Science and Technology of Ceramic Fuel Cells

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N G U Y E N Q U A N G MINH

ALLIEDSIGNAL, INC. AEROSPACE EQUIPMENT SYSTEMS

TORRANCE, CALIFORNIA, U.S.A.

T A K E H I KO T A K A H A S H I

PROFESSOR EMERITUS NAGOYA UNIVERSITY

NAGOYA, OAPAN

1995

ELSEVIER

AMSTERDAM �9 LAUSANNE �9 NEW YORK �9 OXFORD �9 SHANNON �9 TOKYO

ELSEVIER SCIENCE B.V. Sara Burgerhartstraat 25 P.O. Box 2II, IOOO AE Amsterdam, The Netherlands

L i b r a r y of Congress C a t a l o g i n g - i n - P u b l i c a t i o n Data

Nguyen, Quang Minh. Science and technology of ceramic fue l c e l l s / Nguyen Quang Minh,

Takehiko Takahashi . p. cm.

Inc ludes b i b l i o g r a p h i c a l r e f e r e n c e s and index. ISBN 0-444-89568-X 1. Sol id oxide f ue l c e l l s . I . Takahashi , Takehiko. I I . T i t l e .

TK2931.N48 1995 621 .31 '2429- -dc20 95-21137

CIP

ISBN: o 444 89568 X

�9 I995 Elsevier Science B.V. All rights reserved.

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Dedicated to my late father, Mr. Nguy~n Du~

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Dedicated to my wife, RiO Takahashi

T.T.


PREFACE

Ceramic fuel cells, commonly known as solid oxide fuel cells (SOFCs),

have been under development for a broad spectrum of electric power generation applications. The most attractive feature of the SOFC is its clean and efficient

production of electricity from a variety of fuels. The SOFC has the potential to

be manufactured and operated cost-effectively and, thus, promises to be an

important alternative source for generating electric power in the future. At present, SOFC technology is still in its development stage, and several

technical challenges remain to be resolved before the fuel cell becomes a full- fledged, practical power system. In the past several years, SOFC technology has

received much attention. Development effort in this area has expanded significantly, and the number of conferences and publications on SOFCs has

increased dramatically. The widening interest in this technology arises, in part,

from the continuing need to develop cleaner and more efficient means of

converting energy sources into useful forms. Recent advances in ceramic technology, especially in synthesizing fine powders, engineering material

compositions, tailoring composition/property relationships, and processing intricate structures, have also contributed to the increased interest in SOFCs. These technological advancements have led to improved designs, better performance, and cost-effective manufacturing methods. As a result, significant

progress toward practical applications of the SOFC has been made. This book has been written to provide a comprehensive treatise on

SOFCs and to fill the need for a reference book on the technology. It is directed toward scientists, engineers, and technical managers working with SOFCs,

ceramic devices based on conducting materials, and in related fields such as

solid-state ionics and electronic ceramics. The book can also be used as a text

for science and engineering courses at the senior or graduate level. As the field of SOFCs is still evolving, this book has been prepared with an emphasis on the

discussion of known facts and reported experimental data. Dated information has

been kept to a minimum; however, in some instances, this type of information

was included to give insights into work now in progress and to provide guidance

viii Preface

for future work. It is possible that some ideas and explanations discussed in the

book may change as the technology progresses. We greatly appreciate the encouragement and support of many colleagues

during the writing of this book. We are grateful to Ms. Lynn Silver for her

editorial assistance.

Torrance, California, USA Nagoya, Japan

�9 NGUYEN QUANG MINH TAKEHIKO TAKAHASHI

December 1994

CONTENTS

Chapter 1

Chapter 2

Chapter 3

INTRODUCTION

1.1 Scope, 1 1.2 General Characteristics of Ceramic Fuel Cells, 3

1.2.1 Types of Ceramic Fuel Cells, 4 1.2.2 Cell Components, 6 1.2.3 Comparison with Other Types of Fuel Cells, 10

1.3 Historical Background of Ceramic Fuel Cells, 10 References, 14

PRINCIPLES OF OPERATION

2.1 General, 15 2.2 Thermodynamic Principles, 16 2.3 Fuel Cell Efficiency, 20

2.3.1 Electrochemical Efficiency, 21 2.3.2 Other Efficiencies, 24

2.4 Power Generation, 25 2.5 Characteristics of Ceramic Fuel Cells, 28

2.5.1 Features, 28 2.5.2 Effect of Electronic Conduction in Electrolyte, 29

2.6 Types of Fuel and Oxidant, 36 2.6.1 Fuel, 37 2.6.2 Oxidant, 38

2.7 Fuel-Processing System, 38 2.8 Power-Conditioning System, 39 References, 40

ELECTRICAL CONDUCTION IN CERAMICS

3.1 3.2

General, 41 Defects in Fluoride-Type Oxides, 43 3.2.1 Defect Structure of Doped MO2, 44 3.2.2 Conductivities of Oxygen Ions, Electrons,

and Electron Holes, 50 3.2.3 Defect Domains, 52 3.2.4 Defect Associations and Clusters, 54

15

41

x Contents

Chapter 4

Chapter 5

Chapter 6

3.3 Defects in Perovskite-Type Oxides, 56 3.4 Conduction Processes and Transference Numbers, 60

3.4.1 General Transport Equations, 60 3.4.2 Electronic, Ionic, and Total Current, 61 3.4.3 Transference Number Measurements, 62

References, 67

ELECTROLYTE

4.1 Requirements, 69 4.2 Stabilized Zirconia, 70

4.2.1 Preparation, 71 4.2.2 General Properties and Phase Transformation, 74 4.2.3 Stability, 76 4.2.4 Electrical Conductivity, 78 4.2.5 Chemical Interaction, 87 4.2.6 Thermal Expansion, 87 4.2.7 Mechanical Properties, 90 Doped Ceria, 92 Stabilized Bismuthsesquioxide, 96 Other Oxygen-Ion Conductors, 101 Protonic Conductors, 102

References, 107

4.3 4.4 4.5 4.6

CATHODE

5.1 Requirements, 117 5.2 Lanthanum Manganite, 118

5.2.1 Preparation, 118 5.2.2 General Properties, Phase Transformation,

and Stoichiometry, 120 5.2.3 Stability, 126 5.2.4 Electrical Conductivity, 127 5.2.5 Chemical Interaction, 132 5.2.6 Thermal Expansion, 135 5.2.7 Other Properties, 137

5.3 Lanthanum Cobaltite, 138 5.4 Other Materials, 139 References, 140

ANODE

6.1 6.2

Requirements, 147 Nickel/Yttria-Stabilized Zirconia Cermet, 148

69

117

147

Contents xi

Chapter 7

Chapter 8

Chapter 9

6.2.1 Preparation, 149 6.2.2 Stability, 153 6.2.3 Electrical Conductivity, 156 6.2.4 Chemical Interaction, 158 6.2.5 Thermal Expansion, 159

6.3 Other Materials, 160 References, 161

INTERCONNECT

7.1 Requirements, 165 7.2 Lanthanum Chromite, 166

7.2.1 Preparation, 166 7.2.2 General Properties, Phase Transformation,

and Stoichiometry, 168 7.2.3 Stability, 171 7.2.4 Electrical Conductivity, 172 7.2.5 Chemical Interaction, 180 7.2.6 Thermal Expansion, 181 7.2.7 Sinterability, 183 7.2.8 Gas Permeability, 188

7.3 Other Materials, 188 References, 191

165

ELECTRODE REACTION 199

8.1 General, 199 8.2 Reactions at Anode, 200

8.2.1 Electrochemical Oxidation of Hydrogen, 200 8.2.2 Electrochemical Oxidation of Carbon Monoxide, 208 8.2.3 Reaction of Sulfide Impurities, 209 8.2.4 Reforming of Hydrocarbons, 210

8.3 Reactions at Cathode, 212 8.3.1 Oxygen Reduction at Metal Electrode, 214 8.3.2 Oxygen Reduction at Oxide Electrode, 216

References, 225

STACK DESIGN AND FABRICATION

9.1 9.2

General, 233 Sealless Tubular Design, 235 9.2.1 Design Features, 235 9.2.2 Advantages and Disadvantages, 239 9.2.3 Fabrication, 240

233

xii Contents

9.2.4 Performance and Technological Status, 252 9.3 Segmented-Cell-in-Series Design, 255

9.3.1 Design Features, 255 9.3.2 Advantages and Disadvantages, 259 9.3.3 Fabrication, 260 9.3.4 Performance and Technological Status, 266

9.4 Monolithic Design, 268 9.4.1 Design Features, 268 9.4.2 Advantages and Disadvantages, 270 9.4.3 Fabrication, 272 9.4.4 Performance and Technological Status, 280

9.5 Flat-Plate Design, 282 9.5.1 Design Features, 282 9.5.2 Advantages and Disadvantages, 287 9.5.3 Fabrication, 287 9.5.4 Performance and Technological Status, 293

References, 297

Chapter 10 MODELING AND ANALYSIS

10.1 General, 307 10.2 Stress Analysis, 308 10.3 Electrical Analysis, 314 10.4 Modeling of Current and Temperature Distribution, 321 References, 328

307

Chapter 11 SYSTEM AND APPLICATION

11.1 General, 331 11.2 Electric Utility, 332 11.3 Cogeneration, 335 11.4 Transportation, 338 11.5 Space and Other Applications, 343 References, 347

331

Appendix. Selected References Relevant to Solid Oxide Fuel Cell Technology 351

Reprint Permission 355

Index 357

Chapter 1

INTRODUCTION

1.1 SCOPE

Fuel cells are a radically different way of making electrical power from

a variety of fuels. A fuel cell is an energy conversion device that produces electricity (and heat) directly from a gaseous fuel by electrochemical combination of the fuel with an oxidant. Such a device bypasses the conversion of chemical energy of fuel into thermal and mechanical energy, and thus achieves theoretical

efficiency significantly higher than that of conventional methods of power generation. In addition to the high conversion efficiency, fuel cells have the characteristics of environmental compatibility, modularity, siting flexibility, and multifuel capability.

(i) High conversion efficiency: The primary feature of a fuel cell is its

high fuel-to-electricity conversion efficiency (45 to 60%). A fuel cell converts the chemical energy of fuel directly into electrical energy. Thus, the usual losses involved in the conversion of fuel to heat, to mechanical energy, and then to electrical energy are avoided. The efficiency of a fuel cell is further improved

when the byproduct heat is fully utilized (in cogeneration or bottoming cycles). (ii) Environmental compatibility: Fuel cells are capable of using practical

fuels as an energy source with insignificant environmental impact. Emissions of key pollutants from fuel cells are several orders of magnitude lower than those produced by conventional power generators. Production of undesirable materials

such as NOx, SOx, and particulates is either negligible or undetectable for fuel cell systems (for examples, see Figures 11.1 and 11.6, Chapter 11).

(iii) Modularity: Fuel cells have the characteristic of modularity, i.e.,

cells can be made in modular sizes. Thus, fuel cell size can be easily increased

or decreased. Since the efficiency of a fuel cell is relatively independent of size,

fuel cells can be designed to follow loads with fast response times without significant efficiency loss at part-load operation.

2 Chapter 1

(iv) Siting flexibility: Because fuel cells can be made in a variety of sizes, they can be placed at different locations with minimum siting restrictions.

Fuel cell operation is quiet because a fuel cell has no moving parts; the only noises are those from auxiliary equipment. Consequently, fuel cells can be easily located near points of use such as urban residential areas.

(v) Multifuel capability: Certain types of fuel cells have multifuel capability. High-temperature fuel cells can process (reform) hydrocarbon fuels internally and do not need expensive subsystems to process conventional fuels into simple forms.

Ceramic fuel cells having the attributes discussed above are among the several fuel cell technologies being developed for a broad spectrum of electric

power generation applications. The key characteristic of this type of fuel cell is its ceramic electrolyte. The use of a solid electrolyte in ceramic fuel cells eliminates material corrosion and electrolyte management problems and permits unique cell designs with performance improvements. The conductivity requirement for the ceramic electrolyte necessitates high operating temperatures (600 ~ to 1000~ High operating temperatures promote rapid reaction kinetics, allow reforming of hydrocarbon fuels within the fuel cell, and produce high- quality byproduct heat suitable for use in cogeneration or bottoming cycles. On

the other hand, high operating temperatures impose stringent material and processing requirements. The present key technological challenge facing ceramic fuel cells is the development of suitable materials and fabrication processes to incorporate materials into required structures.

To date, although ceramic fuel cell technology is still evolving, it has made excellent technical progress. Multikilowatt fuel cells incorporating various features of a practical power generation system have been operated for thousands of hours and have shown excellent performance. Recently, ceramic fuel cell

research and development has received much attention, reflecting widening interest in this technology.

The objectives of this book are to provide a comprehensive treatise on both the fundamental and technological aspects of ceramic fuel cells and to serve

as a reference source on this emerging technology. This book consists of eleven chapters. Chapter 1 introduces the general characteristics of ceramic fuel cells and gives a brief historical perspective on the development of this type of fuel cell. Chapter 2 provides an overview of the operating principles, with brief

discussions on thermodynamic aspects of cell operation, key features of the fuel cell, types of fuel and oxidant, and other elements of a fuel cell power system.

Introduction 3

Since the operation of a ceramic fuel cell is based fundamentally on electrical

processes in the ceramic components, it is instructive to outline some of the

relevant theoretical considerations on electrical conduction in ceramics. Chapter

3 thus includes discussion on the general principles of electrical conduction, the

relationships between conduction and defect structure, and transference numbers

of ions and electrons in ceramics. Chapters 4, 5, 6, and 7 cover the principal

components of a fuel cell stack: the electrolyte, the cathode, the anode, and the

interconnect, respectively. Emphasis is given to the discussion on the preparation

of each component material, its stability, and its chemical, electrical, and thermal

properties under cell fabrication and operation conditions. Chapter 8 reviews

various aspects of the electrode reactions, including reforming and contaminant

reactions, in a ceramic fuel cell. The discussion in this chapter focuses on the

reaction mechanisms of the hydrogen oxidation and the oxygen reduction.

Chapter 9 is devoted to a thorough discussion of stack design and fabrication.

Detailed description of design characteristics, gas manifolding, and fabrication

processes, along with a summary of the technological status of each stack design,

are presented. Chapter 10 provides a treatment of modeling and analysis used

in ceramic fuel cell design, especially thermal stress analysis, electrical analysis,

and performance modeling of various cell and stack configurations. Finally, the

applications of ceramic fuel cells are discussed in Chapter 11.

1.2 GENERAL CHARACTERISTICS OF CERAMIC FUEL CELLS

A ceramic fuel cell is an all-solid-state energy conversion device that produces electricity by electrochemically combining fuel and oxidant gases across

an ionic conducting ceramic. A ceramic fuel cell consists of two electrodes (the anode and cathode) separated by a solid electrolyte. Fuel is fed to the anode,

undergoes an oxidation reaction, and releases electrons to the external circuit. Oxidant is fed to the cathode, accepts electrons from the external circuit, and

undergoes a reduction reaction. The electron flow (from the anode to the

cathode) produces direct-current electricity (Figure 1.1) [1.1]. The solid

electrolyte conducts ions between the two electrodes.

Present ceramic fuel cells use exclusively hydrogen as fuel, and oxygen

as oxidant. In theory, any gases capable of being electrochemically oxidized and

reduced can be used as fuel and oxidant in a fuel cell. However, hydrogen is

currently the most common fuel, since it has high electrochemical reactivity and

can be derived from common fuels such as hydrocarbons, alcohols, or coal.

4 Chapter 1

FUEL~

, ox,o..

ANODE

ELECTROLYTE

CATHODE

EXTERNAL LOAD AND HEAT /

Figure 1.1. Schematic diagram of fuel cell operation [1.1]

Oxygen is the most common oxidant, since it is readily and economically

available from air. For the hydrogen/oxygenreaction, to date, only oxides are being considered for use as ceramic fuel cell electrolytes. Since fuel cells are

commonly identified by the type of electrolyte used, ceramic fuel cells are referred to as solid oxide fuel cells (SOFCs). Due to the conductivity requirement for the oxide electrolyte, current SOFCs operate in the temperature range

of 600 ~ to IO00~

1.2.1 Types of ceramic fuel cells

A fuel cell electrolyte must ionically conduct one of the elements present

in the fuel or oxidant. Thus, a solid electrolyte for SOFCs based on the

electrochemical reactions of hydrogen and oxygen must conduct either oxygen

ions or hydrogen ions (protons). (Although hydroxide-ion conduction is also

possible, it has been shown to be a proton conduction with oxygen-ion carrier

species. It is a special case and, for simplification,it will be considered as proton

conduction here.) The present generation of ceramic fuel cells can be classified

into two types [1.2]: (i) those based on oxygen-ion-conducting electrolytes and

(ii) those based on proton-conducting electrolytes. Figures 1.2 and 1.3 show the reactions in an oxygen-ion-conductor SOFC and a proton-conductor SOFC,

respectively. An oxygen-ion-conductor SOFC can be considered as an oxygen

concentration cell, and a proton-conductor SOFC as a hydrogen concentration

Introduction 5

FUEL .... ~ / H2, CO. H20,C;2

0 H2'CO / 1, / ~2 / / ANODE H~O. H20 + C02

+H20~C02 + H2 ~ / e ' . ~ "~ .... C O ~ ~ e - ~ "

INTERFACE ,0 = ,,0: 0 = + H 2 ----~H20+ 2e-

~EXCESS FUEL~ "1 TO BURNE~/

e-

OUTER CIRCUIT

,,THO0, ~ 0 2 + 4e- )20 = 132-t 02/ | / e"

/ / /

AIR OR 02 " \ 02 CATHOD EXHAUST/

REACTION H2 + 1/202 ~- H2O CO + 1/202 = C02

Figure 1.2. Schematic diagram of reactions in SOFCs based on oxygen-ion conductors [1.1]

FUEL H2 " ~

ANODE H 2 ~ 2H + + 2e-

/

I / H2

H~ ..... -e~''~' !N_TERFACE_---~

2H + +I/202 + 2e--*H20

CATHOOE

OXIDANT AIR OR 02

,H +

, \r

02 H20

H2

H+~ j e 1 I H OUTER ~, _ [, CIRCUIT

, ! HHo ,,..-

OVERALL REACTION H2 + V202 ~ H20

EXCESS FUER~ "-I TO BUR.NE

J CATHOOE~ "-I EXHAUST

Figure 1.3. Schematic diagram of reactions in SOFCs based on proton conductors [1.1]

6 Chapter 1

cell. The major difference between the two SOFC types is the side in the fuel

cell in which water is produced (the fuel side in oxygen-ion conductor cells and

the oxidant side in proton-conductor cells). Also, certain gases, such as CO, can

be used as fuel in oxygen-ion conductor SOFCs but not in proton-conductor

SOFCs. To date, almost all of the development work on ceramic fuel cells has

focused on SOFCs with oxygen-ion-conducting ZrO2 electrolytes. Work on

proton-conductor SOFCs is limited to material studies, clarification of conduction

mechanisms, and testing of small, laboratory-scale cells.

1.2.2 Cell components

A SOFC single cell consists of an oxide electrolyte sandwiched between

an anode and a cathode. Under typical operating conditions (with hydrogen fuel

and oxygen oxidant), a single cell produces less than 1 V. Thus, practical

SOFCs are not operated as single units; rather, they are connected in electrical

series to build voltage. A series of cells is referred to as a stack. A component,

variously called an interconnect or a bipolar separator, connects the anode of one

cell to the cathode of the next in a stack (Figure 1.4). SOFC stacks can be

configured in series, parallel, both series and parallel, or as single units,

depending on the particular application.

REPEATING ELEMENTS

INTERCONNECT

ANODE

ELECTROLYTE

CATHODE

Figure 1.4. Fuel cell component

Introduction 7

The principal components of a SOFC stack are the electrolyte, the anode,

the cathode, and the interconnect. Each component serves several functions in the fuel cell and must meet certain requirements. Each component must have the

proper stability (chemical, phase, morphological, and dimensional) in oxidizing

and/or reducing environments, chemical compatibility with other components,

and proper conductivity. The components for ceramic fuel cells must, in

addition, have similar coefficients of thermal expansion to avoid separation or

cracking during fabrication and operation. The electrolyte and interconnect must

be dense to prevent gas mixing, while the anode and cathode must be porous to

allow gas transport to the reaction sites. The requirements for the various cell

component are summarized in Table 1.1.

In addition to the requirements listed in Table 1.1, other desirable properties for the cell components from practical viewpoints are high strength

and toughness, fabricability, and low cost. Also, for certain cell designs, the

components for a ceramic fuel cell must be amenable to limited fabrication

conditions since the process conditions cannot be selected independently for each

component. For example, if the components are built up one by one, the

temperature of sintering for each successive component must be lower than that

of the preceding component to avoid altering the microstructure of the preceding

component. If the components are formed in the green state, then all components

must be sintered under the same firing conditions. Furthermore, the components

of a ceramic fuel cell must be compatible not only at the operating temperature but also at the much higher temperatures at which the ceramic structures are fabricated.

Cell components are connected (in electrical series) in proper order in a

stack. The height or number of single cells (thus, voltage) and footprint or active

area (thus, current) of a stack can vary, depending on the particular design and

power output required. Because all the components are solid, the SOFC stack

can be configured into unique shapes unachievable in other types of fuel cells.

At present, four common stack configurations have been proposed and fabricated

for SOFCs: the sealless tubular design, the segmented-cell-in-series design, the

monolithic design, and the flat-plate design (for more details, see Chapter 9).

Each design may have several different versions and is presently at a different

stage of technology development. Figure 1.5 shows, as an example, the

schematic diagrams of the various SOFC stack designs [1.3].

8 C

ha

pter

1 TABLE 1.1

Requirements for Ceramic Fuel Cell Components

- - - - - --

Requirements Component

Conductivity Stability Compatibility Porosity Thermal Expansion

Electrolyte High ionic conductivity Chemical, phase, morphological, No damaging chemical Fully Thermal expansion Negligible electronic con- and dimensional stability in fuel interactions or interdif- dense match with adjoin- ductivity and oxidant environments fusion with adjoining cell ing components

components

Cathode High electronic conductivity Chemical, phase, morphological, No damaging chemical Porous Thermal expansion and dimensional stability in oxi- interactions or interdif- match with adjoin- dant environment fusion with adjoining cell ing components

components

Anode High electronic conductivity Chemical, phase, morphological, No damaging chemical Porous Thermal expansion and dimensional stability in fuel interactions or interdif- match with adjoin- environment fusion with adjoining cell ing components

components

Interconnect High electronic conductivity Chemical, phase, morphological, No damaging chemical Fully Thermal expansion Negligible ionic conductivity and dimensional stability in fuel interactions or interdif- dense match with adjoin-

and oxidant environments fusion with adjoining cell ing components components

Introduction 9

INTERCONNECTION

ELECTRODE

, ~ . I POROUS S ~

AIR FLOW ~ - ~ ~ FUEL ELECTRODE

ELECTROLYTE

ELECTROLYTE CATHODE INTERCONNECT'~~

/ ~ i l i \i : : Po.ous s0PPo.T i : :::i: : : : : i : /

~U )) ) ) ' ~ /) /)_ ) --" OXIDANT -----e,

Seal-less Tubular Design Segmented-Cell-in-Series Design

ELECTROLYTE

Monolithic Design Flat-plate Design

Figure 1.5. SOFC stack designs [1.3]

10 Chapter 1

1.2.3 Comparison with other types of fuel cells

The SOFC is one of several types of fuel cells currently under develop-

ment for clean and efficient electric power generation from a variety of fuels.

Besides the SOFC, the other major types of fuel cells are polymer membrane,

alkaline, phosphoric acid, and molten carbonate fuel cells. Among these fuel

cells, the phosphoric acid fuel cell is presently at the initial stage of commer-

cialization for electric utility and cogeneration uses. The molten carbonate is the

next most likely candidate for commercialization, whereas the SOFC is

considered as the third-generation technology. The polymer membrane fuel cell

is being developed mainly for space and transportation applications, and the

alkaline fuel cell is an important power source for space flights. Typical features

and operational characteristics of the SOFC and other types of fuel cells are listed

in Table 1.2.

1.3 HISTORICAL BACKGROUND OF CERAMIC FUEL CELLS

The principles of fuel cell operation were first reported by Sir William

Grove in 1839 [1.4]. His fuel cell used dilute sulfuric acid as the electrolyte and

operated at room temperature. Ceramic fuel cells came much later and began

with Nernst's discovery of solid-oxide electrolyte in 1899 [1.5] and the operation

of the first ceramic fuel cell at 1000~ by Baur and Preis in 1937 [1.6].

Nernst discovered solid oxygen-ion conductors when he invented the so-

called glower in the end of the 19th century [1.5]. Nernst proposed to use solid

compositions such as ZrO2 with 15 wt% Y203 addition (called the Nernst mass)

as a glower to replace carbon filaments in electric lamps. The Nernst glower

was operated for hundreds of hours on direct current, though electrolysis was

found to occur. It was explained that any loss of oxygen liberated at the anode

was balanced by an equal amount of oxygen taken into the glower at the cathode.

This phenomenon was the reverse of fuel cell operation. In 1935, Schottky

published a paper suggesting that the Nernst mass could be used as a fuel cell

solid electrolyte [1.7].

In 1937 Baur and Preis demonstrated the operation of the first ceramic

fuel cells [1.6]. They used mainly ZrO2-based ionic conductors (e.g., ZrO2 with

10 wt % MgO or 15 wt % Y203 addition) in the form of a tubular crucible as the

electrolyte, with iron or carbon as the anode and Fe304 as the cathode. Observed

open-circuit voltages were between 1.1 and 1.2 V at 1000 ~ to 1050~ Baur and

Introdu

ction

11 TABLE 1.2

Typical Features and Operational Characteristics of SOFC and Other Types of Fuel Cells

- -- - -- - -

Type of fuel cell

Solid oxide Molten carbonate Phosphoric acid Alkaline Polymer membrane

Electrolyte Solid Y20,-stabilized Molten Li,CO,- H,PO,

zrOz (YSZ) K2C03

KOH solution Perfluorosulfonic acid membrane

Electrolyte support None LiAlO, S ic Asbestos None

Cathode Sr-doped LaMnO, Li-doped NiO PTFE'-bonded Pt on C Pt-Au PTFE-bonded Pt on C

Anode NiIYSZ Ni PTFE-bonded Pt on C Pt-Pd PTFE-bonded Pt on C

Interconnect/Bipolar Doped LaCrO, SS" clad with Ni Glassy carbon Ni Graphite

Operating temperature 1000°C 650°C 200°C 100°C 80°C

Operating pressure 1 atm"' 1 to 3 arm 1 to 8 arm 1 to 10 atm 1 to 5 atm

Fuel H,. CO Hz, CO H? Hz H2

Oxidant 0 2 O? + CO, 0, 0, 0,

Contaminant tolerance < 10-100 ppm sulfur < ppm sulfur < l t 0 2 % C O No C02, CO < 50 ppm CO

< 50 ppm sulfur No sulfur No sulfur

'PTFE = Polytetrafluoroethylene; "SS = stainless steel; "'atm = 1.01 x 1@ Pa

12 Chapter 1

Preis constructed a ceramic fuel cell battery consisting of eight ZrO2-Y203 crucibles filled with coke and immersed in a common magnetite bath. With

hydrogen, CO, or town gas as fuel, the open-circuit voltage was 0.8 V per cell

(0.2 V lower than the theoretical value). At a current density of approximately 0.3 mA/cm 2, the cell voltage was 0.65 V, corresponding to an imernal resistance

of 1.8 to 2.6 ~. Although operation was demonstrated, the current outputs of these cells were too low to be practical.

Initial development work on practical ceramic fuel cells began in the early 1960s. The cell configuration in this time period was either a flat-plate

design using the electrolyte in the form of a disk, or a segmented-cell-in-series design (bell-and-spigot configuration) using short tubular segments of the

electrolyte joined together with conducting seals. These designs used very thick electrolytes, thus suffering significant internal resistance losses. This led to the

development of the thin-wall concept to improve cell performance. In 1970s the banded configuration (a segmented-cell-in-series design) was proposed, which

made use of the thin-wall concept in which a number of thin-film cells were deposited on a porous support. Development of fuel cells based on this

configuration is still going on; kilowatt-size stacks of banded SOFC cells have been tested. In 1980 the sealless tubular design was proposed, with several

advantages over the segmented-cell-in-series design. The key features of the sealless tubular design include individual thin cells formed on a tubular support

and electrically connected into a bundle in a fuel-reducing atmosphere. This design is presently the most advanced; multikilowatt sealless tubular SOFC generators have been fabricated and operated for thousands of hours. In 1982 the monolithic design, in which cells are configured in a honeycomb structure (resulting in extraordinarily high power density), was advanced. At the same time, interest in the fiat-plate design has been renewed, and due to many

advances in ceramic forming and processing technologies, various advanced fiat-

plate concepts have been proposed.

Early SOFC stacks used noble metals (e.g., platinum) as electrode and

interconnect materials. In the early 1970s nickel/YSZ, doped In203, and

CoCr204 were used as anode, cathode, and interconnect, respectively. CoCr204

was later replaced by LaCrO3, and in 1980 LaMnO3 and LaCoO 3 were proposed

for cathode use. Recently, high-temperature alloys have been tested as interconnect material for flat-plate SOFCs. Figure 1.6 summarizes key

historical events in the development of the SOFC technology.

Introdu

ction

13

ADVANCED FLAT-PUTE SOFC I

MONOLITHIC SOFC

I I SEALLESS TUBULAR SOFC I

SEGMENTED-CELL-IN-SERIES SOFC (BANDED CONFl URATION) '1 I

SEGMENTED-CELL-IN-SERIES SOFC [BELL-AND-SPIGOT)

I I I CELL OPERATION FLAT-PLATE SOFC

DISCOVERY OF SOLID ELECTROLYTE IBAUR AND PREIS) I I I I

I YSZ ELECTROLYTE

INTERCONNECT

Ni/YSZ ANODE

In,O, CATHODE

CoCr,O, INTERCONNECT ' I

LaCrO, INTERCONNECT I

b M n O , CATHODE

LaCoO, CATHODE

t METALLIC INTERCONNECT

Figure 1.6. Key historical events in SOFC technology development

14 Chapter 1

References

1.1

1.2 1.3

1.4 1.5 1.6 1.7

Morgantown Energy Technology Center, Fuel Cells -- Technology Status Report,

Report No. DOE/METC-87/0257, Morgantown Energy Technology Center, Morgantown, WV, 1986. N.Q. Minh, J. Am. Ceram. Soc., 76 (1993) 563. N.Q. Minh, in Science and Technology of Zirconia V, S.P.S. Badwal, M.J. Bannister, and R.H.J. Hannink (eds.), Technomic Publishing Company, Lancaster, PA, 1993, p. 652. W.R. Grove, Philos. Mag., 14 (1839) 127. W. Nernst, Z. Elektrochem., 6 (1899) 41. E. Baur and H. Preis, Z. Elektrochem.,43 (1937) 727. W. Schottky, Wiss. VerOff. Siemens Werken, 14 (1935) 1.

Chapter 2

PRINCIPLES OF OPERATION

2.1 GENERAL

A solid oxide fuel cell (SOFC) is a ceramic device that converts the

chemical energy of a fuel gas and an oxidant gas directly to electrical energy

without combustion as an intermediate step. The operating principles of a SOFC

and other types of fuel cells are similar to those of a battery, i.e., electrochemical

combination of reactants to generate electricity. However, unlike a battery, a

SOFC does not run down or require recharging; the fuel cell employs gases

(from an external source) as reactants, and operates as long as both fuel and

oxidant are supplied to the electrodes. In fuel cells, the electricity generation

mechanism is based on the electrochemical combustion of the fuel; i.e., the

overall reaction is the same as that of the combustion; however, the reaction is

made of two separate electrochemical reactions. For example, for hydrogen fuel

and oxygen oxidant, the electrochemical combustion reaction consists of the

oxidation of hydrogen at the anode and the reduction of oxygen at the cathode.

The overall reaction in this case, like combustion, yields water as the reaction

product. In the operation of a SOFC, fuel (e.g., hydrogen) is fed to the anode,

where it is oxidized and electrons are released to the external (outer) circuit.

Oxidant (e.g., oxygen) is fed to the cathode, where it is reduced and electrons

are accepted from the external circuit. The electron flow (from the anode to the

cathode) through the external circuit produces direct-current (DC) electricity.

The electrochemical transformation of the fuel and oxidant in a SOFC is

isothermal; i.e., the fuel cell directly uses the available free energy in the fuel at

the operating temperature.

Direct conversion of fuel energy to electricity is the key characteristic of

fuel cell operation. In a conventional thermal power system, the chemical energy

of the fuel is transformed first to thermal energy, then to mechanical energy, and

finally to electrical energy. Other' energy systems such as magnetohydrodynamic

(MHD) generators and thermionic converters involve the conversion of chemical

16 Chapter 2

energy to thermal energy and then to electricity. The efficiency of the thermal-

to-mechanical and thermal-to-electrical energy conversions is subject to the

Carnot limitation. The Carnot efficiency of a combustion-type system operated

between high-temperature (Th) and low-temperature (T/) heat sources is given as

e = 1 Tt (Eq. 2.1)

where e is the efficiency and T is the temperature in K. To obtain a high value

of E, a T h as high as possible and a TI as low as possible are desirable. However,

there is a practical limit on this efficiency due to high limits on Th (temperature

of material stability) and low limits on ~ (room temperature). The operation of

a fuel cell is not Carnot-limited. A fuel cell converts chemical energy directly

to electricity and thus can yield a higher efficiency than a combustion-type

conversion system.

For practical power generation applications, a SOFC must use commer-

cially available fuels and produce alternate-current (AC) electricity. Therefore,

besides the fuel cell power section, a SOFC system has two other main

components: a fuel processor and a power conditioner. The fuel processor

converts a practical fuel to a suitable gas that is then fed to the fuel cell stack.

The power conditioner uses solid-state technology to efficiently convert DC

electricity to AC.

2.2 THERMODYNAMIC PRINCIPLES

The operation of a SOFC involves the reduction of the oxidant at the

cathode and the oxidation of the fuel at the anode. At present, the most common

fuel and oxidant for use in SOFCs are hydrogen and oxygen (electrochemically

active). Thus, for SOFCs having an oxygen-ion-conducting electrolyte, the

reactions in the fuel cell involve the oxidation and reduction of oxygen at the

electrodes. Similarly, for SOFCs having a proton-conducting electrolyte, the

reactions involve the oxidation and reduction of hydrogen.

At the cathode, the reduction of oxygen in a SOFC based on an oxygen-

ion-conducting electrolyte is given as

2- (Eq. 2 2) O2(c) + 4e- = 2 0 ( e )

Principles of Operation 17

where the subscripts (c) and (e) represent the states at the cathode and in the electrolyte, respectively. At the anode, the reverse reaction of Eq. 2.2 can be regarded thermodynamically as the primary electromotive reaction, i.e.,

2 - 20(~) = 02(a) -I- 4e- (Eq. 2.3)

where the subscript (a) represents the state at the anode. Consequently, the overall cell reaction (which determines cell voltage) is represented by the

following equation:

O2(c) : O2(a) (Eq. 2.4)

The SOFC is therefore considered to be an oxygen concentration cell, and the electromotive force (emf) or reversible (thermodynamic) voltage, Er, is

given by the Nernst equation

P E , - R T ln o~) (Eq. 2.5)

4F P o2(a)

where R is the gas constant, T the temperature, F the Faraday, and P o 2 the partial pressure of oxygen at the electrode.

For a certain oxygen partial pressure at the cathode, the magnitude of Er

depends on the anode oxygen partial pressure, thus on the type and composition of the fuel fed to the anode. For example, when CO is fed to the anode, the following reaction takes place at the anode:

C O ( a ) + 1/~ O2(a ) -- CO2(a) (Eq. 2.6)

The oxygen partial pressure at the anode is given by

�9 Po2(a) PCO(a)K(2.6)

(Eq. 2.7)

where K(2.6 ) is the equilibrium constant of Eq. 2.6. Substituting the equation for

the anode oxygen partial pressure (Eq. 2.7) into Eq. 2.5 yields

R T R T gco(a) E r = E 0 + ~ LrlPo2~c ) + In (Eq. 2.8)

4F 2F Pc02<o)

18 Chapter 2

where E ~ is the reversible voltage at the standard state and is given as

EO - RT 2F INK(26)

Thus, Eq. 2.8 gives the reversible voltage of the cell

(Eq. 2.9)

CO2(a) , CO(a)I Oxygen-ion conductor [ O2(c) (Eq. 2.10)

with the following cell reaction

2CO(a ) -+- O2(c) = 2CO2(a) (Eq. 2.11)

Similar equations can be obtained when other gases are used as the fuel

in the anode. For example, for hydrogen fuel, the cell reaction is

H2(a) d- 1/~O2(c)= H20(a ) (Eq. 2.12)

The reversible cell voltage is given as

RT RT PI'I2(a) E~ = E ~ + lnP%,) + - - l n ~ (Eq. 2.13)

, 4F 2F P~o<o)

where E ~ is the standard cell voltage.

At the standard state, Er equals E ~

established for every fuel:

and the following equation is

EO AG O AH ~ - TAS ~ - - (Eq. 2.14) zF zF

where AG O is the standard Gibbs free energy change of the combustion reaction

of the fuel, AH ~ the standard enthalpy change, AS ~ the standard entropy change,

and z the number of moles of oxygen required to oxidize one mole of fuel

multiplied by four. The maximum energy obtained in this case is given by -AG O ,

and the ideal thermodynamic efficiency eT is represented by AG~ Table 2.1

lists AG ~ A//~ E ~ and eT for the combustion of several fuels (when the fuel cell

reaction products are in gaseous state). (The efficiency is referred to as low

heating value (LHV) when the H20 product is in the form of steam and as high

heating value (HHV) when the H20 is produced in the liquid state.)


TABLE 2.1

Thermodynamic Data and Thermodynamic Efficiencies for Several Cell Reactions

Reaction T, K AG ~ kJ ~ , kJ E ~ V cv

H2 + 1AO2 = H20 1000 -192.5 -247.3 0.997 0.78

1250 -178.2 -249.8 0.924 0.71

CO + 1AO2 = CO2 1000 -195.4 - 2 8 3 . 3 1.013 0.69 1250 -173.2 - 2 8 3 . 3 0.898 0.61

CH 4 + 202 = C O 2 -Jr- 2H20 1000 -802.5 -800.4 1.039 1.00

1250 -802.9 -801.2 1.039 1.00

C + 02 = CO2 1000 -396.6 -396.2 1.027 1.00 1250 -396.6 -396.6 1.027 1.00

The effect of temperature and pressure on the reversible voltage of a fuel

cell can be analyzed based on changes in the free energy with temperature and

pressure. Thus, the following equations can be derived:

OEr AS (Eq. 2.15) ( - ~ ) P - zF

OEr A V (Eq. 2.16) ( - - ~ ) r - zF

where A V is the volume change. For the hydrogen/oxygen reaction, the entropy

change is negative; therefore, the reversible or thermodynamic voltage of the fuel

cell decreases with increasing temperature. For the same reaction, the volume

change is negative; therefore, the reversible voltage increases with increasing

pressure.

An increase in the operating temperature is beneficial to fuel cell

performance because of enhanced mass transfer, increased reaction rate, and

usually lower material resistance (thus reduced cell polarization). Increase in

operating temperature, on the other hand, may limit material choice and

20 Chapter 2

accelerate material-related problems such as interaction, degradation, and

sintering. An increase in the operating pressure improves fuel cell performance because of increased reactant partial pressure and increased mass transport rate.

On the other hand, increase in pressure also imposes limits on material selection and causes other problems or concerns such as integrity of the gas seal and cell structure.

In reversible operation, the heat absorbed by the fuel cell at constant

temperature equals the change in entropy of the reactants

Q, = TAS = A H - A G (Eq. 2.17)

If the change in entropy is negative, as is the case in most common fuel cell

reactions, heat is generated by the cell. In practical fuel cell operation, this heat

increases the operating temperature. Performance losses due to higher temperatures are generally avoided by cooling the cell to remove the heat

generated.

2.3 FUEL CELL EFFICIENCY

The overall efficiency of a SOFC, EFC , is the product of the electro-

chemical efficiency 6E and the heating value efficiency ell. The electrochemical efficiency is, in turn, the product of the thermodynamic or thermal efficiency er, the voltage efficiency ev, and the Faradaic or current efficiency ej of the fuel cell. Thus,

%c = e~ea = exeveJea (Eq. 2.18)

A fuel cell system consists of, in addition to the fuel cell, the fuel processor and the power conditioner. Thus, the efficiency of a SOFC system without byproduct

heat utilization, es, is given as

r =r e'av eve (Eq. 2.19)

where EFp and Epc are the efficiencies of the fuel processor and power conditioner,

respectively. If the byproduct heat is used in a bottoming cycle and/or for

cogeneration, the system efficiency 6s/H is then given by the ratio of the total AC power plus bottoming cycle and cogeneration heat to the high heating value

(HHV) of the fuel fed to the fuel processor. The various efficiencies mentioned

above are generally independent of each other and can be optimized separately.


2.3.1 Electrochemical efficiency

The electrochemical efficiency of a SOFC includes three elements: the

thermodynamic efficiency, the voltage efficiency, and the current efficiency

(often referred to as the fuel utilization). (i) Thermodynamic efficiency: In a SOFC and other types of fuel cells,

the Gibbs free energy change of the cell reaction may be totally converted to

electrical energy. Thus, a fuel cell has an intrinsic (maximum) thermodynamic

efficiency given by

A G TAS er - AH - 1 AH (Eq. 2.20)

For common fuels such as hydrogen, CO, and hydrocarbons (negative AH and

negative AS), the thermodynamic efficiency of a SOFC is < 1. (eT of >_ 1 is

conceptually possible for cell reactions with positive AS.)

(ii) Voltage efficiency: In an operating SOFC, the cell voltage is always less than the reversible voltage. As the current is drawn from the fuel cell, the cell voltage falls, due to various losses. The reduction in the cell voltage under current load depends on current density and several factors such as temperature,

pressure, gas flow rate and composition, and cell material. (This reduction is not

characteristic of high-temperature SOFCs, but is common not only to all types

of fuel cells but also to all electrochemical cells.) The voltage efficiency, ev, is

defined as the ratio of the operating cell voltage under load, E, to the equilibrium

cell voltage, Er, and is given as

E e v - (Eq. 2.21) e,

(It should be noted that the equilibrium voltage, commonly referred to as the open-circuit voltage, may be different from the reversible voltage if there are side

reactions, gas cross leakage, etc.)

The difference between the operating cell voltage and the expected

reversible voltage is termed polarization, overvoltage, or overpotential and is

presented as r/. The total polarization of a cell, r/, is the sum of four types of

polarization: charge transfer or activation polarization rlA, diffusion or

concentration polarization rio, reaction polarization r/~, and resistance or ohmic

polarization r/a:

TI = TIA + ~D + ~R + ~O (Eq. 2.22)

22 Chapter 2

Polarization cannot be eliminated but can be minimized by material modification

and cell design. Temperature, pressure, electrolyte composition, and electrode material naturally influence cell polarization. For example, increasing

temperature enhances mass transfer, increases the reaction rate, and usually decreases cell resistance, thus reducing cell polarization and increasing voltage

efficiency.

(a) Charge transfer or activation polarization" Chemical reactions including electrochemical reactions involve an energy barrier that must be

overcome by the reacting species. This energy barrier, called the activation energy, results in activation or charge transfer polarization, r/A. The activation

polarization may be regarded as the extra potential necessary to reduce the energy barrier of the rate-determining step of the reaction to a value such that the

electrode reaction proceeds at a desired rate. Activation polarization is related to current density, j, by the following equation:

J =joexp RT J - j ~ ~-~ (Eq. 2.23)

where a is the transfer coefficient,, and Jo is the exchange current density. The

transfer coefficient is considered as the fraction of the change in the polarization

which leads to a change in the reaction rate constant. The exchange current density is the (equal) forward and reverse electrode reaction rate at the

equilibrium potential. High exchange current density means high electrochemical reaction rate and, thus, good fuel cell performance might be expected. The exchange current density can be determined experimentally by extrapolating plots of log j versus r/to r/ = 0. When the irreversibility of the electrode reaction is small, the second term on the right hand side of Eq. 2.23 may be neglected, and

by taking the common logarithms for both sides of Eq. 2.23, the Tafel equation

is obtained

rlA = a • b logj (Eq. 2.24)

where a and b are constants which are related to electrode material and type of

electrode reaction.

Charge transfer or activation polarization is generally due to one or more

slow rate-determining steps in the electrode reaction. The slow step could be

related to adsorption of reactant onto the surface of the electrode, electron

transfer, desorption of product, or any other step in the reaction. The electrode

reaction rate is a function of temperature, pressure, and electrode material. At


high temperatures as in the case of SOFCs, reaction rate is rapid, and as a result,

charge transfer or activation polarization is usually small.

(b) Diffusion or concentration polarization: Diffusion or concentration

polarization, riD, appears when the electrode reaction is hindered by mass

transport effects, i.e., when the feeding velocity of the reactant and/or the

removing velocity of the reaction product from the electrode is slower than that

corresponding to the discharge currentj. When the electrode process is governed

completely by diffusion (because of low concentration of reactant in the feed

gases or because of reactant conversion approaching 100 %), the limiting current,

JL, is reached (characterized by a rapid drop in cell voltage). The limiting

current can be calculated from the diffusion coefficient of the reacting ions, D,

the activity of the reacting ions, aM~§ and the thickness of the diffusion layer, di,

by applying Fick's law as

zFDaMz +

JL - (Eq. 2 25) 8

For an electrode process free of activation polarization, the diffusion or

concentration polarization is expressed as

llD _ R T l n ( I _ j_" ] (Eq. 2.26) zF t

The diffusion polarization is dependent on the mass transport properties of the system. Mass transport is a function of temperature, pressure, concentra-

tion, and the physical properties of the system. In SOFCs, the reactants must diffuse through the porous anode and cathode so the electrode structure is

important. Since the electrode reaction rate is a function of the concentration of

the reactant gases, the diffusion polarization becomes more severe as the degree

of conversion increases.

(c) Reaction polarization: The reaction polarization, r/R, appears when

the rate of the reaction to supply cell reactants or to remove products (in the

vicinity of the electrode before or after the cell reaction) is slow. This polariza-

tion is similar to the concentration polarization. At high operating temperatures,

the reaction polarization is usually small.

(d) Resistance or ohmic polarization: The ohmic polarization is caused

by resistance to conduction of ions (through the electrolyte) and electrons

(through the electrodes and current collectors), and by contact resistance between

24 Chapter 2

cell components. The ohmic polarization, r/,, is given as

TI~ - jR~ (Eq. 2.27)

where R i represents the total cell resistance, including both ionic and electronic resistances. (The resistance polarization is commonly separated from other types

of polarization and referred to as ohmic loss.)

(iii) Current efficiency: The efficiency of a SOFC falls if all of the reactants are not converted to reaction products, or if some electrons are involved

in an alternative reaction such as corrosion. For 100% conversion of fuel, the

amount of current density, j'~, produced is given as (Faraday's law)

-- zF df (Eq. 2.28) Jp dt

where f represents the amount of fuel and t the time (df/dt is the molar flow rate of the fuel). For the amount of fuel actually consumed, the current density

produced is given by

(If (Eq. 2.29)

The current efficiency, ej, is the ratio of the actual current produced to the current available from complete electrochemical conversion of the fuel

J (Eq. 2.30) e j = m

In the case of fuel cells, current efficiency is commonly expressed as fuel

utilization. Since high utilization of fuel in a fuel cell in general and a SOFC in

particular results in increased diffusion polarization, fuel cells are often operated

at less than 100 % fuel utilization.

2.3.2 Other efficiencies

In addition to the electrochemical efficiency, other efficiencies considered

in a SOFC system are the heating value efficiency and the system efficiency.

(i) Heating value efficiency." The heating value efficiency must be

considered in cases where the fuel contains inert gases, impurities, and other

combustibles in addition to electrochemically active species (hydrogen) [2.1]. In general, gases produced from practical fuel processors contain varying amounts


of inerts, particles, and combustibles. A heat engine is capable of converting all

the chemical energy of combustible materials into heat energy. Some of these may not be utilized in a fuel cell, even at the high operating temperatures of the

SOFC. The heating value efficiency, ~,, is defined as the ratio of the amount of

heat energy of fuel species available in the fuel cell to generate electricity, AH ~

to the amount of heat energy included in all combustible species in the fuel gases

fed to the fuel cell, AHco m (both at the low heating value because water product

is generally not condensed in the fuel cell)

A H ~ e a = (Eq. 2.31)

aH~om

This efficiency may be close to unity for high-temperature SOFCs operating on

simple fuels such as methane.

(ii) System efficiency: As discussed earlier, the fuel cell efficiency (~Fc)

considers only the fuel cell power section in a fuel cell system. The efficiency

of the system should include the efficiencies of the fuel processor, power

conditioner, and byproduct heat utilization. The fuel cell system efficiency, Es/.,

is expressed as

Fuel Cell(A C) +Bottoming Cycle(A6") + Cogeneration(Heat) Csm = Raw Fuel into FueIProcessor(HHlO (Eq. 2.32)

In this case, a bottoming cycle is used to generate additional electricity using fuel cell byproduct heat in a gas turbine or steam turbine. Cogeneration uses the fuel

cell byproduct heat for space heating, to supply hot water, or to generate steam

for industrial process.

2.4 P O W E R G E N E R A T I O N

A typical voltage/current plot of an operating SOFC is presented in

Figure 2.1. The figure illustrates the regions in which various types of voltage

losses predominate. From Figure 2.1, it can be seen that at low current

densities, the major contribution to the cell voltage losses is from the activation

polarization, as indicated by the sharp drop in cell voltage with increasing

current. As the current increases, the resistance polarization or ohmic loss

dominates, as exhibited by the linearity in Figure 2.1. At high current densities,

the cell resistance is controlled by mass transport limitations, resulting in a rapid

decrease in cell voltage. Under ideal conditions (i.e., all the polarization losses

26 Chapter 2

REVERSIBLE CELL VOLTAGE

~ - - f VOLTAGE LOSSES DUE TO

V o , r ~ E ,oss~s DuE To \ OHMIC DROP / , ~

. . . . . =VOLTAGE LOSSES DUE TO O,FFU sloN o VEn~TENTIAL

CELL CURRENT

Figure 2.1. Typical voltage~current relation for an operating fuel cell

are zero), one would expect constant voltage, i.e., a straight line parallel to the

current axis, as indicated by the dotted line in Figure 2.1.

The power output Pw of a fuel cell is the product of the voltage and

current

Pw -- E1 (Eq. 2.33)

where I is the current. Qualitatively, one can predict the nature of the

power/current relation by examining the behavior at low and high currents. At

low currents, I ~ 0, and at high currents, E ~ 0. Hence, the power values

approach zero at both ends. A current at which the power is at a maximum lies

between these two extremes. Figure 2.2 is a qualitative presentation of the

power/current relation in a hypothetical fuel cell. The hypothetical curve,

corresponding to no overpotential losses in the fuel cell, is also shown. It should

be stressed that the power/current plot in Figure 2.2 refers to ideal single-cell

fuel cells in which planar electrodes are used. In most typical fuel cells, owing

to the sophisticated construction of the porous gas diffusion electrodes, the

limiting current density is not reached, even at high currents, and the cell

voltage/current is approximately linear. In such cases, the power/current relation

tends to be parabolic.

It is difficult to obtain a general expression for maximum power when

all forms of overpotential are present. However, this may be done for some

limiting cases, e.g., the case where the cell voltage varies linearly with the


/ / IDEAL POWER/CURRENT

RELATION WITHOUT / o .OWE,

/ / / ' / DUE TO 0HMIC LOSSE s

/ BETWEEN ELECTRODES

Y APOcWECALOoSNS EoSvDU EpoTT~ NT, A L

MAXIMUM

POWER SHUTDOWN DUE TO MASS TRANSPORT LIMITATION~

CURRENT

Figure 2.2. Typical power~current relation for a hypothetical fuel cell

current. In this case, the equation for the cell voltage as a function of current is

given by

E = E , . - Rt/ (Eq. 2.34)

where R t represents the total polarization loss (essentially ohmic loss). Hence,

Pw = I(E~ - Rtl) (Eq. 2.35)

Eq. 2.35 shows that the power/current variation is parabolic. The conditions for

the maximum power can be deduced as

E I - ~ (Eq. 2.36)

2R t

e - Er (Eq. 2.37) 2

Then, the maximum power is given as

P _ E~ (Eq. 2.38) w(m~x) 4 R t

Thus, the maximum power is realized when the cell voltage is equal to one-half

the reversible voltage. At maximum power, voltage efficiency is 50%.

28 Chapter 2

2.5 CHARACTERISTICS OF CERAMIC FUEL CELLS

Compared with the other types of fuel cells, ceramic fuel cells, or

SOFCs, characteristically have all-solid-state construction (mainly ceramic), along with high operating temperatures (highest among the present generation of fuel cells). These two characteristics give rise to a number of distinctive features in the SOFC. For example, because all the cell components are solid, the SOFC

can be configured into compact and lightweight structures unachievable in fuel cell systems having a liquid electrolyte. In addition, the SOFC uniquely has the

characteristic of the possibility of mixed (ionic and electronic) conduction in the solid oxide electrolyte. The presence of electronic conduction in the electrolyte

can significantly influence fuel cell efficiency.

