable
standards of scholarly presentation and is fully adequate, in scope and quality, as a
dissertation for the degree of Doctor of Philosophy. / /
David P. Norton
Professor of Materials Science and
Engineering
I certify that I have read this study and that in my opinion it conforms to acceptable
standards of scholarly presentation and is fully adequate, in scope and quality, as a
dissertation for the degree of Doctor of Philosophy.
Engineering

This dissertation was submitted to the Graduate Faculty of the College of
Engineering and to the Graduate School and was accepted as partial fulfillment of the
requirements for the degree of Doctor of Philosophy.
December 2004 |
Pramod P. Khargonekar
Dean, College of Engineering
Winfred M. Phillips
Dean, Graduate School




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