2.5.1 Features

Solid oxide fuel cells (and other types of fuel cells) offer several advantages over conventional methods of power generation: higher conversion efficiency, modular construction, high efficiency at partial loads, minimal siting restrictions, and production of far fewer pollutants. As compared with the other

types of fuel cell, SOFCs exhibit the following features [2.2]: (i) At the high operating temperatures of the SOFC, electrode reactions

are rapid; therefore, catalysts are not needed for the electrodes. On the other hand, because of these high operating temperatures, material selection is limited

and material compatibility can be a serious problem. (ii) SOFCs do not suffer from CO poisoning. In fact, CO can be used

directly as fuel in SOFCs. The SOFC can tolerate a relatively high impurity

content in the fuel. (iii) SOFCs have multifuel capability. In addition to hydrogen, suitable

fuels include gasoline, alcohol, natural gas, synthetic gas made from coal or plant

matter, and a number of other possibilities. Because of the high operating temperatures, hydrocarbon fuels can be reformed internally within the fuel cell.

Thus, expensive external fuel reforming equipment is not necessary. (iv) The SOFC produces useful high-temperature, quality heat suitable for

use in cogeneration or bottoming cycles. The overall efficiency of the system

can be significantly increased when this byproduct heat is fully utilized.


(v) Because the SOFC electrolyte is solid, electrolyte management (to prevent electrolyte loss and composition change) is not a problem. Unlike fuel

cell systems having a liquid electrolyte, SOFCs need not store excess electrolyte.

(vi) Because of its all-solid state, the SOFC can be configured into

geometries impossible in fuel cells having a liquid electrolyte. However,

fabrication of SOFC cell components into the required structure is often a technical challenge.

(vii) SOFC components exhibit relatively low electrical conductivity. Fabricated cell components must be very thin to reduce internal resistance losses.

2.5.2 Effect of electronic conduction in electrolyte

In SOFCs, the electrolyte is an ionic oxide. Because the electrolyte is

an ionic solid, its electronic conductivity cannot be absolutely zero (see Chapter

3). In general, the SOFC electrolyte is selected such that the electronic transference number is as small as possible under normal operating conditions to

minimize electronic conduction losses. However, in some cases, depending on the partial pressure of oxygen at the electrodes, the electronic conduction in the

electrolyte cannot be neglected. Even if an electrolyte having some electronic

conduction is used, a higher-efficiency output may possibly be obtained if the

absolute value of the ionic conductivity of the mixed conducting electrolyte is

higher than that of a purely ionic conducting electrolyte [2.3]. When a SOFC is constructed with a mixed oxygen-ion and electronic

conductor, the electronic current flows through the electrolyte even at open- circuit conditions. In this case, the terminal voltage, E/, is somewhat lower than the theoretical voltage of the cell reaction, E. The transference number of the oxygen ion, ti, is given as

t~ - ' (Eq. 2.39)

where ty i is the oxygen-ion conductivity and a is the total conductivity (the sum

of the oxygen-ion conductivity ai and the electronic conductivity ae). Under this

condition, the equivalent circuit of the ceramic fuel cell in operation can be

represented as shown in Figure 2.3. The cell can be considered to be short

circuited by the resistor, the resistance of which is L/a e where L is the thickness of the electrolyte.

30 Chapter 2

ANODE I CELL

= =p- -

I

L m . . . . ,..,,,

E /

CATHODE I I 1 I

, - - 1 - - . I I I I

. _J

J J

OUTER CIRCUIT

Figure 2.3. Equivalent circuit of an operating SOFC [2.3]

Thus,

E / = je L-~- = E - j i L (Eq. 2.40) II e O i

where j~ and Ji represent the electronic and ionic current density in the electrolyte, respectively. At the open circuit of the cell,

j , = j~ (Eq. 2.41)

The following equation is established:

E o . g / = ' (Eq. 2.42)

O e+ff i

When the cell is discharged through the external circuit, the electronic

current in the electrolyte decreases as the terminal voltage of the cell falls. The

following equation can be derived for j~:

Je = (1 - ti) ( o E t i - J) (Eq. 2.43)

where j is the current density output, which is given by Ji -J~.

The total energy generated by the fuel cell reaction consists of the Joule

heat due to Ji (Pw(i)) and j~ (Pw(e)) and the power output Pw. These quantities can


be represented by the following"

Pw(i) = ( E - E/)ji (Eq. 2.44)

Pw(,) = E(i, (Eq. 2.45)

Pw = E(] = t~Ej j2 (Eq. 2.46) 12

From these equations, the total power, Pw~o), is

Pw(o) = Pw(o + Pw(~) + Pw = E(j + j~) (Eq. 2.47)

The energy efficiency, e, of the cell is the ratio of power output to total

power

P e - w (Eq. 2.48)

Pw(o)

Then, the relation between energy efficiency and current output is expressed by

the following equation:

oEti j _ j2 e = (Eq. 2.49)

oE[ti(1 - t i)oE + tiJ ]

Figure 2.4 shows the relation between e and j, for constants E and o, with ti as

a parameter. From this figure, it is evident that the efficiency of the cell with

1 , 0

I O.B

m z 0 . 6 u

!!1

0 .4

0 .2

0.0, 0.0 0.2 0.4 0.6 o.e --i .0

CURRENT OUTPUT, aE

Figure 2.4. Relation between energy efficiency and current output for a constant total conductivity [2.3]

32 Chapter 2

a mixed-conducting electrolyte is lower than that of the cell with a purely ionic

conductor in the whole range of current output (for a constant value of total conductivity). However, if the total conductivity of the mixed-conducting

electrolyte is greater than that of the purely ionic conductor, the efficiency of the

cell may be greater than that of a cell with a purely ionic conductor above a certain current output (Figure 2.5).

Under the conditions in Figure 2.4, the current providing maximum energy efficiency can be calculated from O~/aj = 0 and is given as

Je== = ~ - (1-ti)] (Eq. 2.50)

Substituting Eq. 2.50 for j in Eq. 2.49, the maximum energy efficiency, 6max, can be obtained as

1-~/(1-ti) e=~ x = (Eq. 2.51)

1 +~/(1-ti)

As can be seen from Eq. 2.51, the maximum energy efficiency is dependent only on the oxygen-ion transference number and is independent of total conductivity

and theoretical emf. This is an important criterion for the selection of solid oxygen-ion conductors for SOFC electrolytes. The relationship between 6ma x and

ti is shown in Figure 2.6.

1.0

0.8

I-//// U 0.6 E .~eo.

0.4 e "%. ,,z ~o

0.2 |l/ '='6, ~ %,

o or , L - % 0.0 0.2 0.6 0.8 1.0 i . 0.4

CURRENT OUTPUT, aoE

Figure 2.5. Relation between energy efficiency and current output for various conductivities [2.3]

1.0

0.8

0.6

0.4

0.2

I 1 I I

0.4 0.6 0.8 1.0 0 . 0

0.0 0.2

TRANSFERENCE NUMBER OF OXYGEN ION


Figure 2.6. Relation between maximum energy efficiency and oxygen-ion transference number [2.3]

The power output at maximum energy efficiency can be obtained as

Pw(emx,=oE2/ ( l_ t i ) [ l_ / ( l_ t i )~ (Eq. 2.52)

This equation is shown graphically in Figure 2.7, indicating that a maximum

value of Pw~, max) is obtained at ti = 8/9. The relation between the power output and the energy efficiency is an

important consideration in designing the construction of the ceramic fuel cell.

This relation can be derived by eliminating j from Eqs. 2.46 and 2.49 and is

given as

Pw = loE2eti{[(1-e)zt~-4r (1- r ( l - Q } (Eq. 2.53) 2

As ti --" 1, Pw approaches the following formula:

Pw = ~162 1 - e) (Eq. 2.54)

The relation between Pvr and E for various values of ti, with constants E and a,

is shown in Figure 2.8. In these curves, the arrows show the direction of the

increase in current output.

The power output at maximum energy efficiency does not coincide with

the maximum power output. Figure 2.8 shows that higher power output can be

34 Chapter 2

0.15

0 .10

0 .05

O.OO I ! 1 1 0.0 0.2 0.4 0.6 0.8 1.0

T R A N S F E R E N C E N U M B E R OF O X Y G E N ION

Figure 2.7. Relation between power output at maximum energy efficiency and oxygen-ion transference number [2.3]

0.25

0 .20

O.15

,,=, 0.10

0 .05

0 .00 i . 0.0

tu tO LL 14. ILl

X

I-

I--

I - 0 re. l,iJ

0

"o

4 I

0.2 0.4 0.6 0.8 1.0

E N E R G Y EFFICIENCY

Figure 2.8. Relation between power output and energy efficiency for various oxygen-ion transference numbers [2.3]

obtained for a slight loss of efficiency (since the slope of the Pw-versus-e curve is very steep in the vicinity of maximum efficiency).

As shown in Eq. 2.53, the power output of the cell is a complicated

function of e and ti, but proportional to tr. For the purpose of discussing the

relation among these four quantities, it is convenient to add a a-Pw relation to Figure 2.8 (Figure 2.9). Using Figure 2.9, the characteristics of various cells

Principles of Operation 3=3

"u . . . . . 0

O

" ~ \ IX / i \ ! I g 0.10 z o E7 C 0

! 1 ! I , _o.oo r 0.6 0.5 0.4 0.31 0.2 0.1 0.0 0.2 0.4 I 0.6, 0.8 1.0 o~j

POWER OUTPUT (TOTAL CONDUCTIVITY o'), o'k ..= ENERGY EFFICIENCY

Figure 2.9. Relation between energy efficiency and power output with respect to conductivity and oxygen-ion transference number [2.3]

can be compared with respect to their energy efficiencies and power outputs, if

the conductivities and oxygen-ion transference numbers are known. For example, consider two cells, I and II, the electrolytes of which have total conductivities of

a /and a~/(a/ < %) and oxygen-ion transference numbers of ti,~ and ti,~/(ti,~ > ti, z/), respectively. When cells I and II are operated in the vicinity of ~ = ~K, cell

II has a higher output power than cell I. In this case, the cell having the

electrolyte with the smaller oxygen-ion transference number shows the higher

output. In the vicinity of ~ = ~L (EL > EK), however, the power output of cell I

is superior to that of cell II. Table 2.2 lists the ratio of total conductivity of the

mixed conducting electrolyte to that of a purely ionic electrolyte giving the same

efficiency. For example, in order to get an efficiency of 0.60 using an electrolyte

with the oxygen-ion transference number of 0.94, the conductivity of the

electrolyte needs to be 1.5 times that of the purely ionic conductor.

36 Chapter 2

TABLE 2.2

Ratio of Total Conductivity of Mixed-Conducting Electrolyte to That of Purely Ionic Electrolyte for Various Values of t i and e [2.3]

Efficiency

0.35 0.40 0.45 0.50 0.55 0.60 0.65 0.70 0.75 0.80 0.85

1.000 1.0 1.0 1.0 1.0 1.0

0.995 1.1 1.1 1.1 1.1 1.1 0.990 1.1 1.1 1.1 1.1 1.1 0.980 1.1 1.1 1.1 1.1 1.1 0.960 1.1 1.1 1.1 1.1 1.2 0.940 1.1 1.2 1.2 1.2 1.3 0.920 1.2 1.2 1.2 1.3 1.5 0.900 1.2 1.2 1.3 1.4 0.850 1.3 1.4 0.800 1.6

1.0 1.0 1.1 1.1 1.1 1.1 1.1 1.2 1.2 1.5 1.5

1.0 1.0 1.0

1.1 1.1 1.1 1.1 1.2 1.4 1.3 1.5

1.0 1.5

It is clear from Eq. 2.54 that an oxygen-ion transference number of

greater than 0.89 is necessary to yield an energy efficiency greater than 0.5.

Therefore, for SOFCs, the oxygen-ion transference number in the electrolyte

must be greater than 0.9. However, it is clear that, even if the electrolyte shows

some electronic conduction, the cell can give high power with fairly high

efficiency when its oxygen-ion transference number is not smaller than 0.9,

provided that its conductivity is sufficiently high.

2.6 TYPES OF FUEL AND OXIDANT

In theory, any gases capable of being electrochemically oxidized and

reduced at the fuel cell operating temperature can be used as fuel and oxidant.

However, hydrogen is currently the most common fuel, and oxygen the most

common oxidant for use in SOFCs. Hydrogen can be derived from many

practical fuels such as natural gas, alcohol, or coal. Oxygen is readily available

from air.


2.6.1 Fuel

Present ceramic fuel cells use exclusively hydrogen as the fuel.

Hydrogen can be obtained from a number of hydrogen-containing fuels, e.g.,

natural gas, or from synthesis gases obtained by gasification of carbon sources,

e.g., coal. Among the practical fuels for the SOFC, hydrocarbons are probably

the most common. Hydrocarbon fuels are thermodynamically unstable at the

high operating temperatures of the SOFC. For example, pipeline natural gas

(which consists primarily of methane with small quantities of the higher

hydrocarbon gases such as ethane and propane, and trace amounts of a sulfur-

bearing odorant such as mercaptan, aliphatic sulfides, and thiophene) decomposes

mainly into carbon and hydrogen at fuel cell operating temperatures. An

appropriate amount of steam is often added to prevent carbon deposition and to

reform the hydrocarbons. The steam reforming of hydrocarbons can be carried

out in an external reformer (external reforming) or within the fuel cell (internal

reforming). Sulfur impurities, often present in common fuels, tend to cause cell

performance to degrade and to retard or poison the reforming activity of cell

materials even when steam is added to the fuel. Therefore, if undesirably high

levels of sulfur are present, the fuel must be desulfurized before fed to the fuel

cell. One common fuel for use in SOFCs is natural gas. The thermodynamic

feasibility of reforming natural gas in the SOFC depends on sufficient heat being

available at a suitable temperature for the endothermic reforming reaction. In

general, the electrochemical reactions produce enough heat for use in internal

reforming. For example, in the case of methane, the heat produced by the

electrochemical oxidation reaction at the SOFC operating temperature is about

twice that required by the reforming reaction.

Potential liquid fuels for SOFCs are naphtha, gas oil, or kerosene. In

this case, external reformers are often required. Also, sulfur must be removed

from the fuel before steam reforming. In general, the heavier the liquid fuels,

the higher the sulfur content, and the more difficult the fuel is to desulfurize.

Alcohol can also be used as the fuel; it can be reformed internally to supply

hydrogen for the SOFC.

Another potential source of fuel for the SOFC is coal. Coal may be

gasified externally through steam/carbon reaction to produce hydrogen and

carbon monoxide for SOFCs. Usually, coal is reacted with steam and oxidant

(air or oxygen) in a pressurized, fluidized-bed gasifier to produce hot synthesis

38 Chapter 2

gas. Coal can also react with hot CO2 and H20 (which are the products from ceramic fuel cells) to form hydrogen and CO. The heat produced by the fuel cell

can be used to sustain the gasification process [2.4,2.5]. Biogas and gases from biomass and landfill wastes are also potential fuel

gases for ceramic fuel cells. The composition of such gases is determined by its

respective origin and its generation process [2.6]. For example, gas mixtures produced by anaerobic digestion of manure are rich in methane and contain a

certain amount of sulfur and a small amount of halides. Landfill gases have the composition similar to that of a biogas; however, they contain ammonia, some

chlorides, and carbon-fluorine compounds. Biomass can be gasified either by high-temperature pyrolysis [2.7,2.8] or by partial oxidation [2.9]. Sulfur removal is usually necessary for biogases [2.10] but direct utilization of biogases in the fuel cell is possible when contaminants do not exceed a certain limit

[2.11].

2.6.2 Oxidant

The performance Of ceramic fuel cells is improved when pure oxygen is used as the oxidant; however, air is usually used because of its availability. The difference between cell voltage obtained with pure oxygen and that with air increases as the current density increases, suggesting that diffusion polarization plays an important role during the reduction of oxygen in air. Because of the

high operating temperature, a SOFC system needs an air supply blower, air preheater, and air recirculator.

2.7 FUEL-PROCESSING SYSTEM

The function of a fuel-processing system is to convert practical fuels to

a gas mixture suitable for use in the fuel cell. For natural gas and liquid fuels, the fuel processor consists mainly of sulfur removal and reforming systems. For

coal, the fuel-processing system consists of gasification, cleanup, and desulfuri-

zation subsystems.

The fuel-processing system for natural gas is often designed to remove sulfur compounds to below 0.1 ppm, using active carbon, zinc oxide, or

hydrogen (hydrodesulfurization). The desulfurization method selected depends on the sulfur compounds involved. For example, when the fuel contains

thiophane, a hydrodesulfurization reactor is usually installed in the fuel-


processing system in series before the zinc oxide reactor. After desulfurization, natural gas is reformed, either externally or internally. For internal reforming, natural gas is fed to a prereformer before being reformed within the boundaries

of the SOFC (reforming partially in prereformer tubes and partially in the fuel cell), or is reformed within the cell without a prereformer (100% reforming in

the fuel cell). In both cases, a portion of the spent f,,el stream is recirculated, to guarantee a supply of water vapor adequate to support the reforming reaction.

For heavier liquid fuels, because of difficulties in sulfur removal and cost

concerns, reformers need high sulfur tolerance. This requires higher operating

temperatures for the reformer and/or the use of reforming catalysts that have less overall reforming activity but are also less sensitive to sulfur poisoning. Heavier liquid fuels also contain higher proportions of aromatic hydrocarbons. This leads to another severe technical p r o b l e m - carbon formation in the reformer. These

two factors, sulfur tolerance and carbon formation, have been the key challenges in the development of fuel processors for heavy hydrocarbon liquid fuels. For

the conversion of heavier liquid fuels, high-temperature steam reforming, auto-

thermal reforming, hybrid reforming, and cyclic reforming have been considered [2.12].

For coal, a pressurized, fluidized-bed gasifier (aair- or oxygen-blown) is used to produce coal gases. Carbon conversion reaches 98.5 %, and examples

of gas compositions produced (in vol %) [2.13] are

Air mode

Oxygen mode

H2 CO C H 4 H20 CO2 N2 H2S COS N H 3

16.8 27.9 1.8 3.8 2.8 46.0 0.68 0.02 0.16

29.4 38.0 4.0 17.3 9.5 0.54 1.07 0.02 0.24

The gas from the gasifier is cooled in a steam generator after removal of

entrained particulates. The cooled gas is treated in order to remove sulfur compounds by absorption on beds of zinc ferrite or zinc oxide. The clean gas

is then passed through an expander before it flows into the fuel cell.

2.8 POWER-CONDITIONING SYSTEM

A fuel cell produces DC currem at a certain voltage. If the DC current

is used directly (e.g., supply power to an electrolyzer), a chopper is required for

controlling the voltage output. For other applications, a power-conditioning

system is required in order to convert the DC current to an AC current consistent

with, for example, the electric utility interface standards. A power-conditioning

40 Chapter 2

system may have a self-commutated inverter or an externally commutated

inverter for commutation of DC to AC, and a self-controlled inverter or an

externally controlled inverter for frequency control. For a fuel cell power-

conditioning system, the self-commutated inverter is commonly used. When

connecting to a commercial AC line system, the self-controlled inverter is used,

and in the case of a simple AC system, the externally controlled inverter is often

used.

References

2.1

2.2

2.3 2.4

2.5

2.6

2.7 2.8

2.9 2.10

2.11 2.12

2.13

E.H. Camara and L.G. Marianowski, Handbook of Fuel Cell Performance, Institute of Gas Technology, Chicago, IL, 1980, p. 83. J.L. Bates, in Proceedings of the 16th Energy Technology Conference, February 28- March 2, 1989, Washington, DC, Government Institutes, Rockville, MD, 1989, p. 205. T. Takahashi, K. Ito, and H. Iwahara, Electrochim. Acta, 12 (1967) 21. R.L. Zahradnik, L. Elikan, and D.H. Archer, in Advances in Chemistry Series 47, Fuel Cell Systems, American Chemical Society, Washington, DC, 1965, p. 343. Westinghouse Electric, 1970 Final Report Project Fuel Cell, Research and Development Report No. 57, U.S. Government Printing Office, Washington, DC, 1970. H. Wendt, V. Plzak, and B. Rohland, in Proceedings of the Second International Symposium on Solid Oxide Fuel Cells, July 2-5, 1991, Athens, Greece, F. Grosz, P. Zegers, S.C. Singhal, and O. Yamamoto (eds.), Commission of the European Communities, Luxembourg, 1991, p. 343. W. Kaminsky, Chem.-Ing. Tech., 61 (1989) 775. D.S. Scott, J. Piskorz, M.A. Bergougnou, R. Graham, and R.P. Overend, Ind. Eng. Chem. Res., 27 (1988) 8. W. Adlhoch and N. Briingel, Braunkohle, 41 (1989) 42. E. Pilarczyk and K. Knoblauch, in Proceedings of the Engineering Foundation Conference on Separation Technology, April 27-May 1, 1987, Elmau, Germany, Engineering Foundation, New York, 1987, p. 522. H.H. M6bius and B. Rohland, U.S. Patent No. 3402078, 1964. J.E. Young, in Proceedings of the Symposium on Fuel Cells: Technology Status and Applications, November 16-18, 1981, Chicago, IL, Institute of Gas Technology, Chicago, IL, 1982, p. 153. M. Krumpelt, V. Minkov, J.P. Ackerman, and R.D. Pierce, Fuel Cell Power Plant Designs: A Review, Report No. ANL-85-39, Argonne National Laboratory, Argonne, IL, 1985, p. 31.

Chapter 3

ELECTRICAL CONDUCTION IN CERAMICS

3.1 GENERAL

The operation of a SOFC is based fundamentally on electrical conduction

in the ceramic components. The electrolyte conducts ions between the anode and cathode. The electrodes carry electrons to and from the reactive sites, where the

electrochemical reactions take place. The interconnect conducts electrons from the anode of one cell to the cathode of the next in electrical series. Thus,

electrical conduction processes in ceramics in general and SOFC materials in

particular are of paramount importance to the operation of the SOFC. A

discussion of electrical conduction in ceramics is given here to provide a background on the basic electrical processes in fuel cell operation. Emphasis is

placed on the discussion of electrical conduction in fluorite-type oxides and perovskite oxides, the two common types of oxide used in SOFCs.

In general, electrical conduction in ceramics or crystalline solids depends on the material lattice defects. Lattice defects can be classified into two groups:

(i) stoichiometric defects, in which the crystal composition is unchanged when defects are formed in the lattice and (ii) nonstoichiometric defects, which are formed as a consequence of a change in the crystal composition. For example, AgC1 crystals may contain Ag + ion interstitials (silver ions in interstitial sites)

and Ag + ion vacancies (missing silver ions from their normal sites), thus having stoichiometric defects. Y203-stabilized ZrO2, strontium-doped LaMnO3, and doped LaCrO3 (commonly used as SOFC electrolyte, cathode, and interconnect, respectively) are doped crystalline oxides, thus having nonstoichiometric defects.

Alternatively, defects can be classified into three groups by the size and

shape of the defect: (i) point defects which involve interstitials or vacancies,

(ii) line defects (dislocations) which are effectively point defects in two

dimensions, and (iii) plane defects, in which the whole layer in a crystal structure

is defective. Although the concept of isolated point defects has become less

attractive in recent years (being replaced by microdomain and cluster models),

42 Chapter 3

it is convenient to discuss (and easy to understand) electrical conduction in

crystalline oxide ceramics in terms of the three classical types of point defect.

These three classical types are the Frenkel defect and the Schottky defect

(intrinsic point defects), and the Koch-Wagner defect (extrinsic point defect).

The Frenkel-type defect was introduced by Frenkel in 1926 [3.1]. According to this model, some of the ions in an ionic crystal move from their

normal lattice points to interstitial positions due to thermal fluctuation, forming

interstitial ions (or interstitials). Interstitials vibrate at their points and move to

other interstitial sites by thermal fluctuation, chemical diffusion, or an applied

electric field. The vacant lattice points or vacancies left by the interstitials can

be occupied by other lattice point ions. As a result, vacancies also move in the

crystal. These interstitials and vacancies are called the Frenkel defects. The

number of interstitials in a unit volume (which is equal to that of vacancies), NF, is given as

N e = NiNexp(- Ee 2KT ) (Eq. 3.1)

where N and N/are the total number of ions and interstitial positions in a unit

volume, respectively, K the Boltzmann constant, T the temperature, and EF the formation energy of the Frenkel defect.

The Schottky-type defect was introduced by Schottky in 1935 to explain

simultaneous cationic and anionic conduction in ionic crystals [3.2]. According

to this model, the same number of cationic and anionic vacancies exists in ionic

crystals, and the appearance and disappearance of vacancies are considered to

occur mainly at the crystal surface or dislocation face. The number of the

Schottky ionic pair in a unit volume, Ns, is given as

N s = Nexp(--2KT) Es (Eq. 3.2)

where N is the total number of ionic pairs in a unit volume and Es the formation

energy of the Schottky ionic pair. (In these equations, the thermal expansion of

the crystal and the change of the frequency of lattice vibration are not considered.

The number of defects increases somewhat if these two factors are taken into

account; however, this increase in the number of defects is usually small.)

The existence of the Frenkel or Schottky defect in a crystal can be

estimated by comparing the ion radius, van der Waals energy, and dielectric

constant (i.e., comparing EF and Es). The Frenkel defect usually appears when

Electrical Conduction in Ceramics 43

the radii of the ions of the crystal differ considerably, and the van der Waals

energy and the dielectric constant are somewhat large. The Schottky defect often occurs when the differences between the radii of the cation and anion and their

abilities to be polarized are small, and the van der Waals energy and the

dielectric constant are both relatively small [3.2, 3.3].

The Koch-Wagner type defect was introduced by Koch and Wagner in

1937 [3.4]. This type of defect occurs when impurities or dopants are introduced

into the lattice of the crystal. For example, when CaO is dissolved in ZrO2

crystal to form a solid solution, the substitution of the divalent Ca 2+ ion for the

tetravalent Zr 4+ forms oxygen-ion vacancy defects in the ZrO2 lattice (in order

to maintain the electroneutrality condition in the crystal). Similarly, oxygen-ion

vacancies are also formed when a trivalent metal oxide, such as Y203 or Yb203, is dissolved in ZrO2.

3.2 DEFECTS IN FLUORIDE-TYPE OXIDES

Fluorite-type oxides are commonly ionic conductors and have been

considered for use as electrolytes in SOFCs. The fluorite structure (CaF2) is

adopted by a number of oxides of the general formula MO2, where M is a large

tetravalent cation, e.g. Th 4+ and Ce 4+, as well as by a multitude of halides,

sulfides, hydrides, and intermetallic compounds of composition AX2. The unit

cell of the fluorite-type oxide has the so called M408 structure. This structure

is schematically shown in Figure 3.1. In this fluorite structure, each metal ion

A A

"/// �9 U �9 CATION 0 OXYGEN ION

Figure 3.1. Crystal structure of a fluorite-type oxide

44 Chapter 3

is surrounded by eight oxygen ions, forming a body-centered cubic structure, and

each oxygen ion is surrounded by four metal ions, forming a tetrahedral

arrangement. To form the fluorite structure in MO2, the limiting (minimum)

ionic radius ratio (the ratio of metal-ion radius to oxygen-ion radius) is 0.732. Under normal conditions of temperature and pressure, certain MO2

oxides do not have the fluorite structure because the ionic radius ratio condition

is not satisfied; one example is ZrO2. At room temperature, ZrO2 has a monoclinic crystal structure. The monoclinic structure changes to a tetragonal form above 1170~ The fluorite structure only exists at temperatures above

2370~ However, the addition of certain aliovalent oxides stabilizes the fluorite

structure of ZrO2 from room temperature to its melting point of 2680~ The

fluorite structure of ZrO2 is stabilized by direct substitution of divalent or trivalent cations of appropriate size for the host lattice cation Zr 4§ In this case,

lattice defects are created to preserve the electroneutrality condition in the solid solution. The probable models for structural defects in such case are: (i) an oxygen-ion vacancy model with all metal ions being fixed at their lattice points, (ii) a cation interstitial model with all oxygen ions being fixed at their lattice sites

(Frenkel type), and (iii) a mixed model of (i) and (ii) (Schottky type). It is well

established that the oxygen-ion vacancy model applies to stabilized ZrO2. For

example, calculated and pycnometer density measurements confirm the presence

of oxygen-ion vacancy defects in ZrO2-CaO systems [3.5-3.7]. The oxygen-ion vacancy model has been verified by X-ray [3.8, 3.9], neutron diffraction [3.10- 3.12], and measurements of the diffusion coefficient of oxygen ions in ZrO2-CaO solid solutions [3.13]. Similar results have been obtained for other stabilized zirconias, cerias, and other fluorite-type oxygen-ion conductors. The presence of a high oxygen-vacancy concentration in stabilized ZrO2 gives rise to a high

oxygen-ion mobility, resulting in high oxygen-ion conductivity. Oxygen-ion

conduction takes place in stabilized ZrO2 by movement of oxygen ions via

vacancies.

3.2.1 Defect structure of doped MO~

As mentioned above, doped M O 2 oxides such as stabilized Z r O 2 exhibit

oxygen-ion conduction via oxygen-ion vacancies. However, the electrical

conductivity in an oxygen-ion conductor can appear not only via oxygen-ion

vacancies but also via mobile electronic charge carriers; i.e., both types of

conduction (ionic and electronic) can occur simultaneously. The ionic transfer-

ence number of an oxygen-ion conductor (the ratio of the ionic conductivity to


the total conductivity) can vary depending on the surrounding gaseous environment (the oxygen partial pressure). Thus, it is possible to modify the defect structure (thus, conductivity) of an oxygen-ion conductor by changing the oxygen

partial pressure. The dependence of defect structure on oxygen partial pressure is considered here for a general system, MO2-AO, assuming that the intrinsic defects in pure MO2 are of the Schottky type. Kr6ger-Vink notation is used to

describe the various defects (V for vacancy, bracket for concentration, e / and h"

for effective charges of electron and hole, and n and p for electron and hole concentrations, respectively). The following discussion largely follows those

which have appeared in the literature [3 .14 -3 .18] .

The incorporation of AO into MO2 is described by the following equation:

// x AO = A M + Vo + O O (Eq. 3.3)

This equation describes a situation where the M site is occupied by A, generating a doubly charged negative site, and it is necessary to generate two anion sites to maintain the crystal structure, one site being occupied by an oxygen ion and the other vacant, leaving a doubly positive charge.

The incorporation of oxygen from the environment into the solid MO2 is described as

ZO . . . X �9

2 2 + Vo + 2e/ Oo (Eq. 3 4)

and the equilibrium constant K for the reaction is given by

-2 K - n p-1/2 o~ (Eq. 3.5)

[Vo]

The intrinsic Schottky equilibrium of the oxide MO2 is given by the following reaction"

i l l ~ l ~ .. n i l - , M + 2 V o (Eq. 3.6)

The equilibrium constant K s for Eq. 3.6 is

K s - [V~//I[Vo]2 (Eq. 3.7)

46 Chapter 3

The equations describing the intrinsic electronic equilibrium and its equilibrium constant Ki are as follows:

nil = h ' + e / (Eq. 3.8)

K~ = np (Eq. 3.9)

From these equations, the electroneutrality condition is given as

n + 4[I/~ a] + 2[A~] = p + 2[Vo] (Eq. 3.10)

(i) Low oxygen partial pressure region: In the low oxygen partial pressure region, as the oxygen pressure decreases, the concentration of oxygen-

ion vacancy will increase (to maintain the equilibrium constant K). This increase

causes the metal vacancy concentration to decrease (to maintain the Schottky constant Ks). In this situation,

[V o] , [I/~//] (Eq. 3.11)

Eventually, the concentration of oxygen-ion vacancy exceeds that of Ar~, which concentration is fixed by the dopant level. Thus, n must increase to maintain the electroneutrality condition, and accordingly, p must decrease. In this case, the electroneutrality condition given in Eq. 3.10 is reduced to

n = 2[V o] (Eq. 3.12)

From Eqs. 3.5 and 3.12, the oxygen partial pressure dependence of oxygen-ion vacancy is obtained as

[l/'O] = (4f)-IDp -I/6 02 (Eq. 3.13)

The same dependence on oxygen pressure can be derived for the electron concentration. From Eqs. 3.7 and 3.13, the oxygen partial pressure dependence

of the metal vacancy concentration is expressed as

Ks(4K)2taP~C (Eq. 3.14)

From Eqs. 3.9, 3.12, and 3.13, the dependence of the hole concentration on oxygen partial pressure is given by

p = Ki(K)IDP 1/6 (Eq. 3.15) 2 "2


(ii) Intermediate oxygen partialpressure region: Under some intermediate oxygen pressure range, it is possible to approximate the electroneutrality

condition as

[Vo] = [A~] (Eq. 3.16)

This relationship indicates that over this intermediate oxygen pressure range, the

concentration of oxygen-ion vacancies is not dependent on oxygen partial

pressure, but fixed by the dopant level. By the same procedures described earlier, the oxygen partial pressure

dependencies of the defects can be calculated and are given as follows:

[I/Mm] - KS (Eq. 3.17) A//12 ~XMl

// -1/2p-114 n = (K[AM]) 02 (Eq. 3.18)

// p = (Eq. 3.19)

These equations indicate that as the oxygen partial pressure increases, the

concentration of the electrons decreases and that of the holes increases. On the other hand, the metal-ion vacancy concentration is independent of oxygen partial pressure and determined solely by K s and the oxygen-ion vacancy concentration.

(iii) High oxygen partial pressure region: At some sufficiently high

oxygen pressures, the electroneutrality condition can be approximated as

p = 2[A~] (Eq. 3.20)

In this oxygen partial pressure range, the oxygen pressure dependencies of the

anion and cation vacancies are given as

[Vo] - ~4[AM]2po~I2 (Eq. 3.2 I) KK .2

l

viaJ = rail14 16t- -Mj ~ P o 2 (Eq. 3.22)

48 Chapter 3

and the electron concentration is given by

/q n - (Eq. 3.23)

2[A~I

As can be seen from Eq. 3.23, the electron concentration is constant in this

oxygen pressure region.

(iv) Very high oxygen partial pressure region: At some very high oxygen pressures (although this condition will not occur in operating ceramic fuel cells),

the concentration of metal-ion vacancy becomes very large, i.e.,

[I,'a~//] > [ A ~ ] (Eq. 3.24)

The electroneutrality condition can be expressed approximately as

p = 4[I,;M///] (Eq. 3.25)

For this very high oxygen pressure region, various defect concentrations are

described as follows"

(4Ks) 2 1/Sp-l/lO [Vo] = [ KK~] o~ (Eq. 3.26)

K s ( )215Po~ ( E q . 3.27) 16

(Ale" "~ 115(]['F2"~2/5 D 1/5 P = v ~ - - s J ~ . . . . i J - o 2 (Eq. 3.28)

/ : -2 /5 t Ki ,~1/5p-115 /1 = .,~ ~,-"7-~__-J �9 O

4 K s 2 (Eq. 3.29)

The various defect concentrations in a MO2-AO system for various

oxygen partial pressure regions are shown schematically in Figure 3.2 [3.16]. The defect structure of a fluorite oxide MO: doped with B 2 0 3 c a n be

considered by the same procedures described above. The overall electroneu-

trality condition in this case can be expressed as shown below by assuming low

defect concentrations so that Henry's law is obeyed

n + 2[O~/] + [B~] = p + 2[Vo] (Eq. 3.30)


I i S . . . . . . . . . . .

f p-TYPE n-TYPE EMICONDUCTION IONIC CONDUCTION SEMICONDUCTION '

Z IVERY HIGl~l HIGH I INTERMEDIATE I LOW o_t--e w . -p ~i ~ ,, ----v-- p I..- I o , I o , . t o , I o ,

/ i '. i

Z / I I I i l l

I I / / -

, [ ~ " ] (9 I I I 0 I I - J I

I I I I I

. . I . .___LOG O X Y G E N P A R T I A L PRESSURE

Figure 3.2. Variation of defect concentration as a function of oxygen partial pressure for a MO2-AO system [3.16]

Figure 3.3 shows a schematic illustration of the variation of defect concentration

as a function of oxygen pressure for a fluorite-type oxide doped with trivalent

metal oxide [3.17]. When the MO2 originally has the fluorite-type structure

(e.g., CeO2), it has been generally accepted that the predominant intrinsic defects

in the oxide are anion Frenkel defects i e Vo and 0/ / , �9 �9 i �9

�9 . I / . . / " u .I " = 2 [ v ~ i , [B.] = 2[v~l p = tB.i p = 2[0, ] l , 9 I I I -'@

"<~o'"~ I , I s ,,~ ! ~ _ ~ - . . _ i [ ~ ] I I J ~

. . . . . . . . . . . . . .

l.U

0 1.1

<~ r " 1 ;- . . . . " 0 -J P ' ~ / I I o ~ . ~ , ~ . ~

. . . . .

LOG O X Y G E N P A R T I A L PRESSURE

Figure 3.3. Variation of defect concentration as a function of oxygen partial pressure for a MO2-B203 system [3.17]

50 Chapter 3

3.2 .2 Conductivities of oxygen ions, electrons, and electron holes

The formation of oxygen-ion vacancies in doped MO2 oxides gives rise to oxygen-ion conduction in these solid solutions. The conduction occurs by

diffusion of oxygen ions via vacancies. In general, the conduction process in an

oxide ionic conductor is mixed ionic and electronic. The total electrical

conductivity, a, of a fluorite-type oxide is given as

o ~

o = o~ + o n + or, = 2e[Vo]lXr6 + e n i d , + e p l x p (Eq. 3.31)

where /z refers to the mobilities of the subscript species in which i, n, and p indicate ions, electrons, and electron holes, respectively. The ionic conductivity

due to the migration of cations of the dopants and the host is neglected because

the mobilities of the cations have been shown by diffusivity measurements to be

several orders of magnitude lower than the mobility of oxygen-ion vacancy. The following discussion is given for a MO2-B203 solid solution; similar

treatment can be used for MO2-AO systems. At low oxygen partial pressures, the hole concentration is relatively small, the total electrical conductivity becomes

0 = 2 e [ V o l t X v o + e n , , (Eq. 3.32)

However, since the mobility of electrons is generally 104 to 108 times higher than

that of oxygen-ion vacancies even at elevated temperatures, the fluorite oxide is

approximated to be a pure n-type semiconductor. Therefore, the conductivity for a doped MO2 under this condition is given by the following equation:

= p-116 o = o n = en tx n e(2K)-l/31x n 02 (Eq. 3.33)

At intermediate oxygen partial pressures, the total electrical conductivity

can be approximated as

_ p-114 eKi(2K[B~])lr21.tpp1/42 o = 2e[Vol . vo + e(2K~l)1/21.t, , 02 + (Eq. 3.34)

where the electroneutrality condition is

2tVol- 03 ,1 (Eq. 3.35)


At high oxygen partial pressures, the fluorite oxide is approximated to be a pure p-type semiconductor, and the conductivity is given as

o = e [~] i . tp (Eq. 3.36)

Since the mobility of oxygen-ion vacancy is generally much lower than

that of electrons or electron holes, MO2 doped with a trivalent metal oxide can

only exhibit an appreciable ionic conductivity over a wide range of oxygen partial

pressures when the concentration of oxygen-ion vacancy is considerably larger

than n and p. Figure 3.4 shows a schematic diagram of the variation of electrical

conductivity as a function of oxygen pressure for a fluorite-type oxide doped with

trivalent metal oxide. For this system, the ionic conductivity is at the maximum

when o, = %. The electronic equilibrium constant Ki (Eq. 3.9) can be expressed as

E - - -~T ) (Eq. 3.37) K i np NcNvexp(-

where Eg is the band gap energy between the valence band and conduction band,

and Nc and Nv the effective density of the electronic state in the conduction band and in the valence band, respectively. From Eq. 3.37, the minimum electronic conductivity, O'e(min.), is given as

E N )1r2 e o~(,,,~..) = 2e(p,,pp cNv exp(--2KT) (Eq. 3.38)

.,..... n-TYPE - ' l I p---- p-TYPE L SEMICONDUCTION I I I SEMICONDUCTION J

I .... ~ I I I ~, I o: I _ , ~ " - , , , ~ i o I I ~ ,-~ _ ~ o , , I ~,(~ _ _ p , ~ . ,o i"o . , . I

_, ', ,

I , MIXED' ~ ~,,"1. MIXED I I COND. I I

I i I t - . I i I I

I I .-I IONIC "1. . I I I I ~'---| D O M A I N I " - ' ~ I I I I I I I I I

I I I~LECTROLYTiC I I I I ' J ' I I I P . / D O M A I N . P , I I I ~ I I I I , I

L O G O X Y G E N P A R T I A L P R E S S U R E

Figure 3.4. Variation of electrical conductivity as a function of oxygen partial pressure for a MO2-B203 system [3.17]

52 Chapter 3

3.2.3 Defect domains

The ionic domain boundaries Pp and P~ of a fluorite-type oxide are the

oxygen partial pressures at which Po2 = Pp when O" i : t ip, and Po2 = P~ when

o i " - O n. Within this oxygen pressure range and away from the boundaries, the

oxide is considered to be an ionic conductor. (The electrolytic domain boundaries are the oxygen partial pressures where o- i "-- 1 0 0 0rp and a i = 100 cry.

Within this oxygen partial pressure range, the conductivity of the oxide is exclusively ionic.) The ionic domain boundaries of a fluorite oxide can be

estimated by the following procedure.

The mobilities of each conducting species (ions, electrons, and holes) can

be expressed by the following equations:

An. A exp(- ) (Eq. 3 39) r

B E - ~ exp(- ~" ~" T KT ) (Eq. 3.40)

E.._ exp( - ~e) (Eq. 3.41) It, - T ~:T

where A~, B,, and C~, are constants that include entropy terms, AH m the migration

enthalpy of oxygen-ion vacancies, and E~,,n and E~,,p the migration energy of

electrons and electron holes, respectively. Then, the conductivities are given as

aHm ~ = ~176 KT ) (Eq. 3.42)

o ~ En~p-114 0 n = 0 n T- exp(--~---Tj o~ (Eq. 3.43)

o ~ Ep pl/4 op = opT- exp(---~) 02 (Eq. 3.44)

where each cr ~ represents conventional constants, which are assumed to be

independent of Po2 and T, and

AHlro En - 2 + E~m (Eq. 3.45)

Ep = Eg + Er,~, + AHv-o

(Eq. 3.46)


where z~v. ~ is the oxygen vacancy formation enthalpy. At the oxygen partial

pressure Po where a, = ap, the expression for Po can be obtained by solving

Eqs. 3.43 and 3.44

oO 2(E,-Ep lnP o = 21n n (Eq. 3.47)

o o Ir p

This equation defines the boundary separating predominantly n-type from p-type

conductivity. The expressions for P, and Pp are obtained by setting a, = o i and

(Tp = O" i

0

lnP n = 41n(-~) - 4 O i

(aHVo/2) +e..,,-AH,,, rT

(Eq. 3.48)

(ant'J2) + aH' -E*-e" ' " lnPp = 41n( ) - 4 (Eq. 3.49)

o KT p

By combining Eqs. 3.34, 3.37, 3.39 to 3.46 with Eqs. 3.48 and 3.49, the ionic

domain width is given as (written for MO2-B203)

/ 2 2 In Pp = 4( [BM] A , Ee,+E..,,+E.~,-2AH,.

P,, N, NvB C ) + 4 ~r (Eq. 3.50)

It can be seen from Eq. 3.50 that an increase in [B~I and a large Eg, E~,, or E,,p will increase the ionic domain, and an increase in Z ~ m and T will decrease the ionic domain. This suggests that oxygen-ion conducting ceramics exhibiting low

values of Z~/m can be potential electrolyte candidates for SOFCs. It should be

noted that oxygen-ion conductors having relatively large '~ /m are less sensitive

to decrease of the ionic domain with increasing temperature. In addition to the diagram shown in Figure 3.4, there is also a Patterson-

type map suitable for illustrating the ionic and electrolytic domain boundaries. Figure 3.5 is an example of the Patterson map showing the domain boundaries

of stabilized ZrO2 [3.19-3.21].

54 Chapter 3

t UJ re'

t~ I.M t r ix. ...i <

tw < ix. z UA

>- X O (9 O _J

up D O M A I N / ___ " / IONIC DOMAIN 0, r = o" S O ' /

\ -t, = o.aa t , = o.5 \

1 / T E M P E R A T U R E --,,

Figure 3.5. Electronic, ionic, and electrolytic domains of a M O 2 - B 2 0 3 system [3.17]

3.2.4 Defect associations and clusters

The discussion regarding doped MO2 systems given above considers a random distribution of noninteracting defects. However, at low temperatures and even at low dopant concentrations, oppositely charged oxygen-ion vacancies and dopant cations may associate to form randomly distributed pairs, and the concentration of free oxygen-ion vacancy is determined by the association equilibrium �9

�9 " I I x V o + A~ = (VoAM) (Eq. 3.51)

o r

/

Vo + = (Eq. 3.52)

At high temperatures, the oxygen-ion vacancies are most likely to be free at

certain defect concentration ranges. Figure 3.6 shows the calculated temperature

for the break between associated and free vacancy behavior for CeO2 doped with CaO, as a function of dopant concentration for various association energies, EA [3.22]. However, the formation of defect pairs also depends on the magnitude of the association energy. For example, the conductivity data at high tempera-

tures for calcium-doped ThO2 in the range of dopant concentrations similar to that of CeO2 appears to indicate the association of vacancies and does not show


Figure 3.6.

3 4 0 0 , , , , , , , - , , , , - , ooo{ , v / _ / / EA-0.5

, / / y y 2 0 0 0 . . /

n..< 0 . 2 e V

~- 1 5 0 0

,,~ 1 0 0 0

I l l 5OO

0 I I I

0 2 4 6 8 10

M O L % C a O

Calculated break temperature between associated and free vacancies in CaO- doped Ce02 for various association energies [3.22]

break temperatures as CeO 2 does [3.23]. This can be explained by the difference

in the association energy between the two compounds: the association energies

(1.16 to 1.42 eV or 111.9 to 137.0 kJ/mol) of ThO2 solid solutions with calcium

dopant concentrations of 1 to 7 mol % are much higher than those (0.20 to 0.50

eV or 19.3 to 48.2 kJ/mol) of CeO 2 solid solutions.

At higher defect concentrations and even at high temperatures, a random

distribution of defects and defect pairs may disappear. Defects may aggregate

into defect clusters that become ordered or form two- or three-dimensional

defects (dislocation loops, shear planes, or voids, as well as changes in the local

crystal structure). As the binding energy of defect clusters has been known to

increase with the number of dopant cations, the vacancies are expected to be

more deeply trapped with increasing dopant concentration. Even with a random

distribution of the dopant cations, local fluctuations can result in the formation

of defect clusters. In these cases, the system may be described as a relatively

homogeneous matrix with defect clusters or as a vacancy-free, ordered matrix

with microdomains of another ordered structure. For example, in the case of

stabilized zirconias, the ordered matrix is ZrO 2 and the microdomains could be

sluggish long-range ordering forms of CaZr409 in the ZrO2-CaO system [3.24],

and Zr3Y4012 in the ZrOz-Y203 system [3.25-3.27].

The experimentally observed decrease of oxygen-ion conductivity with

increasing vacancy concentration following doping with large amounts of dopants

56 Chapter 3

is probably due to trapping of vacancies in aggregates or clusters. Trapped vacancies found in the clusters, however, are not immobile, but must overcome an energy barrier to move (by dissociation or rearrangement of clusters). This

energy barrier is higher than that present in systems having only single vacancies.

3.3 DEFECTS IN PEROVSKITE-TYPE OXIDES

Oxides of the general formula ABO3 (A is a divalent or trivalent cation

and B is a tetravalent or trivalent cation) belong to the class of compounds having

the perovskite structure. The ideal perovskite structure is cubic. A large number of perovskite oxides are orthorhombic, rhombohedral, or tetragonal but can be

approximated as cubic. The rare-earth perovskite oxides are commonly used as cathode and interconnect materials for ceramic fuel cells. For fuel cell applications, the electrical conductivity of the oxides is often enhanced by substituting acceptor- or donor-type cations for either the A or B sites. At

present, the perovskite oxides commonly used in SOFCs are LaMnO3 for the

cathode and LaCrO3 for the interconnect. The electrical conductivity of LaCrO3 (or LaMnO3) is considered to be due to the 3d electrons of the Cr 3+ (or Mn 3§

ions. An increase in conductivity is expected when a lower-valence ion is substituted for either the La 3§ or Cr 3+ (or Mn 3+) sites, resulting in the formation of Cr 4+ (or Mn 4+) ions. A proposed defect structure of these perovskite-type

oxides is discussed here as applied to a p-type perovskite REBO3 (RE is a trivalent rare-earth ion) with acceptor dopant A 2§ [3.28]. This type of defect behavior has been found to apply to magnesium-doped LaCrO3 and strontium-

doped LaMnO3 [3.29,3.30] (see Chapters 5 and 7). For simplicity, it is assumed that the p-type disorder prevails in

nonstoichiometric perovskite oxide. The intrinsic Schottky equilibrium of REBO 3

is represented by

nil = I / ~ + I/a a + 3 V o (Eq. 3.53)

and the equilibrium constant Ks is given by

K s = [l/R~[l/aa][Vo] 3 (Eq. 3.54)

The p-type nonstoichiometric reaction is given as

-3o - v 'g § § 3 0 0 § 6h. (Eq. 3 .55) 2 2


and as the cation stoichiometry must remain constant, the equilibrium constant

for Eq 3.55 is

(Eq. 3.56)

When REBO 3 is doped with AO, the substitution of A 2+ for the RE 3+ site

results in the formation of ARE, which possesses one effective negative charge. This charge will be compensated for either by C r 3+ ---> Cr 4+ or by the formation

of oxygen-ion vacancies. Then, the electroneutrality condition is expressed as

2[Vo] + p = 6[I/d/] + [A~j (Eq. 3.57)

The defect behavior of REBO3 doped with AO can be divided into five

oxygen partial pressure regions. (i) Region I: In this very high oxygen partial pressure region, native

nonstoichiometry prevails, and the electroneutrality becomes

p = 6[VCB "] (Eq. 3.58)

From Eqs. 3.56 and 3.58, the hole concentration is given as

p = (36K)l/SPo/~n (Eq. 3.59)

Since the equilibrium constant of intrinsic electronic defects, K/, equals np, the electron concentration is

n = Ki p-3/,6 (Eq. 3 60) 02 ~ (36K) l/s

and the oxygen-ion vacancy concentration is

K s /3 r, -1/s (Eq. 3 61)

[Vo] = 2.5 K 1/12"O2

/

[ARE] (ii) Region H: In this region, as both [Vo] and [1/~/] are smaller than

, the electroneutrality condition becomes

p = [A~] (Eq. 3.62)

58 Chapter 3

Thus,

[~B//] : g 1]2 p3/4 02 (Eq. 3.63)

[ A ~ 3

From Eqs. 3.54, 3.56, and 3.62, the concentration of oxygen-ion vacancy is given as follows:

t Ksxl tar A / 12p-112 [Vo] = , - - - ~ , t - ~ R E J - 0 2 (Eq. 3.64)

and the electron concentration becomes

n = K. (Eq. 3.65) [A~E]

(iii) Region IlL" As the oxygen partial pressure decreases, oxygen is lost to form oxygen-ion vacancies, and the electroneutrality condition is represented

by

p + 2[Vo] = [ A ~ (Eq. 3.66)

The equations for the various defects are as follows:

[A/RE] rM ] ' /2K ''1/6 _ tXXREJ ""S p-l~4 (Eq. 3.67)

[Vo] = 2 ~ o~

p __ A ]1/2 Tr ,'1/6 "XREJ "xS p1/4 (Eq. 3.68)

o,

[ARE]-I/2Ks 1/6 p-x/a (Eq. 3.69) n = K i 2-1/2K -1/6 -o~

(iv) Region IV: In this low oxygen activity region, the electroneutrality

condition and the relations for p and n are given by the following equations:

2[Vo] - [A~E] (Eq. 3.70)

K 1/3 D1/4 P ~ ~1"1/3 �9 02 ( E q . 3 . 7 1 )

JtX S

Electrical Conduction in Ceramics 5 9

and

Ks /3 l)-1/4 n = K K 1/3" 02

(Eq. 3.72)

(v) Region V: Eventually, with decreasing oxygen partial pressure, B 3+

may be reduced to B 2+. In this case, the reduction is the dominant reaction (with

equilibrium constant KR), and the defect concentrations are given as follows"

n = [ V o ] = (4KR)l/3Po~/6" (Eq. 3.73)

2

p Ki D1/6 = �9 o, (Eq. 3.74) 2(4Ke) 1/3

Figure 3.7 summarizes the defect behavior in various oxygen partial

pressure regions for a REBO3 perovskite doped with AO acceptor dopant. The above defect model does not take into account association of defects.

However, point defects in REBO3 perovskites may not be single, unassociated, and randomly distributed. Defect models assuming formation of defect pairs

have been proposed [3.31].

V IV III II I e �9

�9 " , �9 t V o J = �9 " P = [ A ~ ] " r

. , I.o . .

�9 o O r o~ , ~ - ~ . ' ~ * , . �9

�9 �9 �9 n - ~ �9

L O G O X Y G E N P A R T I A L P R E S S U R E

Figure 3.7. Defect concentration as a function of oxygen partial pressure for a rare-earth perovskite oxide

60 Chapter 3

3.4 CONDUCTION PROCESSES AND TRANSFERENCE NUMBERS

The total conduction process in an oxide solid conductor is generally

mixed ionic and electronic. Therefore, it is necessary to characterize the partial current density and transference number of each conducting species as a function

of oxygen partial pressure. This section discusses general transport processes, ionic and electronic currents, and transference numbers in oxide electrical

conductors.

3.4.1 General transport equations

Let the partial current density carried by conducting species X be jx, the

charge of X be zxq where Zx is the valence of X, q be the absolute electronic

charge, the concentration of X be nx, the mobility be #x, and the diffusion

coefficient associated with X be Dx. Then, jx is expressed by

Jx = - n x IZx Iq~txV~ - Dx IZx IqVnx (Eq. 3.75)

where 9 is the electrical potential. T, he electrochemical potential for X, /Xx, is given as

~tx = ~x + Izxlq~ (Eq. 3.76)

where ~x~iS the chemical potential for X as represented approximately by the following equation:

! 0 Ix x - I~ x + KTlnn x (Eq. 3.77)

(nx is used instead of the activity of X.) The Nernst-Einstein equation for Dx is

given by

K/'~ x D x - (Eq. 3.78) IZxlq

and the gradient in the electrochemical potential is expressed as

_ K/Vn x V~x = + [zxlqV~ (Eq. 3.79)

nx

Combining Eqs. 3.75 and 3.78 yields

Jx = -nx IZx Iq~txV* - ~xKrVnx (Eq. 3.80)


From Eqs. 3.79 and 3.80, the partial current density of species X is given by

Jx = -nxl-txVlXx (Eq. 3.81)

The conductivity of species X is given as

o x = nx~xlZxlq (Eq. 3.82)

Thus, from Eqs. 3.81 and 3.82, the partial current density can be expressed as

oxV~t x Jx - - ~ (Eq. 3.83)

I z x l q

From this equation, the partial current density for a particular species is shown to be proportional to the gradient of the electrochemical potential for that species.

3.4.2 Electronic, ionic, and total current

The total current density carried by electronic species (electron or

electron hole), Je, is given by

From Eq. 3.83,

j , = j,, - jp (Eq. 3.84)

= + t, V~ (Eq. 3.85) j , "VI~ ,, p q q

Eq. 3.85 can be rearranged to

(O n + O p) Je = - Vp. n + Op( -.~n + -~p) (Eq. 3.86)

q q

Assuming that the thermodynamic equilibrium between electrons and holes is not disturbed by the flow of particles, the following equation is established:

Thus,

V~ n Vlsp + =0 (Eq. 3.87)

(O +Op j~ = _ n 7-~',, (Eq. 3.88)

q

62 Chapter 3

Since tr~ = tr~ + trp, the electronic current density is given as

m

o -u o .Vi.t (Eq 3.89) ~V_,, - p q q

Similarly, the ionic current density is given by

j, = - (~176 = (~176 V~I (Eq. 3.90) Iz lq lq

where the subscripts 1 and 2 represent the cations and anions, respectively.

The total current density is given as

f f - - 0 i j = j , + j , = "Vl~p V-~x (Eq. 3.91)

q Iz lq

where X is one of the various species in the specimen by which the thermody-

namic equilibrium is expressed as

x + [z lh" = x (Eq. 3.92)

3.4.3 Transference number measurements

Solid oxide electrical conductors are generally mixed electronic and ionic

conducting materials. Under specified conditions, whether an oxide is ionic,

electronic, or mixed conducting depends on the relative magnitude of the ionic

transference number ti (including cation and anion contributions) and the

electronic transference number t~ (the sum of the electron and electron hole

transference numbers, t~ + tp). Evaluation of transference numbers of an oxide

conductor is important in the selection of a material for use as a SOFC

electrolyte, electrode, or interconnect component.

For fuel cell applications, the transference number of ions must be above

0.99 and that of electrons below 0.01 for electrolyte, and a large transference

number of electrons is desired for electrode materials. For example, (ZrO2)0.8:

(CaO)0.15 is an ionic conductor suitable for SOFC electrolyte application. With P02 = 1 atm (105 Pa) on one side and Po2 = 2 • 10 -~ to 10 -17 atm (2 • 10 4 to

10 -~2 Pa) on the other side, the material has an ionic transference number of

0.998 to 0.994 at 1000 oc [3.32]. Ionic transference numbers can be determined

using oxygen gas concentration cells. For example, in the concentration cell

shown below


Fe, FeO J (ZrO:)o.85-(CaO)o.,, ] Cu:O,Cu (Eq. 3.93)

the oxygen pressures at the anode and cathode are determined by the equilibriums

FeO ~ Fe + %02 CuO ~ 2Cu + 1/202

(Eq. 3.94)

(Eq. 3.95)

By this oxygen concentration method, the ionic transference number of

(ZrO2)0.85(CaO)0.15 has been found to be 0.997 at 700 ~ and 0.996 at 800~

[3.33]. Figure 3.8 lists some fluorite-type oxygen-ion conductors which exhibit

oxygen-ion transference numbers above 0.99 at 1000~ as a function of composi-

tion and oxygen partial pressure [3.34].

In an oxygen gas concentration cell, the electromotive force, dE, of the

cell is given by

d E / = ti(-~)dlnPo~ q

(Eq. 3.96)

COMPOSITION

MOL%

ZrOz-CaO (85:15)

ZrO2-Y016 (80:20)

i ThO2-Y016 (99:1 ) i

ThO2-Y01.5 (96:4)

ThO2-Y016 (90:10)

ThOz-YOl.s (85:15) ThO2-La01.6 (90:1 O}

ThO2"La01.6 (85:15]

Tho2-LaO, 5 (75i2r~) I

; T h o , - c " o ]

02. TRANSFERENCE NUMBER > 0.99 ( 1000 o C)

~. ~ ~ 0 ~, 0 =, ~ o o ~ o o z = o II 0 7 ~7 7, ~ '7, ,.;~o-

- , I , ~ _ . I {

I [ I [

" , t I I I ' = I I I I I I I I

i171l i f

I I [ I 1 I I I

I , . I I

0 5 10 15 20 25 -LOG OXYGEN PARTIAL PRESSURE, atm (1.01 x 10 ~ Pa)

Figure 3.8. Oxygen partial pressure range for fluorite-type conductors exhibiting oxygen-ion transference numbers > 0.99 at IO00~ [3.34]

64 Chapter 3

Then, the measurable electromotive force of the concentration cell is given as

= KTi~~ (Eq. 3.97) E / 4q J tta'o~,~

Differentiation of Eq. 3.97 with respect to lnPo2 yields the following relation for

the ionic transference number:

4q aE / (Eq. 3.98) ti = KT 01nPo2

This equation shows that the ionic transference number can be obtained

experimentally at any oxygen partial pressure from the slope of the curve of the

electromotive force versus lnPo2 at that oxygen pressure. Thus, in the case of

fuel cells, by plotting E / (obtained at the same cathode oxygen pressure, Po2c~),

and different anode oxygen pressures, PO2(a), ) against lneo2(a), t i can be obtained

as a function of PO2(a). By this type of measurement, t i has been found to be

almost unity over wide oxygen pressure ranges for stabilized zirconia at elevated

temperatures [3.5, 3 . 10 , 3 . 35 , 3 .36] .

In certain cases, it may be necessary to determine the value of t~ directly

(instead of evaluating it by 1 - O, especially when the value of t e is small (as in

the case of SOFC electrolytes). In SOFCs, even under open-circuit conditions,

currents continue to flow so long as their sum is zero. In such conditions, the

particles move via chemical diffusion. According to ambipolar diffusion theory

[ 3 . 3 7 , 3 . 3 8 ] , the chemical or ambipolar diffusion coefficient, Da, in mixed

conductors is given as

D a = teD i + t i d e (Eq. 3.99)

where D i and D e are the oxygen ion and electronic diffusion coefficients,

respectively. When ti ~ te (i.e., for predominantly ionic conductors),

D a = t iD e = D e (Eq. 3.100)

The two experimental methods most commonly used to evaluate small

electronic contributions in an oxide conductor are the permeation [3.35] and

polarized cell [3.39] techniques. (i) P e r m e a t i o n t e c h n i q u e : When a fixed gradient of oxygen pressure is

maintained across a specimen, a steady state flux of neutral oxygen flows from


the high oxygen pressure side to the low oxygen pressure side. The oxygen flux,

Jo, is given by the following equation:

Jo = -Da~Co (Eq. 3.101)

where Co represents the concentration of oxygen.

By using Eq. 3.100, expressing the Wagner scaling equation in terms of

conductivities [3.36], and introducing KTInP for the chemical potential of neutral

species, the permeation current density, j , is given by

KT o t~ VInP (Eq. 3 102) J = 4q ~ o~

Integration over the sample thickness, L, leads to

KT tin/,// j = ~O4qL ~J~,, tedlnPo2 (Eq. 3.103)

where P / and P H are oxygen pressures at the high-pressure side and a fixed low-

pressure side of the permeation cell, respectively. For the evaluation of the

transference number, t~, at a certain oxygen pressure of P, it is useful to

differentiate Eq. 3.103 with respect to the upper integration limit to obtain

OlnPo~ eo, -ell -- o i t~ (Eq. 3.104)

From the relation above, te at P can be evaluated in terms of the slope of the plot

of experimental permeation current densityj versus lnP". It should be noted that

this analysis has no restrictions and can be used even when the dependence of or e

on oxygen partial pressure (electronic conductivity being n- or p-type) is not known.

(ii) Polarized cell technique: Partial electronic conductivities can also be

determined by using ion-blocking electrodes to suppress ionic transport so that

only electrons and electron holes can pass through the system. This is accom-

plished by using a chemically inert material, like Pt or Au, as a blocking

electrode (the cathode in this case), so that no ions are supplied to the specimen

when a DC voltage, E, (below the decomposition voltage of the sample material)

is applied. An electrode that is reversible to ionic species is used as the counter

electrode (the anode) to fix the chemical potential of the components. In this

case, conduction takes place only by the migration of electronic carriers.

66 Chapter 3

Assuming that the electronic carriers are dilute enough to obey ideal solution la~ and to have a constant mobility [3.39], the following equation has been derived:

i _ exp(_~T) ~[ P ~:lJ (Eq. 3.105)

where a~ / and a/represent the electron and hole conductivities in chemical

potential determined by the reversible electrode. Then, a plot of the left side of

Eq. 3.105 versus exp(qE/KT) should give a straight line. The electron and hole conductivities are obtained from the slope of the plot and from the intercept with

the ordinate, respectively [3.40]. The key technical challenge in this technique is to have satisfactory

blocking conditions. However, it has been demonstrated that even though blocking of the ionic flow is incomplete, the polarized cell treatment discussed above describes well observed experimental data [3.41].

The techniques described above are static methods; in addition, the application of these techniques under dynamic conditions using transient measurements has been demonstrated. For example, if the applied voltage is removed after attaining a steady state, the concentration gradients will decay until equilibrium concentrations are reached. These transients will show measurable voltage relaxations. From the time dependence of these voltage relaxations, expressions for diffusion coefficients of electrons and electron holes can be obtained [3.42-3.44].

(iii) Other techniques: There are also many other techniques for determining transference numbers [3.35,3.40,3.45-3.47]. For instance, the electron and electron hole contributions are evaluated by subtracting the oxygen pressure independent component from the total conductivity at a certain oxygen

partial pressure [3.40,3.45,3.46]. The transference number of a conducting

species may be determined by measuring its self diffusion coefficient, calculating its conductivity, and comparing the calculated conductivity with the total

conductivity [3.35]. The individual contributions of anions, cations,and electrons

to the total conductivity may be evaluated by electrolysis experiments. The thermoelectronic emission from oxides at elevated temperatures has been found to be useful to identify electronic contributions, since the emission flux is sensitive only to the electron density [3.47].


References

3.1

3.2

3.3 3.4

3.5

3.6

3.7 3.8 3.9 3.10

3.11 3.12 3.13

3.14

3.15

3.16

3.17

3.18

3.19

3.20

3.21

3.22

3.23 3.24

3.25

3.26

3.27

J. Frenkel, Z. Phys. Chem., 35 (1926) 652.

W. Schottky, Z. Phys. Chem., 29B (1935) 335. W. Jost, J. Chem. Phys., 1 (1933)466. E. Koch and C. Wagner, Z. Phys. Chem., 38B (1937) 295.

F. Fund, Z. Elektrochem., 55 (1951) 363.

F. Fund, Z. Phys. Chem., 199 (1952) 142.

C. Wagner, Naturwissenschaften, 31 (1943) 265. T.Y. Tien and E.C. Subbarao, J. Chem. Phys., 39 (1962) 572. H. Schmalzried, Z. Elektrochem., 66 (1962) 572. R.E. Carter and W.L. Roth, G.E. Research Report No. 63-RL-3479M, General

Electric, Schenectady, NY, 1963. J. Faber, Jr., M.A. Seitz, and M.H. MOiler, J. Phys. Chem. Solids, 37 (1976) 903. J. Faber, Jr., M.A. Seitz, and M.H. MOiler, J. Phys. Chem. Solids, 37 (1976) 909. W.D. Kingery, J. Pappis, M.E. Doty, and D.C. Hill, J. Am. Ceram. Soc., 42

(1959) 393.

C.B. Choudhary, H.S. Maiti, and E.C. Subbarao, in Solid Electrolytes and Their Applications, E. C. Subbarao (ed.), Plenum Press, New York, 1980, p. 1.

E.C. Subbarao and H.S. Maiti, Solid State lonics, 11 (1984) 317.

C.A.C. Sequeira and J.M.B.F. Diniz, Phys. Status Solidi A, 165 (1988) 123.

O.T. Sorensen, 0. Johannesen, and K. Clausen, in Transport-Structure Relations in Fast Ion and Mixed Conductors, F.W. Poulsen, N.H. Andersen, K. Clausen, S. Skaarup, and O.T. S~rensen (eds.), Ris~ National Laboratory, Roskilde, Denmark, 1985, p. 93. 0. Johannsen, in Proceedings of International Energy Agency Workshop on Mathematical Modelling of Natural Gas Fueled Solid Oxide Fuel Cells & Systems, July 2-6, 1989, Charmey, Switzerland, Swiss Federal Office of Energy, Bern, Switzerland, 1989, p. 99.

L. Heyne and D. Engleson, J. Electrochem. Soc., 19 (1977) 727.

H.L. Tuller, in Nonstoichiometric Oxides, O.T. Sorensen (ed.), Academic Press, New York, 1981, p. 271.

J.W. Patterson, in Electrical Conductivity in Ceramic and Glass, Part B, N.M. Tallan (ed.), Marcel Dekker, New York, 1974, p. 453.

A.S. Nowick and D.S. Park, in Superionic Conductors, G.D. Mahan and W.L. Roth

(eds.), Plenum Press, New York, 1976, p. 395.

H.S. Maiti and E.C. Subbarao, J. Electrochem. Soc., 123 (1976) 1713. B. Hodson and P.T. Moseley, J. Solid State Chem., 19 (1976) 383.

S.F. Bartram, Inorg. Chem., 5 (1966) 749.

M.R. Thornber, D.J.M. Bevan, and J. Graham, Acta Cryst. B, 24 (1968) 1183.

S.P. Ray and D.E. Cox, J. Solid State Chem., 15 (1975) 333.

68 Chapter 3

3.28

3.29

3.30 3.31 3.32

3.33

3.34

3.35

3.36

3.37

3.38

3.39

3.40

3.41 3.42

3.43

3.44

3.45

3.46

3.47

H.U. Anderson, in Proceedings of the 14th Riso International Symposium on Materials Science, High Temperature Electrochemical Behaviour of Fast Ion and Mixed Conductors, September 6-10, 1993, Roskilde, Denmark, F.W. Poulsen, J.J.

Bentzen, T. Jacobsen, E. Skou, and M.J.L. OstergArd (eds.), Riso National Laboratory, Roskilde, Denmark, 1993, p. 1.

H.U. Anderson, J.H. Kuo, and D.M. Sparlin, in Proceedings of the First International Symposium on Solid Oxide Fuel Cells, October 16-18, 1989, Hollywood, FL, S.C. Singhal (ed.), Electrochemical Society, Pennington, NJ, 1989,

p. 111. H.U. Anderson, Solid State Ionics, 52 (1992) 33. P. Kofstad and A. Petrov, see Ref. 3.28, p. 287. J. Weissbart and R. Ruka, in Fuel Cells, Vol. II, G.J. Young (ed.), Reinhold, New York, 1963, p. 37. S.F. Pal'guev and A.D. Neuimin, in Transactions (Trudy) of the Institute of Electrochemistry No. 1, Electrochemistry of Molten and Solid Electrolytes, Consultants Bureau, New York, 1961, p. 90. B.C.H. Steele and C.B. Alcock, Trans. Metall. Soc. AIME, 233 (1965) 1359.

L.A. Simpson and R.C. Carter, J. Am. Ceram. Soc., 49 (1966) 139.

K. Kiukkola and C. Wagner, J. Electrochem. Soc., 104 (1961) 379.

W. van Roosbroeck, Phys. Rev., 91 (1953) 282.

L. Heyne, in Solid Electrolytes, S. Geller (ed.), Springer-Verlag, New York, 1977,

p. 169. C. Wagner, in Proceedings of the 7th Meeting of C.LT.C.E., 1955, Lindau, Germany, Butterworths Scientific, London, U.K., 1957, p. 361. J.W. Patterson, E.C. Bogren, and R.A. Rapp, J. Electrochem. Soc., 114 (1967)

752. D.Y. Wang and A.S. Nowick, J. Electrochem. Soc., 127 (1980) 752.

W. Weppner, Z. Naturforsch. A, 31 (1976) 1336.

L. Heyne and N.M. Beekmans, Proc. Br. Ceram. Soc., 19 (1971) 229.

K. Kitazawa and R.L. Coble, J. Am. Ceram. Soc., 57 (1974) 360.

H.L. Tuller and A.S. Nowick, J. Electrochem. Soc., 122 (1975) 255.

T. Takahashi, H. Iwahara, and T. Esaka, J. Electrochem. Soc., 124 (1977) 1563.

J.C. Rifflet, P. Odier, A.M. Anthony, and J.P. Loup, J. Am. Ceram. Soc., 58

(1975) 493.

Chapter 4

ELECTROLYTE

4.1 REQUIREMENTS

The main function of the SOFC electrolyte is to conduct ions between the anode and cathode. The electrolyte carries the ions produced at one electrode to the other electrode to balance the charge from the electron flow and complete the electrical circuit in the fuel cell. The conducting ion must be or contain one of the elements present in the fuel and/or oxidant. The electrolyte also separates the fuel from the oxidant in the fuel cell. Thus, the electrolyte material must be stable in both the reducing and oxidizing environments, impermeable to the

reacting gases, and sufficiently conductive (ionically) at the operating conditions. Since the SOFC operates at high temperatures (600 ~ to 1000~ the electrolyte

must be chemically and thermally compatible with the other cell components from room temperature to those operating temperatures, and to even higher temperatures at which the fuel cell is fabricated. The key requirements for the electrolyte in the SOFC are discussed below. This discussion is qualitative because the specific requirements depend on selected materials and cell and stack designs [4.1].

(i) Stability: The electrolyte must be chemically, morphologically, and dimensionally stable in the dual atmosphere (reducing atmosphere on one side and oxidizing atmosphere on the other). It may not experience any disruptive phase transformation (involving large changes in molar volume) between room

temperature and fabrication temperature.

(ii) Conductivity: The electrolyte must possess adequate ionic conductivity in the dual atmosphere (at the fuel cell operating temperature); its ionic

conductivity must be as high as possible to minimize ohmic losses. The electrolyte must also have negligible electronic conductivity to prevent voltage

losses due to the electronic current flowing through the electrolyte. The electrolyte conductivity must not change (age) significantly over prolonged periods.

70 Chapter 4

(iii) Compatibility: The electrolyte must be chemically compatible with the other cell components, not only at the operating temperature, but also at the much higher temperature at which the fuel cell ceramic structure is fabricated. The electrolyte is typically the first component considered in the construction of

the SOFC. Therefore, other cell materials must be selected to limit chemical interaction and elemental diffusion between the electrolyte and the other

components in order to minimize unacceptable occurrences such as insulating phase formation, stability reduction, change in thermal expansion, introduction of significant electronic conductivity in the electrolyte, etc.

(iv) Thermal expansion: The thermal expansion of the electrolyte must match (from room temperature to operation and fabrication temperatures) that of other cell components to avoid cracking and delamination during fabrication and operation, including thermal cycling. The coefficient of thermal expansion of the electrolyte material must remain unchanged despite changes in oxygen partial pressures of the fuel and oxidant atmospheres during operation. The electrolyte

is typically selected as the baseline material, and the thermal expansion characteristics of other materials are tailored to match those of the electrolyte.

(v) Porosity: The electrolyte must be dense (or contain no connected porosity) to prevent gas cross leakage. The electrolyte material must be impervious to both oxidant (oxygen) and fuel (hydrogen) gases between room temperature and operating temperature.

In addition to these requirements, other desirable properties for the SOFC electrolyte are high strength and toughness, fabricability, and low cost.

Present SOFCs use, almost exclusively, stabilized zirconia (ZrO2), especially yttria (Y203)-stabilized ZrO2 (YSZ), as the electrolyte. Other oxygen- ion conductors (such as doped CeO2, stabilized Bi203, and other conducting oxides) have also been proposed as electrolyte materials for the SOFC, especially for reduced-temperature operation (600 ~ to 800~ In addition, a number of proton-conducting oxides have been studied for use as SOFC electrolytes.

4.2 STABILIZED ZIRCONIA

Stabilized ZrO 2 has been used almost exclusively as the electrolyte in SOFCs because the material possesses an adequate level of oxygen-ion

conductivity and exhibits desirable stability in both oxidizing and reducing atmospheres. Among the various stabilized zirconias, the Y203-stabilized

material is the most common.

Electrolyte 71

4.2.1 Preparation

In SOFCs, the stabilized ZrO2 electrolyte is generally fabricated as a polycrystalline dense film or layer. Various processes have been developed to make thin ZrO2 electrolyte layers for the SOFC (see fabrication sections of Chapter 9 for more details). These processes are based on either the particulate approach or the deposition (including coating) approach. The particulate approach involves compaction of ZrO2 powder into the desired shape, and densification at elevated temperatures. Examples of the particulate approach are tape casting and tape calendering. The deposition approach involves the formation of a thin layer (on a substrate or support) by a chemical or physical process. Examples of the deposition approach are electrochemical vapor deposition (EVD) and plasma spraying.

The most important aspect in the fabrication of the electrolyte is the production of a fully dense layer (as thin as possible). In the particulate approach, the densification of powder at elevated temperatures is dependent on material and processing factors, such as powder characteristics (reactivity, purity, morphology), particle packing (green density), and processing conditions (temperature, time, atmosphere). Ideally, fine (high surface area) spherical particles with a narrow size distribution are desirable because they result in high reactivity and high packing density, and therefore, enhanced densification of powders (into sintered bodies with uniform microstructure) at low sintering temperature. At present, ZrO2 powders with such characteristics can be prepared and have been processed into dense layers under conditions suitable for SOFC fabrication. Recently, further advancements have been made in powder synthesis and processing of ZrO2 materials. For example, YSZ powders of submicrometer size (nanocrystalline size) have been formed into a green body (about 50 % green density) and fired to 95 % theoretical density in air at 1125~ [4.2]. Microwave sintering has been demonstrated to densify ZrO2 powders at reduced temperatures [4.3] and shorter hold times [4.4]. For SOFC applications, common particulate methods are tape casting [4.5,4.6] and tape calendering [4.7,4.8]. Tape calendering is also suitable for making pinhole-free, micrometer-thick YSZ

electrolyte layers for reduced-temperature operation [4.9]. Figure 4.1 shows an example of the microstructure of YSZ electrolyte made by tape calendering, a particulate method.

72 Chapter 4

Figure 4.1. Microstructure of YSZ electrolyte made by tape calendering

In the deposition approach, the electrolyte is commonly formed as a thin layer on a substrate or support. The key consideration in this approach is the ability of the method to produce a dense and pinhole-free layer. Therefore, process parameters of selected methods are often tailored to achieve this requirement. At present, the common deposition methods for the SOFC are EVD [4.10] and plasma spraying [4.11]. Many other deposition methods have also been proposed and investigated, including chemical vapor deposition (CVD) [4.12], rf sputtering [4.13], rf ion plating [4.14], spray pyrolysis [4.15], slurry coating [4.14], vapor-phase electrolytic deposition [4.16], CO2 laser evaporation [4.17], laser spraying [4.18], pulse laser deposition [4.19], sol-gel coating [4.20], plasma enhanced metal organic CVD [4.21], jet vapor deposition [4.22], electrophoretic deposition [4.23], vacuum evaporation [4.24], and hybrid plasma spraying [4.25]. Several of these techniques have been studied for making electrolyte layers thinner than 10/zm (see Chapter 9, section 9.5.3 for further details and more references).

In the particulate methods discussed above and certain deposition techniques, ZrO2 precursor powders containing a desired amount of stabilizing oxide are required. The ideal powder should be homogeneous (having the stabilizing additive uniformly distributed on an atomic scale), yet large enough to handle, while the surface area should be sufficiently high to allow sintering at the lowest possible temperature [4.26]. A wide range of stabilized ZrO2 powders have been produced by various methods (Table 4.1) [4.27]. At present, highly reactive and uniform ZrO2 powders are commercially available. Figure 4.2 shows the steps of a manufacturing process based on hydrolysis of ZrOC12 and

Electrolyte 73

TABLE 4.1

Characteristics of Powders Produced by Various Preparation Methods [4.27]

Method

Calcining Agglomerate Tap Surface

temperature size density area

(~ (~m) (% theor.) (mZ/g)

Hot kerosene

Citrate Sol-gel microsphere

Peroxide Acetone-toluene

Alkoxide

Chloride

600 2-20 18 4

650 5 58

650 1-20 26 32

600 <__50 38 82 650 ___ 100 32 16

650 12 90

650 10 123

ZIRCON SAND (ZrO2SiO2)

Na2ZrO3Na2SiO3

ZrOCI2-HCI

ZrOCI2"8H20 SOLUTION

Y203

CHLORIDE SOLUTION

HYDROLYSIS

CALCINATION

MILLING

SPRAYING

YSZ POWDER

Figure 4.2. A manufacturing process for producing YSZ powder

74 Chapter 4

YC13 to produce fine YSZ powders. As ZrO2 powders are generally prepared

from chloride precursors, these materials may contain chloride residue. The

presence of a significant amount of chloride (1 wt%) can raise the sintering

temperature of the material as much as 150~ [4.28,4.29]. If necessary,

laundering techniques can be used to remove the chloride impurity [4.28, 4.29].

4.2.2 General properties and phase transformation

ZrO2, in its pure form, exhibits three well-defined polymorphs. At room temperature, ZrO2 has a monoclinic crystal structure. The monoclinic structure changes to a tetragonal form above 1170~ and to a cubic fluorite structure

above 2370~ The monoclinic/tetragonal transformation in ZrO2 is thermodynamically reversible but associated with a large volume change (3 to 5%)

(contraction on heating and expansion on cooling). The cubic phase exists up to the melting point of 2680~ However, the addition of certain aliovalent oxides

can stabilize the cubic fluorite structure of ZrO2 from room temperature to its melting point. Table 4.2 lists several properties of ZrO2. The properties of stabilized ZrO2 have been extensively studied [4.30-4.33], and several reviews

on the subject are available [4.34-4.36]. The common stabilizing oxides for ZrO2 are CaO, Y203, MgO, Sc203,

and certain rare-earth oxides. These oxides exhibit a relatively high solubility

in ZrO2 and are able to form various solid solutions with ZrO2, including cubic fluorite solid solutions which are stable over wide ranges of composition and

temperature. Figure 4.3 shows the equilibrium phase diagram for the ZrO2-CaO

system (where M~, Ts~, and C~s indicate monoclinic, tetragonal, and cubic solid

solution, respectively) [4.37]. The system has a eutectoid at about 17 mol % CaO in the composition range CaO-CaZrO3. From 6 to 17 mol% CaO, the material

(known as partially stabilized ZrO2) consists of tetragonal solid solution and monoclinic solid solution phases above 1140 ~ Slow cooling from 1140 ~ to

1000~ results in the tetragonal solid solution phase and CaZr409 (the eutectoid decomposition product). Further cooling below 1000~ causes the martensitic

transformation of the tetragonal solid solution to monoclinic solid solution. The cubic phase is thermodynamically unstable at low temperatures. The eutectoid

decomposition of cubic solid solution occurs at 1140~

Electrolyte 75

TABLE 4.2

Properties of ZrO2

Melting point, ~

Density, g/cm 3

Undoped ZrO 2 (monoclinic)

YSZ (8 mol% Y203) Electrical conductivity, f~-lcm-~

YSZ (9 mol% Y203) (1000~

YSZ (9 mol% Y203) (600~ Thermal conductivity, W/cm.K

Thermal expansion coefficient, cm/cm.K

Undoped ZrO2 (20 ~ to 1180~

YSZ (8 mol% Y203) (100 ~ to 1000~ Standard enthalpy of formation (25~ kJ/mol

Standard entropy of formation (25~ J/mol.K

Bend strength, MPa

YSZ (8 mol% Y203) at 25~ YSZ (8 mol% Y203) at 1000~

Fracture toughness at 25~ (YSZ), MN'm -3/2

2680 [4.34]

5.56 [4.34] 5.90 [4.34]

0.12 [4.36] 0.006 [4.36] 0.02 [*]

8.12 [4.106] 10.8 [4.110] -1,097.5 [**]

50.4 [**]

300 [***]

225 [4.118] 3 [***]

[*] from brochure of Vesuvius Mc Danel, Beaver Falls, PA

[**] from JANAF Thermochemical Tables, Third Edition, American Institute of Physics, New York [***] from brochure of Tosoh Corporation, Tokyo, Japan

"" ... 2 2 5 0 + 20 "C ~ - - "

0 o z o o o u~ n,- T =) lss ss Css Css t- CaZr03 I--

< C s s r r

u J n 1 5 0 0

gJ Tss '* 0 0 0 0 ' �9 �9 1310 -4- 4 0 " C I-- M s s �9 . o o , �9 �9

Css + CaZr 4 0 9

�9 _ , , . . . . . . . . i o o o [ , . . �9 o ~ c.z,o .

t" "

M s s "P C e Z t 4 0 9 i / m

0 I0 20 30 4 0 ZrO 2 CeZr 4 0 9

M o I % C a O

50 C0ZrO$

Figure 4.3. Phase diagram of a ZrOz-CaO system [4.37]

76 Chapter 4

For the ZrO2-Y203 system, several equilibrium phase diagrams have been

reported (and certain discrepancies still exist). Figure 4.4 shows a phase diagram

of the ZrO2-Y203 system [4.38]. The diagram of the low Y203 region is given in Figure 4.5 [4.39, 4.40]. It can be seen from Figure 4.5 that the addition of Y203 to ZrO 2 reduces the temperature of the tetragonal/monoclinic transformation, and the transformation temperature decreases with increasing Y203 content (in the composition range 0 to 2.5 mol % Y203). In this composition range, the tetragonal solid solution is transformable, i.e., the tetragonal phase will transform on cooling to the monoclinic phase. At higher Y203 contents, a mixture of

nontransformable tetragonal and cubic solid solutions exists. Further increase in Y203 content results in a homogeneous cubic solid solution. The minimum Y203

amount required to fully stabilize the cubic phase of ZrO2 is about 8 to 10 mol % at 1000~ [4.38,4.39,4.41]. The ZrO2-MgO has a eutectoid at 13 mol% MgO [4.37]. A cubic solid solution exists at this composition above 1400~ The cubic solid solution decomposes into a tetragonal solid solution and MgO below 1400~ and again into a monoclinic solid solution and MgO below 1240~

Other ZrO2-M203 systems (where M is ytterbium, scandium, neodymium,

samarium, or gadolinium) have also shown stabilized solid solutions in a certain

M203 range [4.42-4.51]. The minimum amount of M203 necessary to stabilize ZrO2 in cubic structures is close to the composition which gives the highest

conductivity (approximately 8 mol % for Yb203, 10 mol % for Sc203, 15 mol % for Nd203, 10 mol % for Sm203, and 10 mol % for Gd203).

4.2.3 Stability

ZrO 2 is chemically stable in the SOFC oxidizing and reducing atmospheres. Only under highly reducing conditions, e.g., below 10 .25 to 10 .30 atm

(10 .20 to 10 .25 Pa) at 1000~ (not normal SOFC conditions), ZrO2 is blackened

(reduced) by forming a zirconia suboxide, ZrO2_~. The black coloration has been

ascribed to entrapped electrons [4.52, 4.53]. The reduction or blackening of ZrO2

can also be produced if a current density higher than the limiting current densities

(associated with the gas electrode reactions) is forced through the electrolyte. Pure ZrO2 does not serve as a good electrolyte because its ionic

conductivity is too low, and it has disruptive phase transformations on heating and cooling. Doping ZrO2 with a certain divalent or trivalent oxide can stabilize

the cubic fluorite phase and, at the same time, increase its oxygen vacancy concentration. This enhances the ionic conductivity and extends the oxygen partial

Electrolyte 77

3 0 0 0 r ~ ~ 1

i r F

20001 U o ui er"

I-- < 11= W

'" i I--

1 0 0 0

" ~ ~ ~ . . ~ ... ' ~ -~ ~

oo t Fr~

| { 13!! oc I I

I | o I ! \ / Fs=

\ / + r,,,, +o~\o /z,.~.o,. s= | ~r \\ i i

\ /

Z r 0 2

F = + C =

1377:1: 5=C

Z%Y4Otz + Cs=

- ! i I

] i

i , ~s= ~

\ \

\ \

\ \ \ - \

\

\

\ ,

'i I

Ms= + Z%Y,Otz h ~ , ~ ,

2"0 4 0 6 0 8 0 Y203

Mol % Y203

Figure 4.4. Phase diagram of a ZEO2-Y203 system [4.38]

ooo t 2500

I T I LIQUID

{-) 2000 o

er

1500

I.U

I.U I-- 1000

500

0 ~ - . ( 6 % ) TET ( 1 1 % _ ) _ _ ~ ,.

0 5 10 15 20

z~o~ Mol % YO 1.s ~12(Y2O3)---

Figure 4.5. Phase diagram of a ZFO2-Y203 system in the low }'203 region [4.39]

78 Chapter 4

pressure range of ionic conduction, making stabilized ZrO2 suitable for use as an electrolyte in SOFCs. (This extended oxygen partial pressure range covers the conditions, 1 to 10 -17 atm (105 to 10 -~2 Pa), to which a SOFC electrolyte is exposed in the operating fuel cell.) In general, fully stabilized ZrO2 is preferred as SOFC electrolyte material in order to maximize conductivity. The use of fully stabilized ZrO2 also avoids the problems of phase transformation and conductivity aging associated with partially stabilized materials during cell operation.

4.2.4 Electrical conductivity

ZrO2 stabilized with a cation of appropriate size is an oxygen-ion conductor. ZrO2 is stabilized by direct substitution of divalent or trivalent cations for the host lattice cation Zr 4§ This substitution creates a large concentration of oxygen vacancies by charge compensation according to following equation (written for Y203 stabilization using Kr6ger-Vink notation):

ZrO2 y~o3 _. 2y/z, + Vo + 300 (Eq. 4.1)

The high oxygen vacancy concentration gives rise to high oxygen-ion mobility. Oxygen-ion conduction takes place in stabilized ZrO2 by movement of oxygen ions via vacancies. The general features of the defect structure and oxygen-ion conduction in stabilized ZrO2 are discussed in detail in Chapter 3.

The ionic conductivity of stabilized ZrO2 has been extensively studied and measured. The ionic conductivity behavior of the material is influenced by several factors such as dopant and dopant concentration, temperature, atmosphere, grain boundary, and time.

(i) Conductivity measurement: Complex impedance measurements are

commonly used to determine the ionic conductivity of stabilized ZrO2. It is well established that ionic conductivity of polycrystalline ZrO2 (and ceramics in general) depends on the microstructures, especially grain boundaries, of the material. The complex impedance method has also been used to separate the

contributions of grain boundaries to the total conductivity (resistivity) [4.54- 5.57]. Analysis of impedance data often employs equivalent circuits to elucidate the contributions of grain boundaries. The individual components of the equivalent circuit are related to parameters of the resistivity such as the bulk or grain interior resistance (Rgi) and capacitance (Cgi) and the grain-boundary resistance (Rgb) and capacitance (Cgb). Figure 4.6 shows, as an example, two equivalent circuits for polycrystalline ZrO2 [4.58]: one showing impedance

Electrolyte 79

Rgi

Cgi

] ~gb

c~o

Rb Cb

R~

c~

R~

c~

Re

Ce

t r fe

. fg FREQUENCY IN

F - - - Rgi - - - 4 " - - - Rgl, ~ - 4 - ' ~ R e "-I

Z / , ~ ---"-

Figure 4.6. Equivalent circuits and schematic complex impedance plot of polycrystalline Zr02 [4.581

components arising from grain interiors (gO, grain boundaries (gb), and electrodes

(e) [4.59] and one showing parallel connection of geometric capacitance (Cg) and

nonblocked (R~) and blocked (R b, Cb) ionic paths [4.60]. A complex impedance plot of the imaginary (Z//) vs. the real (Z t) components for these equivalent circuits (over a wide range of frequencies) results in three semicircles (Figure 4.6). The diameter of each semicircle is the resistance of the corresponding circuit element. The high-frequency semicircle provides the bulk resistance and capacitance of the interior of the grains; the intermediate-frequency semicircle provides the grain-boundary resistance and capacitance; and the low-frequency semicircle provides information on the oxygen-ion transfer at the electrodes. (In

practice, circular arcs rather than semicircles are often observed because the structure of actual grain boundaries is obviously more complex than that envisaged in the equivalent circuit.) It should be noted that the influence of grain boundaries on conductivity may vary depending on temperature. As a result, the

complex impedance plot (number and size of semicircles) may also vary with temperature. Figure 4.7 shows an example of a complex impedance plot for

( Z r O 2 ) o . 9 ( Y 2 0 3 ) o . 1 obtained at 800~ in air [4.61]. (ii) Influence of dopant and dopant concentration: The ionic conductivity

of stabilized Z r O 2 obviously depends on dopant and dopant concentration. Table

4.3 lists the conductivity data for stabilized Z r O 2 doped with various rare-earth oxides.

80 Chapter 4

100

80

60

N 40

20

Z I, fl

480 500 520 540 560 580 600 I i I i I I I

~o | - ~

1 2 k 2 k , r '~ 5OO

/ ~ I ! I I I

Figure 4.7. Complex impedance plot of polycrystalline (ZrOz)o.9(Y2Oz)o.1 at 800 o C in air [4. 61]

TABLE 4.3

Conductivity Data for Stabilized ZrO 2 Doped with Rare-Earth Oxides [4.35]

D o p a n t C o m p o s i t i o n Conductivity (1000~ Activation energy (M203) (mol % M203) (10 -2 fl-lcm-l) (kJ/mol)

Nd203 15 1.4 104 Sm203 10 5.8 92 Y203 8 10.0 96 Yb203 10 11.0 82 Sc203 10 25.0 62

With the same dopant concentration, the conductivity of stabilized ZrO 2

depends on the size of the dopant cation. Thus, the conductivity increases from

neodymium to scandium dopant (in the order shown in Table 4.3) since the ionic radius decreases from Nd 3+ to Sc 3+. (The radii of Nd 3+, Sm 3+, Gd 3+. y3+,

Yb 3+, and Sc 3 § are 0.104, 0.100, 0.097, 0.092, 0.086, and 0.081 nm,

respectively.) One explanation for this trend is that when Zr 4§ (the ionic radius

is 0.079 nm) is substituted by a cation with a more similar ionic radius, the

Electrolyte 81

lattice has less strain and fewer association of the cation with oxygen-ion vacancy, e.g., (MzrVo)". This results in greater mobility for oxygen ions, thus higher ionic conductivity. Another explanation for the decrease in conductivity with increasing dopant size is the steric blocking effect of larger cations

[4.56, 4.62-4.64]. As seen from Table 4.3, stabilization of ZrO2 with Y203 does

not yield the highest conductivity. However, Y203-stabilized ZrO2 or YSZ is most frequently used as SOFC electrolyte because of availability and cost.

The conductivity of stabilized ZrO2 varies with dopant concentration.

Figure 4.8 shows, as an example, variation of conductivity with dopant

concentration for various doped ZrO2 compounds at 1080 K [4.42,4.65].

Isothermal plots of conductivity as a function of dopant concentration exhibit a

maximum. Above the maximum conductivity dopant concentration, the

conductivity decreases with increase in dopant content, and this trend is

accompanied by an increase in activation energy for conduction. For the ZrO2-

Y203 system, the Y203 composition at which the conductivity maximum occurs has not been unambiguously defined, although it appears to fall between 8 and

10 mol % [4.66]. This value is close to or near the minimum quantity of dopant

required to fully stabilize the cubic fluorite phase [4.67-4. 76].

10-1

% 10 2 (,,}

c~ >: I,- > I- u a z 10-3 0 u

10 -4 l I l I 1

4 8 12 16 2O

MoI% M203 or MO

Figure 4.8. Variation of ionic conductivity of stabilized Z r O 2 with dopant concentration at 1080 K [4.65]

82 Chapter 4

The decrease in conductivity at higher dopant concentrations is believed

to be due to defect ordering, vacancy clustering, or electrostatic interaction [4.41]. One such mechanism is discussed here to explain the conductivity

behavior of the ZrO2-Y203 system [4.49]. At low Y203 concentrations, the average distance between defect complexes (considered as a one-fold associate

defect complex, YZr Vo) is large, and each oxygen-ion vacancy is trapped and immobilized inside the defect complex, resulting in low oxygen-ion conductivity. With increasing Y203 concentration, the defect complexes begin to overlap one another. The effective carrier concentration and the migration path for oxygen-

ion vacancies (through these defect complexes) thus increase with accompanying increase in conductivity. Further increase in the Y203 concentration leads to the

Zr VOYzr) [4. 77]. This appearance of two-fold associate defect complexes (Y / " /

decreases the effective carrier concentration and the effective oxygen-ion

migration path. The result is a decrease in conductivity, thereby producing a maximum conductivity as a function of dopant concentration. The appearance of a maximum in the oxygen-ion conductivity of stabilized ZrO2 is thus due to

increased trapping of vacancies as the concentration of Y203 dopant increases; i.e., as the dopant concentration increases, the distorted region decreases in size, and the disorder becomes dominated by aggregates of clusters.

(iii) Influence of temperature: The electrical conductivity of fully stabilized ZrO2 as a function of temperature typically follows Arrhenius-type

behavior. Figure 4.9 shows several Arrhenius conductivity plots for YSZ having

different Y203 contents [4.44]. In general, Arrhenius plots for YSZ tend to show two regions (low-temperature and high-temperature) of different slopes (or conduction activation energies) [4. 78]. The change in the slope of Arrhenius

plots has been attributed to the formation of dopant cation/vacancy complexes at

low temperatures [4. 79]. The conductivity/temperature relationship for YSZ is

given by the following equation [4.80]:

oT = Aoexp[-(tz+13T-~r] ~ (Eq. 4.2) leT

where o is the conductivity, T the temperature, r the Boltzmann constant, Ao the

preexponential constant, and a and B the positive constants. This equation has

been found to hold for both single crystal and polycrystalline YSZ. It should be noted that in a certain temperature range, c~ + BT l can be approximated as constant (thus, giving the activation energy for conduction, E,), and, in this case,

the more familiar Arrhenius equation is obtained, i.e., aT = A o exp(-Eo/KT).

Electrolyte 83

I 0 s

10 4

E U

>2 1 0 3 F- > ,,=,. I-. 0 ' } ,.=.

t ,Li

10 2

I- 1 1 I 1 I

oo6 10 tool% Y=,Oz l~x + 15 mol% v / / , 2o.,o1~ / p

- ///

/

/

1 l 1 I l _

0.6 0.7 0.8 0.9 1.0 1.1 1.2 1.3 lIT, 10"3K "1

Figure 4.9. Arrhenius resistivity plots for Y203-doped ZrO 2 materials [4.44]

(iv) Influence of atmosphere: In general, the ionic conductivity of cubic stabilized ZrO2 is independent of oxygen partial pressure over several orders of magnitude (Figure 4.10) [4.81]. Under these conditions, the ionic transport number is very close to unity (i.e., negligible electronic conduction). As discussed in Chapter 3, electronic conductivity in an ionic conductor cannot be absolutely zero. For example, the ionic (ai) and electronic conductivities (ae and

oh) of (ZrO2)o.92(Y203)0.08 as a function of temperature (800 ~ to 1050~ and partial pressure of oxygen (0.21 to 10 17 atm or 0.21 x 105 to 10 .2 Pa) are given

by the following empirical equations [4.82]:

o~(~-lcm -1) = 1.63x102exp( -0.79 eV ) (Eq. 4.3) KT

o �9 (~-lcm-l) = 1.31 x 10 7 exp( -3.88 eV) p-1/4o, (Eq. 4.4) KT

84 Chapter 4

lOO

10-1

10-2

18.18 mol% Y01.5

33.33 mol% Y01.5 / 15 mol% CaO /

46.16 mol% Y01.5

10-3 _

_

10-4 0

66.67 mol% Y01.5

i I i I J 1 l l J I l 1 l

4 8 12 16 20 24 2~

-LOG OXYGEN PARTIAL PRESSURE, atm (1.01 x 10 s Pa)

Figure 4.10. Ionic conductivity of stabilized Zr02 at IO00~ as a function Of oxygen partial pressure [4.81]

o h (~-xcm-1) = 2.35 x 102 exp( - 1.67 e V ) p 114 ~:T o~ (Eq. 4.5)

where eV = 1.6021 • 10 -19 J. Under typical SOFC oxygen partial pressures (0.21 to 10 -17 arm), as shown by the equations above, the electronic (electron and hole) conductivities are negligible, compared to the ionic conductivity. At very low oxygen partial pressures, the electronic conductivity then becomes

significant, and the total conductivity starts to increase with decreasing oxygen

partial pressure (Figure 4.11) [4.83]. The oxygen partial pressure at which the electronic conductivity becomes significant is higher at higher temperatures, as

shown in Figure 4.11. In a highly reducing atmosphere, the increase in conductivity occurs mainly in the bulk material. The grain-boundary conductivity

varies very little with atmosphere [4.84].

(v) Influence of grain boundary: In SOFCs, YSZ is used in the form of polycrystalline thin films or layers. The conductivity of polycrystalline YSZ is known to consist of the bulk conductivity and the grain-boundary conductivity.

The grain-boundary conductivity of stabilized Z r O 2 is mainly due to the presence of impurities or second phases introduced via the raw materials or during

Electrolyte 85

| 700~ 800~ 900~ ~ :

~ --- o , / / /

700o C -5

-7 (-.) / / \ x

-9

- 1 1 ! / / , !, '/ : \~ ~-""-- t ! .

0 - 10 -20 -30 -40 -50

LOG OXYGEN PARTIAL PRESSURE, atm (1.01 x 10 s Pa)

Figure 4.11. Conductivities of (ZrO2)o.9(Y203)o. 1 as a function of oxygen partial pressure [4.831

fabrication processes. (In general, grain boundaries are considered as highly disordered regions with possible differences in chemical composition compared to the interior of the grains.) Figure 4.12 shows, as an example, the grain- boundary and bulk conductivities of YSZ prepared from two commercial powders

[4.851. Compounds such as SiO2 and AI203 are commonly present as impurities

in commercial YSZ powders or added to starting powders as sintering aids. SiO2 and A1203 tend to form thin imergranular glassy phases (composed of SiO2, A1203, and Y203) in YSZ ceramics at high temperatures [4.86-4.88]. These glassy phases tend to be segregated in large isolated pockets along the boundaries [4.89,4.90] or to migrate to the external surface of the ceramics [4.91]. As a result, intergrain (grain-boundary) conductivity is quite different from the intragrain (bulk) conductivity.

The grain-boundary conductivity is strongly influenced by several factors such as grain size and impurity level. For small grains (< 2 to 4/~m), grain- boundary conductivity is independent of the grain size and is 100 times lower than that of the bulk. For large grains (> 2 to 4 ~tm), grain-boundary conductivity decreases with increasing grain size [4.92-4.94]. The grain- boundary conductivity decreases with increased impurity level [4.95,4.96]. In general, for high-density, pure polycrystalline materials, the contribution from

86 Chapter 4

r,,.) "7 c >: I.-- > I-- C3 :3 a Z 0 fO

TEMPERATURE, ~

900 800 700 J 6 0 0 500 4 0 0

10-1

1 0 -2

10 .3

10 .4

i i

~,, , GRAIN BOUNDARY CONDUCTIVITY

N ~ o ~, ------ POWDER 1 >r162 . . . . . . POWDER 2 ~,

, o . . , c o . o u c . , v , . Y .

;-, COMM .C,. owo ., -'COMMERCIAL POWDER 2 \ \ \ ,

1 0 5 �9 ' ' , , , ,

9 10 11 12 13 14 15

lIT, 104K "1

Figure 4.12. Bulk and grain-boundary conductivities of YSZ samples (prepared from two commercial powders) [4.85]

the grain-boundary conductivity is relatively small. Also, the influence of grain boundaries is typically considerable at low and intermediate temperatures (e.g., < 700~ At higher temperatures, grain boundaries have a lesser influence on the conductivity, especially when the sample is prepared from powders without the presence of significant amounts of impurities or additives.

(v) Influence of time (aging): Stabilized ZrO2 can exhibit a time dependence of its conductivity [4.97-4.99]. This conductivity aging effect has

been shown to be especially important in polycrystalline ZrO2 [4.100].

Typically, the bulk conductivity of fully stabilized ZrO2 reaches a steady value in a relatively short time (as shown by measurements of single crystals) [4.80]. This relatively rapid attainment of steady state has been attributed to the reorganization of the crystal [4.100]. The grain-boundary conductivity, on the other hand, tends to change with time. This change mainly involves gradual enhancement of the blocking effect as a result of the segregation of oversaturated impurities at the boundaries [4.100]. For partially stabilized ZrO2, the conductivity aging effect is considered to be due to the precipitation of tetragonal phase from the cubic matrix [4.101].

Electrolyte 87

The time dependence of the resistivity of polycrystalline YSZ has been measured and is given by the following general equation:

P(O = A - B x e x p ( - K l t ) - B2exp(-K2t) (Eq. 4.6)

where o is the resistivity, t the time, and A, B~, K~, B2, and K2 positive constants with the subscripts 1 and 2 designating relations to the bulk and grain-boundary

resistances, respectively [4.102]. In SOFC applications where the fully stabilized

YSZ electrolyte is typically prepared from pure materials, the conductivity aging effect is often insignificant.

4.2.5 Chemical interaction

At the fuel cell operating temperature (600 ~ to 1000~ stabilized ZrO2 exhibits little or no chemical interaction with other components (LaMnO3 cathode, nickel/YSZ cermet anode, and LaCrO3 interconnect). Chemical interaction of stabilized ZrO2, especially with LaMnO3, becomes more important at higher temperatures (see chemical interaction sections of Chapters 5, 6, and

7 for more details). ZrO2 reacts with LaMnO3 to form insulating phases such as

La2Zr207 at the interface at temperatures above 1100 oC [4.103-4.105]. These insulating phases are undesirable in SOFCs and must be minimized because they cause cell performance to degrade significantly.

4.2.6 Thermal expansion

The thermal expansion coefficient of undoped ZrO 2 single crystals is about 8.12 x 10 -6 cm/cm-K in the temperature range of 20 ~ to 1180~ [4.106].

D o p e d Z r O 2 materials typically have higher thermal expansion coefficients. For example, the thermal expansion coefficient of Z r O 2 crystals doped with 4 wt% CaO is about 10.08 x 10 -6 cm/cm.K [4.106]. Table 4.4 lists thermal expansion

data for various YzO3-stabilized Z r O 2 materials [4.107]. In general, the thermal expansion of partially stabilized Z r O 2 is very similar to that of fully stabilized material and is essentially unaffected by the presence of tetragonal precipitates

[4.106]. The thermal expansion coefficients of Z r O 2 (doped with different Y203

concentrations) at various temperatures are given in Figure 4.13 (from a Tosoh

zirconia powder technical bulletin).

T A B L E 4 .4

Dopant content

Thermal Expansion Data for YzO3-Stabilized ZrO2 Materials

Temperature (oc)

Thermal expansion coefficient (10 .6 cm/cm. K)

Ref.

==

Single-crystal

5 wt% Y203 8 wt% Y203 12 wt% Y203 20 wt% Y203

Polycrystalline

3 mol% Y203 6 mol % Y203

7.5 mol% Y203 8 mol% Y203 9 mol% Y203

20- 1500 10.99 [4.106] 20- 1500 10.92 [4.106] 20- 1500 10.23 [4.106] 20- 1500 11.08 [4.106]

1000 10..~ 1000 10.2

25- 1000 10.0 100- 1000 10.8

960 9.8

[4.108] [4.108] [4.109] [4.110] [4.111]

11 E (3

=, 10 o

9

z 8 U.I

U 7 u

U .

uJ 6 O U Z 5 O N 4 z < 3 G .

x m 2

<

W

-r 0 k-

88 Chapter 4

�9 2 m o l % Y 2 0 3

A 3 " "

0 4 " "

�9 6 . . . .

�9 1 | i i 1 i J 1 i i

2o0 ,00 600 8o0 000 200 ,00 600

TEMPERATURE, ~

Figure 4.13. Thermal expansion coefficients of YSZ at different temperatures (from Tosoh Zirconia Powder Technical Bulletin)

Electrolyte 89

As discussed in Chapter 10, a small difference in the thermal expansion coefficients of cell components in the SOFC can produce large stresses during fabrication and operation and cause cracking in the ceramic structure. Therefore, matching thermal expansion in the cell components is critical. Figure 4.14 is an example of thermal expansion curves for several stabilized ZrO2 and perovskite oxide materials to illustrate that without material tailoring or modification,

significant thermal expansion mismatch can exist [4.112]. In SOFCs, the YSZ electrolyte is typically selected as the baseline material, and the thermal

expansion of other cell materials is modified to match that of the electrolyte. The thermal expansion coefficient of perovskite cathode and interconnect may be adjusted by tailoring the dopant element and concentration (see thermal expansion sections of Chapters 5 and 7). The thermal expansion coefficient of nickel

cermet anode may be tailored by modifying the nickel content and ZrO2 concentration or by using additives (see thermal expansion section of Chapter 6).

1.0

0.5

0.0 0

1 !

o g , , .

- - . b

/ /

a. (ZrO~lo.91(Sc~O3)o.04slYb~O3)o.04s b. (ZrO2)o.91(Y203)o.09 c. Lao.TCao.3 MnOz.~ d. Lao.TSro.3MnO ~-6 e. Yo.2Sro..FeO3_ ~ f. Lao.3Sro.TFeO3. J g. Ndo.eSro.4CoO3.~ h. Lao.4Sro.sCoO3.~

. . . . . L___ . . . . ! 500 1000

TEMPERATURE, ~

Figure 4.14. Thermal expansion curves for several stabilized ZrO 2 and perovskite oxide materials [4.112]

90 Chapter 4

4.2.7 Mechanical properties

At room temperature, YSZ (8 mol% Y203) typically has a bending strength of about 300 to 400 MPa and a fracture toughness of about 3 MN-rff/2. The mechanical properties of a YSZ electrolyte layer obviously vary, depending

on the characteristics of starting powders used in the fabrication (such as particle size, particle size distribution, and agglomerate strength) and fabrication route

and fabrication conditions [4.113-4.117]. For example, YSZ powders having

strong and large agglomerates (up to 100/zm in diameter) are difficult to break down and tend to cause defects in the prepared electrolyte, resulting in poor

strength of the component [4.114]. YSZ sheets produced by tape calendering have superior mechanical properties, with a mean strength about 15 % higher than

that of the material made by tape casting [4.117]. At present, few data have been reported on the mechanical properties of YSZ at elevated temperatures.

One measurement indicates a mean strength of about 280 MPa for the YSZ at 900~ as compared to the value of 368 MPa at room temperature [4.117]. A bending strength of about 225 MPa has also been reported for YSZ at 1000~

[4.1181. Using YSZ having high strength and toughness as the electrolyte in

SOFCs is desirable. A strong and tough electrolyte is less sensitive to the presence of flaws and imparts better fracture resistance to the fuel cell during fabrication and operation. Several approaches have been taken to improve the mechanical properties of the YSZ. These approaches are based on improvement of fracture resistance of YSZ by toughening. The YSZ has been toughened by introducing inclusions of monoclinic ZrO: [4.119]. However, the addition of

monoclinic ZrO2 has been found to cause the conductivity of YSZ to degrade to a level unacceptable for SOFC application. Fine particles of partially stabilized

ZrO 2 [4.119], A1203 [4.118,4.120-4.126], and MgO [4.127] have also been

added to YSZ. These additives increase the fracture toughness and strength of the YSZ without deleteriously affecting the electrical conductivity of the material.

For example, the composition of YSZ with 30 wt% partially stabilized ZrO2 is attractive as the SOFC electrolyte. The fracture toughness of this composition

is about 2.95 MPa-rri/2 (200% higher than that of YSZ fabricated under similar

process conditions), and the ionic conductivity is about 0.15 fl-~cm -~ (17 % lower

than that of YSZ) [4.118]. The composition of YSZ with 20 wt % A1203 has

been recommended as the SOFC electrolyte. This composition has a bending strength of 33 kgf/mm 2 or 323 MPa (compared with 24 kgf/mm 2 or 235 MPa for

Electrolyte 91

YSZ) and an ionic conductivity of about O. 10 f}-lcm-1 at 1000~ (compared with 0.12 ~-~cm -~ for YSZ) [4.121]. Figure 4.15 shows an example of the variation of the bending strength of YSZ with temperature and A1203 content [4.118].

Because of high toughness and strength, tetragonal ZrO2 (TZP) has been

proposed as a SOFC electrolyte material [4.128-4.131]. For comparison, the toughness of YSZ ranges between 1 and 3 MPa.m--1/2, whereas that of TZP

ranges between 6 and 9 MPa.m la At temperatures below 600~ the electrical

conductivity of TZP is greater than that of YSZ (Figure 4.16) [4.128,4.129]. This suggests the possibility of its use as electrolyte in SOFCs operated at low

temperatures. However, there are two main concerns regarding SOFCs based on TZP electrolyte: mechanical integrity and aging effect. In general, TZP

electrolytes can suffer mechanical degradation after long-term exposure to fuel cell operating temperatures. This degradation results from the growth and subsequent transformation of the metastable tetragonal phase [4.132]. This

transformation proceeds as a function of time, and the contribution of dislocations and internal stresses is believed to be significant to the tetragonal-to-m0noclinic

phase transformation. TZP also shows pronounced conductivity aging [4.133,

4.134]. The aging behavior in TZP is a result, in part, of enhancement of the

grain-boundary blocking effect. The segregation of impurities to the grain

Figure 4.15.

4 0

I.- (9 z 30 IL l n "

(/3 n

Z m

Z ' ~ 2 0

~ E Z ~

Q.

'"' 1 0

"r I--

25oc . . ~

t... . . . . . . . . / - - " ~ 500~

1 1 1 1 0 10 20 30

AI20 ~ CONTENT, wt%

Variation of the bend strength of YSZ with temperature and Al20s content [4.]]8]

92 Chapter 4

v

o

i.c

o .,.J

1 0 3

1 0 2

101

1 0 0

10-~ 1 . 0

i i [ f I ! i ; i

-I

YSZ

boundary \

I ! . l ] , t [ J 1 _ 1 . 2 1 . 4 1 . 6 1 . 8

1/7", 1 0 3 K 1

Figure 4.16. Bulk and grain-boundary conductivities of TZP (bulk conductivity of YSZ included for comparison) [4.128]

boundaries during aging is one of the mechanisms responsible for this effect.

Furthermore, TZP has been found to have poor phase stability when undergoing thermal cycling between room temperature and 800~ This phase instability results in an increase in monoclinic phase in bulk TZP and at the grain boundary and, accordingly, causes a decrease in both bulk and grain-boundary conductivi-

ties [4.135]. Incorporation of CeO2 into TZP appears to reduce the conductivity

aging effect [4.136-4.138].

4.3 DOPED CERIA

Pure C e O 2 has the cubic fluorite structure up to its melting point.

Therefore, CeO2, unlike ZrO2, does not need any stabilization. Depending on temperature and oxygen partial pressure, the material exhibits a large oxygen

deficiency with the formula CeO2_~, where t5 may be as large as 0.3. For small oxygen deficiencies ((5 < 10-3), doubly ionized oxygen vacancies are the principal

ionic defects with compensating electrons [4.139]. For large oxygen deficien-

cies, a transition toward singly ionized vacancies has been observed. Pure CeO2

has n-type electrical conduction, and the conduction takes place by a small-

polaron hopping mechanism. Pure CeO2 has negligible ionic conductivity.

The radius of the Ce 4§ ion in CeO2 is sufficiently large that a variety of

dopant can be incorporated to form solid solutions. CeO2 doped with a divalent

Electrolyte 93

or trivalent oxide shows relatively high oxygen-ion conductivity at elevated temperatures. Doped CeO2, however, has the tendency to undergo reduction (Ce 4+ ion to Ce 3+ ion) at low oxygen partial pressures (with the consequent

introduction of electronic defects). This reduction restricts the range of oxygen partial pressure over which the ionic transference number remains close to unity.

Thus, at 800~ the oxygen partial pressure range for predominant ionic conduction in CeO2 is limited to about 10 -~2 atm (10 -7 Pa) [4.56]. Figure 4.17 shows, as an example, the oxygen-ion transference number for La203-doped

CeO2 as a function of temperature and dopant content [4.140]. In spite of the tendency to be reduced in a reducing environment, doped

CeO2 has been considered for use as SOFC electrolyte, especially at reduced temperatures, because of its high ionic conductivities. Compared to stabilized ZrO2, doped CeO2 shows a higher conductivity and a lower conduction activation

energy. Doped CeO2 has also been considered as an electrolyte for SOFCs operated directly on methanol.

Various dopants have been used with CeO2, including La203 [4.140, 4.141], Y203 [4.79,4.142-4.148], Sm203 [4.149,4.150], Gd203 [4.148,4.151- 4.159], other rare-earth oxides [4.160], Gd203+Pr203 [4.161], CaO [4.145- 4.147, 4.162, 4.163], and SrO [4.164]. The Arrhenius plots for the ionic conductivities of several CeO2 solid solutions are shown in Figure 4.18 [4.165]. Table 4.5 lists conductivity data for various doped CeO2 materials.

1.0 CELL . , ~ ,,'5, o

- x . x

0.9

z 0.8 F,,

0.7

, / 6oooc _~ 0.6 / ~ 1 7 6

X 0.5 " t~ 1000~

0 . 4 1 , 1 1 1 .1 1 _1 . 0.2 0.3 0.4 0.5

CONTENT OF LaOl.s, mol

Figure 4.17. Oxygen-ion transference number of La203-doped Ce02 as a function of temperature and dopant content [4.140]

94 Chapter 4

T, E 0

"7 cI I

b

0 ._I

-I

3

2 -

, . =

3 -

0.~

TEMPERATURE, ~ 900 800 700 600 500

I i i i - i

0 ~70 ~4o

e e

I I e

0 (CeOJo.e(SmO~.6)o.2 A (CeOJo.a(GdO~.s)o.2 ? (CeOJo.a(YO~.5)o.2 [ ] (CeOJo.e(CaO)o.2 �9 Ce02 �9 (ZrO2)o.es(YO ~.s)o.~ 5

O El

o

O l

�9 3 7 I I

i 0 1 . I , ! l 1 0.8 .9 1.0 1.1 1.2 1.3 1.4

1/T, 10"3K "1

Figure 4.18. Arrhenius plots of ionic conductivities of doped Ce02 compounds [4.165]

TABLE 4.5

Ionic Conductivity of Doped CeO2

Dopant Content Ionic conductivity at 800~ Activation energy Ref. (mol%) (10 .2 0lcm -1) (kJ/mol)

La203 10 2.0 - [4.140] Y203 20 5.5 26 [4.160] Gd203 20 8.3 44 [4.160] Sm203 20 11.7 49 [4.160] CaO 10 3.5 88 [4.142] SrO 10 5.0 77 [4.164]

Electrolyte 95

For the same dopant concentration, the ionic conductivity of M203-doped CeO 3 (where M is a rare-earth ion) increases with increasing ionic radius (in the order of yttrium, ytterbium, holmium, dysprosium, gadolinium, samarium) to a

maximum at a radius of 0.109 nm and then decreases with further increase in ionic radius (in the order of samarium, neodymium, lamhanum) [4.160]. This trend has been attributed to the binding energy between dopant and host cations

[4.160, 4.166]. The electrical conductivity of doped CeO2 increases at low oxygen partial pressures (Figure 4.19). This conductivity increase is due to the presence of n-type electronic conduction under those conditions. For SOFC applications, it is important to know the electrolytic and ionic domain boundaries of the electrolyte. The main factors defining the electrolytic and ionic boundaries of doped CeO2 are dopant and temperature. An example of the electrolytic and

ionic boundaries as a function of temperature for (CeO2)0.95(Y203)0.05 and (CeO2)0.8(Gd203)0.2 is shown in Figure 4.20 [4.153].

The reduction, thus the electronic conductivity, of doped CeO2 under reducing atmospheres can be minimized by modifying the dopant. A dopant concept has been attempted to improve the ionic domain of CeO2 [4.154]. Certain dopant modifications have been found to improve the ionic boundary to 10 -2~ atm (10 -16 Pa) at 700~ For example, the replacement of 3 mol% gadolinium by

praseodymium in Ce0.8Gdo.202_ ~ (to form Ce0.8Gdo.lvPr0.0302_a) improves the electrolytic domain of the material by nearly two orders of magnitude without significantly affecting the ionic conductivity [4.161].

|1 , , l i

10 -0"5 !- O (CeO2)o.e(SmOl.s)0.2 ! A (CeO2)~176

~___ 10 "1"0 _

(J

Z

~ t i0-I.5

l I 1 I 1

0 5 10 15 20

-LOG OXYGEN PARTIAL PRESSURE, atm (1.01 x 105Pa)

Figure 4.19. Conductivity of doped CeO 2 as a function of oxygen partial pressure [4.160]

t~ I1.

m o I-.-

TEMPERATURE, ~ 12001000 800 600

i "1 I ! I I

i - - o.

z

(.9 0 ..J

0

-10

-20

-30

96 Chapter 4

450 400 i i l

0"5 t. I a: (Ce02)o ~s(Y203)o.os I b: (CeOz)o i , (Gd203) o.,

- 4 0 ! J I , I ! ! I _! 0.6 0.7 0.8 0.9 1.0 1.1 1.2 1.3 1.4 1

1/7", lO'aK "1

Figure 4.20. Electrolytic and ionic domain boundaries of doped CeO 2 [4.153]

Another approach to suppress the reduction of CeO 2 under reducing

atmospheres is to coat CeO2 solid solutions with a film of a more stable ionic

conducting compound. For example, (CeO2)0.8(Sm203)0.2 has been coated with an Y203-stabilized ZrO2 thin layer (e.g., 2/xm thick) on the fuel side to produce a stable SOFC electrolyte [4.160,4.165,4.167-4.170]. In this case, the stability of CeO2 is dependent on the oxygen partial pressure at the ZrO2/CeO2 interface [4.171]. The oxygen partial pressure at the interface, in turn, depends on the transport characteristics, especially the electronic conductivity of CeO2. High electronic conductivity (in addition to high ionic conductivity) of the CeO2 side ensures high oxygen partial pressure at the interface, and, thus, high thermody-

namic stability.

4.4 STABILIZED BISMUTHSESQUIOXIDE

Pure Bi203 has two crystallographic polymorphs: t5 phase with a fluorite

cubic structure above 730~ and ct phase with a monoclinic structure below 730~ di-li203 (having 25 % of the anion sites vacant) exhibits high oxygen-ion conductivity (about 1 fl-lcm-I near the melting point of 825 ~ ot-Bi203 exhibits

p-type conduction. The (5 phase of Bi20 a can be stabilized at low temperatures

by doping with a certain metal oxide [4.172-4.178].

Electrolyte 97

Among the known oxygen-ion conductors, stabilized Bi203 shows the

highest ionic conductivity at comparable temperatures. For example, in certain composition ranges, the conductivities of stabilized Bi203 exceed 10 -I ~-lcm-I at 700~ and 10 .2 f/-~cm -I at 500~ one to two orders of magnitude higher than

that of stabilized ZrO2. This greater ionic conductivity of stabilized Bi203 offers

the possibility of its use as electrolyte in SOFCs operated at reduced temperatures. The main drawback of this material is its small oxygen partial pressure

range of ionic conduction. Stabilized Bi203 is easily reduced under low oxygen partial pressures and decomposes into bismuth metal at an oxygen partial pressure of about 10 -13 atm (10 .8 Pa) at 600~ [4.179]. Therefore, practical use of

stabilized Bi203 as SOFC electrolyte requires protection of the material from

direct exposure to reducing atmospheres [4.180]. Stabilized Bi203 has a face-centered cubic (fcc) or rhombohedral single

phase or both face-centered cubic and rhombohedral phases, depending on the ionic radius of the dopant cation. The fcc phase is obtained when Bi203 is doped

with a trivalent, pentavalent, or hexavalent metal oxide, and the ionic radius of the dopant cation is relatively small. The fcc phase is stable over a wide

temperature range and shows no phase transformation up to 800~ The rhombohedral phase is obtained when Bi203 is doped with a certain divalent or trivalent metal oxide, and the ionic radius of the dopant cation is relatively large. In this case, at high temperatures, the stabilized material consists of a high-

conductivity rhombohedral phase,/31. This rhombohedral phase is transformed into another rhombohedral phase,/32, at low temperatures. The rhombohedral

phase/32 also exhibits high conductivity, although lower than that of/3~. Both fcc and rhombohedral phases are obtained (depending on composition) when Bi203

is doped with a metal oxide having a metal cation of medium-size ionic radius. Several stabilized Bi203 systems are briefly discussed below.

Bi203-MO systems (where M is calcium, strontium, or barium) have been studied. Bi203-MO solid solutions tend to form rhombohedral phase that

shows high oxygen-ion conductivity. Figure 4.21 shows the conductivity of

Bi203 doped with CaO, SrO, and BaO as a function of temperature [4:172]. These plots exhibit a sharp jump in the conductivity at 600 ~ to 650~ This

jump has been attributed to the 131"~132 transition within the rhombohedral

structure. Among the various stabilized Bi203 materials, the BaO-dopedBi203

has one of the highest conductivities. For example, at 600~ the conductivity

of (BiEO3)0.a4(BaO) 0.16 is about 0.88 f/-lcm-I [4.181].

98 Chapter 4

T E M P E R A T U R E , ~ 8 0 0 6 0 0 4 0 0

1 '1 1 "I �9 BaO

1 0 ~ x SrO �9 CaO

? ~ 10 "!

~ 1 0 . 2 -

0

10"3 --

1 0 .4 I 1 1 0 . 8 1 . 0

I I 1 1 ! 1

1.2 1.4 1.6 1.8

lIT, 10"3K "1

Figure 4.21. Conductivity of Bi203 doped with CaO, SrO, and BaO [4.172]

Bi203-M203 systems (where M is a rare-earth metal) have been most

extensively studied. In these systems, the range of fcc and rhombohedral phase formation depends on the ionic radius of the rare-earth cation and the rare-earth oxide content (Figure 4.22) [4.182]. Figure 4.23 shows the conductivity of

BiEOa-Y203 as a function of dopant concentration and temperature to illustrate the conductivity behavior of BiEOa-M203 systems [4.183]. It can be seen from Figure 4.23 that solid solutions containing more than 25 mol % Y203 show no jump in conductivity over the whole range of temperature examined. The solid solutions having less than 25 mole % Y203, on the other hand, show abrupt

increase in conductivity (due to phase transition), and exhibit thermal hysteresis

in the Arrhenius plot (as represented by dotted lines in the figure). The

minimum dopant content without jump in conductivity corresponds to the lowest

content of added oxides for forming the solid solution in the low-temperature

region. Detailed studies on the electronic conductivity of (Bi203)0.73(Y203)0.27 indicate that the lower limit of oxygen partial pressure under which the ionic transference numbers are higher than 0.99 is about 10 -13 atm (10 -8 Pa) at 600~ Figure 4.24 shows the oxygen partial pressure dependence of the hole and

electron conductivities of Bi203-M203, together with the oxygen-ion conductivity

at 500 ~ 600 ~ and 700~ [4.179]. Table 4.6 summarizes the conductivity of

various Bi203-M203 materials.

Electrolyte 9 9

1 , 2 i �9 �9 i - , , , ,

E

, t A i

0 . 0

La

'Nd

Sm

Gd

Yb

0.5 1.0

x IN (gi203)l.x(M203)x

Figure 4.22. Range of face-centered cubic and rhombohedral phase formation of Bi203-M203 systems [4.182]

TEMPERATURE, ~

800 700 600 500 4()0 i . . 1 I I i

" : 1 o tool% Y2O3 100 .~ 2 5 mol%

:: "~ 2 3 20 mol% . 0 . . i 4 2 5 m o l %

: "*'-,t ~ 5 33 mol% 10 "1 6 42.5 tool%

.~3~ 7 50 mol% - 8 60 mol%

5 1 0 - 2 ! 6 ~,

3 4 -

g ~ 10.3

0.9 1.0 1.1 1.2 1.3 1.4 1.5 1/T, 10-3K 1

Figure 4.23. Conductivity of Bi203-Y203 as a function of Y203 concentration and temperature [4.183]

100 Chapter 4

, , �9 = . . , ,

at ' 7oo~ ~, 600oc E

2 5~176176

1 1"6

-LOG OXYGEN PARTIAL PRESSURE, atm (1.01 • 10 s Pa)

Figure 4.24. Dependence of conductivities of Bi203-M203 on oxygen partial pressure [4.179]

TABLE 4.6

Conductivity Data for Bi203-M203

Conductivity Dopant Dopant content (10 .2 fllcm-1) Ref.

(mol % M203) 500~ 700~

Dy203 28.5 0.71 14.4 [4.184] ErzO 3 20 0.23 37.0 [4.184] Y203 20 0.80 50.0 [4.183] Gd203 14 0.11 12.0 [4.185] Nd203 10 0.30 85.0 [4.182] La203 15 0.20 75.0 [4.182]

Electrolyte 101

Other Bi203-based systems of interest include Bi203-MO2 (M is tellurium;

e.g., Bi203-TeO2 [4.186]), Bi203-MO3 (M is tungsten or molybdenum; e.g., Bi203-WO3 [4.187]), Bi203-M205 (M is vanadium, niobium, or tantalum; e.g., Bi203-V205 [4.172,4.188,4.189]), and Bi203-M60~ (M is praseodymium; e.g., Bi203-Pr6Ol~ [4.186]). Bi203 has also been stabilized by two or more oxides.

Some examples are Bi203-Y203-Nb205 [4.190], Bi203-Ln2Oa-TeO2 (Ln is lanthanum, samarium, gadolinium, or erbium) [4.191], and Bi4V2_xMxOll_~ (M is a divalent to pentavalem cation) (or the so-called BIMEVOX) systems [4.192]. The BIMEVOX systems being investigated include BICUVOX (copper doped), BINIVOX (nickel doped), BIMOVOX (molybdenum doped), BIWVOX (tungsten doped), BIPBVOX (lead doped), BIZNVOV (zinc doped), and BICOVOX (cobalt

doped) [4.193-4.196].

4.5 OTHER OXYGEN-ION CONDUCTORS

Pyrochlore oxides (of the general formula A2B207 with a defect fluorite superstructure) have been proposed as SOFC electrolytes. The key features of these materials are their relatively high intrinsic oxygen-ion conductivity (thus no doping is required) and their equilibrium pyrochlore phase (thus no long-term aging effects are expected). Work on pyrochlores has been focused on the (A3+)2(B4+)207 system, especially the Ln2Zr207 where Ln is a lanthanoid such as

Gd2(ZrxTil_x)207 (GZT) and Y2(ZrxTil_x)207 (YZT) [4.197-4.202]. Of these pyrochlore compounds, GZT has been suggested as the electrolyte material for SOFCs. When x < 0.2, the compound shows mixed ionic and electronic conduction. However, when x > 0.2, the electronic component of the conductivity decreases markedly. At x > 0.4, the oxygen-ion conductivity dominates and exceeds 10 .3.5 91cm -~ at 800~ Increases in ionic conductivity

with increasing x in GZT are related to increasing levels of oxygen disorder. The ionic conductivity of the titanate end member, Gd2Ti207, can be increased further by doping, for example, with an acceptor such as calcium on the A site.

In general, higher dopant level and closer host/dopant size match lead to high ionic conductivities (with low activation energies) and low electronic conductivities. The GZT system shows a broad ionic domain, e.g., P~ is about 10 -21 atm

(10 -16 Pa) for Gd2(Zro.6Tio.4)207 at 1000~ The ionic domain of the material falls rapidly with increasing temperature.

Perovskite oxides of the form ABO 3 (A is a divalent or trivalent cation and B is a tetravalent or trivalent cation) have been considered as SOFC

102 Chapter 4

electrolytes, especially for reduced-temperature applications. When substituted with an appropriate dopant, certain perovskite oxides exhibit high ionic conductivity [4.203]. Typically, doped perovskites are pure oxygen-ion conductors at low oxygen partial pressures but become mixed conductors

(acquiring p-type electronic conductivity) at high oxygen partial pressures. For

example, BaCeO3 doped with 10 mol% gadolinium shows an oxygen-ion conductivity of about 1.1 x 10 .2 ~-lcm-I and an electronic conductivity of about 5 x 10 .3 f~-lcm-1 in oxygen at 600~ [4.204,4.205]. An empirical approach has

been developed to identify perovskite ionic conductors [4.206-4.209]. The approach is based on the relationship between the activation energy for oxygen- ion conduction and the free volume of lattice structures. The observed trend between the activation and the free volume has been used to develop a free volume maximization algorithm which permits the identification of cations in the perovskite A and B sites to maximize the free volume for ionic transport through the lattice. As an example, some of the perovskites identified are BaCe0.gGdo.~O3,

CaAlo.7Tio.303, and SrZr0.9Sc0.103. In ionic-conducting perovskites, the migrating ion may be 02. or H +, depending on the nature of the perovskite and several

factors such as water vapor and temperature. For example, the conducting species in DyA103 and BaCeO3 are considered as oxygen ions, whereas in

KTaO3, CaHfO3, BaZrO3, and SrCeO3, they are considered as protons. In BaCe0.9Gdo.103_~, the ionic transport mechanism varies with temperature and atmosphere. The conducting species change from protons at temperatures below 600~ to oxygen ions above 900~ under fuel cell operating conditions. Some BaCeO3-based oxides also show mixed oxygen-ion and protonic conduction under SOFC conditions [4.210-4.213].

4.6 PROTONIC CONDUCTORS

Some perovskite-type materials have been known to exhibit protonic

conduction in hydrogen and/or wet atmospheres at high temperatures [4.214- 4.216]. For example, when A in ABO3 is substituted with an appropriate lower- valency cation, oxygen-ion vacancies are formed to maintain the electroneutrality

condition in the crystal. During processing at high temperatures, oxygen-ion vacancies react with oxygen to produce electron holes as shown by the following

equation:

10 x. 2 2+ V ~ 1 7 6 (Eq. 4.7)

Electrolyte 103

In this case, the material shows r~ixed oxygen-ion and electron hole conduction.

However, in a hydrogen atmosphere, electron holes react with hydrogen to

produce protons

-~I~ + h ' = H § (Eq. 4.8)

Or, in a wet atmosphere, the following reaction takes place

H20 § 2 h ' = !O + 2 + 2 H (Eq. 4.9)

o r

•

t-120 + V o = O o + 2H + (Eq. 4.10)

o r

x

1-12o § Oo + V o - 2(OH)o (Eq. 4.11)

Protons produced in this case (Eqs. 4.9 and 4.10) are considered to be

interstitials [4.215,4.217]. On the other hand, hydroxide ions produced (Eq.

4.11) are considered to migrate between sites adjacent to oxygen ions [4.218] or

migrate via vacancies [4.204, 4.207]. Hydroxide-ion conduction is considered to

occur when water is present in the environment. In doped BaCeO3_~, the activation energy for conduction (typically below 0.52 eV or 50 kJ/mol) is lower

than that for the motion of OH-via vacancies (about 0.78 eV or 75 kJ/mol)

[4. 218, 4.219], indicating migration of protons as the conduction mechanism. Several proton-conducting perovskites have been studied for use as SOFC

electrolytes. The most widely studied high-temperature protonic conductors are

the solid solutions based on BaCeO3 and SrCeO 3 [4.220, 4.221]. (i) Doped BaCe03: Doped BaCeO3 has been shown to possess significant

proton conduction at high temperatures. For example, BaCe0.85Ca0.~503_~ exhibits

a protonic conductivity of about 2 • 10 -3 ~-lcm-1 at 9 0 0 ~ with an activation

energy of 0.54 eV (52 kJ/mol) [4.222]. Figure 4.25 shows the structure of

BaCe0.9Gd0.102.95 as determined by powder neutron diffraction [4.215,4.223]. The structure is orthorhombic with the lattice constants of a = 0.8773 nm, b =

0.6244 nm, and c = 0.6222 nm. Though this structure does not contain any

proton and the chemical formula does not include hydrogen, protonic conduction

in the material is possible [4.215]. This is because the compound incorporates

water vapor from the atmosphere during preparation and subsequent cooling, and

so acquires a concentration of mobile protons. However, under typical SOFC

104 Chapter 4

0 - - -

(....,

0 0 2-

cea+(Gd 3"~)

Ba2. +

a

Figure 4.25. Structure of BaCeo.9Gdo.102.95 [4.215]

conditions, doped BaCeO3 may show mixed proton and oxygen-ion conduction. According to Eqs. 4.7 and 4.10, the p-type electronic conductivity of

doped BaCeO3 depends on oxygen partial pressure and the protonic conductivity on water vapor pressure as shown by the following equations:

(Eq. 4.12)

(Eq. 4.13)

where K7 and gl0 are the equilibrium constants of Eqs. 4.7 and 4.10, respectively. Figure 4.26 shows the conductivity of BaCe0.gGdo.~O2.95 as a function of oxygen partial pressure at two water vapor pressures [4.215]. According to Eq. 4.13, the protonic conductivity should increase by a factor of 10, since the water partial pressure is increased by a factor of 100. The independence of the conductivity of water vapor partial pressure (as shown in Figure 4.26) suggests that protonic conduction is a minor component in this temperature range. Conductivity measurements for Y203-doped BaCeO 3 treated in H20- and D20- saturated gas indicate that the conductivity ratio, tr(H+)/o(D§ is greater than the classical value of ~ and the difference in activation energy is about 0.05 eV (4.8 kJ/mol). Thus, nonclassical proton hopping may occur in doped BaCeO3. The carrier concentration (obtained from analysis of the conductivity preexponential for proton-saturated samples) is nearly an order of magnitude lower than the hydrogen concentration (determined from isotopic exchange), suggesting that a large fraction of protons are in bound sites [4.224].

Electrolyte 105

WATER PARTIAL PRESSURE

2 x 10 .4 atm (OPEN SYMBOLS) 0.25 SYMBOLS)

0.20

~- 0.15 >

o D 0.10 7 a

o o 0 . 0 5 ~ =

0.00 ,. i i i i 25 20 15 10

TEMPERATURE

o 1 2 0 0 " C v 1 1 0 0 * C / / o 1 0 0 0 * C ~ / -, 90OoC / o 8 0 0 0 C

I i l

5 0

-LOG OXYGEN PARTIAL PRESSURE, atm(1.01 x 10 s Pa)

Figure 4.26. Conductivity of BaCeo.9Gdo.102.95 as a function of oxygen partial pressure and water partial pressure [4.215]

The temperature dependence of the conductivity of BaCel_xMxO3_~ in hydrogen gas is shown in Figure 4.27 [4.225]. As discussed earlier, doped

BaCeO3 cam be a proton and oxygen-ion mixed conductors. Some examples are

BaCe0.gNdo.lO3_~ [4.226], BaCe0.8Dy0.203_~ [4.212], and BaCel_xSmxO3_ ~ (x = 0.05 to 0.15) [4.227]. In order to reduce the oxygen-ion contribution to the conductivity, partial substitution of barium with calcium has been suggested. This type of substitution lowers the degree of crystal symmetry and restrains the

motion of the oxygen ion [4.228]. Gadolinium-doped BaCeO3 has also been found to exhibit primarily proton conduction. The Gd 3+ ions have been suggested as substitutes not only for C e 4+ but also for Ba 2+ [4.229-4.331].

(ii) Doped SrCe03: Doped SrCeO3 is another common material being considered as a protonic conductor for SOFC applications. In hydrogen

atmospheres, doped SrCeO 3 exhibits significant protonic conductivity. Examples of the conductivities of several doped SrCeO 3 compounds in hydrogen are given

in Figure 4.28 [4.232]. The conductivities of SrCel_xMxO3_ ~ have been

determined in a hydrogen atmosphere where M is scandium, zinc, manganese,

yttrium, indium, neodymium, samarium, dysprosium, or ytterbium, and x =

0.05 to 0.10 [4.232, 4.233]. The compound SrCe0.95Yb0.0503_ ~ has a protonic conductivity of 2 x 10 -3 i ] - lcm -l in NE/5 % H2 at 600~ with an activation energy

of 0.59 eV (57 kJ/mol) [4.234]. The conductivity increases proportionally to the

106 Chapter 4

10 -1

u "7

>.: 1 0 .2 t- > I-

E3 z O

1 0 .3

TEMPERATURE, ~ 1000 900 800 700 600

I I I I I

DOPANT [] Yz03 (x = 0.1) O Nd203 (x = 0.1) A LazOz (x = 0.1) �9 Nd203 (x = 0.05) V' CaO (x = 0.05)

"', ~. . . .~

% .

- _ _ SrCeo.~gbo.~03.6 ~ -

O' ' i i 7 .8 0.9 1.0 1.1 1.2 l IT, 103K "1

Figure 4.27. Conductivities of BaCel_~lx03. ~ in a hydrogen atmosphere [4.225]

TEMPERATURE, ~

1000 900 800 700 600 i ~ 1 i 1 i j

~-~ 1 SrCeossYboosO3.6 1 0-2 ~---'~ 2 SrCeo9oYo loO3-6

~'E I ~ ~ ~ o ~ 3 3 SrCe~176176 1 ~ 0

g ~ 10-3

0.8 0.9 1.0 1.1 1.2 1/T, 103K "1

Figure 4.28. Conductivities of doped SrCe03.~ in hydrogen [4.232]

Electrolyte 107

square root of the water vapor pressure (in contrast to BaCeO3) and is indepen-

dent of the oxygen partial pressure [4.235]. Secondary ion mass spectrometry

(SIMS) measurements for SrCe0.95Ybo.05034 indicate protons as the probable

conducting species for material annealed both in dry air and in wet hydrogen gas

at 800~ [4.236]. Furthermore, the protonic conductivity of SrCeOa-based

oxides has been found to increase with increasing content of oxygen-ion vacancy,

but it is almost independent of the dopant concentration [4.237]. The activation

energy of proton conduction in ytterbium-doped SrCeO 3 is about 0.62 to 0.63 eV

(60 to 61 kJ/mol) and independent of dopant coment [4.238-4.240]. The proton

concentration in SrCeo.95Ybo.osO34 has been determined to be about 2 mol% at

600~ and 1 mol % at 1000~ in water vapor [4.241, 4.242]. These results

appear to support the application of Eq. 4.13 for doped SrCeO3.

In addition to doped BaCeO3 and SrCeO3, several other oxides have been

investigated as proton-conducting electrolytes for use in SOFCs. Examples of

these materials include doped zirconates (CaZrO3, BaZrO3, and SrZrO3)

[4.243,4.244], BaTho.9Gdo.~O34 [4.208,4.245], Ba2Gdlnl_xGaxO 3 [4.246], and

doped KTaO3 [4.247]. To date, work on protonic-conductor SOFCs has been

limited to material studies, clarification of the conduction mechanism, and testing

of small, laboratory-scale single cells.

References

4.1

4.2

4.3

4.4 4.5

4.6

4.7

N.Q. Minh, in Science and Technology of Zirconia V, S.P.S. Badwal, M.J. Bannister, and R.H.J. Hannink (eds.), Technomic Publishing Company, Lancaster, PA, 1993, p. 652. M. Ciftcioglu and M.J. Mayo, paper presented at The Spring Meeting of the Materials Research Society, April 16-21, 1990, San Francisco, CA, Report No. DE90-015700, SAND-90-1639C, CONF-900466-78, National Technical Information Service, Alexandria, VA, 1990. M.A. Janney, C.L. Calhoun, and H.D. Kimrey, J. Am. Ceram. Soc., 75 (1992) 341. S.A. Nightingale, R.H.J. Hannink, and S. Street, see Ref. 4.1, p. 299. C.C. McPheeters and T.D. Claar, in 1986 Fuel Cell Seminar Abstracts, October 26- 29, 1986, Tucson, AZ, Courtesy Associates, Washington, DC, 1986, p. 64. S.V. Phillips and A.K. Datta, in Proceedings of the Institute of Energy Conference on Ceramics in Energy Applications, April 1990, Sheffield, U.K., Hilger, Bristol, U.K., 1990, p. 183. N.Q. Minh, C.R. Horne, F.S. Liu, D.M. Moffatt, P.R. Staszak, T.L. Stillwagon, and J.J. Van Ackeren, in Proceedings of the 25th IECEC, August 12-17, 1990,

108 Chapter 4

4.8

4.9

4.10

4.11

4.12 4.13

4.14 4.15

4.16 4.17

4.18

4.19

4.20

4.21

4.22

4.23 4.24

4.25

Reno, NV, Vol. 3, American Institute of Chemical Engineers, New York, 1990, p.

230. N.Q. Minh, C.R. Horne, F. Liu, P.R. Staszak, T.L. Stillwagon, and J.J. Van Ackeren, in Proceedings of the First International Symposium on Solid Oxide Fuel Cells, October 16-18, 1989, Hollywood, FL, S.C. Singhal (ed.), Electrochemical Society, Pennington, NJ, 1989, p. 307.

N.Q. Minh and C.R. Horne, in Proceedings of the 14th Riso International Symposium on Materials Science, High Ten~erature Electrochemical Behaviour of Fast Ion and Mixed Conductors, September 6-10, 1993, Roskilde, Denmark, F.W. Poulsen, J.J. Bentzen, T. Jacobsen, E. Skou, and M.J.L. OstergArd (eds.), Riso National Laboratory, Roskilde, Denmark, 1993, p. 337. A.O. Isenberg, Solid State lonics, 3/4 (1981) 431. Y. Ohno, S. Nagata, and H. Sato, in Proceedings of the 15th IECEC, August 18-22, 1980, Seattle, WA, American Institute of Aeronautics and Astronautics, New York, 1980, p. 881. M.F. Carolan and J.N. Michaels, Solid State Ionics, 25 (1987) 207. N. Nakagawa, C. Kuroda and M. Ishida, in Proceedings of the International Sympo- sium on Solid Oxide Fuel Cells, November 13-14, 1989, Nagoya, Japan, O. Yamamoto, M. Dokiya, and H. Tagawa (eds.), Science House, Tokyo, Japan, 1989,

p. 58. H. Arai, see Ref. 4.13, p. 12. T. Setoguchi, M. Sawano, E. Eguchi, and H. Arai, Solid State Ionics, 40/41 (1990) 502. Z. Ogumi, Y. Tsuji, Y. Uchimoto, and Z. Takehara, see Ref. 4.13, p. 203. N. Nakagawa, S. Kosuge, H. Tsuneizumi, E. Matsuda, H. Mihara and Y. Sato, see Ref. 4.8, p. 71. K. Tsukamoto, F. Uchiyama, Y. Kaga, Y. Ohno, T. Yanagisawa, A. Monma, Y. Takahagi, M. J. Lain, and T. Nakajima, Solid State Ionics, 40/41 (1990) 1003. N. Nicoloso, B. Leibold, and H.U. Habermeier, in Laser Ablation of Electronic Materials, E. Fogarassy and S. Lazare (eds.), Elsevier, Amsterdam, The Nether- lands, 1992, p. 385. T.W. Kueper, S.J. Visco, and L.C. DeJonghe, Solid State lonics, 52 (1992) 251.

E.-T. Kim, J.W. Lee and S.-G. Yoon, J. Electrochem. Soc., 140 (1993) 2625.

B.L. Halpern, J.J. Schmitt, J.W. Golz, and Y. Di, in Proceedings of the Fourth

Annual Fuel Cell Contractors Review Meeting, July 14-15, 1992, Morgantown, WV, W.J. Huber (ed.), Report No. DOE/METC-92/6127, U.S. Department of Energy, Washington, DC, 1992, p. 102. T. Ishihara, K. Sato, Y. Mizuhara, and Y. Takita, Chem. Lett., (1992)943. T. Namikawa, Y. Yamazaki, I. Saitoh, T. Kanai, S. Sumiya, and M. Satoh, Denki Kagaku, 52 (1984) 714. H. Hamatani, T. Okada, and T. Yoshida, see Ref. 4.13, p. 197.

Electrolyte 109

4.26

4.27

4.28 4.29

4.30

4.31

4.32

4.33

4.34

4.35

4.36 4.37 4.38 4.39 4.40 4.41 4.42

4.43 4.44

4.45 4.46

4.47

4.48

4.49

4.50

4.51

R.T. Brook, in Science and Technology of Zirconia, A.H. Heuer and L.W. Hobbs (eds.), American Ceramic Society, Columbus, OH, 1981, p. 272. M.A.C.G. van de Graaf and A.J. Burggraaf, in Science and Technology of Zirconia

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Columbus, OH, 1984, p. 744.

C.E. Scott and J.S. Reed, Am. Ceram. Soc. Bull., 57 (1978) 741.

C.E. Scott and J.S. Reed, Am. Ceram. Soc. Bull., 58 (1979) 587.

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116 Chapter 4

4.223 4.224

4.225

4.226

4.227

4.228 4.229 4.230

4.231

4.232

4.233 4.234 4.235 4.236 4.237

4.238 4.239 4.240 4.241 4.242

4.243

4.244

4.245

4.246

4.247

K.S. Knight, M. Soar, and N. Bonanos, J. Mater. Chem., 2 (1992) 709. T. Scherben, Yu.M. Baikov, and E.K. Shalkova, Solid State Ionics, 66 (1993) 159.

H. Iwahara, H. Uchida, K. Ono, and K. Ogaki, J. Electrochem. Soc., 135 (1988)

529. H. Iwahara, H. Uchida, and K. Morimoto, J. Electrochem. Soc., 137 (1990)462.

H. Iwahara, T. Yajima, T. Hibino, and H. Uchida, J. Electrochem. Soc., 140

(1993) 1687. T. Yajima, H. Iwahara, and H. Uchida, Solid State Ionics, 47 (1991) 117. K.D. Kreuer, E. Sch6nherr, and J. Maier, see Ref. 4.9, p. 297. H. Iwahara, T. Yajima, T. Hibino, K. Ozaki, and H. Suzuki, Solid State lonics, 61 (1993) 65. J.F. Liu and A.S. Nowick, in MRS Syn~osium Proceedings, Solid State Ionics, Vol.

210, G.A. Nazri, D.F. Shriver, R.A. Huggins, and M. Balkanski (eds.), Material Research Society, Pittsburgh, PA, 1991, p. 675. H. Iwahara, T. Esaka, H. Uchida, and N. Maeda, Solid State Ionics, 3/4 (1981)

359. H. Iwahara, Solid State Ionics, 28-30 (1988) 573. N. Bonanos, B. Ellis, and M.N. Mahmood, Solid State Ionics, 28-30 (1988) 579. H. Uchida, N. Maeda, and H. Iwahara, Solid State lonics, 11 (1983) 117. T. Ohgi, T. Namikawa, and Y. Yamazaki, see Ref. 4.9, p. 357. T. Yajima, H. Iwahara, H. Uchida, and K. Koide, Solid State Ionics, 40/41 (1990)

914. T. Yajima and H. Iwahara, Solid State Ionics, 50 (1992) 281. T. Yajima and H. Iwahara, Solid State lonics, 53-56 (1992) 983. T. Scherban and A.S. Nowick, Solid State Ionics, 35 (1989) 189. H. Uchida, H. Yoshikawa, and H. Iwahara, Solid State Ionics, 35 (1989) 229. H. Uchida, H. Yoshikawa, T. Esaka, S. Ohtsu, and H. Iwahara, Solid State Ionics,

36 (1989) 89. T. Yajima, H. Kazeoka, T. Yogo, and H. Iwahara, Solid State Ionics, 47 (1991) 271.

H.H. Huang, M. Ishigame, and S. Shin, Solid State Ionics, 47 (1991) 251. R.L. Cook, J.J. Osborne, J.H. White, R.C. MacDuff, and A.F. Sammells, J.

Electrochem. Soc., 139 (1992) L19. M. Schwartz, B.F. Link, and A.F. Sammells, J. Electrochem. Soc., 140 (1993)

L62.

T. Scherdan, S.Q. Fu, and A.S. Nowick, see Ref. 4.231, p. 663.

Chapter 5

CATHODE

5.1 REQUIREMENTS

The main function of the cathode is to provide reaction sites for the

electrochemical reduction of the oxidant. Thus, the cathode material must be

stable in the oxidant oxidizing environment and have sufficient electronic

conductivity and catalytic activity for the oxidant gas reaction at the operating

conditions. Since the SOFC operates at high temperatures (600 ~ to 1000~ the

cathode must be chemically and thermally compatible with the other cell

components, from room temperature to those operating temperatures and to the

even higher temperatures at which the fuel cell is fabricated. The key require-

ments for the cathode in the SOFC are discussed below. This discussion is

qualitative because the specific requirements depend on selected materials and

cell and stack designs [5.1]. (i) Stability: The cathode must be chemically, morphologically, and

dimensionally stable in the oxidant environment. The cathode material must have

no disruptive phase transformation (involving large changes in molar volume)

between room temperature and fabrication temperature. The cathode must

maintain its desired microstructure in long-term operation; significant microstruc-

tural changes can cause degradation in cell performance.

(ii) Conductivity: The cathode must possess sufficient electronic

conductivity to support electron flow in the oxidizing environment at the

operating temperature. In general, maximum possible cathode conductivity is

desirable to minimize ohmic losses.

(iii) Compatibility: The cathode must be chemically compatible with other

components, not only at the operating temperature, but also at the much higher

temperature at which the fuel cell ceramic structure is fabricated. Chemical

interaction or elemental interdiffusion between the cathode and adjoining

components must be limited in order to minimize unacceptable occurrences such

118 Chapter 5

as second phase formation, stability reduction, change in thermal expansion,

introduction of electronic conductivity in the electrolyte, etc.

(iv) Thermal expansion: The thermal expansion of the cathode must match (from room temperature to operation and fabrication temperatures) that of

other cell components to avoid cracking and delamination during fabrication and

operation, including thermal cycling. (v) Porosity: The cathode must have sufficient porosity to allow gas

transport to the reaction sites. The lower limit on porosity is set by mass transport considerations. (The porosity limit may be less critical for mixed conducting materials.) The upper limit is based on consideration of mechanical

strength of the component. (vi) Catalytic activity: The cathode must have sufficient catalytic activity,

thus low polarization, for the electrochemical reduction of the oxidant. In addition to these requirements, other desirable properties for the SOFC

cathode are high strength and toughness, fabricability, and low cost. Because of the high operating temperature of the SOFC (600 ~ to

1000~ only noble metals and electronic or mixed conducting oxides can be used as cathode materials. Noble metals such as platinum and palladium are

unsuitable for practical applications because of prohibitive cost. Many doped oxides are available, but only a few meet the requirements of thermal expansion match and compatibility with the electrolyte. At present, doped lanthanum

manganite (LaMnO3) is most commonly used.

5.2 LANTHANUM MANGANITE

Doped LaMnO3 has been extensively used as cathode material in the SOFC. This selection has been based primarily on three factors: high electrical

conductivity in oxidizing atmospheres, adequate compatibility with Y 2 0 3 -

stabilized ZrO2 (YSZ) electrolyte, and acceptable thermal expansion match with

other cell components. Other properties of LaMnO3 are tailored to meet the

requirements for SOFC applications.

5.2.1 Preparation

Various methods have been used for the fabrication of LaMnO3 cathodes (for references, see fabrication sections of Chapter 9). In the sealless tubular

design, LaMnO3 cathode layers are deposited on the porous support by ' slurry

Cathode 119

coating. In the segmented-cell-in-series design (banded configuration), the cathode is made by flame spraying ceramic powders onto the surface of the electrolyte. In the monolithic design, the cathode layer is formed by tape calendering (with the other components) and then cofired at elevated temperatures. In the flat-plate design, the cathode is applied on the electrolyte by various

techniques such as casting, printing, spraying, and depositing. Figure 5.1 shows an example of the microstructure of LaMnO3 cathode made by tape calendering.

In general, the processing conditions of each fabrication method are tailored to produce the required cathode structure. Each method is also designed such that no conditions in any preparation step destroy desired characteristics of the LaMnO3 material. Recently, several new techniques have been investigated for the fabrication of the cathode, especially as thin films. For example, thin LaMnO3 layers have been formed on YSZ by electroless deposition from a

solution containing manganese ions, lanthanum ions, and an oxidizing agent [5.2]. A pyrosol derived method has been applied to produce strontium-doped

LaMnO3 films 5 to 10 ~m thick for SOFC cathode applications [5.3]. LaMnO3

layers have also been prepared by techniques such as oxidation plating [5.4], electrochemical deposition [5.5], vacuum evaporation [5.6], and film coating

[5.71. Many fabrication methods for LaMnO3 use fine powders as starting

materials. LaMnO3 powders are commonly prepared by the amorphous citrate (Pechini or liquid mix) process [5.8]. The process involves dissolving appropriate proportions of salts of desired metal elements in a citric acid/ethylene

Figure 5.1. Microstructure of LaMnO~ cathode made by tape calendering

120 Chapter 5

glycol solution. The solvent is evaporated, and a transparent amorphous resin

containing a homogeneous mixture of the metals is formed. The resin is then decomposed by heating to remove the organics and produce the desired oxide.

The amorphous citrate or liquid mix process involves two basic chemical reactions: (i) complexation or chelation between metal ions and citric acid, and (ii) polyesterification of complexes with ethylene glycol. The complexation and polyesterification reactions preserve the homogeneity of the metal salt solution in a gel. As a result, powders produced have high surface area and precise compositional stoichiometry. Polymerization and decomposition conditions of the

amorphous citrate process are often tailored or modified to achieve the desired

characteristics of the oxide powder product [5.9-5.12]. The process has also been modified by changing the citric acid/ethylene glycol ratio of the preparation

solution [5.9, 5.11, 5.13, 5.14]. In addition to the Pechini process, several other methods have been

developed for the synthesis of LaMnO3 powders. The freeze drying method

produces LaMnO3 with a very high surface area (14 to 32 m2/g) by spraying a suitable nitrate solution into a cryogenic medium [5.15]. The spray pyrolysis process prepares fine oxide powders by thermally decomposing aqueous nitrate

solution droplets in a hot reaction chamber [5.16-5.19]. The sol-gel process

produces strontium-doped LaMnO3 powder (particle size of 0.3 to 0.7/zm and surface area of 17.5 to 23.5 m2/g) using polyacrylic acid to make a gel [5.20]. The glycine/nitrate combustion method prepares extremely fine LaMnO3 materials by evaporating and autoigniting an aqueous metal nitrate/glycine

solution [5.21].

5.2.2 General properties, phase transformation, and stoichiometry

LaMnO3 belongs to the class of perovskite oxides of the general formula

ABO 3. The cubic perovskite structure comprises three interpenetrating three-

dimensional networks" two consisting of the separate A and B cations and the

third of corner-sharing 06 octahedra. Thus, the ideal structure of LaMnO3 can

be described as a MnO6/2 framework of corner-shared octahedra that contains lanthanum cations within 12-coordinate sites (Figure 5.2). The cubic structure

of LaMnO3 may undergo atomic distortion leading to orthorhombic or rhombohe-

dral unit cells. LaMnO3 melts at about 1880~ [5.22]. The phase diagram of

the La203-Mn203 system is shown in Figure 5.3 [5.22]. Table 5.1 lists the

properties of LaMnO3.

Cathode 121

Q L a

Mn06

Figure 5.2. Meal perovskite structure of LaMnO 3

v

tt.

I - <I: t r u..i 13.

UJ

I---

2 6 0 0

2 2 0 0

1 8 0 0

1 4 0 0

o

1000 0 .0

La203

" P - perovskite-type LaMn03 phase 1 - liquid o - observed

La203", - - - estimated

, - - ........ liquid liquid , - ' ' P + l '. ' " ' . ,- 11 x x

p ", 4- x x x

, liquid ', - - " - f + ,~" M,,,_q,

P + c-Mn304

/ p ~ P + t Mn304

o L__i _~..1]~___~ ~ Mn~ 03~

La203 + P

o

_ _ . l . . . . . . t . . . . . . . . . .

0.2 0 .4 0 .6 0 .8 1.0

Mn20s

Mn/ (La + Mn)

Figure 5.3. Phase diagram of a Mn203-La203 system [5.22]

122 Chapter 5

TABLE 5.1

Properties of LaMnO 3

Melting point, ~ Density, g/cm 3 Thermal conductivity, W/cm.K Thermal expansion coefficient, 10 .6 cm/cm.K

(25 o to 1100 ~ C) Standard enthalpy of reaction, kJ/mol

(from LazO3(s), MnO(~), and O2(g), 1064 to 1308 K) Standard entropy change, J/mol.K

(from La203(~), MnO(s), and O2(g), 1064 to 1308 K) Strength at 25 ~ (30 % porosity), MPa

1880 [5.22] 6.57 [5.221 0.04 [est.] 11.2 [5.231

-168 [5.24]

-65 [5.24]

25 [est.]

Undoped stoichiometric LaMnO3 is orthorhombic at room temperature

[5.25,5.26] and shows an orthorhombic/rhombohedral crystallographic

transformation at about 600~ [5.26]. This transformation has been attributed

to the oxidation of some Mn 3§ to Mn 4§ ions [5.26]. Thus, the orthorhombic/

rhombohedral transition temperature is dependent on the Mn 4§ content, and

therefore, very sensitive to the stoichiometry of the material, especially that of

oxygen. For example, various transformation temperatures (285~ [5.27], 300~ [5.28], 387~ [5.29], 450~ [5.25]) have been reported for LaMnO3

samples having different degrees of oxygen nonstoichiometry. At high Mn 4§

content or high oxygen-excess level (i.e., LaMnO3+~ with t5 > 0.1), the

compound is rhombohedral at room temperature [5.30-5.32]. Doping such as

substituting a lower-valence cation for lanthanum and manganese sites increases

the Mn 4§ concentration in LaMnO3, thus affecting the transformation tempera-

ture. For example, strontium doping and calcium doping change the structure

of LaMnO3 from orthorhombic to rhombohedral at room temperature [5.31,5.33, 5.341.

LaMnO 3 can have oxygen excess, stoichiometry, or deficiency depending

on the preparation conditions such as firing atmosphere, temperature, and time.

For example, LaMnO3.~5 (30% Mn 4§ concentration) is obtained when heated at

1100~ in oxygen for 6 days [5.26]. On the other hand, LaMnO:.99 is obtained

when quenched from 1300~ in air [5.31]. At high temperatures, the oxygen

stoichiometry of LaMnO3 varies as a function of oxygen partial pressure and

Cathode 123

temperature. For example, at 1200~ the oxygen stoichiometry of LaMnO3 ranges from 3.079 to 2.947 under oxygen partial pressures of 1 to 10 -116~ atm

(10 5 tO 10 -6.6o Pa) [5.35]. In oxidizing atmospheres, LaMnO3 has oxygen excess

and the amount of excess oxygen depends on temperature. In reducing

atmospheres, the material becomes oxygen deficient. Under highly reducing

conditions, LaMnO3 dissociates into La203 and MnO; however, the dissociation

is reversible [5.36]. For doped LaMnO3, the level of oxygen excess-decreases

with increasing dopant content [5.36,5.37]. Thus, LaMnO3 containing 20 mol %

calcium dopant has no oxygen excess even at high oxygen activities [5.38].

Figures 5.4, 5.5, and 5.6 show how the oxygen content of undoped and

strontium-doped LaMnO3 varies as a function of oxygen partial pressure,

temperature, and dopant concentration [5.36]. For SOFC applications,

significant changes in oxygen stoichiometry of LaMnO3 (especially during

fabrication) must be avoided to minimize undesired dimensional changes [5.39].

In addition to oxygen nonstoichiometry, LaMnO3 can also exhibit

lanthanum deficiency or excess. LaMnO3 with lanthanum excess may contain

La203 second phase, which tends to be hydrated to La(OH)3. This hydration is

undesirable for SOFC applications because it can cause disintegration of the

sintered LaMnO3 structure. LaMnO3 can have up to about 10% lanthanum

deficiency without second phase formation [5.31]. Above this level, Mn304 is

present as the second phase. In general, it is very difficult to prepare stoichio-

metric LaMnO3; the undoped material synthesized under normal conditions often

possesses about 0.2 mol % lanthanum vacancies [5.36].

3.10

3.08

3.06 Z

I.-tUZO 3'04 1 0 3.02

3.00 >" I X

0 2.98 i "

-18

zx 1000oc �9 1100~ a 1200~ y y / . . ~ .

=,.. = .p , . / - , ,= r'l

/2

' l ' l

-16 -14 -12 -110 ]8 -6 -4 ]2 0 LOG OXYGEN PARTIAL PRESSURE, atm (1.01 X 105 Pa)

Figure 5.4. Oxygen content of undoped LaMnO~ as a function of oxygen partial pressure and temperature [5.36]

124 Chapter 5

3.10

"6 E , 3.00 - k-

Z I l l F- Z 0 0 Z tu 2.90 (9 >- • 0

2.80 . -18

. . . . . . . . . . . _j; .= .=~A-.,=- ~ , ~ = = f

f ~.=r .- o

/ / o-

/= o a I O 0 0 o c o / /~a ~ �9 1100~

. o 1 2 0 0 ~

D

I I "1 I -16 -114 -12 -10 -8 -6 -4 :2 0

LOG OXYGEN PARTIAL PRESSURE, atm (1.01 X 106 Pa)

Figure 5.5. Oxygen content of Lao9SroiMnO 3 as a function of oxygen partial pressure and temperature [5.36]

3.10 ,

~. 3 . 0 0 -

Z

Z 0 2 . 9 0 - 0 Z ILl (9 >- X 2 . 8 0 - 0

2.70 . ! 8

_ , + A - _ |i, IiIll~-II~I''~ - -

~=~ , + , � 9

j L @ ,v A LaMn0a

�9 Lao.agSro.ol MnOa

o Lao.e6Sro.o~MnOa

�9 Lao.ooSro.+oMnOa

�9 o Lao eoSro.2oMn0a

-;6 -1'4 -1'2 -;o ;8 .... -'6 -~, -~ LOG OXYGEN PARTIAL PRESSURE, atm (1.01 x 106 Pa)

Figure 5.6. Oxygen content of Lal.xSrxMnO s as a function of oxygen partial pressure at 1000~ [5.36]

Cathode 125

The defect chemistry of undoped and doped LaMnO3 has been studied [5.29, 5.36, 5.38, 5.40-5.46]. A defect model based on unassociated and randomly distributed point defects (metal and oxygen vacancies) has been proposed. In this conventional defect structure model, the formation of metal vacancies in the high oxygen activity region (the oxide has oxygen excess) can be expressed by the

following equation (using Kr6ger-Vink notation):

6MnMn +2230= CGa + CMn + / / // 6Mndn + 30 o (Eq. 5.1)

Assuming equal concentration of vacant lanthanum and manganese sites, the

concentration of metal vacancies is thus proportional to the 3/4 root of the oxygen

partial pressure (Region I). As the oxygen activity decreases, the oxide becomes stoichiometric (Region II). At low oxygen activities, the oxide is oxygen deficient, and oxygen vacancies are the predominant point defects. The

formation of oxygen vacancies can be written as

10 2Mnr~n + Oo = 2Mnr~n + Vo + ~- 2 (Eq. 5.2)

Under these conditions, the concentration of oxygen vacancies is inversely proportional to the 1/2 root of oxygen partial pressure (Region III). At

sufficiently low oxygen activities, the concentration of oxygen vacancies is constant; thus, the oxygen nonstoichiometry is independent of the oxygen partial

pressure. Under these conditions, the concentration of electron holes is proportional to the 1/4 root of oxygen partial pressure (Region IV). Figure 5.7 summarizes the variation of the oxygen stoichiometry of LaMnO3 as a function of oxygen partial pressure according to the proposed model.

Figure 5. 7. Oxygen content as a function of oxygen activity as predicted from defect model

126 Chapter 5

In general, LaMnO3 exhibits a defect chemistry very close to that shown

in Figure 5.7 for high oxygen activity regions (Region I, II, and III) (see Figures 5.4, 5.5 and 5.6). Other experimental results support the basis of the model:

lanthanum and manganese vacancies are randomly distributed in equal amounts in oxygen-excess LaMnO3 [5.47,5.48]. However, the material behavior deviates significantly from the model predictions for low oxygen activity regions. (No region IV is observed.) This deviation has been explained by a change in the manganese valency such as the thermally excited disproportionation of Mn 3§ into

Mn 2§ and Mn 4§ ions [5.36,5.41-5.43,5.48]. Other models based on defect

clusters have been suggested to explain the behavior of oxygen-deficient LaMnO3

in the low oxygen activity regimes [5.41,5.45]. For example, association of

oxygen vacancies and dopant ions has been proposed [5.45].

5.2.3 Stability

At the SOFC operating temperature, LaMnO3 is stable in oxidizing atmospheres but decomposes under highly reducing conditions. The lowest

oxygen partial pressure before LaMnO3 dissociates into multiple phases is termed the critical oxygen partial pressure. At 1000~ the critical oxygen partial pressure for undoped LaMnO3 is about 10 -14 to 10 -~5 atm (10 .9 to 10 -1~ Pa)

[5.24,5.49-5.52]. The critical oxygen partial pressure for LaMnO3 depends on temperature, shifting to higher values at higher temperatures. At the same temperature, this critical pressure increases with increasing dopant concentration [5.36]. Thus, a high dopant content generally results in reduced stability for the LaMnO3 compound. Figure 5.8 shows, as an example, the critical oxygen partial

pressure of La~_xSrxMnO3_ ~ as a function of temperature and dopant content

[5.40]. At the SOFC operating temperature, LaMnO3 decomposes directly to

La203 and MnO at the critical oxygen partial pressure [5.35,5.49,5.53]. However, at lower temperatures (350 ~ to 600~ the material tends to transform

to other phases such as LazMnO4, La8Mn8023, LanMn4Oll [5.54,5.55]. Nonstoichiometry can influence the stability of LaMnO3. Lanthanum

excess tends to precipitate as La203, causing hydroxide formation and subsequent

disintegration of sintered LaMnO3 structures at room temperature [5.56]. Lanthanum-deficient LaMnO3 is more stable; however, the lanthanum deficiency

must be less than 10% to prevent Mn304 formation [5.31]. LaMnO3 with lanthanum deficiency is recommended for use in SOFCs.

Cathode 127

o

X

0 -5 v

E IO - 1 0 u~ fie ::=) (/) u) - 15 I l l iv , o. _1 _~ -2o I-- < O.

z -25 I l l

(9 >. X 0 - 3 0 (9 0 . . . I

T, ~

1 2 0 0 1 0 0 0 BOO i T ,

600

- t - : Lao.6Sro.4Mn03.6

- 0 - : Lao.eSro.2Mn03_ 6

- 0 - : Lao.oMn03.6 zx :, LaMn03_6

= i i i | l |

6 7 8 9 10 11 12 13

I/T, 10 .4 K -I

Figure 5.8. Critical oxygen partial pressure of LavxSr~l/ln03 [5.40]


LaMnO3 has intrinsic p-type conductivity due to the formation of cation vacancies. The material has an electrical conductivity of about 10 -4 f]-lcm-l at

room temperature and about 0.1 ~2-1cm -1 at 700~ The electrical conductivity

of LaMnO3 has been enhanced for SOFC applications by substituting a lower-

valence cation on either the A or B sites.

LaMnO3 has been substituted with various cations such as barium [5.57- 5.59], calcium [5.16,5.42,5.57,5.60-5.64], chromium [5.65-5.67], cobalt

[5.25,5.68-5.72], copper [5.73], lead [5.74], magnesium [5.75], nickel

[5.28,5.30], potassium [5.76,5.77], rubidium [5.77], sodium [5.77,5.78],

strontium [5.29,5.31,5.34,5.78-5.87], titanium [5.58], and yttrium [5.88]. Besides electrical conductivity, many of these substituents are used to modify

other properties (e.g., thermal expansion) of LaMnO3. Presently, strontium and

calcium dopants are most commonly used in SOFCs because LaMnO3 doped with

these cations has high electronic conductivity in oxidizing atmospheres and

relatively matches the thermal expansion of other cell components.

128 Chapter 5

Strontium doping enhances the electronic conductivity of LaMnO3 by increasing the Mn 4+ content by the substitution of La 3§ by Sr 2+"

x S r O x _ 3 + ~ 2+-= ,r 3+ ,= . 4+,--, LaMnO 3 __, La~_x~rx ~vm~_xMnx u 3 (Eq. 5.3)

The electronic conductivity of strontium-doped LaMnO3 takes place via the small

polaron conduction mechanism. Figure 5.9 shows the temperature dependence

of the conductivity of both undoped and strontium-doped LaMnO3 [5.29]. For undoped LaMnO3, the break in the conductivity plot seen in Figure 5.9 can be explained by the orthorhombic/rhombohedral crystallographic transition. At

temperatures below 1000 ~ plots of In aT (where tr is the conductivity and T is the temperature) vs 1/T are linear, as predicted by the following equation derived for the small polaron conduction:

tr = (AJT)exp(-E/KT) (Eq. 5.4)

where A~ is the preexponential factor, K the Boltzmann constant, and E~ the

activation energy for conduction. The Eo values calculated from the slopes of the plots are 18.3, 18.3, 15.4, and 8.7 kJ/mol for undoped, 5, 10, and 20-mol%-

strontium-doped LaMnO3. For LaMnO3 doped with < 20 mol% strontium

dopant, the conductivity of the material increases with increasing temperature and increasing strontium concentration at temperatures below 1000~ The conductivity becomes nearly constant at temperatures above 1000 oc, suggesting

a semiconducting-to-metallic transition [5.29,5.31]. The material may show the metallic conduction behavior at a lower transition temperature depending on the

1 4

"7, E l l

'7, C

~ 8 0 _d

~1~11~ I I - I I I . . ii1..._ B...... _ -

o LaMn03

A Lao.96Sro.o~Mn03 "o. 0 ~ " I Lao.9oSro.loMnO 3 - .

" Lao.eoSro.2oMn03 " - .

" 0

9 12 15 18 21 24

lIT, 10 .4 K ~

Figure 5.9. Electrical conductivity of undoped and strontium-doped LaMnO~ [5.29]

Cathode 129

lanthanum stoichiometry [5.82]. As the strontium content is > 20 to 30 mol %,

the metallic-type conduction is then observed for LaMnO3 over the entire

temperature range [5.31,5.84,5.86]. The variation of the electrical conductivity

of La~_xSrxMnO 3 with strontium content appears to exhibit a maximum. At the

SOFC operating temperatures (600 ~ to 1000~ the strontium level where the

maximum conductivity is observed is about 55 mol % [5.86]. Similar to strontium doping, calcium doping significantly enhances the

electronic conductivity of LaMnO3, and the conduction of the calcium-doped

compound also occurs via a small polaron hopping mechanism. At a calcium

content of < 50 mol%, the conductivity of La~_xCaxMnO3 increases with

increasing temperature and shows a metallic-type conduction transition at about

600~ to 700~ [5.16, 5.89]. Above 60 mol % calcium, the conduction behavior

of the material is entirely metallic (negative temperature dependence) [5.89]. At

the SOFC operating temperature, the maximum conductivity of La~_xCaxMnO 3

occurs at 50 mol % calcium [5.89]. Table 5.2 lists conductivity data for several doped LaMnO3 compounds.

It should be noted that the values given in Table 5.2 are very much dependent

on preparation conditions and material stoichiometries; other values have been

reported in the literature.

TABLE 5.2

Conductivity Data for Doped LaMnO3

Dopant Composition Conductivity, 1000 ~ C Activation energy (MO) (mol % MO) (fl-lcml) (kJ/mol)

Ref.

SrO 10 130 15.4 SrO 20 175 8.7 SrO 50 290 4.5 CaO 25 165 11.6 CaO 45 240 7.9 NiO 20 100 18.6

SrO,Cr203 10,20 25 13.5 SrO,Co203 20,20 150 -

[5.291 [5.291 I5 151 [5.75J [5.751 [5.281 [5.6o] [5.7Ol

130 Chapter 5

The electrical conductivity of undoped and doped LaMnO3 decreases significantly in the low oxygen partial pressure region. The dependence of electrical conductivity of LaMnO3 on oxygen partial pressure has been correlated to the material defect structure [5.29,5.38]. For example, the defect reaction for

La~_xSrxMnO3_ ~ may be represented by

1 o 2MnMn + O~ = 2MnMn + V~ + 7 2 (Eq. 5.5)

The equilibrium constant in terms of mole fraction is given as

(1 -x+2~)2(~)Po/~ K = (Eq. 5.6)

(x-2tS)2(3 -~i)

This equation can be approximated as

g (x-2tS) 2

(Eq. 5.7)

The electrical conductivity a of the material is given by

tr = e ~ p (Eq. 5.8)

where e is the electron charge, /z the mobility, and p the carrier concentration (p = x - 26). By solving Eq. 5.7 and combining it with Eq. 5.8, the conductivity becomes

I "1' D 1/2 [ [ ~ ~. r , , " D - 1 / 2

a = e 4 K , o, twx~xro, + 1) 1/2 - 1] (Eq. 5.9)

In the high oxygen activity region, x approaches zero; Eq. 5.9 then reduces to Eq. 5.8. Thus, the conductivity is independent of oxygen partial pressure.

However, in the low oxygen activity region, Eq. 5.9 becomes

_ ,, X \ 1 / 2 r ~ 1 / 4 = e g ~ - ~ - g ) ro~ (Eq. 5.10)

The computed conductivity for strontium-doped LaMnO3 as a function of oxygen

activity shows excellent agreement with the experimental data (Figure 5.10).

Cathode 131

2.4

2.2 o

2.0

1.8 I- L) z~ 1.6 z 0 (9 1.4 (9 0 1.2

1.0 .

~~~ , 1000~ P/ = 1100~

/ / * 1200~

// . . 6 - 14 - ; 2 - 1 0 -'8 -6 -4 :2 0

LOG OXYGEN PARTIAL PRESSURE, atm (1.01 x 106 Pa)

Figure 5.10. Conductivity as a function of oxygen partial pressure for LaosSro~/lnO 3 at various temperatures [5.29]

Similar results are also obtained for undoped LaMnO3 (Figure 5.11). As seen from Figures 5.10 and 5.11, the electrical conductivity of both undoped and strontium-doped LaMnO3 shows little dependence on oxygen partial pressure in the range of high oxygen activity, and this range narrows with increasing temperature. As the oxygen activity decreases, the conductivity decreases according to Eq. 5.10 and then drops abruptly (due to dissociation of LaMnO3 phase) at highly reducing conditions (critical oxygen partial pressures).

2.0

u O 1.9

m > 1.8 o

z O o 1.7 (9 o

1.6

~ \ - .

~//~= = / ~ a 1000~ = 1 1 0 0 ~ o 1 2 0 0 o c

-16 -14 -12 -10 -8 -6 -4 -2 0 LOG OXYGEN PARTIAL PRESSURE, atm (1.01 x 105 Pa)

Figure 5.11. Conductivity as a function of oxygen partial pressure for undoped LaMnO 3 at various temperatures [5.29]

132 Chapter 5


The high fabrication and operation temperatures of the SOFC raise

concerns regarding the chemical compatibility of LaMnO3 cathode with other cell

components, especially the YSZ electrolyte. Manganese is known to be a mobile

species at high temperatures and can easily diffuse into the electrolyte, changing the electrical characteristics or the structure of both the cathode and the

electrolyte [5.90]. Fabrication temperature is generally limited to below 1400~

to minimize this migration. At temperatures below 1000~ manganese

migration is negligible [5.91,5.92], and there is no report on any significant manganese effects for SOFCs operated up to 10,000 h.

The chemical interaction between LaMnO3 and YSZ has been studied at SOFC operating and fabrication temperatures. No significant reactions between

LaMnO3 and ZrO2 have been observed at temperatures below 1100 ~ to 1200~

above 1200~ the manganite reacts with ZrO2 to produce La2Zr207 [5.93-5.97]. Other reaction products (e.g., SrZrO3 and other zirconate phases) may also form

with LaMnO3 containing high dopant content [5.94, 5.97-5.101]. The interaction

is significant at high temperatures and long times. For example, a layer of

La2Zr207 up to 5 /~m thick can form at the LaMnO3/YSZ interface if fired at 1450~ for 48 h [5.97]. The formation of La2Zr207 and other phases at the LaMnO3 cathode/YSZ electrolyte interface is undesirable in SOFCs: these compounds can deleteriously affect the performance of the cell because they can both act as an insulating layer and create thermal stresses at the interface. The electrical conductivity of La2Zr207 and SrZrO3 is several orders of magnitude

lower than that of YSZ and the thermal expansion coefficients of the materials

are significantly lower than that of YSZ [5.102-5.106]. The LaMnO3/YSZ interaction appears to proceed via the unidirectional

diffusion of manganese, lanthanum, or dopant cation (e.g., strontium, calcium)

into the YSZ [5.107,5.108]. A proposed mechanism for the interaction suggests

high manganese diffusion into YSZ, resulting in lanthanum excess at the

interface, which then reacts with the YSZ to form zirconates [5.108]. The

presence of a dopant in LaMnO3 suppresses the manganese migration; thus,

substitution of lanthanum in the manganite with a low dopant concentration

reduces the La2Zr207 formation [5.93,5.108]. A high dopant content results in

the formation of other phases such as SrZrO3, CaZrO3. Figure 5.12 shows, as an example, how the dopant content influences the reactivity of LaMnO3 and the

reaction products formed at different temperatures [5.101]. The interaction

1 7 0 0

La2ZraO~ /

LaaZr20~ SrZr03

E 1 5 0 0 ,.d n r :::) I-- ,a: n r I.,U n

~ 1300

Cathode 133

1100 . . . . . . 0 .0 0.1 0 .2 0 .3 0 . 4 0 .5 0 .6 0 .7

S T R O N T I U M CONTENT, mol

Figure 5.12. Diagram showing reaction products formed from La~_xSr~/InO 3 and YSZ at different temperatures [5.101]

between LaMnO3 and YSZ is also influenced by the stoichiometry of the perovskite material. Lanthanum deficiency reduces the reactivity of LaMnO3 toward YSZ [5.93,5.94,5.109].

The interaction between LaMnO3 and YSZ has been examined by thermodynamic analysis of various equilibria in terms of chemical potentials of all elements involved [5.110-5.119]. Figure 5.13 is an example of a constructed triangle diagram based on such analysis. The diagram predicts the various processes taking place between strontium,calcium-doped LaMnO3 and YSZ at 1400~ in air. Chemical potential diagrams are also constructed. The chemical

potential diagram of the La-Mn-Zr-O system at 1300 ~ is shown in Figure 5.14. This diagram also takes into account the dissolution of lanthanum and manganese into YSZ and the lanthanum nonstoichiometry in LayMnO3_~ (since manganese is known to diffuse into the YSZ, and the lanthanum deficiency in LaMnO3 affects

its reaction with YSZ). The diagram suggests that the LazZrzO7 formation takes

place for y > 0.86, whereas the manganese dissolution may become significant

for y < 0.86. The LazZrzO7 formation gives rise to the lanthanum depletion in

the LaMnO3 phase, resulting in a decrease in lanthanum stoichiometry y. The manganese dissolution gives rise to the manganese depletion in the perovskite.

When LaMnO3 is doped with a lower-valence cation such as strontium, zirconates formed change with strontium content (from La2Zr207 to SrZrO3) (Figure 5.15).

La~_xSrxMnO 3 can be in equilibrium with La2Zr207 or SrZrO3 or both, depending on the particular composition.

134 Chapter 5

LoMnO 3

y ~ ~_~(Sr,C o) ZrOq

CoMnO 3 SrMnO 3

Figure 5.13. Diagram showing possible reactions between strontium, calcium-doped LaMnO 3 and YSZ at 1400 ~ C [5.115]

10

6 r~

' 4

(.9 0 -- 2

- 2

10

o 0 0 J - - - ~ - - + - ' -

I ~ ~ ~ - ~ - ~ I _ ~ .... II "-- i I i I I o ~ . / 41o Mn3u4 H -~- d c~do " l - J j ~ .., , .

12 14 16 18 20 22 L O G aMn/azr

Figure 5.14. Chemical potential diagram of a La-Mn-Zr-O system at 1300~ [5.117] (a is the activity)

Cathode 135

_J

V)

(.9 O _.J

15

lO

. . Z . . ^ / / : I I - J ~ . - - - ~ 1 / / i p"o12"" % , / I

i SrO / .'~," ,,.[-SrZrO 3 I .b

,, [ fo rmot ion!_ / / , �9 - - , / .O'b" I AIB stockingl ." ," .,~/~ e~

I ~ t ~I..'.'" JA .~'

. . . . . . . . . .

5 ; [L~ i ~_ JA-_Si!e ii '" "i / f~176 i ~ "

Lo203 (Lo, S")Mn% i M n30/. ,

5 10 15 20 LOG aMn/aLa

Figure 5.15. Chemical potential diagram of a La-Sr-Mn-Zr-O system at 1300~ [5.118]

In addition to reaction with YSZ, LaMnO3 may react with doped LaCrO3

when in contact with this interconnect material at high temperatures. Strontium and calcium dopants in LaCrO3 tend to migrate to the interface to form compounds such as (Sr,La)3Mn207 and (Ca,La)3Mn/O7 [5.117,5.120]. The interaction may also form solid solutions between LaMnO3 and LaCrO3 at the interface [5.121].


The thermal expansion coefficient of undoped LaMnO3 (La0.99MnO3) from 25~ to 1100~ is 11.2• x 10 -6 cm/cm.K [5.23]. Lanthanum deficiency

and oxygen nonstoichiometry appear to lower the thermal expansion of the material [5.31,5.122].

Strontium doping increases the thermal expansion coefficient of LaMnO3,

and the coefficient increases with increasing strontium dopant content [5.23]. Substitution for lanthanum with smaller cations, such as calcium or yttrium,

lowers the thermal expansion coefficient of LaMnO3 [5.16,5.88,5.89]. For

136 Chapter 5

example, the average thermal expansion coefficient of Lao.9Cao.lMnO 3 from room temperature to 1000~ is about 10.1 x 10 -6 cm/cm-K [5.16]. The coefficient

of calcium-doped LaMnO3 is higher at higher calcium concentration [5.16, 5.122] (10.1 to 11.4 x 10 -6 cm/cm.K for calcium content ranging from 10 to 50 mol%

[5.16]). Forming solid solutions of LaMnO 3 with LaCoO3 and LaCrO3 leads to higher thermal expansion coefficients [5.67,5.69,5. 70,5.123]. Figure 5.16 shows

the variation of the thermal expansion of Lao.5Sro.sMnO3-LaCoO3 solid solutions as a function of cobalt content [5.70]. Tables 5.3 and 5.4 list selected thermal expansion data for doped LaMnO3.

Z uJ o U.. LL UJ 0 u u z E o 4

z

~ o x ' - W

_J

-1-

26

22

18

14

10

�9 Powder (X-Ray diffractometry) /

l t = t

II

EJ

i i " "

t I

I

o i

0,0 0.2 0 ,4 0 ,6 0 .8 1.0

COBALT CONTENT, mol

II

Figure 5.16. Thermal expansion coefficients of Lao.sSro.sMnl.xCoxO 3 compounds [5.70]

TABLE 5.3

Thermal Expansion Coefficient of Lal_xSrxMnO 3

(25 ~ to l l00~ [5.23]

Composition Thermal expansion coefficient

(10 .6 cm/cm.K)

Lao.99MnO 3 11.2

Lao.94Sro.osMnO 3 11.7

Lao.s9Sro.loMnO 3 12.0

Lao.79Sro.2oMnO 3 12.4

Lao.69Sro.3oMnO 3 12.8

Cathode 137

TABLE 5.4

Thermal Expansion Coefficients of Doped LaMnO3

Composition Thermal expansion coefficient Ref. (10 .6 cm/cm.K)

Lao.9Sro.lMnO3 12.0 [5.23] Lao.sSro.sMnO3 13.2 [5.75] Lao. 9Cao. l MnO3 10.1 [5.16] Lao 5Cao.sMnO3 11.4 [5.16] Lao.nYo. 1S r0.sMnO3 10.5 [5.88] Lao.vSro.3Mno.7Cro.303 14.5 [5.67] Lao.8Sro.2Mno.7Coo.303 15.0 [5.70]

5.2.7 Other properties

LaMnO3 can be sintered to full density at temperatures above 1250~ in

air. Because the material can be easily densified under normal firing conditions,

starting powders and processing procedures are often tailored to reduce the

sinterability of LaMnO3 to produce a porous cathode structure (with required

stability) for SOFC applications [5.17,5.124-5.126]. The sinterability of

LaMnO3 depends on, among other factors, the dopant, dopant level, and cation nonstoichiometry of the material. For example, calcium-doped LaMnO3 is more

sinterable than strontium- and barium-doped compounds [5.127]; LaMnO3 with

a higher strontium dopant content requires a higher sintering temperature [5.81, 5.85,5.127]; and lanthanum deficiency improves the sinterability of LaMnO3

[5.81,5.127]. The oxygen diffusion in doped LaMnO3 plays an important role in the

oxygen electrode reaction in SOFCs. The chemical diffusion coefficient of

oxygen in Lao.79Sro.2oMnO3_ ~ and Lao 5oSro.5oMnO3_~ has been measured to be about 10 .8 to 10 -6 cmZ/s in the temperature range of 700 ~ to 860~ indicating rapid

oxygen transport [5.128, 5.129]. On the other hand, the oxygen self-diffusion

coefficient has been found to be only about 10 -~5 to 10 ~2 cm2/s for Lal_xSrxMnO3_ ~

at 700 ~ to 900~ [5.130-5.132]. This suggests that for SOFC applications, the

LaMnO3 cathode porosity needs to be optimized to improve oxygen ion transport.

138 Chapter 5

Oxygen permeation in LaMnO3 has been described in terms of a point defect

model [5.133].

5.3 LANTHANUM COBALTITE

Doped lanthanum cobaltite (LaCoO3) is another perovskite of interest as

SOFC cathode material. The material belongs to the same class of oxide

compounds as LaMnO3. LaCoO3 is rhombohedral from room temperature to

1000 oc [5.134-5.136]. The material rhombohedral structure may transform to

a cubic phase; the transformation temperature depends on dopant content

[5.135,5.136]. LaCoO3 has no oxygen excess but shows a large oxygen

deficiency at high temperatures, especially when doped with a lower-valence

cation such as strontium. At 1000~ the oxygen stoichiometry of undoped

LaCoO3_ 6 ranges from 2.975 to 2.750 under oxygen partial pressures of 10 .2 to

10 .4 atm (103 to 10 Pa) [5.13 7]. At 800 ~ C, the undoped compound is essentially

stoichiometric. When doped with strontium, LaCoO3 becomes oxygen-deficient:

the oxygen stoichiometry of La0.7Sr0.3CoO3_~ varies from 2.970 to 2.840 under oxygen partial pressures of 1 to 10 .5 atm (105 to 1 Pa) [5.137]. The oxygen

deficiency of the cobaltite increases with increasing temperature, increasing

dopant content, and decreasing oxygen partial pressure [5.136-5.142]. The oxygen nonstoichiometry in LaCoO3 is inversely proportional to the l/ruth power

of oxygen partial pressure. The values of 1/m have been found to be about 0.45

for undoped LaCoO 3 [5.143] and about 0.13 for Lao.7Sr0.3CoO 3 [5.140]. (i) Stability: LaCoO3 is stable in oxidizing atmospheres but decomposes

in reducing environments. The material is less resistant toward reduction, when

compared with LaMnO3. At 1000~ LaCoO 3 dissociates into other phases at

the critical oxygen partial pressure of less than 10 -7 atm (10 .2 Pa) [5.49,5.52, 5.144, 5.145]. (For comparison, LaMnO3 decomposes at oxygen partial pressure

of 10 -~5 atm or 10 -~~ Pa.) Doping reduces the stability of LaCoO3, and the critical

oxygen partial pressure shifts to higher values with increasing dopant content

[5.146]. (ii) Conductivity: LaCoO 3 has been shown to have intrinsic p-type

conductivity, and the electrical conductivity can be enhanced by substituting a

lower-valence ion on the lanthanum site. Strontium and calcium are the most

commonly used dopants for LaCoO 3. The conductivity of Lao.8Sr0.2CoO3_~ and

Lao.8Cao.2CoO3_~ at 1000~ is 1200 f~-~cm -~ [5.70,5.135,5.136] and 800 fl-~cm -~

[5.147, 5.148], respectively. The conductivity of doped LaCoO3 increases with

Cathode 139

increasing dopant level and exhibits a maximum at 40 mol% for strontium

[.5.136] and 30 mol % for calcium [5.138]. Undoped LaCoO3 shows a semiconducting-metallic conduction transition at about 800~ in an oxygen atmosphere [5.135]. Strontium and calcium doping appears to lower this transition temperature. Above 30 mol % strontium, the conductivity of the doped material

is mainly metallic from room temperature to 1000~ [5.135,5.136,5.147]. (iii) Chemical interaction: LaCoO 3 interacts readily with YSZ to form

La2Zr207 at high temperatures ( > 1100 oc) [5.116, 5.117]. Dopant present in

LaCoO 3 may also react to produce other phases (e.g., SrZrO3 for strontium

dopant [5.98, 5.102, 5.105]). Reaction products tend to form a multilayer reaction

zone [5.116,5.149]. (iv) Thermal expansion: The average thermal expansion coefficient of

LaCoO 3 is about 22 to 24 x 10 -6 cm/cm.K[5.147,5.149]. The coefficient can

be modified by strontium [5.147], calcium [5.138,5.147], manganese [5.70], and

nickel [5.149] substitution for lanthanum and cobalt. For example, the thermal

expansion coefficient for LaCo0.6Nio.403 is about 17 x 10 -6 cm/cm.K[5.149]. In general, the thermal expansion coefficient of LaCoO 3 (even with doping) is

significantly higher than that of YSZ electrolyte.

5.4 OTHER MATERIALS

A number of doped oxides have been proposed and investigated as SOFC cathode materials [5.150,5.151]. The disadvantages of most of these materials are thermal expansion mismatch, incompatibility with YSZ, and lack of

conductivity. Early SOFCs used tin-doped indium oxide (In203) as the cathode

material [5.152-5.1.54]. This material exhibits excellent electrical conductivity under fuel cell operating conditions and can be applied as thin film by the

chemical vapor deposition method [5.155]. The material performed satisfactorily

for up to 5000 h without degradation. However, In203 represents the most costly

and least thermodynamically stable component in the SOFC. Therefore, In203

has been replaced by doped LaMnO3. Other materials containing In203, such as

In203-PrO2-HfO 2, In203-PrO2-ZrO2, In203-ZrO 2 have also been proposed [5.156- 5.159]. Doped YMnO3, CaMnO3, and YFeO 3 compositions have been considered

for possible SOFC cathode application [5.38,5.42,5.160-5.162]. Recently,

compounds based on LaFeO3-LaCoO3 solid solutions have been studied as

cathode materials in reduced-temperature SOFCs [5.163-5.165] due to their

desirable mixed conducting properties [5.166,5.167]. Strontium-doped PrMnO3

140 Chapter 5

compositions show promise as a suitable cathode material for reduced-tempera-

ture applications [5.168]. RuO2/YSZ materials have also been evaluated for

SOFC cathode uses [5.169].

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5.3

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Cathode 143

5.72

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Cathode 145

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Chapter 6

ANODE

6.1 REQUIREMENTS

The main function of the SOFC anode is to provide reaction sites for the

electrochemical oxidation of the fuel. Thus, the anode material must be stable in the fuel reducing environment and have sufficient electronic conductivity and

catalytic activity for the fuel gas reaction at the operating conditions. Since the SOFC operates at high temperatures (600 ~ to 1000~ the anode must be

chemically and thermally compatible with the other cell components from room temperature to those operating temperatures and even higher temperatures at

which the fuel cell is fabricated. The key requirements for the anode in the

SOFC are discussed below. This discussion is qualitative because the specific

requirements depend on selected materials and cell and stack designs [6.1]. (i) Stability: The anode must be chemically, morphologically, and

dimensionally stable in the fuel environment. The anode material must be stable,

not only at the fuel inlet conditions, but also at the more oxidizing conditions of the fuel outlet. The material must have no disruptive phase transformation (involving large changes in molar volume) between room temperature and

fabrication temperature. The anode must maintain its dimensions and desired microstructure in a hydrogen atmosphere and in long-term operation. Significant

structural changes can cause degradation in cell performance and mechanical integrity.

(ii) Conductivity: The anode must possess sufficient electronic conduct-

ivity for electron flow in the reducing environment at the operating temperature.

In general, maximum possible anode conductivity is desirable to minimize ohmic

losses. The anode conductivity must not change significantly due to changes in

oxygen partial pressure in the fuel environment during cell operation.

(iii) Compatibility: The anode must be chemically compatible with other

components, not only at the operating temperature, but also at the much higher temperature at which the fuel cell ceramic structure is fabricated. ChemiCal

148 Chapter 6

interaction or elemental interdiffusion between the anode and adjoining

components must be limited to minimize unacceptable occurrences such as second

phase formation, stability reduction, change in thermal expansion, introduction

of electronic conductivity in the electrolyte, etc.

(iv) Thermal expansion: The thermal expansion of the anode must match (from room temperature to operation and fabrication temperatures) that of other

cell components to avoid cracking and delamination during fabrication and operation, including thermal cycling. The coefficient of thermal expansion of the

anode must remain unchanged despite changes in oxygen partial pressure of the fuel atmosphere during operation.

(v) Porosity: The anode must have sufficient porosity to allow gas

transport to the reaction sites. The lower limit on porosity is set by mass transport considerations. (The porosity limit may be less critical for mixed conducting anode materials.) The upper limit is based on consideration of mechanical strength of the component.

(vi) Catalytic activity: The anode must have sufficient catalytic activity,

thus low polarization, for the electrochemical oxidation of the fuel. The anode must be tolerant to certain levels of contaminants (e.g., sulfur) commonly present

in fuel gas. If the anode is used as a catalyst for internal reforming of hydrocar-

bons, the anode must also maintain its reforming effectiveness over long operating periods.

In addition to these requirements, other desirable properties for the SOFC anode are high strength and toughness, fabricability, and low cost.

A number of ceramic and metallic materials can potentially meet the requirements discussed above. Because of the reducing conditions of the fuel gas, metals can be used as SOFC anode materials. At the high operating

temperature of the SOFC, suitable metals are limited mainly to nickel, cobalt,

and noble metals. Electronic conducting ceramics and mixed conducting oxides

(stable in the fuel reducing environment) are also suitable. Currently, the most

common anode material is nickel metal, often dispersed on an Y203-stabilized

ZrO2 (YSZ) support (cermet).

6.2 NICKEL/YTTRIA-STABILIZED ZIRCONIA CERMET

At present, nickel is used almost exclusively as the SOFC anode material.

This metal is preferred primarily because of its low cost (when compared with

other metals such as cobalt, platinum, and palladium). To maintain the required

Anode 149

porous structure of nickel over long periods at high temperatures and to provide

other desired properties for the anode, YSZ is often incorporated as a support. The functions of the YSZ in the anode are to support the nickel-metal particles,

inhibit coarsening of the metallic particles at the fuel cell operating temperature, and provide an anode thermal expansion coefficient acceptably close to those of

the other cell components. The YSZ support is considered as "inactive" although the support may play an important role in dictating the catalytic activity of the

anode. Some properties of nickel/YSZ cermet are given in Table 6.1.

TABLE 6.1

Properties of Nickel/YSZ Cermet (in Reducing Atmosphere)

Melting point, ~ (melting point of nickel)

Density, g/cm 3

(30 vol% nickel) Conductivity at 1000 ~ fl%m 1

(30 vol % nickel, 30 % porosity) Thermal expansion coefficient, 10 .6 cm/cm.K

(30 vol % nickel, 30 % porosity) Strength at 25 ~ MPa

(30 vol % nickel, 30 % porosity)

1453

6.87

- 500

-12.5

- 100

6 . 2 . 1 P r e p a r a t i o n

A variety of preparation methods are available for making nickel/YSZ

cermet anodes. These methods include conventional ceramic forming techniques

(such as tape casting, calendering), coating techniques (such as screen printing,

slurry coating), and deposition techniques (such as plasma spraying, chemical

vapor deposition (CVD)) (for references, see fabrication sections of Chapter 9).

For example, the anode of the sealless tubular SOFC is fabricated by slurry

coating of nickel and electrochemical vapor deposition of YSZ. The anode of the

banded SOFC (segmented-cell-in-series design) is made by flame spraying of NiO

and YSZ. Tape casting and tape calendering are commonly used to form

NiO/YSZ anode layers for the monolithic SOFC. The nature of the design and

150 Chapter 6

assembly of the flat-plate SOFC permits a variety of methods for anode

fabrication; suitable techniques range from casting, printing to plating, spraying. In most cases, the SOFC anode is first made with NiO and YSZ. The

NiO is then reduced in situ to nickel metal when exposed to the fuel in the fuel

cell. Examples of SOFC anode microstructures before and after NiO reduction are shown in Figure 6.1. The reduction of NiO to nickel increases the porosity of the anode. The porosity increase is caused by volume change as a result of oxygen loss due to the conversion of the oxide to its metallic form. The

relationship between the air-fired porosity and the hydrogen-reduced porosity of nickel/YSZ anodes is presented in Figure 6.2 [6.2].

One of the key factors in the preparation of nickel/YSZ anodes is to tailor and control electrode morphology, because the characteristics and stability of the anode microstructure are known to significantly affect electrode electrochemical performance [6. 3, 6. 4]. In the conventional powder mixing process, the anode morphology can be tailored by controlling the starting powder properties. Figure 6.3 shows the influence of particle size ratio on anode overpotential, thus

cell electrochemical performance [6.5]. A minimum anode overpotential can be

obtained at a certain nickel-to-YSZ particle size ratio.

Figure 6.1. SOFC anode microstructure after air firing (A) and after hydrogen reduction (B)

Anode 151

50

40 CD

o tr o a. 30

W (D D s UJ nr 2O

z W s O tr

I0 >- T

0 10

CALCULATED LINE

~ f

20 30

AIR-FIRED POROSITY, %

Figure 6.2. Relationship between air-fired porosity and hydrogen-reduced porosity of nickel/YSZ anodes [6.2]

0.7

0.6

0.5

0.4

0.3

0.2

0.1

0 - - 1 0 ,4

�9 " ' ' " " I ' ' ' " ' " 5 " ' " ' ' ' " " I ' ' ' " ~ I , . . . i , ,

1000~ ~1 NiO 2 1 / J m

I YSZ 0.35/1m o

I ~ Y S Z 0 . 1 1 # m "

o

o

o

.

�9

L �9 . . . . d . . . . . . J , , , , . . , J , , , , , . . , I , , , "

1 0 .3 1 0 .2 1 0 -~ 1 0 ~ 101

YSZ/NiO PARTICLE SIZE RATIO

Figure 6.3. Relationship between nickel/YSZ anode overpotential and particle size ratio of starting powders [6.5]

152 Chapter 6

Fabrication techniques have also been devised to prepare anode structure such that nickel particle sintering is minimized during long-term operation. In the fabrication of sealless tubular SOFCs, the nickel/YSZ anode is produced by slurry coating, followed electrochemical vapor deposition (EVD) (see Chapter 9). The YSZ deposited by EVD grows around the nickel phase, thus inhibiting nickel sintering. A technique called pyrolysis of metallic soap slurry has been developed to deposit fine YSZ particles on the surface of NiO [6.6, 6. 7]. The process involves preparing a slurry of NiO in a zirconium and yttrium octylate solution and firing to polymerize and decompose the organometallics to form YSZ on the NiO particles (Figure 6.4). This process produces a controlled nickel/YSZ microstructure with improved adhesion and morphological stability. A vapor phase method based on CVD and EVD deposits YSZ on NiO (by reacting ZrC14 and YC13 gases with oxygen released from NiO) [6.8], thus reducing nickel particle coarsening in the anode structure. A preparation method using liquid phase synthesis with YSZ sol produces a well-dispersed nickel on a MgO-YSZ support (Figure 6.5) [6.9]. Anodes fabricated by this technique show long-term stability due to suppressed grain growth.

Figure 6.4. Micrograph of anode prepared by pyrolysis of metallic soap slurry (courtesy of Tokyo Gas)

Anode 153

Figure 6.5. Microstructure of nickel/MgO-YSZ anode prepared with YSZ sol [6.9]

When the nickel/YSZ anode is formed on sintered YSZ electrolyte,

adhesion of the electrode is an important consideration [6.10]. Processing

conditions must be carefully tailored and optimized to produce anodes with good

interfacial bonding. For example, a certain YSZ powder ratio and milling time

have been found to improve adhesion in spraying the NiO/YSZ mixture [6.11].

Excessively high firing temperatures, often beneficial to stronger adhesion

between the electrode and electrolyte, can lead to formation of insulating phases due to interface interactions, resulting in poor electrode performance.

6.2.2 Stability

Both nickel and YSZ phases of the SOFC anode are known to be

chemically stable in the fuel reducing environment. Also, the materials exhibit

no phase transformation between room temperature and the operating tempera-

ture. The main concerns with the stability of the nickel/YSZ anode relate to the

electrode dimensional change and the sintering of nickel particles in long-term

operation at elevated temperatures. These changes are particularly important

when the anode is prepared by firing a mixture of NiO and YSZ powders.

The nickel/YSZ anode may change its dimensions and microstructure

over long periods of time or during NiO reduction if the YSZ does not form a

continuous network to support the nickel particles. The formation of a three-

dimensional YSZ network in the SOFC anode strongly depends on fabrication

154 Chapter 6

conditions and starting material characteristics and compositions. For example,

NiO/YSZ anodes produced by screen printing often have poor YSZ networks

[6.12, 6.13]. Screen-print anodes, therefore, can show measurable volume

change as a function of time at high temperatures. The YSZ particle size and

YSZ content have a notable effect on this change (Figure 6.6) [6.14-6.16]. In

general, greater than 50 wt% YSZ is required to build a continuous ZrO2

network in the nickel/YSZ cermet.

It is well known that fine nickel particles of the nickel/YSZ anode tend to coarsen under the fuel cell operating conditions. Nickel coarsening (or

sintering) results in loss of active surface area and reduced conductivity of the anode, leading to degradation of cell performance. Figure 6.7 shows, as an example, the effect of nickel sintering on the polarization of the anode [6.4,

6.10]. Since nickel particles are high-surface-area solids, there will always be

a thermodynamically driving force to decrease free energy, i.e., to minimize

surface area. Thus, the sintering behavior of the nickel/YSZ anode is strongly

dependent on the wetting properties of the nickel on the YSZ. In general, the rate of anode sintering is dependent on the nickel-particle-size distribution, with

increase in the rate as the width of dispersion increases. The rate of sintering also increases as the nickel content in the anode increases.

YSZ 13

1 . 0 o.6 pm - - O

3 6

~Sr- 0 0.9

O ~ F TEMPERATURE 1000~

/ TIME 15 h . _ ~u DIAMETER OF NICKEL PARTICLE 2.6 pm r r 40 pm

O F I I I 1 I , I . i 0 2 0 4 0 6 0

Y S Z C O N T E N T , w t %

Figure 6.6. Relationship between anode volume change and YSZ content with various YSZ particle sizes [6.15]

Anode 155

O

z" O m

I-- < N n.- < _ J

o

L U

0 z <

6

5

4

3

2

1 ."" a ram"

0 0

Q

a

10000C

CURRENT DENSITY: 222 mAJcm = FUEL UTILIZATION: 50% GAS COMPOSITION: Hz/3%H20

! �9 ! �9 ! �9

1000 2000 3000 4 0 0 0

TIME, h

Figure 6. 7. Effect of nickel sintering on cermet anode polarization [6.4]

The coarsening of the nickel/YSZ cermet anode can be analyzed using

a transmission line analog model [6.4]. The polarization of the anode based on this model is given as

1 (r3+krt) 1/31 1 ~/ = _-=-_[pC' ]Tcoth [ (_.p. r. L ) ] ] (Eq. 6.1)

ro ~ (ro 3 +krt) ~'3

where Np is the number of pores per unit area, o is the electrolyte resistivity, ro

is the initial particle radius, kr is the proportionality constant, t is the time, L is

the electrode thickness, and r is given as

r = ~rZ (Eq. 6.2) 2

r r is the pore radius and Z is the interfacial resistance between nickel and

~'SZ. At initial stages of coarsening (t = 0), Eq. 6.1 becomes

vf~--cothF__O L (Eq. 6.3)

156 Chapter 6

At long periods of time (t very large), Eq. 6.1 reduces to

1

1 ff(krt) ~ (Eq. 6.4)

Thus, the anode polarization will increase rapidly at the beginning and continue to increase as long as the driving force for nickel sintering remains significant.


The electrical conductivity of nickel/YSZ cermet is strongly dependent on its nickel content. The conductivity of the cermet as a function of nickel content shows the S-shaped curve predicted by percolation theory (Figure 6.8) [6.17, 6.18]. The theory as it applies to electrical conductivity of composites has been discussed elsewhere [6.19]. The percolation threshold for the conductivity

1 0 4

1 0 3

10 2

N 101 I-- o

z 0 1 0 ~ o

10 -1

10 .2

LOWER SURFACE AREA YSZ ~ ~ "" " - - ' -

I

r AREA YSZ

I I

0 10 20 3 0 4 0 5 0 6 0

V O L % NICKEL

Figure 6.8. Conductivity of Ni/YSZ cermet at I O00~ as a function of nickel content [6.17]

Anode 157

of the cermet is at about 30 vol % nickel. This percolation behavior can be explained by the presence of two conduction mechanisms through the cermet: an electronic path through the nickel phase and an ionic path through the YSZ

phase. Below 30 vol% nickel, the conductivity of the cermet is similar to that of YSZ, indicating an ionic conduction path through the YSZ phase. Above 30

vol% nickel, the conductivity is about three orders of magnitude higher,

corresponding to a change in mechanism to electronic conduction through the

nickel phase. This is supported by the fact that the conductivity of a nickel/YSZ cermet containing more than 30 vol % nickel decreases with increasing tempera-

ture, and the activation energy for conduction is similar to that of pure nickel

(5.38 kJ/mol) [6.17]. Above 30 vol% nickel, the conductivity of the cermet is also dependent on its microstructure (YSZ surface area). At the same nickel content, a YSZ support with lower surface area has better nickel coverage, resulting in improved nickel particle-to-particle contact, thus higher conductivity

for the cermet.

The dependence of the conductivity of the nickel/YSZ anode (nickel

content > 30 vol%) on temperature (700 ~ to 1000~ follows Arrhenius behavior. Figure 6.9 shows an example of the linear Arrhenius plots of

logarithm conductivity of nickel/YSZ versus reciprocal temperature.

10 4

"7, E

'7,

>: i - ra > l-- ~O

n Z 0 ~O

103

102

7.5 8.0

40 VOL% NICKEL _________---,,---

3 2 VOL% N I C K E L

8.5 9.0 9.5 10.0 10.5

TEMPERATURE, 1 04 /K

Figure 6.9. Temperature dependence of conductivity of nickel/YSZ cermet with nickel content greater than 30 vol % [6.17]

158 Chapter 6

In general, hydrogen reduction of NiO in the anode is fast at high

temperatures. Although most of the reduction occurs in the first several minutes,

it may take longer for the anode conductivity to reach a steady state due to

continuing reduction and rearrangement of nickel particles as the reduction

proceeds. During the reduction, the conductivity of the anode usually reaches

a maximum very quickly, then falls off slowly until a steady state is achieved

[6.17]. The maximum occurs when enough NiO is reduced to form a conducting nickel-metal matrix, and the fall-off corresponds to loss of nickel particle contact

as the particles shrink due to further NiO reduction. Thus, this conductivity fall- off varies depending on the nickel content in the anode (Figure 6.10)


The nickel/YSZ cermet anode has negligible chemical interaction with

YSZ electrolyte and LaCrO3 interconnect at temperatures below 1000~

However, at higher temperatures, NiO may react with LaCrO3 to form poor

conducting phases such as NiCrO4 [6.20]. In cofiring NiO/YSZ anode laminated

with LaCrO3 interconnect, liquid phases present in the LaCrO3 tend to migrate into the electrode, forming a reaction layer at the electrode/interconnect interface

(Figure 6.11) [6. 21]. For example, a dense interfacial region rich in calcium and

10 3

-E o 10 z

>': 10' I-- > I - (.) 10 0 :3 (3 Z 0 10 -~ (..)

1000oc

".... ""~

~176176

.o

26 VOL% Ni

1 04'.

~EO 103,.

~..~102.

~ 1 0 ~,.

~ 1 0 o , s

10-). 0

1000oC

50 VOL% Ni

1 0 .2 0 200 400 600 800 200 400 600 800

TIME, s TIME, s

Figure 6.10. Anode conductivities as a function of time during NiO reduction [6.17]

Anode 159

100

7 5

z i i i

i i i - - 5 0 i i i

2 5

0

R E A C T I O N A N O D E L A Y E R I N T E R C O N N E C T C A T H O D E

! " ' O "-" O " - Cl

A ~Z~.... A t A '~ /X

-.....

4~ o o o , -o-o.__ ~/ - r

,j,,,

0 5 0 1 0 0 1 5 0

D I S T A N C E , pm

2 0 0

~ O Ni ~ E I Zr - - Z ~ La

~ 0 Cr V DOPANT �9 DOPANT

~ & S r ~ V M n

Figure 6.11. Elemental distribution in cofired anode (NiO/YSZ)/interconnect (doped LaCrO ycathode (Sr-doped LaMnO 3) [6.21]

chromium has been observed in cofiring NiO/YSZ with calcium, cobalt-doped LaCrO3 [6.22]. At 1400~ for 1 h, calcium and chromium diffuse more than 100/zm (from the anode/interconnect interface) into the porous electrode [6.20].

To date, no effective methods have been found to prevent liquid phase migration from LaCrO3 in cofiring the anode with the interconnect.


The thermal expansion of the nickel/YSZ anode varies with the cermet composition, increasing with increased nickel content [6.18, 6.23, 6.24]. Figure

6.12 shows a plot of the average coefficient of thermal expansion of cermet anode (from room temperature to 1200~ as a function of NiO (or nickel)

volume percent. The nickel/YSZ anode generally has a higher thermal expansion coefficient than YSZ and other cell components. A significant degree of

mismatch between the thermal expansion coefficient of the anode and those of other cell components can result in large stresses, causing cracking or delamination during fabrication and operation (see Chapter 10). Various means have

160 Chapter 6

V O L % Ni

0 15 30 51 72 1 O0 v a 15 I i I I I I I;.,

E

b 1 4 -

z

i i i1

0

z _o 12- z < D .

x w 1 1 - _.1

er

"1" 10 I l I / I I

0 20 40 60 80 1 O0

VOL% NiO

Figure 6.12. Thermal expansion coefficient of cermet anode as a function of NiO (or nickel) content [6.23]

been developed to tolerate and minimize anode thermal expansion mismatch. For example, improving the fracture toughness of the electrolyte through additives has been attempted to provide sufficient tolerance of stresses generated by thermal expansion mismatch [6.25]. Control of critical processing flaws is

another effective means to increase fracture toughness of the electrolyte. Varying the thickness and thickness ratio of the cell components can be used to increase tolerance of thermal expansion mismatch [6.23]. Minor constituents have been

added to the anode formulation to improve the anode thermal expansion match

with those of the other cell components. Additives can not cause any deleterious

effects on other anode properties such as unacceptable degradation in electrical

conductivity and dimensional stability.

6.3 OTHER MATERIALS

Cobalt is another suitable SOFC anode material, since the metal can

withstand the fuel environment and remains non-oxidized. Cobalt/calcium-doped

ZrO2 anodes have been used in SOFCs [6.26]. Compared with nickel, cobalt has the advantage of high sulfur tolerance; however, cobalt is not commonly used

Anode 161

because of its high cost. Also, the oxidation potential of cobalt is higher than

that of nickel metal, thus requiring a less complete fuel combustion. Recently, ruthenium/YSZ cermets have been tested as SOFC anodes

[6.27-6.29]. Ruthenium has a higher melting point (2310~ than nickel

(1453~ thus providing better resistance to particle coarsening. Ru/YSZ cermets have been shown to have minimum sintering at fuel cell operating

temperatures, high catalytic activity for steam reforming, and negligible carbon deposition under reforming conditions. Ruthenium/YSZ anodes are fabricated

by slurry coating of ruthenium metal particles, followed by EVD of the YSZ

phase. Mixed conducting oxides have also been investigated as SOFC anode

materials. In mixed conductors in which both oxygen ions and electrons are

mobile, the electrochemical reactions occur over the entire electrode/gas interfacial area. Thus, polarization losses with a mixed conducting electrode are

expected to be significantly less than with electrodes exhibiting only electronic

conductivity. ZrO2-Y203-TiO2 solid solutions are of particular interest for SOFC

applications because of their compatibility with the YSZ electrolyte [6.30-6.36].

Up to 15 mol % TiO2 can be dissolved in ZrO2 (stabilized with 12 mol % Y203)

to form a mixed conducting fluorite structure phase. The optimum composition

for a mixed conducting ZrO2-Y203-TiO2 anode is that at which the transference numbers or percentages of electronic and oxygen-ion conductivity are equal. For

example, the optimum TiO2 composition in 10 mol % Y203-stabilized ZrO2 would

be 9.3 mol% at 1000~ [6.30]. CeO2-based materials have been evaluated for SOFC anode applications

[6.37-6.41]. Doped CeO2, exhibiting mixed conduction in the fuel reducing

environment, has been studied both as anode material and as support, replacing the ZrO2 in nickel/YSZ cermet. The material has also shown considerable

promise as electrode material for direct oxidation of CH 4. Anodes made of doped CeO2 particles with highly dispersed metal catalysts on the surface have

shown significantly improved catalytic activity, especially at reduced tempera-

tures [6.42].

References

6.1 N.Q. Minh, in Science and Technology of Zirconia V, S.P.S. Badwal, M.J. Bannister, and R.H.J. Hannink (eds.), Technomic Publishing Company, Lancaster, PA, 1993, p. 652.

162 Chapter 6

6.2

6.3

6.4

6.5 6.6

6.7

6.8

6.9

6.10

6.11 6.12

6.13

6.14

6.15

6.16

N.Q. Minh, C.E. McPheeters, and J.E. Brule, Monolithic Solid Oxide Fuel Cell Technology Development, Phase 1A, Final Report, February 1987-March 1989, Report No. GRI-89/0267, Gas Research Institute, Chicago, IL, 1989. T. Kawada, N. Sakai, H. Yokokawa, M. Dokiya, M. Mori, and T. Iwata, J. Electrochem. Soc., 137 (1990) 3042.

S. Elangovan and A. Khandkar, in Proceedings of the First International Symposium on Ionic and Mixed Conducting Ceramics, October 16-17, 1991, Phoenix, AZ, T.A. Ramanarayanan and H.L. Tuller (eds.), Electrochemical Society, Pennington, NJ, 1991, p. 122. T. Hikita, see Ref. 6.1, p. 674.

Y. Matsuzaki, M. Hishinuma, T. Kawashima, I. Yasuda, T. Koyama, and T. Hikita, in 1992 Fuel Cell Seminar Abstracts, November 29-December 2, 1992, Tucson, AZ, Courtesy Associates, Washington, DC, 1992, p. 237. T. Hikita, M. Hishinuma, T. Kawashima, I. Yasuda, T. Koyama, and Y.

Matsuzaki, in Proceedings of the Third International Symposium on Solid Oxide Fuel Cells, May 16-21, 1993, Honolulu, HI, S.C. Singhal and H. Iwahara (eds.), Electrochemical Society, Pennington, NJ, 1993, p. 714. T. Ogawa, T. Ioroi, Y. Uchimoto, Z. Ogumi, and Z. Takehara, see Ref. 6.6, p. 479. Fuel Cell RD & D in Japan, Fuel Cell Development Information Center, Tokyo, Japan, 1992, p. 66.

Ceramatec, Inc., Development of Planar Geometry Solid Oxide Fuel Cell Technolo- gy, Phase II, Annual Report, October 1987-October 1988, Report No. GRI-89/0161, Gas Research Institute, Chicago, IL, 1989. C. Bagger, see Ref. 6.6, p. 241.

Ceramatec, Inc., Development and Optimization of Planar Geometry Solid Oxide Fuel Cells, Annual Report, December 1990-December 1991, Report No. GRI- 92/0132, Gas Research Institute, Chicago, IL, 1992. A.C. Khandkar, S. Elangovan, M. Liu, and M. Timper, in Proceedings of the Symposium on High Temperature Electrode Materials and Characterization, May 5- 10, 1991, Washington, DC, D.D. MacDonald and A.C. Khandkar (eds.), Electrochemical Society, Pennington, NJ, 1991, p. 175. S. Murakami, Y. Miyake, Y. Akiyama, N. Ishida, T. Saito, and N. Furukawa, in

Proceedings of the International Symposium on Solid Oxide Fuel Cells, November 14-14, 1989, Nagoya, Japan, O. Yamamoto, M. Dokiya, and H. Tagawa (eds.), Science House, Tokyo, Japan, 1989, p. 187.

S. Murakami, Y. Akiyama, N. Ishida, T. Yasuo, T. Saito, and N. Furukawa, in

Proceedings of the Second International Symposium on Solid Oxide Fuel Cells, July 2-5, 1991, Athens, Greece, F. Grosz, P. Zegers, S.C. Singhal, and O. Yamamoto

(eds.), Commission of the European Communities, Luxembourg, 1991, p. 105. T. Saito, Y. Akiyama, N. Ishida, T. Yasuo, S. Taniguchi, S. Murakami, and N. Furukawa, Denki Kagaku, 61 (1993) 228.

Anode 163

6.17

6.18

6.19

6.20 6.21

6.22

6.23 6.24

6.25

6.26

6.27 6.28

6.29

6.30

6.31 6.32

6.33

6.34

6.35

D.W. Dees, T.D. Claar, T.E. Easier, D.C. Fee, and F.C. Mrazek, J. Electrochem. Soc., 134 (1987) 2141. E. Ivers-Tiff6e, W. Wersing, and M. SchieB1, Ber. Bunsenges. Phys. Chem., 94

(1990) 978. D.S. McLachlan, M. Blaszkiewicz, and R.E. Newnham, J. Am. Ceram. Soc., 73

(1990) 2187. T.R. Armstrong, L.A. Chick, and J.L. Bates, see Ref. 6.7, p. 632. N.Q. Minh, T.R. Armstrong, J.R. Esopa, J.V. Guiheen, C.R. Horne, F.S. Liu, T.L. Stillwagon, and J.J. Van Ackeren, see Ref. 6.15, p. 93. N.Q. Minh, A. Amiro, T.R. Armstrong, J.R. Esopa, J.V. Guiheen, C.R. Horne, and J.J. Van Ackeren, see Ref. 6.6, p. 607. S. Majumdar, T. Claar, and B. Flandermeyer, J. Am. Ceram. Soc., 69 (1986) 628. T.E. Easier, B.K. Flandermeyer, T.D. Claar, D.E. Busch, R.J. Fousek, J.J.

Picciolo, and R.B. Poeppel, in 1986 Fuel Cell Seminar Abstracts, October 26-29, 1986, Tucson, AZ, Courtesy Associates, Washington, DC, 1986, p. 72.

J.P. Singh, A.L. Bosak, D.W. Dees, and C.C. McPheeters, in 1988 Fuel Cell Seminar Abstracts, October 23-26, 1988, Long Beach, CA, Courtesy Associates,

Washington, DC, 1988, p. 145. T.L. Markin, R.J. Bones, and R.M. Dell, in Conference on Superionic Conductors, General Electric Research and Development Center, Schenectady, NY, G.D. Mahan and W.L. Roth (eds.), Plenum Press, New York, 1976, p. 15. M. Suzuki, H. Sasaki, S. Otoshi, and M. Ippommatsu, see Ref. 6.15, p. 585. H. Sasaki, M. Suzuki, S. Otoshi, A. Kajimura, and M. Ippommatsu, J. Electro- chem. Soc., 139 (1992) L12. M. Ippommatsu, H. Sasaki, A. Hirano, S. Otoshi, M. Suzuki, and A. Kajimura, in Proceedings of the 1992 International Gas Research Conference, November 16-19, 1992, Orlando, FL, H.A. Thompson (ed.), Government Institutes, Rockville, MD, 1993, p. 2062.

S.S. Liou and W.L. Worrell, in Proceedings of the First International Symposium on Solid Oxide Fuel Cells, October 16-18, 1989, Hollywood, FL, S.C. Singhal (ed.), Electrochemical Society, Pennington, NJ, 1989, p. 81. S.S. Liou and W.L. Worrell, Appl. Phys. A, 49 (1992) 25. K.E. Swider and W.L. Worrell, see Ref. 6.15, p. 593.

T. Lindegaard, C. Clausen, and M. Mogensen, in Proceedings of the 14th Riso International Symposium on Materials Science, High Temperature Electrochemical Behaviour of Fast Ion and Mixed Conductors, September 6-10, 1993, Roskilde,

Denmark, F.W. Poulsen, J.J. Bentzen, T. Jacobsen, S. Skou, and M.J.L. OstergArd (eds.), Riso National Laboratory, Roskilde, Denmark, 1993, p. 311.

J.L. Bates, L.A. Chick, G.E. Youngblood, and W.J. Weber, Advanced Materials and Electrochemical Processes in High-Temperature Solid Electrolytes, Report No. DE91-005181, U.S. Department of Energy, Washington, DC, 1990.

R.M.C. Marques, J.R. Frade, and F.M.B. Marques, see Ref. 6.7, p. 513.

164 Chapter 6

6.36

6.37

6.38

6.39 6.40

6.41

6.42

M.T. Colomer, J.R. Jurado, R.M.C. Marques, and F.M.B. Marques, see Ref. 6.7, p. 523.

T. Takahashi, H. Iwahara, and Y. Suzuki, in Proceedings of the Third International Symposium on Fuel Cells, June 16-20, 1969, Brussels, Belgium, Presses Acad6mi- ques Europ6ennes, Brussels, Belgium, 1969, p. 113.

I.S. Metcalfe, P.H. Middleton, P. Petrolekas, and B.C.H. Steele, SolM State Ionics, 57 (1992) 259.

M. Mogensen and J.J. Bentzen, see Ref. 6.30, p. 99.

M. Mogensen, B. Kindl, and B. Malmgren-Hansen, in 1990 Fuel Cell Seminar Abstracts, November 25-28, 1990, Phoenix, AZ, Courtesy Associates, Washington, DC, 1990, p. 195. M. Mogensen, see Ref. 6.15, p. 577.

M. Watanabe, H. Uchida, M. Shibata, N. Mochizuki, and K. Amikura, J. Electrochem. Soc., 141 (1994) 342.

Chapter 7

INTERCONNECT

7.1 REQUIREMENTS

The primary function of the SOFC interconnect is to connect the anode of one cell to the cathode of the next cell in electrical series. The interconnect also separates the fuel from the oxidant in adjoining cells of a stack. Thus, the interconnect material must be stable in both the reducing and oxidizing environments, impermeable to gases, and sufficiently conductive to support

electron flow at the operating conditions. Since the SOFC operates at high temperatures (600~ to 1000~ the interconnect must be chemically and thermally compatible with the other cell components from room temperature to those operating temperatures, and to even higher temperatures at which the fuel cell is fabricated. The key requirements for the interconnect in the SOFC are

discussed below. This discussion is qualitative because the specific requirements depend on selected materials and cell and stack designs [7.1].

(i) Stability: The interconnect must be chemically, morphologically, and dimensionally stable in the dual atmosphere (reducing atmosphere on one side and oxidizing atmosphere on the other). The interconnect must have no disruptive phase transformation (involving large changes in molar volume) between room temperature and fabrication temperature. The interconnect must maintain its dimensions (particularly, no major expansion or contraction) when exposed to the dual atmosphere.

(ii) Conductivity: The interconnect must possess adequate electronic

conductivity in the dual atmosphere to conduct electrons between electrodes of

adjacent cells. In general, maximum possible interconnect conductivity is

desirable to minimize ohmic losses. The interconnect conductivity must not

change significantly due to changes in oxygen partial pressures in the fuel and oxidant environments during cell operation.

(iii) Compatibility: The interconnect must be chemically compatible with the other cell components, not only at the operating temperature, but also at the

166 Chapter 7

much higher temperature at which the fuel cell ceramic structure is fabricated.

Chemical interaction or elemental interdiffusion between the interconnect and adjoining components must be limited in order to minimize unacceptable

occurrences such as second phase formation, stability reduction, change in thermal expansion, sinterability loss, etc. The interconnect material must tolerate certain levels of contaminants (e.g., sulfur) commonly present in fuel gas.

(iv) Thermal expansion: The thermal expansion of the interconnect must match (from room temperature to operation and fabrication temperatures) that of other cell components to avoid cracking and delamination during fabrication and operation, including thermal cycling. The coefficient of thermal expansion of the interconnect material must remain unchanged despite changes in oxygen partial pressures of the fuel and oxidant atmospheres during operation.

(v) Porosity: The interconnect must be dense (or contain no connected

porosity) to prevent gas cross leakage. The interconnect material must be impervious to both oxygen and hydrogen gases between room temperature and

operating temperature. In addition to these requirements, other desirable properties for the SOFC

interconnect are high strength and toughness, fabricability, and low cost. The stringent requirements discussed above eliminate all but a few oxide

systems from consideration for the interconnect in the SOFC. High-temperature alloys have also been considered as interconnect material, especially for fiat-plate SOFCs. Currently, LaCrO3 is the most common interconnect material for

SOFCs.

7.2 LANTHANUM CHROMITE

LaCrO3 has been used as SOFC interconnect material since the 1970s.

The particularly attractive features of the material include high electronic conductivity under fuel and oxidant atmospheres, adequate stability in the fuel

cell environment, and reasonable compatibility with other cell components. LaCrO3 is often doped to tailor and control its properties for SOFC applications.

7.2.1 Preparation

LaCrO3 imerconnects have been prepared by a number of fabrication

methods. A key factor in selecting a fabrication technique for the interconnect

is the technique's ability to produce gastight LaCrO3 under acceptable process

Interconnect 167

conditions. Presem fabrication methods for LaCrO3 imerconnects include EVD (sealless tubular design); plasma spraying (segmented-cell-in-series design, banded configuration); tape calendering (monolithic design); and CVD, plasma spraying, pressing, and tape casting (flat-plate design) (for references, see fabrication sections of Chapter 9). Recently, several other deposition techniques have been investigated for making thin interconnect layers [7.2-7. 6].

In many of the fabrication processes mentioned above, fine powders are used as starting materials. Although LaCrO3 can be prepared by solid-state reaction of oxides at elevated temperatures, this method often produces powders that do not meet the required levels of purity, chemical homogeneity, and reactivity. Therefore, LaCrO3 powders are commonly prepared by solution techniques. Powders prepared by these techniques generally exhibit higher purity, improved chemical homogeneity, increased reactivity, and more precise stoichiometry. The most common solution method to prepare LaCrO3 is the amorphous citrate (Pechini) process [7. 7] (for more details, see preparation section of Chapter 5). The process produces LaCrO3 powders of high quality via polymeric precursors made of citric acid and ethylene glycol.

LaCrO3 has also been synthesized by various modified amorphous citrate processes. The modifications mainly involve changing the citric acid/ethylene glycol molar ratio. (The original amorphous citrate process uses a 20/80 ratio.) In one modified version, a 50/50 (equimolar) ratio is used to produce a porous gel [7.8, 7.9]. Pulverization and calcination of this gel yield submicrometer LaCrO3 powders with a desirable narrow particle-size distribution. Other modified versions use either 100 % citric acid or 100 % ethylene glycol. In the citric acid process, citric acid is mixed with metal ion solution to produce a foamy gel precursor (after dehydration) [7.10, 7.11]. Pyrolysis at relatively low temperatures transforms the precursor into a high-surface-area oxide. In the ethylene glycol process, ethylene glycol is mixed with metal ion solution and nitric acid to produce a gelatinous liquid [7.12]. Calcination at temperatures above 700~ converts the liquid to powders having submicron particles.

Recently, a combustion method, called the glycine/nitrate process, has been developed to synthesize fine LaCrO3 powders [7.13, 7.14]. The process involves dissolving glycine and metal nitrates in an aqueous solution and heating the solution to evaporate water until the solution thickens and self-ignites to produce oxide product ash. In this process, the glycine forms complexes with the metal ions to increase solubility and prevent selective precipitation (during water removal), and also functions as fuel for the combustion. The oxide ash is

168 Chapter 7

generally composed of very fine (25- to 100-nm) particles linked together in chains. This process produces oxide powders of very high surface area and compositional homogeneity and with low levels of carbon residue. Other combustion methods (drip pyrolysis and spray pyrolysis) have also been used to synthesis fine LaCrO3. The drip pyrolysis process uses nitrate and acetate solution as starting material and glucose as combustion fuel [7.15]. The spray pyrolysis process uses metal salt solution along with carbohydrate [7.16].

Coprecipitation techniques have also been used to make LaCrO3 [7.17].

These methods entail the precipitation of oxalates or carbonates from a solution of metal nitrates. The methods produce precipitated solids with cation composition within experimental error of the target formulation.

7.2.2 General properties, phase transformation, and stoichiometry

LaCrO3 is a perovskite oxide (ABO3) with highly refractory properties. The compound structure consists of the rare-earth ion (La 3§ occupying the A site in 12-fold coordination with the oxygen ions and the B cation (Cr 3+) in

octahedral coordination. LaCrO3 melts congruently at 2510 ~ +20 ~ C [7.18]. The phase diagram of the Cr203-La203 system is shown in Figure 7.1 [7.18]. Table 7.1 lists the properties of LaCrO3.

I 1 I 1 1 2 5 0 0 1 2 5 ~ 0 " 2 0 ~ _

V z 300 ~ / I k z

" L,Q ~ 2210_20 _ _ . j \ f~>. o, \

, .#~s:

o ?-; . , o o _ 2060"_ 20* .,.v ~/!

1900 --

Cr203

P ' L o C r 0 3 + clCr203

1 1 20 4 0

X=H

P'LaCr03 + H-LaeO 3 (87~ _ _ H-"

P - L o C r 0 3 + A -Lo203

60 80 toz03

MoI %

Figure 7.1. Phase diagram of a Cr203-La203 system (A, H = hexagonal, P = perovskite, X = cubic) [7.18]

Interconnect 169

TABLE 7.1

Properties of LaCrO3

Melting point, ~ Density, g/cm 3

Thermal conductivity, W/cm.K 200~ 1000~

Thermal expansion coefficient, 10 -6 cm/cm.K 25 ~ to 240~ 25 ~ to 1000~

Standard enthalpy change (from La:O3 and Cr:O3), kJ/mol Standard entropy change (from La:O3 and Cr:O3), J/mol.K Bend strength, MPa

25~ 1000~

2510 [7.181 6.74 [7.751

0.05 [7.21] 0.04 [7.21]

6.7 [7.33] 9.2 [7.331 -67.7 [7.107] 10 [7.1071

200 [est.] 100 [est.]

LaCrO 3 is orthorhombic at room temperature and undergoes a crystallo-

graphic transformation from orthorhombic to rhombohedral at about 240 ~ to

290~ [7.19-7.28]. The rhombohedral form of the oxide changes to a hexagonal

structure at about 1000 ~ [7.18, 7. 29-7. 32]. A further increase in temperature

to about 1650~ brings about a transition to a cubic phase [7.18, 7.26, 7.30, 7.31] (although a transition temperature as low as 1030~ has been reported

[7.29, 7.32]). In general, the unit-cell volume of the LaCrO3 phases increases

linearly as a function of temperature, and the volume thermal expansion

coefficient increases from orthorhombic to hexagonal to cubic phases. The phase

transformations in LaCrO3 are accompanied by changes in thermal expansion

coefficient, electrical conductivity, and other properties.

Nonstoichiometry and doping influence phase transformation of the

LaCrO3. The temperature at which the orthorhombic/rhombohedral transforma-

tion occurs is dependent on the lanthanum stoichiometry and increases with

increasing lanthanum/chromium ratio [7.33]. The orthorhombic/rhombohedral

transformation temperature of LaCrO3 is also dependent on dopant. Strontium

substitution for lanthanum in LaCrO3 lowers the transformation temperature

[7.33-7.35]. As little as 10 mol% strontium is sufficient to stabilize the

170 Chapter 7

rhombohedral structure at room temperature. Aluminum [7.23, 7.33] or cobalt

[7. 36] substitution also lowers the transition temperature. Nickel [7. 26, 7. 36],

manganese [7.24], or calcium [7.37, 7.38] substitution raises the transformation temperature. For example, 20 mol % nickel doping increases the orthorhombic/

rhombohedral transition temperature as much as 75 ~ Magnesium substitution,

on the other hand, does not affect the transformation temperature [7.33].

LaCrO3 has oxygen stoichiometry in oxidizing atmospheres but becomes

oxygen deficient at elevated temperatures under highly reducing conditions.

Undoped LaCrO3 exhibits very low oxygen nonstoichiometry [7. 39]; however,

reliable data cannot be obtained for undoped oxide because of its high vaporiza-

tion rates. The oxygen deficiency in LaCrO3 depends on dopant, dopant

concentration, temperature, and oxygen partial pressure (oxygen activity). The

amount of oxygen deficiency in magnesium-doped LaCrO3 as a function of these parameters is shown in Figure 7.2 [7. 40-7. 43] as an example. In general, the oxygen activity at which maximum nonstoichiometry occurs shifts to lower values as the temperature decreases. The degree of oxygen nonstoichiometry increases as the amount of dopant increases at any given oxygen activity.

LaCrO3 can have excess lanthanum, which tends to precipitate as La203,

resulting in hydroxide formation and subsequent disintegration of LaCrO3 at

room temperature [7.34, 7. 44]. The material contains no or very little excess

0.10

0.09

0.08 R 0

E 0 . 0 7

U o.oo i i

tu 0.05

Z tu 0 .04 (9 >.. X 0.03 0

0.02

0.01

0.00

x-- 0 .20

x = 0 . 1 0

3 - 11 -9 -7 -5 -3 - 1 1

LOG OXYGEN PARTIAL PRESSURE, Pa

3 5

Figure 7.2. Oxygen deficiency as a function of oxygen partial pressure and dopant content at 1255~ for LaCr1_xMgx03 [7.41]

Interconnect 171

chromium. (The limit appears to be less than 1.5 mol% [7.34].) On the other hand, LaCrO3 can have lanthanum and chromium deficiencies. However, the maximum allowable deficiency (without second phase formation) also appears to be small in this case (e.g., less than 4 mol% chromium deficit [7.44]).

7.2.3 Stability

LaCrO3 is chemically stable in both oxidizing and reducing atmospheres. The material has insignificant chromium volatilization at the fuel cell operating

conditions. However, at very high temperatures (> 1600~ LaCrO3 appreciably volatilizes chromium oxides in oxidizing atmospheres [7.21,7.45-7.47].

At 1600~ the typical vaporization rate of LaCrO3 is 54 /zg/cm2.hwith an oxygen flow of 13 cm/s and a pressure of 0.2 x 10 s Pa. This high volatilization leads to the low sinterability of the oxide in high oxygen activities (see section on sinterability). Lattice substitution by dopants such as aluminum, strontium, and calcium reduces the volatility of LaCrO3 at those high temperatures [7.18,7.23,7.31,7.48,7.49].

Under reducing environments (e.g., hydrogen atmosphere), LaCrO3

expands as a result of lattice expansion due to loss of oxygen [7.33]. Figure 7.3 shows the relative dimensional change of strontium-doped LaCrO3 as a function

0.40

o -'e- 0.35

0.30 z < "1-"

0.25, _,1 < z 0.20 o O~ z 0.15 Ud

0.10 ILl > C-- 0 .05-

"' 0 . 0 0 -

\ ~

\ \ \ . \ .

\ \ � 9

k

\

-0.05 . . . . . . . . . . . �9

- 2 0 - 1 8 - 1 ' 6 - 1 ' 4 - 1 ' 2 - 1 " 0 -8 - ; -4 "2 0

LOG OXYGEN PARTIAL PRESSURE, arm (1.01 x 10 ~ Pa)

Figure 7.3. Relative dimensional change as a function of oxygen partial pressure for Lao.79Sro.zoCr03 at 1000~ [7.33]

172 Chapter 7

of oxygen partial pressure. The expansion of doped LaCrO3 in hydrogen can

vary depending on the dopant. For instance, magnesium-doped LaCrO3 shows about four times less expansion than does the strontium-doped material under

similar conditions (0.1% vs 0.4%) [7.33]. For SOFC applications, it is

necessary to design the LaCrO3 interconnect in such a way that the expansion at the fuel interface will not result in enough mechanical stress in the fuel cell to

cause cracking or delamination. Dopants and sintering aids used in LaCrO3 can significantly influence the

mechanical strength of the oxide in reducing atmospheres. Gases containing hydrogen and CO2 have been found to regenerate sintering-aid liquid phases present in LaCrO3, and the material becomes embrittled [7.50]. Cobalt in

calcium,cobalt-doped LaCrO3 precipitates as CoO in hydrogen, causing dramatic degradation in the fracture strength of the material [7.51]. In addition, sintering aids can cause chemical degradation on the surface of LaCrO3 interconnects [7.52]. In strontium-doped and calcium-doped LaCrO3, alkali-earth chromates

such as Cam(CrO4) n and SrCrO4 tend to migrate to the exposed surfaces under fuel cell operating conditions, and the chromates decompose at the fuel-side

surface [7. 52].


LaCrO3 can be substituted with a cation on either the lanthanum or chromium sites. Examples of substituents or dopants include strontium [7.35, Z 45, Z 48, Z53-Z58] and calcium [Z18, Z 31, Z 38, Z 44, Z50, 7.59-Z 65] (lanthanum site) and magnesium [7. 23, 7. 40-7. 42, 7. 49, 7. 66-7. 68], cobalt [7. 25, 7. 63], zinc [7. 69], copper [7. 69], iron [7. 70], titanium [7. 71], aluminum

[7.23, 7.49, 7. 72], nickel [7.26, 7.36, 7.50, 7. 73], niobium [7. 74], and manganese

[7. 24, 7. 75-7. 77] (chromium site). The substitution of lanthanum and chromium

in LaCrO3 with another cation significantly influences the electrical properties of

the material. LaCrO3 is a p-type conductor due to holes in the 3d band of the

chromium ions [7.27, 7. 78-7.82]. Under oxidizing conditions, substitution of a lower-valence ion on either the lanthanum or chromium sites of LaCrO3 results

in a charge-compensating transition of Cr 3+ to Cr 4+ ions, thereby enhancing the

electronic conductivity of the material. Under reducing conditions, charge

compensation occurs by the formation of oxygen vacancies; thus, no increase in

the electronic conductivity is anticipated. To have conductivity adequate for use

Interconnect 173

as SOFC interconnect, LaCrO3 is often doped with divalent ions. The most

common dopants are strontium, calcium, and magnesium. The solution limits for strontium [7.35], calcium [7.38], and magnesium [7. 41] in LaCrO 3 are 50, 50, and 15 mol %, respectively.

Figure 7.4 shows the electrical conductivity (in air) of strontium-doped

LaCrO3 as a function of temperature [7.57]. The linearity of the plots of log aT

vs 1/T for strontium-doped LaCrO3 indicates thermally activated hopping of small polarons as the conduction mechanism. Similar behavior is also observed for

other doped chromites. In general, the conductivity of LaCrO3 at elevated temperatures is proportional to the dopant concentration. For example, the

conductivity of calcium-doped LaCrO3 at 1000~ increases from about 20 to 40 to 60 f~-~cm -~ as the calcium content increases from 10 to 20 to 30 mol% [7.62]. It should be noted that the conduction activation energy for doped LaCrO3 is usually independent of dopant concentration. However, the activation energy for LaCrO3 in solid solution with other lanthanum perovskite oxides such as LaCoO3, LaMnO3 (i.e., doped with Co, Mn) varies notably as a function of dopant content

[7.64, 7.75] (Figure 7.5). In this case, the activation energy increases with

dopant content to a maximum, and then decreases as the dopant content

increases. This behavior is attributed to a lower small-polaron site energy at the dopant site as compared to the chromium site. The activation energy for

conduction and the conductivity at 1000~ for several doped LaCrO3 compounds are given in Table 7.2.

1 0 5

1 0 4

'T 1 0 3

I.-." b

10 2

I0 ~

~m~m....u �9 x = 0 .15

"------~..~--~~ / - - , , . . ,_ , . .o . . ,_ " ' - ~ , - ~ t ~ - ~ . _ _ m

_ x = 0 .05 x = 0 .10

1 1

5 10 15

1 IT, 1 0 " 4 / K

Figure 7.4. Conductivity of Lal.xSrxCrO ~ as a function of temperature [7.57]

174 Chapter 7

- - 0.5 "6 E

v L~

~ 0.4

~ 0.3

z

z

- 0.2 t -

>

u

'~ 0 . 1 ~ 1 l t l

0 2 0.4 0.6 0.8 1.0 COBALT CONTENT, rnol

Figure 7.5. Activation energy for conductivity of LaCrl_xCOx03 as a function of cobalt content [7. 64]

T A B L E 7 .2

Conductivity (in Air) Data for Doped LaCrO3

Dopant Composition Conductivity, 1000~ Activation energy (MO) (mol % MO) (fl-lcm -1) (kJ/mol)

Ref.

None 0 1 18 [Z54] MgO 10 3 19 [7.42] SrO 10 14 12 [7.57] CaO 10 20 12 [7.62] CoO 20 15 43 [7.64] MnO 20 0.2 46 [7. 74]

SrO,MnO 10,20 1 50 [ 7. 75] CaO,CoO 10,20 30 19 [7.64]

The conductivity of doped LaCrO3 depends on the equilibrium atmo-

sphere. Equilibration in a reducing atmosphere, such as hydrogen gas, causes

an appreciable decrease in conductivity [7.42, Z54, 7.60, Z65, Z67, 7. 72,7. 77].

For example, the conductivity of magnesium-doped LaCrO3 in hydrogen is an

Interconnect 175

order of magnitude lower than that in air [7. 77]. In a hydrogen atmosphere, the oxygen loss reduces the charge carrier concentration and, thereby, decreases the conductivity of the material. Figure 7.6 shows an example of the difference in the conductivity of doped LaCrO3 in oxidizing and reducing atmospheres [ 7. 83].

Since the LaCrO3 interconnect in a SOFC is subject to a dual atmosphere (with

fuel on one side and oxidant on the other), a conductivity gradient across the material may exist. However, the overall conductivity of the material is still

adequate [7. 84, 7. 85]. For example, under the fuel cell operating environment, the conductivity of 10 mol% magnesium-doped LaCrO3 is about 2 9-1cm -1,

exceeding the conductivity requirement for the interconnect [7.85]. For polycrystalline LaCrO3 exposed to an oxygen potential gradient such as the SOFC

dual atmosphere, the material equilibrates quickly due to its rapid oxygen exchange rate [7. 86].

A defect structure model has been proposed to correlate the dependence of conductivity on oxygen activity for doped LaCrO3 [7. 40-7. 42, 7. 58, 7. 62,

7. 66, 7. 82]. The model is briefly discussed here as applied to magnesium-doped

LaCrO3 [7. 40-7. 42, 7. 66]. In the compound LaCrl_~MgxO3_ ~, assuming p-type

10 2 , , , ; w I -

101

O 100

_> 10-1 I.-

2~ a z 0.2 0 1

10 -3

10 .4

x'-~r X" X - ~ ( X

xx... A IR Oo., o ~ x

N O ~ x ,...

\ %

\ ~ + 2 0 % H 2 0

. . . . . . . . .

8 10 14 18 22 2

1/T, 10"4/K

Figure 7.6. Electrical conductivity of Lao.84Sro.16CrO 3 in oxidizing (air) and reducing (He/HzO) atmospheres [7. 83]

176 Chapter 7

nonstoichiometry, the defect reaction under oxidizing and reducing conditions can

be expressed as (using the Kr6ger-Vink notation)

oo + 2Crcr = 2Crc~r + Vo + 7 2 (Eq. 7.1)

The equilibrium constant for this reaction is as follows:

X 2 "" K = [Crcr] [Vo] ol/2 �9 o, (Eq. 7 .2)

�9 2 x [Crcr] [Oo ]

or (in terms of mole fraction)

g (1-2x+2/i)zt5 ol/2 (x_2/i)2(3_~i1- o,

Eq. 7.3 can be approximated as

(Eq. 7.3)

K = ~ p1/2 o, (Eq. 7.4) (x-2 i) 2

By solving Eq. 7.4, the concentration of oxygen vacancy can be obtained and is given as

Dl/2 x_ 02 (Eq. 7.5) b = - [(8xKPo~/2 + 1)~/2 _ I I 2 8K

On the other hand, the electrical conductivity a of the material is given by

a = e # p (Eq. 7.6)

where e is the electron charge,/z the mobility, and p the concentration of carriers

(p = x - 26). Thus, the conductivity can be expressed as

a = e t~ o l / 2 [ (8xKPo l /2 + 1)1/2 1] 4K-O,

(Eq. 7.7)

In the high oxygen activity regime, 6 is equal to zero, and Eq. 7.7 becomes

o = e/.tx (Eq. 7.8)

Interconnect 177

In the low oxygen activity regime, Eqs. 7.5 and 7.7 become

Dl14 x - 2~i = �9 o~ (Eq. 7.9)

x (2xK) ~/2

X o = ell(~--K)l/lpoi ? (Eq. 7.10)

Figure 7.7 shows the amount of oxygen deficiency in magnesium-doped

LaCrO3 as a function of oxygen partial pressure and dopant concentration

(computed from Eq. 7.9 and given as solid lines). The experimental data

(obtained from thermogravimetric measurements and given as symbols) are also

shown in the figure. The results indicate good agreement between the experi-

mental data and the defect structure model. At low oxygen activities, the

conductivity of LaCr~_xMgxO3_~ is expected to decrease as a function of fourth root

of oxygen partial pressure. Figure 7.8 shows the computed and experimental

conductivity data as a function of oxygen activities. In general, the figure shows

0.0

-0.1

-0.2

-0.3

-0.4 .._..

(..9 o -0.5 _J

-0.6

-0.7

-0.8

//. ",,oO/" o o

/ / = : / o ////~

/l./ :/ , 1 1 i i o

/ I i + ! i o ....... o .~o~. :' ; a . . . . . . . . 0 . 1 0 M g I i i l o i , . _ o . o , M ~

�9 i i . k - / . . . .

3 -1 1 -9 -7 -5 -3 -1 1 3 5


I

Figure 7. 7. Relative carrier concentration as a function of oxygen partial pressure and dopant content for magnesium-doped LaCrO~ at 1300~ [7.41]

178 Chapter 7

0.6

0.4 u (3

0.2 >: I -- > ~- 0 . 0 u

a z o o -0.2 c o _J

-0.4

- 0 . 6 - ' 3

,"

/2) /,'" /

,,/// o --- 1 300~ . . . . . . 1 2 5 0 ~ �9 ~ 1 2 0 0 ~

-11 -9 -7 -5 -3 -1 1 3 5


Figure 7.8. Conductivity as a function of oxygen partial pressure at various temperatures for LaCro gMgo103 [7. 42]

a relatively good fit between the predicted and the observed curves. Thus, the defect model adequately explains the electrical conductivity behavior of doped

LaCrO3. Diagrams can be made from the model to show the regions of oxygen activity and temperature for which stability of conductivity may be expected. Figures 7.9 and 7.10 show how the conductivity and oxygen vacancy concentra-

tion of magnesium-doped LaCrO3 vary with temperature, oxygen activity, and

dopant content.

0.04 m O E

I-: U 0.03 m LL ILl ra Z tu 0.02 (9 >. X O

0.01

- - , x = 0 . 1 0 .4

. . . . . . . . . . . . . . . . . . cr . . . . . o . . . . . -0.2 0 " o o c ---o x = 0 .2 O

,6"

=

-13 -11 -9 -7 -5 -3 -1 1 3

0.2 r - O 6")

0.0 o

z

--I -0.4

- 0 . 6 " D 0

- 0 . 8 3


Figure 7.9. Oxygen deficiency and electrical conductivity for LaCrl_~MgxO 3 at 1200~ [7. 42]

Interconnect 179

0.05

0

E 0.04

O U. ua 0.03 ct z ua 0.02 (9 >. X O 0.01

. ~ o o ~ o

"%,: -~," / ( o o---- 1L~I0*C ";<( . - . ,/ �9 \ I

" 0 0~

,, x , o o i �9 ~ i i .

-13 -11 -9 -7 -5 -3 -1 1 3 5


I-" 0

0.4 0 0 0

0.2 Z 1:7 c C)

o.o -d_ <

-0.2 " D o

-0.4 3

Figure 7.10. Oxygen deficiency and electrical conductivity at different temperatures for LaCro.gMgo.~03 [7. 42]

The applicability of the defect model is consistent with the p-type

conduction in LaCrO3 via a small polaron mechanism. Various Seebeck

coefficient measurements support this conclusion [7. 54, 7. 57, 7. 64, 7. 76, 7. 77].

From the defect model, several thermodynamic properties for the formation of

oxygen vacancy in doped LaCrO3 have been derived. Some of these data are

given in Table 7.3.

T A B L E 7.3

Thermodynamic Data for Oxygen Vacancy Formation in Doped LaCrO3

Compound AH ~ AS ~ Temperature Range (Nominal Composition) (kJ /mol) (J/mol.K) (K)

Ref.

LaogSr0.tCrO3 -302.6 -100.3 1273 to 1573 [7.58] Lao.7Cao.3CrO 3 -259.2 -71.5 1173 to 1323 [Z62] LaCr0.gMg0.tO 3 -272.0 -80.0 1287 to 1640 [7.41]

180 Chapter 7


LaCrO3 does not interact with other cell materials (YSZ, doped LaMnO3, and Ni/YSZ or NiO/YSZ) at the fuel cell operating temperature (_ 1000~ The measured interdiffusion coefficient of chromium and manganese in the LaCrO3-LaMnO3 system indicates insignificant interactions between the two compounds at 1000~ (An interdiffusion zone of less than 2 #m is formed in 50,000 h [7.87].) The chemical potential diagram for the La-Cr-Zr-O system shows no reactions between LaCrO3 and ZrO2 at 1000~ [7.88]. At higher temperatures (> 1300~ mixtures of LaCrO3 and NiO/YSZ react to form NiCrO4 [7. 89], LaCrO3 and LaMnO3 form a solid solution [7. 75], and LaCrO3 and YSZ may react depending on dopant (and dopant content) in the chromite, as well as firing conditions [7.90].

Chemical interactions become significant when the LaCrO3 interconnect is cofired in contact with other components. The most important interaction phenomenon during LaCrO3 cofiring is the migration of liquid phases in the interconnect material into other cell materials [7. 91- 7. 94]. For example, in firing calcium-doped LaCrO3 laminated to NiO/YSZ anode and LaMnO3 cathode, the Cam(CrO4) n liquid formed in the interconnect migrates into the porous electrodes. A dense region containing poor conducting phases such as CaZrO3, NiCr204 often forms at the NiO/YSZ anode interface [7. 89, 7.95, 7.96]. At the LaMnO3 cathode interface, interdiffusion of chromium, calcium, and manganese occurs. Manganese migrates into the chromite via grain boundaries to form compounds such as (La,Ca)3Mn207 [Z97]. As a result of liquid phase migration during cofiring, the LaCrO3 does not densify. Figure 7.11 shows a micrograph of a cofired laminate of doped LaCrO3 interconnect between LaMnO3 cathode and NiO/YSZ anode [7.92]. As can be seen from the figure, the interconnect layer is porous, although this interconnect can be densified to full density when fired alone under similar conditions. In certain cell designs, LaCrO3 interconnects are

also in contact with YSZ electrolyte. Again, the liquid phase Cam(CrO4) n in the calcium-doped LaCrO3 dissolves in the YSZ boundaries, forming (Ca,Y)ZrO3

compounds [7. 98, 7. 99]. Another important chemical interaction of LaCrO3 involves glass sealants

that are often used for gas sealing in the fiat-plate design. The interaction can alter the properties of the glass (e.g., changes in thermal expansion coefficient, softening temperature, etc.), thus degrading its sealing effectiveness. Under the fuel cell operating conditions, LaCrO3 reacts with alkali silicate glasses to form

Interconnect 181

ANODE INTERCONNECT CATHODE

Figure 7.11. Micrograph of doped La Cr0 3 interconnect cofired between LaMnO 3 cathode and NiO/YSZ anode [7.92]

alkali chromates [7.100]. Alkali earth dopants (e.g., calcium) in LaCrO3 tend to dissolve into glasses [7.100, 7.101]. Thus, at present, long-term use of glasses as SOFC sealants is questionable.


From room temperature (25~ to the orthorhombic/rhombohedral transition temperature (240 ~ to 290~ undoped LaCrO3 has a thermal expansion coefficient of about 6.7 x 10 -6 cm/cm.K [7.33]. Above the transition temperature, the rhombohedral LaCrO3 has a higher thermal expansion coefficient, approximately 9.5 x 10 -6 cm/cm.K (Figure 7.12).

12 t 1.0

~ 0 . 8 - Z

~0 .6 - ,l l

~ 0 . 4 - -

"T"

~ 0 . 2

0 . 0 / , , I 1 1 l

200 400 600 800 1000 TEMPERATURE, ~

Figure 7.12. Thermal expansion of Lal_xCrO 3 (-0.1 <_ x < 0.1) [7.33]

182 Chapter 7

The thermal expansion of undoped LaCrO3 is not influenced by the cation stoichiometry over the range of + 10 mol% lanthanum [7.33]. On the other

hand, the thermal expansion of LaCrO3 can be modified by dopant substitution into the chromite structure. The substitution of aluminum ions for chromium

increases the thermal expansion coefficient [7.23, 7.33, 7. 49]. The substitution

of dopants such as calcium [7. 36, 7.102], strontium [7. 33, 7. 36, 7.103], nickel

[ 7. 26], cobalt [7. 25, 7. 36], and manganese [ 7.104] also increases the coefficient. For example, 13 mol % strontium doping raises the thermal expansion coefficient

of LaCrO3 (A-site deficient) to near that of YSZ [7.103]. On the other hand, the substitution of magnesium does not affect the thermal expansion of LaCrO3

[7.33, Z67]. Substitution on both the A and B sites of LaCrO3 can be used to

adjust the thermal expansion of the material. Iron doping of calcium-doped

LaCrO3 lowers the thermal expansion [7.70]. Similarly, calcium doping of

cobalt-doped LaCrO3 reduces the thermal expansion coefficient [7. 63]. Table 7.4 summarizes thermal expansion coefficient data for undoped and

doped LaCrO3 materials. The thermal expansion of LaCrO3 varies with dopant content. An example of such variation is given in Table 7.5 for strontium-doped

LaCrO3 [7. 33].

TABLE 7.4

Thermal Expansion Coefficients of Undoped and Doped LaCrO3

Composition (nominal)

Thermal expansion coefficient (10 .6 cm/cm-K)

Ref.

LaCrO3 9.5 [7.33] Lao.9Sro. l CrO3 10.7 [ 7. 33] Lao.sCao.2CrO 3 10.0 [7.90] LaCro.gMgo.103 9.5 [7.33] LaCro.9Coo.lO 3 13.1 [7.63] LaCro.gNio. 1 O3 10.1 [ 7. 26] LaCro.7Mgo.o5Alo.2503 9.8 [ 7. 23] Lao.aCao.2Cro.9COo.lO 3 11.1 [ 7. 63] Lao.7Cao.3Cro.ssCoo.o5Feo.osNio.osO3 10.8 [ 7.105]

Interconnect 183

TABLE 7.5

Thermal Expansion Coefficient of Strontium-Doped LaCrO3 [7.33]

Composition (nominal)

Thermal expansion coefficient (10 -6 cm/cm.K)

LaCr03 9.5

Lao.98Sro.o2Cr03 10.2 Lao.95Sro.o5 Cr03 10. 9 Lao.9oSro.loCr03 10.7 Lao.ssSro.15Cr03 10.8 Lao.soSro.2oCrO 3 11.1

7.2.7 Sinterability

It is well known that LaCrO3, like other chromium-containing oxides, is

difficult to sinter to high densities under high oxygen activity conditions. At high

oxygen activity levels, the trivalent chromium ion in LaCrO3 can be converted

to oxidation states whose oxides are unstable. As a result, LaCrO3 appreciably

volatilizes chromium oxides in oxidizing atmospheres. Because of the volatiliza-

tion, the predominant form of mass transport is a gaseous process mainly

involving the chromium oxide compounds. This leads to an evaporation/conden-

sation mechanism of sintering in which the chromium ions evaporate in some

form from the surface of the original grains and condense out onto the points of irregular contact, which are the areas of maximum surface energy. Because neck

growth is accomplished without the transport of the material from the bulk of the

grains, little densification or pore removal occurs. The sinterability behavior of

LaCrO3 has been correlated to chemical potential diagrams and vapor pressures

of chromium oxides calculated for La-Cr-O systems [7.106, 7.107]. The poor

sinterability of LaCrO3 in air or oxidizing atmospheres is ascribed to the

formation of a thin layer of Cr203 from CrO 3 gases at the interparticle neck

during the initial stage of sintering:

3 0 = ~La203 + CrO3(g ) LaCrO3(s) + g 2(g) (s) (Eq. 7.11)

184 Chapter 7

followed by

3 0 CrO3(g) = ICrEOa(s) + ~ 2(g) (Eq. 7.12)

To sinter LaCrO3 to high densities, firing temperatures greater than 1600~ under low oxygen partial pressures have generally been used (Figure 7.13) [7.108-7.111]. Control of oxygen partial pressures close to those specified by the Cr/Cr203 phase boundary suppresses grain growth and allows maximum densification.

Several approaches have been investigated to enhance the sinterability of LaCrO3 at firing temperatures below 1600~ in oxidizing atmospheres. These approaches involve use of highly reactive powders, nonstoichiometric materials, dopants, sintering aids, and processing techniques.

(i) Highly reactive powders: Improvements in the densification of LaCrO3 can be obtained when highly reactive (high-surface-area) powders are used. Uniform, high-surface-area powders without agglomeration lower the temperature required to densify LaCrO3. For example, a synthesis method called the nitrate pyrolysis process produces chromite powders (with substantially improved reactivity) that can be sintered at lower temperatures [7.112]. LaCrO3 powders produced by the glycine/nitrate method contain small primary particles and small, soft agglomerates [7.13, 7.113, 7.114]. The powders are highly reactive, sintering to near-full density at 1550~ in air. In general, there appears to be an upper density limit to which high-surface-area powders can be sintered in air.

1 O 0

>- m

z ,,, 90 a l

< ( j m

uJ 80 n- o LM

:E

70

n- i l l a.

60

I ! I I I i O 1 7 2 0 ~

13 1 6 2 5 ~

i / - / b

! l I I ! 1 0 2 4 6 8 10 12 14

--LOG OXYGEN PARTIAL PRESSURE, atm (1.01 x 10 ~ Pa)

Figure 7.13. Sintered density of Lao.s4Sro.16CrO 3 as function of oxygen partial pressure during sintering [7.108]

Interconnect 185

(ii) Nonstoichiometry: Nonstoichiometry can influence sintering of LaCrO3 in air. The sinterability of LaCrO3 has been found to vary significantly with relatively small changes in the material stoichiometry. Chromium nonstoichiometry is particularly effective in enhancing sintering. For strontium-

doped LaCrO3, higher densities have been obtained for samples with chromium deficiency or excess [7.114, 7.115]. Small chromium deficiency markedly increases the densification of strontium-doped LaCrO3 in air (Figure 7.14). For calcium-doped LaCrO3, chromium-deficient samples show high densification in air irrespective of calcium content [ 7. 44, 7. 96, 7.102, 7.116, 7.117]. Chromium-

deficient, calcium-doped LaCrO3, e.g., Lao.TCao.3Crl_yO 3 (0 < y < 0.02), can be sintered to 94% of theoretical density in air at 1300~ whereas poor densifi-

cation is observed for the chromium-stoichiometric compound, La0.7Ca0.3CrO3

(i.e., y = 0). Chromium deficiency in calcium-doped LaCrO3 causes precipitation of La20 (calcium content < 0.15) or CaO (calcium content > 0.15). CaO precipitation has no significant effect on the chemical stability of LaCrO3. Precipitated oxides act as chromium getters, thus minimizing the vapor pressure of chromium oxides [7.106]. However, the most important effect of chromium

deficiency is the formation of liquid phases such as Cam(CrO4) n in calcium-doped materials [ 7. 96, 7.117- 7.120]. Cam(CrO4) n incongruently melts and assists mass transport during firing, resulting in enhanced densification of the LaCrO3. Figure 7.15 shows a schematic diagram of microstructural changes proposed for chromium-deficient calcium-doped LaCrO3 at different firing temperatures.

1.0

0.9 >. I - - O0 7 0.8 LU

UJ > 0.7

w

0.6

! v v ., a,-',,F. ~ % ' - ~ v

! i

r * /

/

/ !

/ / Q.~ STOICHIOMETRIC _

C )_~ 1 i ' ~ / ~D CHOMIUM DEF IC IENT

0 . 5 1 1 1 1 t

0.0 0.1 0.2 0.3 0.4 0.5 0.6

CALCIUM CONTENT, tool

Figure 7.14. Sinterability of stoichiometric, versus chromium-deficient, calcium-doped LaCrO~ (1600~ in air) [7.102]

186 Chapter 7

/ - , ~ 1273 K

\ ~ y ,L_j

~-----J k.______jJ

/ ' - - - -~ ~-> 1273 K

Cam(CrO4) n (rn > n)

' "> 1573 K

Ca rich region Cam(CrO4)" (m > n)

1873 K

CaO

Figure 7.15. Microstructural changes proposed for chromium-deficient, calcium-doped LaCrOs at different firing temperatures [7.96]

(iii) Dopant: One way to enhance the fundamental sinterability of

LaCrO3 is to dope it with an ion that will increase the concentration of vacancies

present, allowing bulk or grain-boundary mass transport. To date, no such dopants have been found for LaCrO3 (which do not also deleteriously affect other properties of the material). Dopants can also improve the sinterability of LaCrO3 by reducing the volatility of chromium from the surface during sintering. For example, the substitution of aluminum for chromium has been found to substantially decrease the volatilization of magnesium-doped LaCrO3 [7. 49]. The most important influence of dopants on the sinterability of LaCrO 3 is the

formation of a transient liquid phase during firing which enhances densification

by liquid-phase sintering. The enhanced sinterability of LaCrO3 doped with

strontium may be due to the formation of SrCrO4 melt [7.13, 7.121]. The

densification of calcium,nickel-doped LaCrO3 can be attributed to the presence

of calcium,chromium-rich melts [7. 50]. Similarly, calcium,cobalt-doped LaCrO3 can be sintered to full density at temperatures below 1400~ in air due to the

presence of a transient liquid phase [ 7. 63, 7.122, 7.123]. Figure 7.16 shows an example of liquid phases formed during sintering

of (La0.6Ca0.a)0.95CrO 3. A semiviscous liquid can be seen spreading along the solid/solid contacts and wetting the grains. In general, at the onset of liquid formation, the chromite sinters rapidly due to the capillary force exerted by the

Interconnect 187

Figure 7.16. Calcium-doped LaCrO 3 surface quenched from sintering temperature of 1300 ~ C (courtesy of University of Missouri-Rolla)

wetting of the solid particles. Both the rate and the degree of densification

depend on solid solubility in the liquid. Thus, a small amount of liquid phase might be insufficient to promote densification. In the Lal_xCaxCr~_yCOyO 3 system, at least 20 mol% calcium and 10 mol% cobalt are required to densify the

material at temperatures below 1400~ in air [7.122]. The composition of a

transient liquid phase, and thus its melting point, are dependent on the amount

of dopant substitution. The liquid tends to disperse along the grain boundaries

and form solid solution with the LaCrO3. As a result, no second phases are observed in the microstructure. As discussed earlier, the main problem with a transient liquid phase is that, in cofiring LaCrO3 with other SOFC components, the liquid phase tends to diffuse into other components, and the chromite does not densify.

(iv) Sintering aids: Introduction of a liquid-phase sintering aid--a second phase with a significantly lower melting point--encourages densification of

LaCrO3 in oxidizing environments. In liquid-phase sintering, the prerequisites

are that a significant amount of liquid phase must be present, the solid must be

soluble in the liquid, and finally the solid particles must be completely wetted by

the liquid. The functions of liquid-phase sintering aids are to pull the particles

closer together by surface tension forces and to enhance diffusion of the solid

phase to the points of particle-to-particle contact to promote material transport via

a solution/precipitation process. The most common sintering aid for LaCrO3 and

many other oxides is B203. Other sintering aids, such as low-melting oxide

eutectics and LaF3, YF3, and MgF2 up to 8 to 10 wt % have been used to densify

188 Chapter 7

LaCrO3 in air at temperatures below 1400~ [7.124-Z126]. Similar to transient liquid sintering, the main difficulty involving sintering aids is that they may

migrate to other cell components, causing elemental migration and morphological changes.

(v) Processing techniques: In addition to the main approaches described

above, the densification of LaCrO3 can be enhanced by modifying or tailoring processing techniques. For example, microwave processing of LaCrO3 has

shown promise of lowering the sintering temperature [Z127]. Firing LaCrO3

between Cr203 plates improves the densification of strontium-doped LaCrO 3

[7.128]. Hot isostatic pressing at 1400~ under 100 MPa for 3 h produces

sintered LaCrO3 of 97 % theoretical density [7.129].

7.2.8 Gas permeability

For efficient operation, LaCrO3 interconnects must be impervious to fuel and oxidant gases, from room temperature to operating temperatures (600 ~ to

1000~ The permeation of oxygen through LaCrO3 has been found to be very

limited [7.67]. At 1000~ with 1 atm (1.01 x 105 Pa) oxygen pressure on one

side and'vacuum on the other side of LaCrO3, the oxygen permeation rate is

about 9.6 • 10 -13 mol/cm2.s. With air on one side and a H20/H2 mixture on the other, the rate is 1.9 • 10 -~ mol/cm2.s. This rate is insignificant in the SOFC

operation. Oxygen permeation through LaCrO3 can be estimated from oxygen vacancy diffusion in the material [7.130-7.132]; for example, with a vacancy diffusion coefficient of 10 .7 cm2/s at 1000 ~ the oxygen permeation current density is estimated to be on the order of 1 mA/cm 2 in an oxygen partial pressure gradient of 0.2 • 105 (oxidant) to 1 x 10 -l~ Pa (fuel) across calcium-doped

LaCrO3 [7.131]. Hydrogen permeation through LaCrO3 has also been measured

[7.72].

7.3 OTHER MATERIALS

As mentioned earlier, only very few oxides can be considered for SOFC

interconnect applications [7. 73, 7.133]. Early SOFCs used doped CoCr204 as

interconnect material [7.134]. This compound was later replaced by LaCrO3.

Interconnect materials based on glass composites have also been investigated

[ 7.135, 7.136]. Recently, YCrO3 has been proposed as an alternative to LaCrO3.

The properties of this compound have been evaluated [Z57, 7.63, Z137-7.139].

Interconnect 189

Although YCrO3 is less refractory than LaCrO3, the yttrium compound has the

advantage of greater stability (less interaction, no hydration) in the fuel cell

environments. To date, only preliminary work has been conducted, and no

SOFC stacks have been tested with YCrO3 interconnects.

In addition to oxide ceramics, metals may be used as SOFC interconnect,

especially for the flat-plate design [7.140-7.148]. The advantages of metallic

interconnects include better chemical compatibility, improved mechanical

properties, and lower cost. The primary concerns regarding metallic intercon-

nects are thermal expansion mismatch with other cell components, and long-term

instability (oxidation, corrosion) under fuel cell operating conditions. Several

high-temperature alloys based on chromium and nickel have been evaluated;

some examples are Inconel 600 [7.146], Inconel 601 [7.148], Hastelloy X

[7.148], and HA-230 [7.142]. Many commercial alloys have proved inadequate

due to incompatibility in thermal expansion coefficients. Also, excessive

oxidation of the alloys has been observed under fuel cell operating conditions,

especially at 1000~ [7.149]. Table 7.6 summarizes the results of the evaluation

of selected alloys for interconnect applications [7.150].

TABLE 7.6

Results of Evaluation of Selected Alloys for Interconnect Applications [7.150]

Physical properties Long-term stability Fabricability (1000~

�9 ~ ~ ~ ~'~ ~ . ~ ~ ~'r. . . . . - o

~- ~ ~ o ~ < ~ z ~ ~.

AC66

HA 230

HA 214

2 1 2 1 4 2 1 3 1 1 1 1

2 1 2 1 2 2 2 2 1 1 1 1

2 1 2 4 4 1 1 1 3 1 1 1

Incoloy M A 9 5 6 1 1 1 4 2

C r - C o alloys 1 3 1 2 1

Cr-Fe alloys 1 2 1 2 1

1 1 1 1 1 1 1

4 2 4 1 3 4 1

2 4 4 1 2 2 1

1 = good; 2 = adequate; 3 -- poor; 4 = inadequate

190 Chapter 7

Certain oxide dispersion strengthened chromium alloys have been developed to provide a better thermal expansion match and better oxidation resistance in fiat-plate SOFCs [7.142,Z149,Z150]. For example, chromium or chromium alloys with the addition of 0.4 to 1 wt% Y203 or La203 have thermal expansion coefficients very similar to that of YSZ (Figure 7.17) [7.149]. Figure 7.18 is a photograph of a metallic interconnect (with machined gas channels) fabricated with Cr5FelY203 alloy (chromium with 5 wt % iron and 1 wt % Y203) .

Over long-term operation at 1000 ~ chromium from chromium-based alloy interconnects may migrate into the perovskite cathode, causing cell performance to degrade. Surface treatment of chromium-based alloys significantly reduces chromium migration [7.151].

When exposed to oxidizing atmospheres at 1000 ~ commercial high- temperature alloys become oxidized, and their conductivity decreases with time. Coating can be used to protect metallic interconnects against oxidation [7.152]. LaCrO3 coatings are of particular interest and can be deposited on metallic interconnects by various methods [7.5, 7. 6, 7.153, 7.154]. Other perovskite materials such as LaCoO3 are also suitable for use as coatings [7.148]. Cermets consisting of an alloy and A1203 have been proposed for SOFC interconnect

applications [Z155].

E o E 16- o

? O t . -

~ 12. 0

Z < a.. 8" • w ..I < ~ 4 uJ "r k-

o o

COMMERCIALLY IAVAILABLE SUPERALLOYS

HA 2 3 0 C r 5 F e i Y 2 0 a ]

I I

..~ I YSZ

J J J / [ CrO.4LazOa I

�9 ' - ~ �9 ' ~ �9 1 '

200 400 600 800 1000 TEMPERATURE, ~

Figure 7.17. Comparison of thermal expansion of commercial and oxide dispersion strengthened chromium alloys with that of YSZ [7.149]

Interconnect 191

Figure 7.18. Photograph of a metallic interconnect plate made of Cr5FelY203 alloy (courtesy of Siemens)

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7.1

7.2

7.3

7.4

7.5

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Interconnect 193

7.29

7.30 7.31

7.32

7.33 7.34

7.35 7.36

7.37

7.38

7.39 7.40

7.41

7.42

7.43 7.44

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7.50 7.51 7.52

7.53

7.54

7.55

7.56

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7.59 7.60 7.61 7.62

7.63

7.64 7.65

7.66

7.67

7.68 7.69 7.70

7.71

7.72

7.73

7.74 7.75 7.76

7.77

7.78

7.79

7.80

7.81

7.82

7.83

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Interconnect 195

7.84

7.85

7.86 7.87

7.88

7.89 7.90

7.91

7.92

7.93

7.94

7.95

7.96

7.97 7.98 7.99

7.100

7.101

7.102

7.103

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196 Chapter 7

7.104 7.105 7.106

7.107

7.108

7.109

7.110 7.111

7.112 7.113

7.114

7.115 7.116

7.117

7.118

7.119 7.120

7.121

7.122

7.123 7.124

7.125

7.126

7.127

J. Palma and C. Pascual, see Ref. 7.2, p. 537. P. Gordes and N. Christiansen, see Ref. 7.3, p. 414.

H. Yokokawa, N. Sakai, T. Kawada, and M. Dokiya, see Ref. 7.87, p. 118.

H. Yokokawa, N. Sakai, T. Kawada, and M. Dokiya, J. Electrochem. Soc., 138 (1991) 1018.

L. Groupp and H.U. Anderson, J. Am. Ceram. Soc., 59 (1976) 449.

M. Berberian, I.B. Cutler, and R.W. Ure, in Proceedings of the Fifteenth Symposium on Engineering Aspects of Magnetohydrodynamics, May 24-26, 1976, Philadelphia, PA, GE Space Sciences Laboratory, 1976, p. IV.3.1.

A.M. George, J.R. Pai, and V.K. Rohatgi, see Ref. 7.109, p. II.1.1. F.J. Brodmann and P.E.D. Morgan, in Conference on High Temperature Sciences Related to Open-Cycle, Coal-Fired MHD Systems, April 4-6, 1977, Argonne National Laboratory, Argonne, IL, Report No. ANL-77-21, Argonne National Laboratory, Argonne, IL, 1977, p. 135. M.R.F. Kuntz, see Ref. 7.3, p. 301.

J.L. Bates, L.A. Chick, and W.J. Weber, in Proceedings of the First Annual Fuel Cells Contractors Review Meeting, May 2-3, 1990, Morgantown, WV, W.J. Huber

(ed.), Report No. DOE/METC-89/6105, U.S. Department of Energy, Washington, DC, 1989, p. 54.

J.L. Bates and L.A. Chick, see Ref. 7.91, p. 159.

L.A. Chick, J.L. Bates, and G.D. Maupin, see Ref. 7.2, p. 621. N. Sakai, T. Kawada, H. Yokokawa, and M. Dokiya, see Ref. 7.2, p. 629.

M. Mori, N. Sakai, T. Kawada, H. Yokokawa, and M. Dokiya, Denki Kagaku, 59 (1991) 314. L.A. Chick, T.R. Armstrong, D.E. McCready, G.W. Coffey, G.D. Maupin, and J.L. Bates, see Ref. 7.3, p. 374. L.A. Chick and J.L. Bates, see Ref. 7.95, p. 563. J.D. Carter, V. Sprenkle, M.M. Nasrallah, and H.U. Anderson, see Ref. 7.3, p. 344.

D.B. Meadowcroft, in Proceedings of the International Conference on Strontium Containing Compounds, June, 1973, Halifax, Nova Scotia, Canada, T.J. Gray (ed.),

Atlantic Industrial Research Institute, Halifax, Nova Scotia, Canada, 1973, p. 119. R. Koc and H.U. Anderson, J. Eur. Ceram. Soc., 9 1992) 285.

R. Koc and H.U. Anderson, see Ref. 7.12, p. 659. B.K. Flandermeyer, J.T. Dusek, P.E. Blackburn, D.W. Dees, C.C. McPheeters,

and R.B. Poeppel, in 1986 Fuel Cell Seminar Abstracts, October 26-29, 1986, Tucson, AZ, Courtesy Associates, Washington, DC, 1986, p. 68. B.K. Flandermeyer, R.B. Poeppel, J.T. Dusek, and H.U. Anderson, U.S. Patent

No. 4749632, June 7, 1988.

C. Milliken and A. Khandkar, see Ref. 7.13, p. 361.

M.A. Janney and H.D. Kimrey, U.S. Department of Energy Report No. CONF- 900546--5, 1990, U.S. Department of Energy, Washington, DC, 1990.

Interconnect 197

7.128 7.129

7.130 7.131

7.132

7.133

7.134

7.135

7.136

7.137

7.138

7.139

7.140 7.141

7.142

7.143

7.144 7.145

7.146

7.147

7.148

7.149

L.W. Tai and P.A. Lessing, J. Am. Ceram. Soc., 74 (1991) 155. S. Song, M. Yoshimura, and S. Somiya, Yogyo-Kyokai-Shi, 90 (1982) 484.

I. Yasuda and T. Hikita, see Ref. 7.3, p. 354. H. Yokokawa, T. Horita, N. Sakai, B.A. van Hassel, T. Kawada, and M. Dokiya,

see Ref. 7.3, p. 364. T. Kawada, T. Horita, N. Sakai, B.A. van Hassel, H. Yokokawa, and M. Dokiya,

ISSI Lett., 4 (2) (1993) 6. W. Baukal and W. Kuhn, J. Power Sources, 1 (1976/1977) 91.

C.C. Sun, E.W. Hawk, and E.F. Sverdrup, J. Electrochem. Soc., 119 (1972) 1433.

P.G. Russell, H.S. Isaacs, A.C.C. Tseung, and S. Srinivasan, in Proceedings ofthe Workshop on High Temperature Solid Oxide Fuel Cells, May 5-6, 1977, Brookhaven National Laboratory, Upton, NY, H.S. Isaacs, S. Srinivasan, and I.L. Harry (eds.),

Report No. BNL 50756, Brookhaven National Laboratory, Upton, NY, 1978, p. 96.

P.G. Russell, H.S. Isaacs, A.C.C. Tseung, and S. Srinivasan, J. Appl. Electro- chem., 11 (1981) 197. T. Negas, W.R. Hosler, and L.P. Domingues, in Proceedings of the Fourth International Meeting on Modem Ceramics Technologies, May 28-31, 1979, Saint- Vincent, Italy, P. Vincenzini (ed.), Elsevier/North Holland, New York, 1980, p.

993. W.J. Weber, J.L. Bates, C.W. Griffin, and L.C. Olsen, in Proceedings of the Symposium on Defect Properties and Processing of High-Technology Nonmetallic Materials, December 2-4, 1985, Boston, MA, Y. Chen, W.D. Kingery, and R.J.

Stokes (eds.), Materials Research Society, Pittsburgh, PA, 1986, p. 235.

G.F. Carini II, H.U. Anderson, D.M. Sparlin, and M.M. Nasrallah, Solid State lonics, 49 (1991) 233. W. Wersing, E. Ivers-Tiffee, M. Schiessl, and H. Greiner, see Ref. 7.87, p. 33.

E. Ivers-Tiff6e, W. Wersing, M. SchieB1, and H. Greiner, Ber. Bunsenges. Phys. Chem., 91 (1990) 978. W. Drenckhahn and H.E. Vollmar, see Ref. 7.95, p. 419. Y. Akiyama, N. Ishida, T. Yasuo, S. Taniguchi, S. Murakami, T. Saito, and N.

Furukawa, see Ref. 7.95, p. 603. Y. Akiyama, T. Yasuo, N. Ishida, S. Taniguchi, and T. Saito, see Ref. 7.3, p. 724. H. Takagi, H. Taira, A. Shiratori, S. Kobayashi, Y. Sugimoto, S. Sakamoto, and

K. Tomono, see Ref. 7.3, p. 738. Y. Sato, H. Nakagawa, H. Mihara, S. Kosuge, H. Tsuneizumi, and E. Morishige,

U.S. Patent No. 5049458, September 17, 1991.

K. Krist and J.D. Wright, see Ref. 7.3, p. 782.

T. Kadowaki, T. Shiomitsu, E. Matsuda, H. Nakagawa, H. Tsuneizumi, and T.

Maruyama, Solid State Ionics, 67 (1993) 65.

W.J. Qudakkers, H. Greiner, and W. K6ck, in Proceedings of the First European Solid Oxide Fuel Cell Forum, October 3-7, 1994, Lucerne, Switzerland, U. Bossel

(ed.), European SOFC Forum Secretariat, Baden, Switzerland, 1994, p. 525.

198 Chapter 7

7.150 7.151

7.152

7.153

7.154

7.155

H. Greiner, private communication. Y. Akiyama, S. Taniguchi, T. Yasuo, M. Kadowaki, and T. Saitoh, J. Power Sources, 50 (1994) 361.

M. Dokiya, T. Horita, N. Sakai, T. Kawada, H. Yokokawa, B.A. van Hassel, and C.S. Montross, see Ref. 7.82, p. 33.

H. Tenmei, H. Michibata, T. Namikawa, and Y. Yamazaki, Denki Kagaku, 58 (1990) 1072.

H. Konno and R. Furuichi, in High Temperature Corrosion of Advanced Materials and Protective Coatings, Y. Saito, B. Onay, and T. Maruyama (eds.), Elsevier Science Publishers, Amsterdam, The Netherlands, 1992, p. 177. H. Seto, T. Miyata, A. Tsunoda, T. Yoshida, and S. Sakurada, see Ref. 7.3, p. 421.

Chapter 8

ELECTRODE REACTION

8.1 GENERAL

The operation of a SOFC involves two primary electrode reactions: the

oxidation of the fuel at the anode and the reduction of the oxidant at the cathode.

In theory, any gases capable of being electrochemically oxidized and reduced at

the SOFC operating temperature can be used as fuel and oxidant. However,

hydrogen is presently the most common fuel for use in SOFCs. Hydrogen has

high electrochemical reactivity and can be derived from common fuels such as

hydrocarbons, alcohol, or coal. Oxygen is the most common oxidant for SOFCs

since oxygen is readily and economically available from air. Thus, the overall

reactions at the SOFC anode and cathode are as follows:

Anode H 2 + O 2--- H20 + 2e- (Eq. 8.1)

Cathode 1/~O 2 + 2e -= 0 2- (Eq. 8.2)

The difference between the thermodynamic (equilibrium) potentials of the

electrode reactions determines the cell reversible (open-circuit) voltage. The

reversible voltage, Er, is the maximum voltage that can be achieved by a SOFC

under specified conditions of temperature and gas composition. The voltage of

an operating cell, E, is always lower than Er. As the current is drawn from the fuel cell, the cell voltage falls, due to internal resistance and polarization losses.

Thus, the voltage of an operating cell is given as

E - E r - I n i - ('Oa + ~ c ) (Eq. 8.3)

In the equation above, I g i is the internal resistance or ohmic loss (I is the current,

R i the internal resistance of the cell), and ~a and ~7~ the anode and cathode

polarization, respectively. Ohmic losses result from the resistance of the

electrolyte and other cell components. Polarization (overpotential) losses are

200 Chapter 8

associated with the electrochemical reactions taking place at the interface between

the electrodes and the electrolyte. The kinetics of the electrode reactions (the

oxidation of hydrogen and the reduction of oxygen) play a critical role in

determining polarization losses in SOFCs. In practical applications, SOFCs may use gaseous mixtures that contain,

in addition to hydrogen, CO, CO2, and H20, e.g., coal gas and natural gas. Because of the high operating temperature (600 ~ to 1000~ of the SOFC cell, the presence of CO and CO2 in the fuel does not poison the anode reaction. (In fact, CO functions as a fuel in YSZ-based SOFCs, and in the presence of H20, the favorable path for the oxidation of CO is via the generation of hydrogen by the shift reaction.) On the other hand, these gas mixtures (natural gas, coal gas)

may contain sulfide impurities. The presence of significant levels of sulfur may cause an unacceptable loss of cell voltage. A thorough knowledge of the level

of sulfide and its influence on the anode reaction is essential in order to operate SOFC systems efficiently.

The high operating temperature of the SOFC also permits the fuel cell to reform conventional hydrocarbon fuels internally. The commonly used anode material, nickel/YSZ cermet, is a suitable catalyst for the reforming reactions. Internal reforming in a SOFC is expected to simplify the overall fuel cell system

design (because of elimination of the external reformer), hence increasing system reliability. An internal reforming SOFC system promises low capital and

operating costs.

8.2 REACTIONS AT ANODE

The main reaction at the SOFC anode is the electrochemical oxidation of hydrogen (and CO) on the electrode material (commonly nickel metal) in contact

with YSZ electrolyte. Other important reactions at the anode are those of sulfur impurities and the reforming of hydrocarbon gases.

8.2.1 Electrochemical oxidation ofhydrogen

The hydrogen oxidation at nickel metal in contact with YSZ produces

water according to the following reaction (using Kr6ger-Vink notation):

H 2 + Oo = H20 + Vo + 2e- (Eq. 8.4)

Electrode Reaction 201

Several mechanisms have been proposed for this reaction; however, the exact

nature of the reaction kinetics has still not been well established [8.1, 8.2]. The

hydrogen oxidation at the nickel/YSZ cermet anode is influenced by two

important factors, namely, gas composition and electrode microstructure. For

example, the presence of H20 and CO in the gas and the morphology of the

nickel and YSZ have been shown to have a dominant effect on the polarization

behavior of the anode. Other factors, e.g., electrode thickness, may also

influence the hydrogen reaction, thus anode performance [8.3, 8.4].

(i) Reaction mechanism: It is generally accepted that nickel metal plays a catalytic role in the oxidation. The influence of electrode materials on the

electrochemical characteristics of the hydrogen reaction has been demonstrated

at SOFC operating temperatures. For example, the catalytic activity of several

metals at 1000~ has been shown to decrease in the order: iron > cobalt >

nickel > molybdenum [8.5]. Overpotential measurements at 800~ have

indicated that the nature of metal electrodes strongly affects the polarization of

the hydrogen reaction [8.6, 8. 7].

There is also evidence suggesting that the YSZ plays an important

electrocatalytic role in the hydrogen reaction. The catalytic effect of the YSZ is

supported by the observation that an increase in electronic conductivity in the

electrolyte enhances the reaction rate [8.8]. The presence of electronic conductivity promotes the spreading of the reaction zone around the three-phase

boundary or triple contact point (gas/electrode/electrolyte), leading to an increase

in available reaction sites [8.9] (Figure 8.1). Thus, modification of the YSZ

! 2 /H2 \ ~ / ' H20

0 2. 0 2 -

Figure 8.1. Spreading of reaction zone around three-phase boundary due to electronic conductivity in the electrolyte [8.12]

202 Chapter 8

electrolyte surface (to introduce electronic conductivity) improves performance

of the anode by reducing its overpotential [8.8,8.10-8.12]. Substitution of the

YSZ in the anode with a mixed conducting material such as CeO2 lowers anodic

polarization resistance [8.6, 8. 7, 8.13-8.15].

At present, a definite mechanism for the hydrogen oxidation at Ni/YSZ

anodes has yet to be established, although the electrochemically active site appears to be the three-phase boundary where the gas, electrode, and electrolyte

meet [8.16-8.18]. To date, several reaction schemes have been suggested. A

mechanism proposed early for the oxidation is based exclusively on the catalytic activity of nickel surface. This mechanism involves the adsorption of hydrogen

on nickel followed by the electrochemical reaction (between adsorbed hydrogen

and oxygen ions). However, the hydrogen reaction appears to be more complex.

For example, AC impedance spectra of the oxidation generally show not a simple

semicircle but two semicircular arcs (Figure 8.2) [8.19-8.22]. Based on AC impedance results, a proposed mechanism suggests possible electrochemical

reaction of hydrogen on the nickel and formation and reaction of hydroxide ions

on the YSZ surface [8.19]

H E = 2Had,N i (Eq. 8.5)

4-

2Haa,N i = 2Had,N i + 2e-

+ 2- 2Had, Ni + 2Oysz = 2OHysz

2 -

2OH~,sz = H20 + O y s Z

(Eq. 8.6)

(Eq. 8.7)

(Eq. 8.8)

"E 0.2 o

z i i i

z 0 n

: ; 0.1 1 kHz 0 (J >-

0 ~ < ? . 1 Hz

0.0 -- 0.0 0.1 0.2 0.3

REAL COMPONENT, Q.cm = 0 . 4

Figure 8.2. Impedance spectrum of a Ni/YSZ anode at I O00~ in H2/H20 atmosphere [8.2]


As shown later, the electrochemical behavior of hydrogen oxidation is significantly influenced by the presence of H20 or oxygen-containing molecules in the fuel. Water may provide adsorbed oxygen species on the nickel, and the

kinetics of the hydrogen oxidation is probably related to the oxygen activity on

the metal surface [8.21]. In this case, the proposed mechanism is as follows:

Dissociative decomposition of H20

H 2 0 = Oad,Ni + H2 (Eq. 8.9)

Dissociative adsorption of hydrogen

H 2 = 2Had,VSZ (on YSZ surface) (Eq. 8.10)

H 2 - 2 H a d , N i _ O (on Oad-Covered nickel surface) (Eq. 8.11)

Charge transfer reaction

Oo = O a 2. vsz § Vo (Eq. 8 12)

2- 2- Oad,YSZ + Oad,N i -- Oad,YSZ + Oad,N i

(Eq. 8.13)

2- Oad,N i = Oad,N i + 2e- (Eq. 8.14)

2Had,VSZ + Oad,YS Z = H20 (Eq. 8.15)

Had, Ni_ O + Oad, Ni = H 2 0 (Eq. 8.16)

In the mechanisms above, both the nickel and YSZ surface comribute to the electrode process, i.e., the triple contact point as the electrochemical reactive site. On the other hand, a mechanism based on the sole catalytic activity of the electrolyte surface has also been advanced. The supporting experimental

observations for this mechanism include the following: the rate of oxidation of

hydrogen at a metal anode with solid electrolyte is independent of the electrode

material; activation enthalpies of the reaction are also independent of the

electrode material; and blackening of the electrolyte by electrolysis leads to

marked enhancement in the reaction rate [8.23-8.27]. It has been concluded that

the major reaction steps occur at active sites on the electrolyte surface. The electrochemical reactive sites (ERSs) are hypothesized to be oxygen vacancies

204 Chapter 8

(VERs) with electrons migrating along the electrolyte surface to or away from these active sites. The following reaction mechanism has been postulated:

H 2 + 2VER s = 2HER s (Eq. 8.17)

O o + 2HER s = H20 + Vo + 2VER s + 2ezro= (Eq. 8.18)

2ezro~ = 2eNi (Eq. 8.19)

In general, the overpotential of the hydrogen oxidation at nickel/YSZ anodes obeys the Tafel equation [8.28-8.30]

RTI s 11 = (Eq. 8.20) 2F Jo

where r/is the overpotential, R the gas constant, T the temperature, j the current

density, and Jo the exchange current density. The overpotential characteristics of the hydrogen reaction are strongly influenced by the formation of oxide on the nickel surface [8. 31-8. 34].

(ii) Effect of gas composition: The overpotential of the electrochemical oxidation of hydrogen is somewhat independent of hydrogen content in the fuel

but strongly influenced by the presence of water. The oxidation of dry hydrogen occurs with significant overpotential or interfacial resistance [8.19, 8. 20], whereas the overpotential is much smaller for hydrogen/water mixtures [8.19, 8.20, 8.35, 8.36]. A few mole percent of H20 in hydrogen can dramatically reduce the anode interfacial resistance (Figure 8.3) [8.19,8.20]. At high H20 concentrations, the interfacial resistance increases with increasing H20/H2 ratio [8.37]. The hydrogen electrode is thus expected to exhibit a minimum interfacial

resistance at a certain H20/H2 ratio. The effect of H20 on the anode polarization may be related to the oxygen partial pressure, and the correlation suggests strong dependence of overpotential on activation of oxygen ions [8. 7, 8.13]. The role

of H20 in the hydrogen oxidation is not clearly understood, although H20 is

believed to adsorb on the surface of the YSZ electrolyte and broaden reactive

sites around the triple contact point, resulting in increased reaction rate [8.8,8.28].

(iii) Effect of electrode microstructure: The anode microstructure can be the dominant factor in determining the overpotential of the hydrogen reaction at

nickel/YSZ electrodes [8.38]. Two microstructure characteristics of the anode,

the surface area and size of the nickel and YSZ particles, play a critical role.


"E o

u~

z

I-- oo I

oo uJ

_.i

1.2

0 .8

s u . rr" I.U p- Z m

0 .4

0 ,0 1 0 .4 1 0 .3 1 0 .2 1 0 1

H20 PARTIAL PRESSURE, atm (1.01 x 10 s Pa)

Figure 8.3. Interfacial resistance of Ni/YSZ anode as a function of 1120 partial pressure at 1000~ [8.20]

This can be seen from Figures 8.4 and 8.5 showing the effect of the surface area

(and particle size) of nickel and YSZ on the anode interfacial resistance [8.39,

8.40]. Thus, the polarization behavior of the hydrogen reaction is strongly

dependent on preparation conditions such as starting material properties and firing temperatures. Figure 8.6 shows, as an example, the influence of firing

temperature on the overpotential of the hydrogen oxidation at nickel/YSZ anodes

prepared by slurry coating [8.29,8.41].

u~ (3 z 2 .0 <

0o I (/) uJ t~ ..J ~ 1.o < u .

z

0 0.0

YSZ DIAMETER: 0 .6 pm

YSZ CONTENT: 10 w t % TIME: 6 h

NICKEL DIAMETER

S l jura 1,,

]}~0.25 pm

i i i

0.5 1.0 1.5

SURFACE AREA OF NICKEL, m2/g

Figure 8.4. Interfacial resistance of Ni/YSZ anode at 1000~ as a function of surface area of nickel [8.40]

206 Chapter 8

30 ELECTRODE AREA: 0.6 cm =

TIME: 6 h

(j z < 20 YSZ DIAMETER F- O3

_1 < -

~ i < 10 LL

l lO,um

5/.m~

0 20 40 60 80

YSZ CONTENT, wt%

Figure 8.5. Relationship between interfacial resistance of Ni/YSZ anode at 1000~ and YSZ particle size and content [8.40]

1 ! I I 11273 K i I H=:H=O = 97:3 I FIRING TEMPERATURE

/ / 1773K

~ -

-1

~ /~" 73 U " o o~ (.9 -2 L ab / / ooo

-' 11 . " ~

~'~ , ,

- 3 0.0 0.1 0.2 0.3

ELECTRODE POLARIZATION, V

Figure 8.6. Influence of firing temperature on polarization of hydrogen oxidation at Ni/YSZ anodes prepared by slurry coating [8.29]


It has been shown that both the ionic path (through YSZ particles) and

the electronic path (through nickel particles) are critical to the performance of the

nickel/YSZ anode [8.41]. Accordingly, anode structures with good nickel-to-

nickel and YSZ-to-YSZ particle contacts are expected to have low overpotentials.

An example of this is the observed decrease in interfacial resistance of the anode

with decreasing nickel particle separation in the electrode microstructure (Figure

8.7) [8.20]. In order to have low overpotentials, a desired anode structure

requires good dispersion of nickel particles on the YSZ in addition to a minimum

nickel content and a certain nickel/YSZ particle ratio. However, fine nickel

particles tend to coarsen at the fuel cell operating temperature (600 ~ to 1000~

[8.42]. This coarsening can be minimized by tailoring the anode microstructure.

For example, a lower interfacial resistances can be obtained for an anode in

which the metal particles are covered with thin film or fine precipitates of YSZ

[8.43-8.45]. Addition of MgO to NiO (to form solid solution) maintains high

specific area of nickel particles in hydrogen at 1000~ [8.46]. This leads to the

low interfacial resistance observed for Ni-MgO/YSZ anodes as compared to that

of Ni/YSZ anodes (Figure 8.8) [8.46].

1.0

"E o

O.8

Z <

0.6

O3 LLI

.~ 0.4

< LI_ ~: o2 F- Z

0.0 i i i

1 2 ! i

3 4 5 6 7 MEAN NICKEL PARTICLE SEPARATION,/.~m

Figure 8. 7. Interfacial resistance of Ni/YSZ anode at 1000~ as a function of mean nickel particle separation [8.20]

208 Chapter 8

"7

E "7

c= 2 u3

z < I -

z) a z 1 o

LLI a o n,,- I -

' " 0 ...I ILl

(9 O ._1

/3 3 2 m o l % Y S Z

/YSZ

Ni /YSZ ~ , i i

8 9 1~0

117", 10 .4 K "1

Figure 8.8. Conductance of Ni/YSZ and Ni-MgO/YSZ anodes [8.46]

8.2.2 Electrochemical oxidation of carbon monoxide

The electrochemical reaction at SOFC anodes can also involve the oxidation of CO. Compared to the hydrogen reaction, the oxidation of CO on metal electrode at SOFC operating temperatures is accompanied by higher

overpotentials [8.47-8.51]. The reaction mechanism for the CO oxidation appears to involve the diffusion of CO and oxygen ions to reactive sites, followed

by slow electrochemical reaction to produce CO2 [8.48, 8.49, 8.52, 8.53]. The

introduction of CO2 into the CO lowers the electrode overpotential. Similar to

the hydrogen reaction, the CO oxidation at a metal electrode exhibits a minimum

interfacial resistance as a function of CO2/CO ratio [8.20,8.51,8.52]. The

minimum interfacial resistance occurs at about 50 to 60 mol % CO [8.20,8.51].

The role of CO2 in the CO oxidation is similar to that of H20 in H20/H2 systems.

One suggestion is that the effect of CO2 is related to oxygen partial pressure

[8.7]. The addition of hydrogen gas to CO reduces interfacial resistance. In

presence of H20, the favorable path for the oxidation of CO is via the generation

of hydrogen by the shift reaction

CO + H/O = CO 2 + H 2 (Eq. 8.21)


Thus, for gas mixtures, such as coal gases that contain H2, H20 , and CO, the

anode is expected to have overpotential comparable to that for H20/H2 mixtures

[8.54].

8.2.3 Reaction of sulfide impurities

Many applications of SOFCs will use gases other than pure hydrogen

(e.g., coal gas and natural gas) as the fuel. Depending on the source of the gases

and the cleanup process used, varying amounts of impurities or contaminants will

be present. The impurity expected to have the greatest impact on SOFC

performance and life is sulfur, present primarily as hydrogen sulfide (H2S).

Even low levels in parts per million (ppm) of H2S may cause significant

performance loss at the anode. Table 8.1 shows an example of the effect of

sulfide on cell voltage at different operating temperatures and current densities

[8.55].

TABLE 8.1

Effect of Hydrogen Sulfide on Cell Voltage [8.55]

Temperature Current density Voltage drop in 100 h (oc) (mA/cm z) (%)

2 ppm H2S 10 ppm H2S

900 160 9.0 > 22.0 1000 160 2.0 10.3 1000 250 _ 15.6

In general, if the sulfur contaminam concentration is low, fuel cell

performance may recover fully upon switching to clean fuel (Figure 8.9) [8.55].

Operation at high sulfur levels (> 100 ppm) can result in severe performance

loss, which is only partially recoverable [8.20]. At those high levels, sulfur may

be incorporated into the electrolyte, and this incorporation could explain why the

sulfur effect is not reversible. Certain modifications (e.g., impregnation of the

anode with nickel and samarium-doped CeO2) improve sulfur tolerance in the

210 Chapter 8

0.70

0.66 >

0 0.62 < I- ._1 o > 0.58 ,_1 _J UJ o

0.54

, ! i ! ,

CURRENT DENSITY: 260 mA/cm =

TEMPERATURE: IOOO~

NO H2S

10 ppm H2S

-q= I--.4---NOH2S "

0.500 700

' I I I I I , : I I

140 280 420 560 TIME, h

Figure 8.9. Effect of H2S on cell voltage [8.551

anode. The limit of tolerance for H2S in the fuel for SOFCs has not been firmly established.

Several possible mechanisms may account for the decrease in the

performance of the anode due to the presence of H2S in the fuel: formation of nickel sulfide on the nickel surface can poison the anode; H2S can poison the

hydrogen reaction by adsorbing on anode active sites; and adsorbed H2S can poison the water gas shift reaction, causing a hydrogen deficiency in the fuel cell. At present, the exact mechanism responsible for SOFC performance loss due to presence of sulfur impurities is not clearly defined.

8.2.4 Reforming of hydrocarbons

Because of its high operating temperature, a SOFC can reform

conventional hydrocarbon fuels internally. Internal reforming in a SOFC

simplifies the overall system design because the external reformer can be

eliminated. A SOFC system with internal reforming has an inherent advantage

on energy efficiency in that the heat required for the reforming reaction is

supplied by the heat generated by the electrochemical reaction.

The feasibility of operating SOFCs directly on hydrocarbons has been

demonstrated [8.56,8.57]. For example, tubular SOFCs have operated on

methane gas (reformed internally by the anode) and have shown performance

comparable to that obtained on fuel mixtures equivalent to prereformed synthetic

gas (Figure 8.10) [8.58]. Flat-plate SOFC stacks have also operated on methane

and have demonstrated performance identical to that of hydrogen fuel [8.59].


1 . 0

0 . 8 >

U J

0 0 . 6 < E-- /

0 > O.4 _ J / I i i

{J 0 . 2 -

O . O - "

I

o CH,t + H20 (STEAM RATIO = 3.0) _

�9 50% Hz + 16.6% CO + 33 .3% H20

FUEL UTILIZATION: 86% OXIDANT: AIR OXIDANT UTILIZATION" 26 % TEMPERATURE" 1000~

, , i I 1 ~ l i ! 1 I

0 120 2 0 360 480 600

C U R R E N T D E N S I T Y , m A / c m 2

Figure 8.10. Voltage~current characteristics of a cell operating on methane and on an equivalent prereformed synthetic gas [8.58]

The steam reforming of hydrocarbon fuels at a SOFC nickel/YSZ anode

involves the following reaction (written for methane):

CH 4 + H20 = CO + 3H 2 (Eq. 8.22)

SOFC anodes has been shown to have sufficient catalytic activity for the

reforming reaction [8.60]. In general, the reforming activity of the anode is improved by high nickel surface area, i.e., small nickel crystalline size. During

the reforming reaction, the water gas shift equilibrium (Eq. 8.21) is also

established. Consequently, a mixture of H2, CO, CO2, and C H 4 is obtained.

The factors which affect the equilibrium are the operating pressure, the

temperature, and the steam ratio (the ratio of moles of steam to moles of carbon

in the feed gas). Eq. 8.22 is favorable at low pressures, high temperatures, and

high steam ratios. If insufficient steam is present on the left side of Eqs. 8.21

and 8.22, carbon may be deposited according to the following reactions:

2CO = CO 2 .~ C

CH 4 -- 2H 2 + C

(Eq. 8.23)

(Eq. 8.24)

212 Chapter 8

Carbon formation is undesirable in SOFCs because deposited carbon can plug gas

flow and cause anode performance to degrade by blocking the active sites. A

steam ratio of 3 or more is usually considered safe, although local variations of

temperature and concentration (both in the gas phase and in the solid phase of the

anode) may increase the tendency toward carbon deposition [8. 61]. The acidity

and water adsorption capability of the support (YSZ) in the anode can also

influence carbon formation. Modification of the support has been investigated

as a means to prevent carbon deposition [8. 62]. The reforming reaction can be poisoned by the presence of sulfur in the

hydrocarbon. The effect of sulfur is well known for nickel catalysts used in

steam reforming. However, the sulfur tolerance of SOFC anodes in the

reforming reaction has not been defined.

The kinetics of hydrocarbon reforming at nickel/YSZ cermet anodes has

not been well studied; on the other hand, the reaction kinetics are expected to

vary, depending on anode fabrication conditions. The rate of the reforming

reaction (Eq. 8.22) depends on methane and water concentration in the gas phase,

thus on the steam ratio. Preliminary results indicate that the reaction is first

order with respect to CH 4 and -1.25 with respect to H20 [8.63,8.64]. The

reforming rate exhibits a maximum at a certain steam ratio. The optimal steam

ratio appears to be influenced by current flow and anode potential [8.65]. A proposed mechanism for the reforming reaction (based on the catalytic

activity of nickel metal) suggests that methane adsorbs dissociatively on nickel

to form carbon and hydrogen atoms, with hydrogen atoms desorbing into the gas

phase [8. 63]. The adsorbed carbon atoms can be gasified by either hydrogen or

steam to form CO, CO2, or CH4. Another mechanism proposes that the major

reaction steps occur at active sites on the YSZ surface rather than on the metal

electrode [8.23-8.27]. In this case, oxygen vacancies in the YSZ with trapped

electrons are hypothesized to be the reactive sites.

Recently, anodes based on conductive CeO 2 have been investigated for

the oxidation of CH4 [8.66-8. 70]. Under appropriate conditions, this type of

electrode can be operated without serious carbon deposition problems.

8.3 REACTIONS AT CATHODE

The reaction at the SOFC cathode involves primarily the reduction of

oxygen. Most of the studies on oxygen reduction have been conducted on

platinum electrodes in the temperature range of 400 ~ to 800~ The oxygen


reaction on oxide electrodes at the SOFC operating conditions has only been

studied in depth recently. The overall reaction for the oxygen reduction at a SOFC cathode can be

written as

02 + 2Vo + 4e- = 200 (Eq. 8.25)

This reaction is made up of a series of bulk and surface processes. One or

several of these processes can be rate-determining steps. To understand the

cathode polarization, identifying the various reaction steps and determining the

rate-controlling step are necessary. Figure 8.11 shows some possible controlling

steps [8. 71]: gas diffusion external to the electrode or within the pores of the

electrode; adsorption and dissociation of oxygen on the electrode surface or the

electrolyte; diffusion of adsorbed oxygen on the electrode, on the electrolyte to

ELECTRODE PORE

02

r i

z o Or) ii ii Oso~. '4 �9 oAos ~ )

z ~z z o I iO .~ GRAIN BOUNDAR' l l ~

0, 0s,0. II- ,,, 0 \ 0 1 , , , , , o o n.-

r" " ~ ~ f f k kt.1 J L ~u R F A C E o / ~ . / ")..31 V D I F F U S I O N ' , '

~ S S e "%

>..Z ~ O o o Z ~ . , ,

<u_._ >o

V ELECTROLYTE ~ . ~ s "

- LJ v~;"

Figure 8.11. Schematic diagram of possible steps for oxygen reduction at an electrode~electrolyte interface [8. 71]

214 Chapter 8

the three-phase boundary (between gas, electrode, and electrolyte), or into the electrode/electrolyte interface; diffusion of dissolved oxygen in the electrode or electrolyte; diffusion of electron holes in the electrolyte; and charge transfer across the electrode/electrolyte phase boundary. The various steps of the oxygen reduction can be as follows (written for platinum electrode and YSZ electrolyte):

O2,bulk "-- O2,Pt pore

O2,Pt pore = 0 2 ad, Pt

0 2 ad,at "-- 2Oad,Vt

2Oad,P t -- 2Oad,ERS

4ept - 4eERS

Vo, sz : Vo, E.s

O ad,ERS + 2eER s + VO,ER S = OoX, ERS x X

OO,ER S -- Oo,YS Z

(ERS: electrochemically reactive site)

(Eq. 8.26)

(Eq. 8.27)

(Eq. 8.28)

(Eq. 8.29)

(Eq. 8.30)

(Eq. 8.31)

(Eq. 8.32) (Eq. 8.33)

8.3.1 Oxygen reduction at metal electrode

The oxygen reduction at metal electrodes, especially platinum electrodes, has been studied extensively; however, considerable discrepancies still exist regarding the reaction kinetics and mechanisms [8.72-8.74]. Various rate- controlling steps have been suggested for the oxygen reaction. The suggested rate-determining steps can be grouped into two major categories" a slow electrochemical reaction (a predominant activation overpotential) and a slow oxygen mass transport (a predominant concentration overpotential) [8. 75]. The following is a summary of the various rate-determining steps proposed for the oxygen reduction at metal electrodes in contact with an oxygen-ion conducting electrolyte:

(i) Diffusion of gaseous oxygen in the pores of the electrode [8.35, 8. 76]: The observed dependence of limiting currents on oxygen partial pressure (proportional to oxygen partial pressure) and gas flow rate supports the diffusion of oxygen as the slow step.


(ii) Adsorption of oxygen on the electrode and the electrolyte [8.77-8. 79]: The oxygen activity in the adsorbed layer determines the electrode potential, and

surface concentration variation is responsible for large capacitive effects.

(iii) Dissociation of oxygen molecules into atoms [8.80-8.85]: The interfacial resistance has been shown to be proportional to the -1A th root of

oxygen partial pressure. (iv) Diffusion of oxygen along the electrode surface [8.80,8.82,8.83,

8.86-8.91]: The activation energy and enthalpy indicate a thermally activated

process that can be attributed to the surface diffusion of adsorbed oxygen. The

electrode conductivity is proportional to the -~/i th and 1/2 th root of oxygen

partial pressure in low and high oxygen partial pressure regimes, respectively.

(v) Diffusion of oxygen atoms through the electrode [8.92-8.94]: The

current that is near zero voltage is dependent on both the 1/2 th root of oxygen

partial pressure and voltage. (vi) Diffusion of electrons in the electrolyte [8.93, 8.95]: The electrode

overpotential behavior as a function of current density and oxygen partial

pressure indicates slow diffusion of electrons in the electrolyte.

(vii) Diffusion of oxygen ions in the electrolyte [8. 76, 8.87, 8.95-8.97]: Polarization curves show purely ohmic behavior, with the current being

determined by the electrolyte resistance.

(viii) Charge transfer reaction [8.98-8.106]: Polarization curves obey the

Butler-Volmer and Tafel equations. It has been well established that the rate-controlling step for the oxygen

reduction on platinum electrodes varies, depending on the experimental conditions such as oxygen partial pressure, temperature range, overpotential regime, and electrode characteristics. For example, in a study of the oxygen reduction mechanism on platinum electrode/stabilized ZrO2 electrolyte at

temperatures above 600~ the proposed rate-determining step is the diffusion

of adsorbed oxygen atoms on the platinum surface [8.83]. Below 500~ the

dissociative adsorption of oxygen molecules on the platinum surface has been

suggested as the slow step. In another mechanistic study of a porous platinum

electrode/scandium-doped ZrO2 system, the proposed rate-limiting steps are mass

transport of oxygen in the gas for the high overpotential region, diffusion of

oxygen ions in the electrolyte for the intermediate overpotential region, and

diffusion of electrons from the platinum to the electrochemical reactive site for the low overpotential region [8. 76,8.87]. In the reversible potential region, the

216 Chapter 8

slow step appears to be the dissociation of absorbed oxygen on a platinum surface.

Electrode morphology also plays a critical role in the oxygen reduction

kinetics [8.107]. Although porous platinum electrodes can be easily prepared by

a variety of methods, the electrode morphology varies from preparation method to preparation method. In addition, those electrodes are subject to drastic, time-

dependent morphological changes under current-carrying conditions at elevated

temperatures [8.100, 8.108]. Different electrode microstructures and morphology changes affect the electrode polarization behavior. For example, platinum

electrodes, made of a small platinum tip, paste, or film, show a different type of

dependence of electrode resistivity and limiting current density on oxygen partial

pressure. Morphological changes appear to affect the reactive site area which,

in turn, influences electrode polarization characteristics [8.107, 8.109-8.111]. The location of the electrochemical reactive site for oxygen reduction at

platinum electrodes appears to be the three-phase boundary where the charge transfer steps occur. However, the electrode reaction is not concentrated on the

triple contact line but may spread over the electrolyte surface. Both the platinum electrode and the solid electrolyte can provide active sites for oxygen adsorption.

The adsorption of oxygen on platinum electrode surfaces has been thoroughly

studied [8. 91, 8.112-8.115]. The surface of the electrolyte has also been shown as sites for the adsorption. Thus, the nature of the electrolyte, especially the electrolyte surface, significantly influences the oxygen reduction. As an example, the electrode resistance for platinum electrodes on doped Bi203 is many times lower than that of doped ZrO2 [8.89,8.103]. Introduction of electronic conductivity at the electrolyte surface broadens the active area, thus increasing the reaction rate [8.8,8.116,8.117]. Blackening of the electrolyte increases the

number of active sites, as well as the area of the reaction zone, resulting in

enhanced reaction rates [8.27]. Addition of ion-conducting oxides to platinum

can improve the electrode reactive areas, thus reducing polarization losses [8.118].

8.3.2 Oxygen reduction at oxide electrode

Compared to oxygen reduction on platinum electrodes, the process on

oxide electrodes has not been studied extensively. The behavior of oxide

electrodes is expected to be different from that of platinum electrodes, since the

polarization process on oxide electrodes is markedly dependent on the electrode


material [8.119, 8.120]. The influence of electrode material on the oxygen reaction kinetics has been demonstrated by the wide variations in specific currents measured for different oxides on YSZ at 1000~ [8.120] and on YSZ and doped CeO2 at lower temperatures [8.121]. For the reduction of oxygen at SOFC operating temperatures, oxide electrodes typically perform better than platinum [8.122,8.123].

Different oxide materials show different catalytic activity for oxygen

reduction. For example, the activity of four doped perovskite electrodes studied

at 800~ can be ranked as follows: LaCoO3 > LaMnO3 > LaFeO3 > LaCrO3 [8.122]. The catalytic activity of oxide electrodes generally decreases with

decreasing temperature, and the degree of activity change due to temperature

effect depends on the electrode material [8.121]. The difference in activity among oxide electrodes has been ascribed to their differing catalytic ability for oxygen-molecule dissociation [8.122] and to their differing defect chemistry [8.122,8.124].

(i) Reaction mechanism: The mechanism of oxygen reduction at oxide electrodes has been studied by several methods such as measurement of electrode conductance as a function of temperature and oxygen partial pressure [8.122], AC impedance spectroscopies [8.125], and polarization techniques [8.126].

Figure 8.12 shows, as an example, the dependence of the electrode conductance

on oxygen partial pressure for strontium-doped LaCoO3 electrode at various

temperatures [8.122]. The slope of the plots (about 1/4) suggests the charge transfer process as the rate-determining step. AC impedance measurements show the presence of three time constants, corresponding to the three possible rate- limiting steps [8.2,8.127]. An example of AC impedance spectra of a strontium- doped LaMnO3/YSZ electrode is given in Figure 8.13 [8.128].

The rate-determining step for the oxygen reduction at oxide electrodes varies, depending on electrode material, electrolyte, and temperature. At 800~ the proposed rate-determining step is the charge transfer reaction for strontium-

doped LaCoO3 electrode in contact with YSZ electrolyte, the dissociation of

adsorbed oxygen molecules for strontium-doped LaFeO3 and LaMnO3, and

oxygen diffusion on the electrode surface for strontium-doped LaCrO3 [8.122]. For LaCoO3 and LaMnO3 electrodes with doped CeO2 electrolyte, the rate- determining step is the charge transfer reaction at low temperatures but switches

to the dissociative adsorption of oxygen at higher temperatures up to 800~ [8.121,8.129].

2 1 8 Chapter 8

ol ,

E o

"7

u3 C) Z <~ I ' - 0

13 Z 0 0 LU 13 0 tr" I-" 0 UJ

-2 -

~ 0 8 0 0 o c

0

7600 C

6 6 0 ~

~ C

0 . 0 - .5 - 1 .0 - 1 .5 - 2 . 0

LOG O X Y G E N P A R T I A L P R E S S U R E ,

a t m ( l . 0 1 x 10 SPa)

Figure 8.12. Dependence of electrode conductance on oxygen partial pressure for strontium-doped LaCoO3 cathode [8.122]

6 0 F: Z i i i Z o 4 o n

o 2 o >.

n,.

z 0

I

I i i 1

1 O0 Hz I ,4. ~ 4 . + 4 .

4 . 4-

4 .

1 kHz + "~ 10 Hz 4 .

1 o kH~1~ o ~ -I ~ I ,~ ~oooc

I 1 0 0 Hz ~ J 1 kHz

10 k H z 1 l l !

0 2 0 4 0 60 8 0 1 O0

REAL COMPONENT, Q

1 0 Hz

4- I - 4"

4- 8 0 0 ~ + 4-

§ 4.

§

1

1 2 0 1 4 0

Figure 8.13. Impedance spectra of strontium-doped LaMnO 3 cathode [8.128]


The oxygen reduction mechanism at the LaMnO3 electrode has not been well studied and appears to be complex. Figure 8.14 shows the polarization curves obtained in air for strontium-doped LaMnO3 at 800~ [8.122]. As shown in the figure, the strontium dopant content in LaMnO3 influences the electrode

overpotential, especially at high current densities. The electrode conductance (obtained from the polarization curves in the low overpotential region) indicates that the dissociation of oxygen molecules is involved in the rate-limiting step. Figure 8.15 shows the steady-state polarization plot at 1000~ in air (normalized

by electrode conductance and thickness) for calcium-doped LaMnO3 cathode

prepared by different techniques [8.130,8.131]. The plot indicates that the reaction kinetics (for small overpotentials) is independent of the electrode microstructure and thickness. Oxygen atoms adsorbed on the electrode surface appear to take part in the rate-determining step [8.130, 8.132]. AC impedance

4 0 0 , , ,

TEMPERATURE = 800~

~ /" x= x = l . 0 3 0 0

tx_o., 200

0 0.3

o

1 0 0 7

0 -4 -3 -2 -1

LOG C U R R E N T D E N S I T Y , A / c m "

Figure 8.14. Polarization curves obtained in air for Lal.xSr~Mn03 [8.122]

220 Chapter 8

I.kl O Z

I - s ::::)

z o o uJ E3 O n,,. I - o i i i ..J LLI >- I ' - (/) Z LLI E:)

I ' - Z LIJ tY" iv" :::) s

(9 O .--I

2 . 5

2 . 0

1.5

1 .0

0 . 5

0 . 0 t 0

~ 121

TEMPERATURE = 1 0 0 0 0 C

I

5 0 1 O 0

C A T H O D I C OVERPOTENTIAL, mV

Figure 8.15. Steady-state polarization plot at IO00~ in air for calcium-doped LaMnO 3 prepared by different techniques [8.130]

study of the oxygen reduction in a Lao.85Sr0.~sMnO3/YSZ system shows three depressed semicircles in the impedance plots [8.127]. These semicircles correspond to the three rate-limiting steps: the high-frequency semicircle has been attributed to the charge transfer reaction of oxygen ions, and the medium- and low-frequency ones attributed to peroxide (022-) dissociation and diffusion. On the other hand, AC impedance measurements using an unbonded interface cell suggest that the oxygen reduction kinetics is limited by the dissociation of oxygen molecules and subsequent migration of adsorbed oxygen atoms to the

triple-phase boundary where the charge transfer takes place [8.133,8.134]. A

similar mechanism is also proposed based on the results of potentiodynamic current/potential experiments and impedance studies [8.135].

An empirical equation has been developed for the oxygen reaction rate

at the LaMnO3 electrode [8.130, 8.132, 8.136]. The reaction rate, given as the steady-state current density j, obeys the following equation:

P j _ j o ( a o _ o~) (Eq. 8.34)

a o

where j0 is the rate constant (exchange current density) and ao the oxygen activity in YSZ at the LaMnO3/YSZ interface. The value of J0 has been found to be


about 2.7 mA/cm 2 at 1000~ and its activation energy is 178 kJ/mol. From the reaction rate, the electrode interfacial resistance, Re, has been calculated and is given as

1 RT[exp(2~IF/RT ) _ exp(-2rlF/RT)] (Eq. 8.35) Re = j S a 4F

where S a is the surface area of the LaMnO3/YSZ interface. Eq. 8.35 holds irrespective of the morphology of the LaMnO3 electrode layer.

(ii) Effect of electrode microstructure: Doped LaMnO3 cathodes exhibit close relationships between electrode microstructures and oxygen reduction overpotentials. The two important microstructural characteristics are the length of the three-phase boundary and the electrode/electrolyte contact area. High triple-phase boundary length reduces electrode overpotential [8.137], whereas increased contact area gives larger electrode capacitance [8.130]. The triple- phase boundary length and the electrode/electrolyte contact area in LaMnO3

electrodes depend on several fabrication parameters, especially firing temperature. Electrodes fired at higher temperatures have smoother particles, and, as a

result, the contact area between the electrode particles and the electrolyte increases, giving larger capacitance. On the other hand, because of these smoother particles, the three-phase boundary becomes shorter, resulting in higher interfacial resistance. Thus, the electrochemical properties of LaMnO3 electrodes are expected to vary with starting material characteristics and processing conditions [8.138, 8.139]. For example, LaMnO3 cathodes prepared from different dispersion media show different microstructures with different triple- phase boundary lengths. The electrode made from n-butyl acetate slurry has a relatively large interfacial area and a long three-phase boundary, and this electrode has been shown to have low interfacial resistance [8.138].

The electrochemical performance of porous LaMnO3 electrodes may degrade with time due to microstructural changes at elevated temperatures

[8.140]. Figure 8.16 shows an example of the changes in the conductance and

capacitance of two LaMnO3 electrodes (fabricated at 1100 ~ and 1200~ as a function of time at 1000~ The observed degradation has been related to the

decrease in the triple-phase boundary caused by electrode sintering. Addition of YSZ to LaMnO3 has been used to minimize sintering, thus maintaining electrode performance over long-term operation [8.140]. Doping LaMnO3 with chromium

is effective in improving the electrode's morphological stability, thus preventing cathode overpotential increases with time [8.141].

222 Chapter 8

1 0 ~

E 0 ,-;.,

CI 8

4 fO

0 2

"o LI.,I ,--I t U

0 0

. / ~ , C ~ . _ F I R E D A T 1 1 0 0 = C r n r " m

- - A , O ; F I R E D A T 1 2 0 0 " C ,~ - 4

o

g m

- 0 - 1 ~ 0

1 1 i o " 5 0 1 O0 1 5 0 2 0 0

T I M E , h

Figure 8.16. Conductance and capacitance of LaMn03 cathode as a function of time at I O00 ~ C [8.140]

The microstructure of the electrode, especially near the electrode/electro-

lyte interface region, can be tailored or modified to increase electrochemical

performance [8.142, 8.143]. An example is strontium-doped LaMnO3 cathodes

fabricated by electrochemical vapor deposition, which produces an extremely

large interfacial area. Electrodes fabricated in this way show only 1 mV

polarization at 3 A/cm 2 (interfacial resistance less than 0.01 9-cm 2 in air)

[8.143,8.144] (Figure 8.17).

~E 0.1 fJ

u~ (J z < I-- or) 0r) ,,, 0.01 n,-

o m u .

o LU

00

W iv" 0.001 <I:

!

AIR A T M O S P H E R E

, I , I

7 8 9 10

117", 10 .4 K 1

"E 0 . 0 3

u~ z

0 .02 D 03 LU rt"

__ 0.01 r UJ

r

cc 0 .00 < 0.01

. . . . . . . | . . . . . . ..,. 1 0 0 0 ~

0.1 1.0 OXYGEN PARTIAL PRESSURE,

atm (1.01 x 10 s Pa)

Figure 8.17. Area-specific resistance (in air) of strontium-doped LaMnO 3 cathode made by E VD [8.143]


(iii) Effect of chemical interaction" It is well known that LaMnO3,

LaCoO3, and other perovskite cathode materials can react with YSZ electrolyte

to form poorly conducting compounds such as La2Zr207. The formation of these

reaction products (at the interface) usually increases electrode overpotential and

causes electrode performance to degrade [8.145-8.151]. The effect of the elec-

trode/electrolyte interaction becomes more significant at higher temperatures and

longer times. Figure 8.18 shows the cathodic polarization curves obtained at

800 o C for LaCoO 3 electrodes annealed at 1000 o C for various times [8.145]. The

observed increase in the overpotential with increasing anneal time is attributed

to the reaction between YSZ and LaCoO 3. The increase in the cathode

polarization observed in long-term operation (Figure 8.19) is partly caused by the

slow chemical interaction at the interface. Lowering fabrication temperatures or

using buffer layers [8.146,8.152] reduces the interaction between the electrode

and the electrolyte. For example, the interfacial resistance of strontium,cobalt-

doped LaFeO3 cathode is significantly reduced when a thin samarium-doped CeO2

buffer layer is placed between the YSZ electrolyte and the cathode to minimize

chemical interaction [8.152].

�9 200 > E J < F-- Z I.iJ t--- O Ix. r 100 LLI > O L) D O "1" !-" < L)

5 h - 1 2 h X

l , 2 o

-2 -1 0

LOG CURRENT DENSITY, A/cm"

Figure 8.18. Polarization curves at 800~ for LaCoO 3 electrodes annealed at lO00~ for various times [8.145]

224 Chapter 8

E

rc"

N

1 0 0

50

10

0 ~ . i 0 2 0 0 4 0 0 6 0 0 8 0 0

Po, = 0.1 a tm

TIME, h

u 0 o

Po= 1 a tm

1 0 1 2 0 0 1 4 0 0

Figure 8.19. Change in polarization of strontium-doped LaMn03 cathode as a function of time at 1000 ~ C [8.148]

To date, the three-phase boundary area is considered as the electrochemically reactive site for oxygen reduction on oxide electrodes. Sufficient evidence has been reported in the literature to support this conclusion. An example is the observed increase in oxygen reaction kinetics of strontium-doped LaMnO3 electrode with increasing triple-phase boundary length [8.133,8.153]. Another example is the dependence of the electrode conductance on the length of the three-phase boundary [8.154]. (The length of the triple-phase boundary depends on electrode preparation conditions and can vary between 102 to 105 cm/cm 2 [8.155].)

The reactive area for the oxygen reduction (oxygen adsorption site) is

often not limited to the triple point but can spread along the electrode surface,

and the spreading of the reaction zone is related to defect chemistry of the cathode material. The concentration of oxygen adsorption sites is directly related

to the concentration of defects in the electrode. In doped perovskites such as

strontium-doped LaMnO3, a certain level of oxygen vacancies is present,

influencing the exchange of oxygen, and thus the reaction rate, at the material

surface [8.156]. The electrocatalytic activity of doped LaMnO3 is greatly

enhanced under high cathodic polarization [8.128,8.157-8.161]. Under those conditions, the electrode material is partly reduced, resulting in a marked

increase in the oxygen vacancy concentration inside the electrode. The oxygen


reaction takes place not only at the three-phase boundary but also on the

electrode surface according to the following equations"

l 0 x "~ 2 + Vo(mang anite ) + 2e- = O0(manganit e )

X X

O0(manganite ) + Vo(electrolyte ) "-Oo(electrolyte ) + Vo(manganite )

(Eq. 8.36)

(Zq. 8.37)

Thus, the enhanced activity of the electrode can be attributed to the presence of

both electrons and oxygen vacancies (mixed electronic and ionic conduction) in

the electrode material. A way to introduce ionic conduction in LaMnO3

electrodes is to add YSZ to the electrode material [8.140,8.162]. In this case,

the reactive site spreads into the electrode and broadens the reaction zone,

significantly reducing electrode polarization as a result (Figure 8.20) [8.162].

E "7,

6

(J Z < I--- r j ::) 4 13 Z 0 0 Z ___ 2 I.-- <

< _.1

0 0 a.. 0 1 2 3

WEIGHT RATIO OF Y S Z / L a M n 0 3

Figure 8.20. Cathodic polarization conductance as a function of the YSZ/LaMnO s weight ratio [8.162]

References

8.1

8.2 N.Q. Minh, J. Am. Ceram. Soc., 76 (1993) 563. M. Mogensen, in Proceedings of the 14th Riso International Symposium on Materials Science, High Temperature Electrochemical Behaviour of Fast Ion and Mixed Conductors, September 6-10, 1993, Roskilde, Denmark, F.W. Poulsen, J.J.

226 Chapter 8

8.3

8.4

8.5

8.6

8.7

8.8 8.9

8.10

8.11 8.12

8.13 8.14 8.15 8.16

8.17

8.18 8.19

8.20

Bentzen, T. Jacobsen, E. Skou, and M.J.L. OstergArd (eds.), Riso National Laboratory, Roskilde, Denmark, 1993, p. 117. G. Maggio, I. Ielo, V. Antonucci, and N. Giordano, in Proceedings of the Second International Symposium on SolM Oxide Fuel Cells, July 2-5, 1991, Athens, Greece, F. Grosz, P. Zegers, S.C. Singhal, and O. Yamamoto (eds.), Commission of the European Communities, Luxembourg, 1991, p. 611. S. Sakamoto, H. Taira, H. Takagi, and K. Tomono, in Extended Abstracts, The

Second Symposium on Solid Oxide Fuel Cells in Japan, December 15-16, 1993, Tokyo, Japan, Solid Oxide Fuel Cell Society of Japan, 1993, p. 99. E.J.L. Schouler and H.S. Isaacs, Solid State Ionics, 5 (1981) 555.

K. Eguchi, T. Setoguchi, M. Sawano, S. Tamura, and H. Arai, see Ref. 8.3, p.

603. T. Setoguchi, K. Okamoto, K. Eguchi, and H. Arai, J. Electrochem. Soc., 139

(1992) 2875. E.J.L. Schouler and M. Kleitz, J. Electrochem. Soc., 134 (1987) 1045. T. Takahashi, H. Iwahara, and Y. Suzuki, in Proceedings of the Third International Symposium on Fuel Cells, Presses Acad6miques Europ6ennes, Brussels, Belgium,

1969, p. 113. E.J.L. Schouler, M. Kleitz, E. Forest, E. Fernandez, and P. Fabry, Solid State

Ionics, 5 (1981) 559. L.J. Olmer, J.C. Viguie, and E.J.L. Schouler, Solid State Ionics, 7 (1982) 23. H. Miyamoto, M. Sumi, K. Mori, I. Koshiro, F. Nanjo, and M. Funatsu, in Proceedings of the Third International Symposium on Solid Oxide Fuel Cells, May 16-21, 1993, Honolulu, HI, S.C. Singhal and H. Iwahara (eds.), Electrochemical Society, Pennington, NJ, 1993, p. 504. K. Eguchi, T. Setoguchi, K. Okamoto, and H. Arai, see Ref. 8.12, p. 494. K. Eguchi, M. Kayano, T. Setoguchi, and H. Arai, see Ref. 8.4, p. 57. T. Inoue, T. Setoguchi, K. Eguchi, and H. Arai, Solid State Ionics, 35 (1989) 285. J. Mizusaki, H. Tagawa, T. Saito, K. Kamitani, T. Yamamura, K. Hirano, S. Ehara, T. Takagi, T. Hikita, M. Ippommatsu, S. Nakagawa, and K. Hashimoto, see

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8.58

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8.132

8.133

8.134

8.135

8.136

8.137

8.138

8.139

8.140

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8.142

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M.J.L. Osterghrd and M. Mogensen, Electrochim. Acta, 38 (1993) 2015.

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M. Suzuki, H. Sasaki, S. Otoshi, A. Kajimura, N. Sugiura, and M. Ippommatsu,

in 1992 Fuel Cell Seminar Abstracts, November 29-December 2, 1992, Tucson, AZ, Courtesy Associates, Washington, DC, 1992, p. 297.

232 Chapter 8

8.145

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8.148 8.149

8.150

8.151 8.152

8.153

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8.157 8.158

8.159 8.160 8.161

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T. Kenjo and M. Nishiya, Solid State Ionics, 57 (1992) 295.

Chapter 9

STACK DESIGN AND FABRICATION

9.1 GENERAL

A SOFC single cell, under typical operating conditions, produces less than 1 V. Therefore, practical SOFCs are not operated as single units; rather, they are connected in electrical series in a stack to build voltage. The power produced by a stack is determined by the number and the area of the individual cells. For practical power generation, the cells of a stack have to be of sufficient size, since the current produced is directly proportional to cell area.

The design of a SOFC stack is governed by the restrictions imposed by

the properties of the selected cell materials. The design must also consider the required physical, chemical, electrical, and electrochemical characteristics of the

cell components at the operating temperature. A SOFC stack must be designed to achieve the desired electrical and electrochemical performance, mechanical

integrity, and manifolding requirements, as discussed below.

(i) Electrical and electrochemical performance: Any stack design must provide for acceptable electrical and electrochemical performance. This requirement means that the design must minimize ohmic and polarization (especially concentration polarization) losses in the stack. To reduce ohmic resistance losses, the cell componems must be made into thin layers, and the current path in the components (especially those having low conductivity) must

be designed to be as short as possible. To reduce concentration polarization

losses, the design choice must provide good access (small diffusion barrier) to

the reaction sites for the reacting gases.

(ii) Mechanical integrity: Any stack must be designed to have adequate mechanical strength for assembly and handling. The design must also satisfy

other structural requirements created by the need to maintain good electrical contact between the individual cell components in a leak-free manner. Thus,

mechanical and thermal stresses must be kept to minimum to prevent cracking, delamination, or detachment of the components under the variety of operating

234 Chapter 9

conditions the stack is expected to experience (e.g., normal operating temperature

gradients, off-design temperature gradients, thermal shock conditions such as sudden power change and cold startup, and mechanical loading expected during installation, moving, and vibration loading).

(iii) Gas manifolding requirements: Any stack design must include suitable means of routing gases (gas manifolding) from a common supply point

to each cell and removing unreacted gases and reaction products. The design must provide even and uniform gas distribution not only across the area of each

cell but also to each cell of the stack. The former is mainly controlled by the configuration and dimension of the gas channels; the latter is mainly controlled

by the gas manifold design. Any manifold system must provide gas sealing between components of separate cells and between manifold and stack to prevent gas leakage.

One important aspect in SOFC stack design is its fabricability. Any

design requires appropriate ceramic fabrication and assembly methods to incorporate the cell materials into the stack configuration. The fabrication and

assembly processes must ensure that no condition or environment in any process step destroys desired material characteristics of any of the components.

Fabrication and assembly methodologies must attain the desired structural integrity, shape, electrical conductivity, and electrochemical performance in the stack.

At present, four common SOFC stack configurations have been proposed

and fabricated" the sealless tubular design, the segmented-cell-in-series design,

the monolithic design, and the flat-plate design [9.1]. The designs differ in the extent of dissipative losses within the cells, in the manner of sealing between fuel and oxidant channels, and in the cell-to-cell electrical connections in a stack of cells. The fabrication costs and ease of assembly vary among the designs. Table

9.1 compares the key characteristics of the four designs [9.2]. The comparison

is relative and qualitative because the differences are hard to quantify at this stage

of development.

The fabrication process selected for each design depends on the configuration of the cells within the stack. Current processes can be classified

into two groups based on the fabrication approach: the deposition approach and the particulate approach. The deposition approach involves formation of thin

layers on a support by a chemical or physical process; the fabrication of the

sealless tubular SOFC and the segmented-cell-in-series SOFC is based mainly on

this approach. The particulate approach involves compaction of ceramic powder

Stack Design and Fabrication 235

into cell components and densification at elevated temperature; the fabrication of

the monolithic SOFC and the flat-plate SOFC is based mainly on this approach.

The details of the current stack designs and fabrication processes for SOFCs are

discussed in the following sections.

TABLE 9.1

Characteristics of SOFC Stack Designs

Design

Feature Sealless Segmented-Cell- Monolithic Flat-Plate Tubular in-Series

Structural support Yes Yes" No No

Internal electrical High High Low Medium resistance

Gas sealing No Yes No b Yes

Power density Low Low High ~ Medium

"Banded configuration (see Section 9.3 on segmented-cell-in-series design) bCoflow configuration (see Section 9.4 on monolithic design)

9.2 SEALLESS TUBULAR DESIGN

The sealless tubular design, first proposed in 1980 [9.3,9.4], is the most

advanced of the several SOFC concepts proposed. Single cells of this design

have been operated for tens of thousands of hours and have shown stable and

excellent performance. Multikilowatt-size stacks have been constructed and

tested with a variety of practical fuels.

9.2.1 Design features

The sealless tubular design consists of the cell components configured as

thin layers on a tubular support closed at one end (Figure 9.1). Each tube of this

design represents a single cell. In this design, the porous support tube is overlaid

with a porous cathode layer. A dense electrolyte layer covers the cathode except

in a strip along the active cell length. This strip of exposed cathode is covered

236 Chapter 9

INTERCONNECT

CATHODE ~ . , . ~

ELECTROLYTE

INTERCONNECT CONTACT

E LE CTRO LYTE INTERCON N ECT

ANODE

7o

CATHODE ..~ SUPPORT TUBE

H2

Figure 9.1. Sealless tubular SOFC design [9.4]

with a gastight interconnect layer. The anode covers the entire electrolyte surface. For cell operation, oxidant is introduced to the cell through an injector

tube, where it traverses and exits from the open annulus between the support and the oxidant injector tube. Fuel flows on the outside of the support tube. A

gastight seal between oxidant and fuel channels is made by the slight overlap of the electrolyte layer over the edge of the interconnect strip. Table 9.2

summarizes the properties of the components of the sealless tubular cell. Early sealless tubular SOFC technology used a ZrO2 support tube 20 cm

long with a 2-mm wall thickness [9.5]. To improve the power output per cell,

the active cell length has been progressively increased (currently up to 2 m), and the thickness of the support tube has been reduced (to 1.2 mm). Work has been

performed to increase the cathode thickness (until it is sufficient to support the

cell) and to eliminate the ZrO2 tube [9.6].


TABLE 9.2

Properties of Sealless Tubular SOFC Components

Component Thickness Other Dimensions Porosity Material

Support tube 1.2 mm Diameter: 1.5 cm 35 % 15 mol% CaO- Length: up to 2 m stabilized ZrO2

Electrolyte 40/~m 0% 10 mol% Y 2 0 3 -

stabilized Z r O 2

Cathode 1.4 mm 35 % 10 mol % SrO- doped LaMnO3

Anode 100 tzm 40% Nickel/Y/O3- stabilized Z r O 2

Interconnect 40/~m Width: 9 mm 0% 10 mol% MgO- doped LaCrO3

Electrical connection in sealless tubular design

In the sealless tubular design, electrical connections are made in the

reducing environment at the anode side (rather than in the oxidizing atmosphere

at the cathode side). This arrangement allows the use of low-cost metals for cell-

to-cell contacts and current collectors. Sealless tubular cells are bundled in series

(anode-to-interconnect) and parallel (anode-to-anode) electrical connection to

form a basic power-generating building block (Figure 9.2). The series and

parallel electrical connection is designed to protect the bundle or stack against

complete failure if an individual cell fails. Nickel felt is used to connect the

anode of one cell to the nickel plating on the interconnect of the next cell for

series connection and the anode of one cell to the anode of the adjacent cell for

parallel connection. The felt provides compliancy, thus minimizing stress placed

on the cells during operation. A sintering temperature of about 1000~ is used

to bond the nickel felt to the cells [9.7]. The cells connected in this manner form

a semirigid structure.

A sealless tubular SOFC bundle is typically an array of three by six cells

[9.8]. A bundle has three parallel strings of six series-connected cells. The

238 Chapter 9

POSITIVE CURRENT COLLECTOR

FUEL

OXIDANT ]==,=~]~I~I::I'-"I OXIDANT

NICKEL FELT

FUEL

= - - " " ' ~ ~ NICKEL NICKEL PLATED ' ~ " ' - FELT

OVER INTERCONNECT ..~

FUEL

OXIDANT ] = = = = ~ ~ OXIDANT

INTERCONNECT

' - - ANODE

SUPPORT TUBE

NEGATIVE "~i"~=~l~ " ~ ~ t , - ~ NICKEL FELT CURRENT COLLECTOR ~ " -

Figure 9.2. Electrical connection in sealless tubular SOFCs [9.5]

individual bundles are connected in series to build voltage and form modules.

The modules are then further combined in either parallel or series arrangements to form a large-scale generator. For example, a 3-kW generator consists of eight

bundles configured in electrical series (144 cells of 36-cm active length) [9.9].

A bundle of 1-m active length cells would generate about 2 kW under design conditions. A conceptual 200-kW generator would comprise fourteen series strings of eight series-connected bundles (2,016 cells total) [9.10].

Gas manifolding in sealless tubular design

Figure 9.3 illustrates the gas manifolding concept in the sealless tubular

SOFC [9.11]. Oxidant is fed through an oxidant plenum (at the top of the

generator) into the oxidant injector tubes, which are suspended (one per cell) from the plenum into the cells. Oxidant flows down the injector tube and

reverses course at the closed end of the cell. As the oxidant flows back up the annular space between the oxidant injector tube and the cell inner diameter, the

electrochemical reaction occurs. Fuel flows from a fuel plenum at the bottom

and up along the outside of the cells, where the electrochemical reaction occurs,

generating electricity and heat. The spent fuel flows through a porous ceramic


OXIDANT

J

,RE.EA U'REACTEO,OEL

DUAL. CELL

FUEL

OXIDANT

Figure 9.3. Gas manifolding concept in sealless tubular design [9.11]

barrier, enters the combustion chamber, and combines with the spent oxidant.

The combustion heat is used to preheat the oxidant entering the cell. Depending

on the operating conditions, the exhaust gases exit the generator from the

combustion chamber at temperatures up to 900~ The temperature distribution

within the cell is maintained by the heat associated with fuel cell power

generation and combustion of the spent fuel in the combustion chamber

(nominally 15 % of the supplied fuel). Excess oxidant flow provides cooling for

the fuel cell.

9.2.2 Advantages and disadvantages

One distinct feature of the sealless tubular design is that it has no gas

seals, so the problems with gastight seals for ceramics at high temperatures are

eliminated. The design has two "leaky" seals. One "leaky" seal at the porous

barrier allows spent fuel to pass through, yet keeps the oxidant out of the fuel

region. Another "leaky" seal is formed from a ball and socket arrangement at

240 Chapter 9

the oxidant plenum/oxidant injector tube juncture. Small leaks (about 5 %) of

oxidant at this "leaky" seal (from the oxidant plenum to the combustion chamber)

do not significantly affect cell performance. In addition, the sealless tubular

design has the advantage of providing a more robust ceramic structure for the

cell, since each thin cell is formed on a thick support. Also, each cell of this

design is built as a unit structure. This allows some freedom of thermal

expansion and minimizes the problem of cracking caused by undue thermally

induced stresses in a monolithic pack of connected cells. Other advantages of the

sealless tubular design include easy connections between cells in the reducing

atmosphere and integration of a high-temperature heat exchanger and the fuel cell

without the need for gastight seals.

On the other hand, the sealless tubular design has a relatively long

current path through the cell [9.12]. During operation, the current travels

through the plane (i.e., across the thickness) of the interconnect and electrolyte,

and in the plane (i.e., along the circumference of the support tube) of the anode and cathode. The current paths are short in the interconnect and electrolyte but

long in the anode and cathode. The long current path in the electrodes results in higher internal cell resistances and greater resistive losses. In addition, the

support tube in the sealless tubular design adds weight and volume to the cell

(thus reducing power density). The support tube is also a limitation on cell

performance. The thick support tube restricts the amount of oxygen which can

be transported to the cathode/electrolyte interface. Thus, gas diffusion through

the tube can become the rate-limiting step and sets the limiting current for the

cell. Even below the limiting current, gas transport represents a loss in cell

performance. Replacing the porous support tube with a cathode tube reduces oxygen diffusion losses and improves cell efficiency and power density. Finally,

the electrochemical vapor deposition (EVD) process used in the fabrication of the

sealless tubular SOFC severely limits the selection of suitable dopants for the

electrolyte and interconnect.

9.2.3 Fabrication

In a sealless tubular SOFC, the cell components are formed sequentially

as thin layers on the porous support tube; thus, the fabrication conditions cannot

be selected independently for each component. As the layers are built up one by

one, the process conditions for each successive layer must be selected such that


no thermal or chemical environment of the step alters the desired properties of

the previously formed layers. At present, the fabrication of the sealless tubular SOFC consists of a

progression of processing steps. Cell construction begins with the porous support tube. The component layers are then built up on the support in the following

order: the cathode, the interconnect, the electrolyte, and the anode. The main

processing techniques employed include extrusion, slurry coating, and EVD.

Figure 9.4 shows a flowchart of the manufacturing process for the sealless

tubular SOFC [9.13].

AREA 1 TUBE MANUFACTURING

�9 SOLIDS PREPARATION �9 TUBE MANUFACTURING �9 TUBE FINISHING

/INSPECTION

A R ~ 6 EVD ANODE

�9 DIP PROCESS �9 EVD FIX �9 CLEANING

AREA 2 CATHODE APPLICATION

�9 MASKING �9 ELECTRODE DEPOSITION

�9 SINTERING

AREA 5 NJ PLATE-INTERCONNECT

�9 CLEANING �9 Ni FLASH AND PLATE

�9 WASHING/DRYING

AREA 3 EVD-INTERCONNECT

�9 CLEANING . MASKING �9 EVD �9 DEMASKING

1' AREA 4

EVD-ELECTROLYTE

�9 CLEANING �9 MASKING �9 EVD �9 DEMASKING

FINISH

CELL TESTING

Figure 9.4. Sealless tubular SOFC fabrication process [9.13]

242 Chapter 9

Fabrication processes

(i) Support tube: CaO-stabilized Z r O 2 has been used as the material for

the support tube because of its thermal expansion match and chemical stability.

At present, the support tube is formed by an extrusion process; extrusion is the

preferred low-cost method of forming large quantities of support tubes [9.14].

The fabrication sequence for the support tube is shown in Figure 9.5 [9.8]. The extrusion process involves forcing a highly viscous, dough-like

plastic mixture of ceramic powder plus binders and additives through a shaped

die [9.15]. In this process, an auger extruder is used. The premix (ceramic

powder plus binders and other additives) is fed into the pug mill section of the

extruder. The pug mill kneads the premix to provide homogeneity, to maximize

plasticity, and to squeeze out excess air. The mixture then enters a deairing

chamber to remove as much air as possible. The mixture finally moves to the

compaction chamber where auger motion or pressure from a piston precompacts

the mixture to remove as much void space as possible prior to extrusion through a shaped die. The resulting long strands of open cylinder are supported on trays

and cut to the desired length. A plug is inserted to close one end. The tube is

then sintered, cut, and ground to final dimensions and shape.

In the extrusion of the support tube, the binder (polyvinyl alcohol) is the

major additive to the ceramic powder (15 mol % CaO-stabilized ZrO2) [9.7, 9.16].

This additive provides a coating over each ceramic particle to allow flow during

extrusion and green strength after extrusion. A small amount of Fe203 is added

as a sintering aid, and cellulose and starch are used as pore formers [9.7].

Extruded tubes are sintered in air at about 1550 ~ to 1650~ An example of a

firing schedule for the support includes a 12-h heatup, a 4-h hold, and a 12-h

cooldown [9.8]. The support tube of the sealless tubular SOFC provides mechanical

integrity and serves as a conduit for the oxidant. Thus, the sintered tube must

possess adequate mechanical strength (low-porosity material required) and allow

high gas diffusion rates (porous material required). In general, the strength, Ss,

of the support tube depends on both porosity, Po, and grain size, Gr, as given by

the following equation:

o -~ S s = Ss Gr e -~" (Eq. 9.1)

where S, ~ is the fracture strength of a theoretically dense, 1-~m grain size

sample and B and ~, are empirical constants (for most ceramics, g = 0.3 to 0.5


ORGANIC BINDER SYSTEM

(CaO)o.16(ZrO2)o.8s

WEIGHING

BALL MILLING

CALCINING

-------~- BLENDING

KNEADING

Fe203

Material input

k---7 r '/Deairtng ~'--'/ ~ Shafts to gears Pug mdl Compaction I c h a m b e ~ and drive motor chamber

Shaft to ~ ~ / ~ d r , v e m o t o r ~

~ " ~ Cutaway show,ng the two rows of pug mill

knives on auger shafts Cutaway of a Cutaway showing auger

simple extrusion die for compaction and extrus=on

EXTRUDING

DRYING

FIRING

INSPECTION

Figure 9.5. Fabrication sequence for support tube [9.8]

244 Chapter 9

and ~ = 4 to 9). Starting material characteristics and processing conditions can

be tailored to obtain a suitable strength, porosity, and gas diffusion coefficient

for the support tube [9.17, 9.18].

Extrusion is often more of an art than a science [9.15]. Therefore, quality is controlled by careful inspection of extruded tubes for defects.

Laminations, longitudinal faults, and dark spots are common defects (macro- defects observable by visual inspection or radiography) in the extrusion of the support tube for the sealless tubular SOFC [9. 7, 9.8]. Laminations are caused by

incomplete deairing of the extrusion mix prior to extrusion. Longitudinal faults are caused by incomplete compaction of the extrusion mix in the die after the mix

passes through the web of the extrusion spider. Dark spots are believed to be due to presence of large pore former particles. Other defects such as cracks in

the tube are determined a candling technique [9.13]. Tube properties (such as

porosity, gas permeability, strength, and dimension) are measured to ensure that each unit meets the specified requirements. For example, for early extruded support tubes, tentative physical property specifications were set at 34 to 35 % porosity, gas permeability of 0.02 Darcy minimum and uniform along the tube

length to within + 15 % of normal values, minimum hoop strength of 34.5 MPa

(5,000 psi), and maximum elastic modulus of 65.5 GPa (9.5 • 10 6 psi) [9.13]. (ii) Cathode: The cathode (next to the support tube in the fabrication

sequence) is currently deposited on the porous support by slurry coating. Early

cathode layers were produced by a spraying technique [9.19]. In the slurry coating method for the cathode, the electrode layer is formed from a dilute aqueous slurry of doped LaMnO3. The porous support tube is suspended in the slurry and a vacuum is applied to the inside of the tube. The cathode material is deposited on the support while the water is drawn through the tube wall to a

collection tank. A uniform-thickness cathode layer can be formed through control of the solids loading, the slurry mixing, the pressure differential, and the

immersion time. A number of support tubes can be deposited simultaneously by

this technique.

The details of the slurry coating process for the cathode are not well

known. Early publications [9. 7, 9.20] indicate that graphite is added to the slurry

as a pore former. With the deposit rate of about 0.12 g/cm2/min, it takes 5 to

10 min to obtain a cathode thickness of 1.2 mm. The deposited tube is dried before firing. Procedures have been developed to prevent radial cracking of the electrode surface during the drying process. The cathode layer is sintered in air


at 1400~ In general, carbon content of several weight percent remains in the

cathode layer after firing. Recently, work has been conducted to replace the porous Z r O 2 support

with a cathode tube. In this case, the cathode tube is fabricated by extrusion.

(iii) Interconnect: The interconnect is made by the EVD technique. The technique has proven to be the only reliable method for producing the thin, dense

interconnect layer for the sealless tubular design. More details of the EVD process will be discussed later. In the interconnect deposition, vaporized chlorides of chromium, lanthanum, and magnesium are passed over the outside (the coated cathode layer side) of the support tube while oxygen and steam are

fed to the interior of the support. The deposition takes place at about 1300~

under vacuum (53 to 267 Pa). Typical deposition time is about 1 h. Powder

masking is used to limit the deposition of interconnect into a strip along the

length of the cathode. The masking material can be easily removed in

postdeposition cleaning operation. A typical micrograph of the EVD interconnect

layer deposited on a porous cathode is shown in Figure 9.6. To obtain optimum quality and speed of the deposition of the interconnect

layer, tailoring and comrolling of certain key process variables (e.g., pressure,

temperature, reacting gas composition, flow distribution, and flow rate) are critical. Pretreatment and fusing of the chlorides are required to ensure a

homogeneous mixture. During the interconnect EVD process, the cathode material can react with the chlorides and may be reduced due to the very low

oxygen activity that exists on the halide vapor side. To prevent harmful

INTERCONNECT

CATHODE

Figure 9.6. Typical micrograph of EVD interconnect on porous cathode (courtesy of Westinghouse)

246 Chapter 9

interactions, interfacial layers of ceramic materials can be used [9.21].

Limitations in the oxygen partial pressure for the stability of the cathode material

can be adequately accommodated in the EVD processing procedures [9.22]. The

interconnect deposition conditions can cause temporary oxygen loss in the

LaMn03 cathode material. However, this oxygen loss is recovered in later processing of the cell.

The interconnect surface is plated with nickel to provide better electrical

contact.

(iv) Electrolyte: The YzO3-stabilized ZrO 2 (YSZ) electrolyte is also made by EVD. The electrolyte is deposited over the entire active area of the cell,

including overlap regions of about 0.5 mm on all sides of the interconnect. The electrolyte overlap at the interconnect edges provides gastight sealing for the cell.

A masking material protects the interconnect during the electrolyte deposition. Figure 9.7 shows a typical EVD electrolyte layer (on a porous cathode).

The electrolyte is formed from the vapors of YC13 and ZrC14 and an oxygen/steam mixture at about 1200~ under vacuum. The deposited electrolyte

shows a fully stabilized ZrO 2 phase. Deposition time and pore closure time for the electrolyte are similar to those of the interconnect. Like the interconnect

EVD, during the initial stages of'the electrolyte deposition process, the halide

vapors can react with the porous LaMnO3 cathode material to form undesirable

reaction products such as LazZr207 (before pore closure can prevent any further chemical reaction). An interfacial layer may be used to prevent possible

interactions; for example, a thin interlayer of CaO-doped YCrO3 has been used for this purpose [9.18]

ELECTROLYTE

CATHODE

Figure 9. 7. Micrograph of typical EVD electrolyte layer on porous cathode (courtesy of Westinghouse)


(v) Anode: The anode is applied on the electrolyte by a two-step process.

In the first step, nickel powder is applied over the electrolyte by dipping in a

nickel slurry. The anode covers the entire electrolyte surface but not in contact

with the interconnect to avoid internal cell shorting. Suitable masking is used to

cover the interconnect. In the second step, the anode is fixed by EVD of YzO3-

stabilized Z r O 2 in the nickel matrix. This zirconia acts as a sintering inhibitor

and maintains a porous, structurally stable anode. The zirconia skeleton grown

by EVD is sufficient strong to support a stable nickel structure during long-term

operation. A representative micrograph of the anode layer prepared by this

procedure is shown in Figure 9.8.

Electrochemical vapor deposition

The EVD process is the key fabrication technique in the sealless tubular

SOFC technology. The process, a modified chemical vapor deposition (CVD),

involves growing a dense layer of electron- or ion-conducting oxide on a porous

substrate at elevated temperatures and reduced pressures [9.23]. In the EVD

process, the oxide layer forms and grows from two gas flows, a metal halide

flow (commonly chloride mixtures) and an oxygen-containing gas flow

(commonly steam/oxygen or steam/hydrogen mixtures), separated by the porous

substrate. This deposition process resembles the growth of an oxide scale on

metal (Wagner oxidation process) [9.24], and the growth mechanism is electro-

chemical in nature. The principles of the EVD process are illustrated in Figure 9.9 [9.25].

ELECTROLYTE

ANODE

Figure 9.8. Typical micrograph of sealless tubular SOFC anode (courtesy of Westinghouse)

248 Chapter 9

Met,2 vAPo, - - - , . / 1' 02- |

POROUS SUBSTRATE

H20 ='~ / " ' ~ H2

STAGE I PORE CLOSURE BY CVD

MeCI2 + H 2 = MeO + 2HCI

STAGE II SCALE GROWTH BY EVD

MeCI 2 + O 2 = MeO + CI 2 + 2e"

H20 + 2 e = H 2 + O 2

Figure 9.9. Principles of electrochemical vapor deposition [9.25]

An oxide layer is deposited by EVD in two stages: a pore closure stage

(stage I) and a scale growth stage (stage II). Stage I involves the formation of the oxide in the pores of the porous

substrate by direct reaction of metal chloride, MeCly, with H20 (where Me is the cation species and y is the valency associated with the cation)

MeCly + YH20 = MeOy/2 + yHC1 (Eq. 9.2)

The oxide, formed by chemical vapor deposition, eventually closes the pores, and

no further direct reaction of MeCly with H20 occurs. Since the porous substrate physically separates the two reactants, any open porosity is a reaction site.

Hence, complete pore closure is assured. Stage II involves the growth of the oxide over the closed pores by a

Wagner oxidation process. The oxide growth arises due to the presence of a large oxygen activity gradient across the substrate. In this stage, H20 is reduced at the water vapor side of the growing oxide scale to produce oxygen ions. The

oxygen ions diffuse through the film to the metal chloride side and react with the


metal chloride vapor to form the oxide

Y "" Y H Yn20 + ~V6 + ye- = 2 + YOo (Eq. 9.3)

Y V 0 + y e - (Eq. 9.4) MeCly + YO0 = MeOy/2 + YC12 +

The scale grows in the direction of the chloride gas phase side since oxygen ions

are generally more mobile than the metal cations. Under the operating conditions

of the EVD process, the oxide (either ion-conducting or electron-conducting)

exhibits both oxygen-ion and electron conductivity. Thus, the oxygen-ion flux

during the oxide growth is balanced by an electron flux, thereby preserving the

electroneutrality of the oxide scale. The scale growth is self-leveling (growth is

fastest where the scale is thinnest) [9.23,9.26], resulting in uniform scale

thickness.

The kinetics and growth mechanism of the EVD process have been

investigated [9.23,9.27-9.41]. A study of the pore closure stage (stage I) in the

EVD of stabilized ZrO2 has indicated that this stage can be modeled as

simultaneous Knudsen diffusion and reaction in a straight pore [9.27]. The

reaction rate is first order in the MeCIy and zero order in H20. In this case,

pores always close at the metal chloride side of the porous substrate because the

metal chloride concentration and reaction are greatest at this side. The time for

pore closure and the final pore radius profile depend on a dimensionless Thiele modulus cI,:

2L2k ,I, - ( E q . 9 . 5 )

D r m p

where L is the thickness of the substrate, k r is the reaction rate constant, D m is

the effective diffusivity of the metal chloride, and r e is the pore radius. Large

values of this parameter should shorten pore closure time and reduce pore

narrowing. A high Thiele modulus can be achieved by increasing the rate

constant (e.g., by increasing the temperature) or by reducing the diffusivity of

the metal chloride through the porous substrate (e.g., by decreasing the pore

diameter).

The growth rate of the oxide scale (stage II) is commonly described by

the classical parabolic rate law. Thus, the film thickness, L, is related to the de-

250 Chapter 9

position time, t, by the following equation:

L 2 = 2kpt (Eq. 9.6)

where kp is the parabolic rate constant. Figure 9.10 shows, as an example, the

plot of the square of the EVD LaCrO3 film thickness vs. time [9.35]. The

linearity of the plot indicates that the film growth is parabolic. The parabolic

rate constant, kp, can be derived using the Wagner oxidation process model. For

example, the parabolic rate constant for the EVD of YSZ has been derived and

is given as

o _1 1

RTVm~ (P/o~ //-~ (Eq. 9 7) - - P o ~ ) "

zF

where R is the gas constant, T the temperature, Vm the molar volume of YSZ, o ~

the electronic conductivity of YSZ at an oxygen activity of unity, z the valency

of oxygen, and F the Faraday. P/ and P// with subscript 02 are the oxygen

partial pressures at the chloride side and water vapor side, respectively. At the

deposition temperatures of 1000 ~ to 1200~ the parabolic rate constant for the

EVD of YSZ ranges from 1.1 • 10 .5 to 3.8 • 103 cmZ/s [9.37]. For the EVD

of MgO-doped LaCrO3, kp has been found to be about 0.18 • 10 .8 cm2/s at

1200

200

"E o 1000

'9, 0

x 800

-.2" if) to 6 0 0 u.I Z v

~ 4oo - r

0 1 r 0

1 1 2'0 40 ,o ,oo

1400

T IME, m in

Figure 9.10. Plot of square of EVD LaCrO 3 layer thickness vs time at 1330~ [9.35]


1330~ [9.35]. The parabolic film growth indicates that the growth process is controlled by the transport of charge species through the oxide film (although mass transport in the pores of the substrate is also proposed [9.30]). Thus, the

rate-limiting step in the EVD growth is the diffusion of electrons through the

oxide (in the case of YSZ) or the diffusion of oxygen ions through oxygen

vacancies (in the case of MgO-doped LaCrO3). In general, the morphology of EVD-deposited films depends on the deposition variables. Films deposited at low temperatures have highly faceted surfaces and show a preferred crystallographic orientation [9.42,9.43]. Films deposited at higher temperatures show a nearly smooth surface and no crystallographic orientation.

The EVD process has been scaled up for large-scale manufacturing [9.44-

9.46]. EVD reactors have been designed based on vacuum furnace technology

and can automatically process batches of as many as 60 cells [9.45]. A

schematic diagram of the scaled-up EVD process in the manufacture of sealless

tubular SOFCs is shown in Figure 9.11 [9.46]. The process relies considerably

VACUUM FURNACE

FILTERS ~ REACTANTS ~ ARGON . _ ~ ~ FUEL CELL

OXIDANT

REACTION CHAMBER PUMPING SYSTEM

CELL INTERNALS BACKUP PUMP PUMPING SYSTEM

Figure 9.11. Schematic diagram of EVD manufacturing process [9.46]

252 Chapter 9

on vacuum pumping systems. Two separate reactor vacuum systems are used,

one for the space surrounding the external cell surface and one for the internal cell space. This process has been successfully used in the production of cells for multikilowatt generators.

9.2.4 Performance and technological status

The sealless tubular SOFC single cell has been fabricated in a number of different sizes (from 20 to 100 cm in length). The performance and life of single

cells of this design have been extensively evaluated. Advancements in materials and processes have significantly improved cell performance over the years

[9.3,9.5,9.11,9.21,9.47-9.49]. Figure 9.12 shows typical performance

characteristics (voltage/current curves) of state-of-the-art single cells [9.11]. Sealless tubular cells have exhibited excellent and stable performance with a variety of fuels, including hydrogen [9.11], synthetic reformate gas [9.47], simulated coal gas [9.49], hydrocarbons [9.50], and natural gas [9.50]. Long- term operation of sealless tubular cells has been tested at various temperatures (from 875 ~ to l l00~ and current densities (from 250 to 450 mA/cm2). Early

cells typically lost 5 to 7 % of their voltages every 1,000 h of operation [9.51].

1.o

0.9

0 . 8

> 0.7 - u3 (9 < 0 . 6 - . _1

o > 0.5, -

0.4 -

0.3 -

0.2 1 oo

l t | ! 1 1 I I l l w t t

Fuel" 8 9 % H 2 + 1 1 % H~O _

O x i d a n t " Air Fuel U t i l i z a t i o n " 8 5 %

o.... O x i d a n t Ut i l izat ion: 2 5 %

"~ : :"-Z.,-,.,. ~ , , ~ ~ ~'t,...~

- - 0 - - 9 5 0 ~

�9 .-e--. 1025~

- - i x - - 1100~

200 300 400 500

CURRENT DENSITY, mA/cm =

Figure 9.12. Voltage~current curves of a sealless tubular SOFC single cell [9.11]


With process improvements, the degradation rate has been reduced to less than 1.5 % per 1,000 h. Several cells have surpassed 30,000 h of operation with

degradation rates in the range of 0.5 to 1.5 % per 1,000 h [9.6]. Sealless tubular SOFC generators of different power output levels have

been constructed and operated [9.9,9.14,9.52-9.59]. The classes of field-tested

generators include 0.4 kW (24 cells, 30-cm cell), 3 kW (144 cells, 36-cm cell),

5 kW (324 cells, 36-cm cell), and 20 kW (576 cells, 50-cm cell). Generators

have been tested on hydrogen, a mixture of hydrogen and carbon monoxide, reformate gas, methane, naphtha, and desulfurized pipeline natural gas. Generator performance has exceeded 5,000 h of operation. For methane and pipeline natural gas operation, the generator design incorporates an integrated internal reformer (Figure 9.13). In this design, a portion of the depleted fuel is recycled to supply steam to the reformer, and the heat from the exhaust gases is used to support the endothermic reforming reaction. Figure 9.14 shows the test

results of a 3-kW generator operated on desulfurized natural gas [9.57]. The generator produced a peak power of 3,039 W at 112 A. The performance of a

OXIDANT

I . . , , . . = . . .

I

T DESULFURIZED NATURAL GAS

EXHAUST

)ANT PLENUM

MBUSTION PLENUM

!PLETED FUEL PLENUM

FUEL CELL

DEPLETED FUEL

REFORMER

Figure 9.13. Sealless tubular SOFC generator design with integrated reformer [9.57]

254 Chapter 9

3 4

32 > ~ ao

~ 28 >

26

24

I FUEL: DESULFURIZED NATURAL GAS M FUEL UTILIZATION: 85% |

T=,: 9930C |

(970 h FROM START OF TEST) ..J,

0 2o

I I i I i

3.0

2.0

~ 1.0

i ~ I l T o . 0 40 60 80 100 120 0 20 40 60 80 100 120 CURRENT, A CURRENT, A

Figure 9.14. Voltage~current curves and power output of a 3-kW tubular SOFC [9.57]

20-kW generator on pipeline natural gas is shown in Figure 9.15. Recently,

25-kW systems (Figure 9.16) [9.60] have been tested [9.61], and 100-kW

generators have been designed as part of a program to move the technology

toward commercialization.

2O0 r ? FUEL = PIPELINE NATURAL GAS f OXIDANT = AIR

TEMPERATURE = 1000~ 1 5 0 >

< 1 0 0 " " " ~ - ~ - - - - - ' ' - I - ..I 0 >

50

i 1 i 1 1 i i l 1 i 1 i i l i i i 1 1 i i 1 i 1

O 500 10OO 1500 2 0 0 0 2 5 0 0 TIME, h

Figure 9.15. Performance of a 20-kW sealless tubular SOFC generator [9.51]


CENTRI . . . . . . BI

AUXILI, AIR Pt

FILTER/SILENCI

A IR IN LET MANIFOLD

SOFC GENERAT(

CONTROL VALVE

RECUPERATOR

ELECTRONICS COMPARTMENT

SU PPORT BASE

E) REFORMERS

. . . . . . . . . . JTLET

Figure 9.16. 25-kW sealless tubular SOFC system [9.60]

9.3 SEGMENTED-CELL-IN-SERIES DESIGN

The segmented-cell-in-series design was proposed in the early 1960s.

Several variations of this design have been developed, and modules of this design

in kilowatt size have been constructed and successfully operated.


The segmented-cell-in-series design consists of segmented cells connected

in electrical and gas flow series. The cells are either arranged as a thin banded

structure on a porous support (banded configuration) (Figure 9.17) or fitted one

into the other to form a tubular self-supporting structure (bell-and-spigot

configuration) (Figure 9.18). The interconnect provides sealing (and electrical

contact) between the anode of one cell and the cathode of the next. In this

design, the fuel flows from one cell to the next inside the tubular stack of cells,

and the oxidant flows outside. In the 'banded configuration where the support

tube (typically 1.5 to 2.5 cm in diameter) is used, cells can be made with

component thickness on the order of 100 to 250/xm. In the bell-and-spigot

configuration, individual cells form short cylinders of about 1.0 to 1.5 cm in

diameter. The cells are about 0.3 to 0.4 mm thick to provide structural support.

256 Chapter 9

ELECTROLYTE CATHODE INTERCONNECT~~

7'~~"~;"~;"~:"~;"~;"r " ' ~"~"~"~"~" ANODE /

FUEL

FUEL

OXIDANT ~.

Figure 9.17. Segmented-cell-in-series design (banded configuration) [9.1]

ELECTROLYTE

/

l OXIDANT

FUEL

CATHODE

INTERCONNECT

Figure 9.18. Segmented-cell-in-series design (bell-and-spigot configuration) [9.1]


Table 9.3 summarizes the typical properties of the components currently used in

the two configurations of the segmented-cell-in-series design.

TABLE 9.3

Properties of Segmented-Cell-in-Series SOFC Components

Banded Configuration Bell-and-Spigot Configuration

Component Material Thickness Material Thickness

Support AIzO 3 2 to 3 mm None

Electrolyte YzO3-stabilized ZrO2 110 to 150/~m YSZ 0.3 mm (YSZ)

Anode Ni/YSZ 80 to 110/zm Ni/YSZ 100/~m

Cathode Doped LaCoO3 150 to 200/~m Doped LaMnO3 300/zm

Interconnect NiA1 200 to 250/~m Doped LaCrO 3 < 1 mm

Electrical connection in segmented-cell-in-series design

Each tube of the segmented-cell-in-series design is a cell stack consisting

of a number of cells connected in electrical series. In an operating stack, the

current flows along the cathode of the first cell, traverses the electrolyte, and

travels along the anode. The current then flows across the interconnect to the

cathode of the second cell, and the process continues. Since the current travels

in the plane of the electrodes in this design, the current path length depends on

the cell size. Thus, each segmented cell is often made as short as possible to

reduce internal resistance losses. For example, present banded-configuration

cells are only about 22 mm wide [9.62]. Cells of the bell-and-spigot configura-

tion are 10 mm long [9.63]. The number of cells per tube is often limited to less

than 20 for practical reasons. At present, each tube of the banded configuration

(12 cells) produces about 35 W [9.64]. The tube of the bell-and-spigot

258 Chapter 9

configuration (10 cells) produces 20 W [9.65]. Segmented-cell-in-series tubes

are connected to form a module for practical power generation.

Gas manifolding in segmented-cell-in-series design

The gas manifolding concept for a module of the banded configuration

is shown in Figure 9.19 [9.66]. The top of the module consists of two

chambers: a fuel distribution chamber and a fuel exhaust chamber. Fuel cell

stacks (tubes) are hung on the bottom plate of the exhaust chamber. Fuel is fed

to the fuel distribution chamber and distributed to each stack through feed tubes.

Spent fuel exiting from each stack flows into the exhaust chamber. Oxidant is

fed at the bottom of the module and preheated in a heat exchanger. Spent

oxidant flows through the pipe at the center of the module and is ducted to the

heat exchanger to heat the incoming oxidant. This gas manifolding concept has

been used in the construction of a 1-kW module.

The gas manifolding concept for a module of the bell-and-spigot

configuration is shown in Figure 9.20 [9.67]. In this concept, fuel cell stacks are

assembled on ceramic support bases which are connected to metal pipes for gas

ducting. Fuel is fed from the metal pipe to the lower bore of the ceramic base

and flows to the stack through a feed tube. The fuel flows down the annular

space between the feed tube and the stack inner diameter. The spent fuel exits

FUE

ii AIR I ~ _

.Ldl~ EXHAUST FUEL

'~' STACK

AIR HEATER

EXHAUST AIR

Figure 9.19. Gas manifolding concept for banded configuration SOFCs [9.66]


~ j STACK

MIC BASE

EXHAUST ~ y FUEL

Figure 9.20. Gas manifoMing concept for bell-and-spigot SOFCs [9.67]

from the stack through the upper bore of the ceramic base. This gas manifolding

concept has been used in a 2-kW bench-scale module.


The segmented-cell-in-series design offers the advantage of improved stack efficiency. Cells connected in electrical and gas flow series waste less

power in resistance losses because the first cell in series has a higher output voltage. For example, four or five cells in fuel flow series can generate about

10 % more power output than a single cell of the same total active area; however,

the benefit of adding cells diminishes rapidly above five or six cells in fuel flow

series. Like the sealless tubular SOFC, the segmented-cell-in-series design improves the structural integrity of the segmented cells: the cells of the banded

configuration are supported on a strong support tube, whereas the electrolyte of

the bell-and-spigot configuration is thick and strong enough to be self-supporting.

Cell internal resistance is an important consideration in the segmented-

cell-in-series design. The long current path in the anode and cathode results in

significant resistive losses. Thus, the cell must be kept short to minimize the

path for current in the electrodes. Similar to the sealless tubular cells, the

relatively thick support tube of the banded configuration is a large diffusional barrier to the fuel, thus limiting cell performance. The self-supporting electrolyte

260 Chapter 9

of the bell-and-spigot configuration reduces the losses arising from gas transport

through a support tube. However, the thick electrolyte layer can significantly increase resistive losses. High-temperature gastight seals are required for the

segmented-cell-in-series SOFC. Fuel and oxidant must be separated by seals on

both ends of the stack, on the feed tube, and between each cell in the stack.

Thus, a large number of such seals is required in a segmented-cell module. In

addition, this design requires the fabrication of many more cells to achieve the

same power as the sealless tubular design.

9.3.3 Fabrication

Several processes have been developed for the fabrication of the

segmented-cell-in-series SOFC. The fabrication of a stack of the banded

configuration involves depositing component layers of a number of cells on the same support tube. The fabrication of a stack of the bell-and-spigot configuration

involves making short-cylinder cells, followed by joining the cylinders with interconnect materials.

Fabrication processes for the banded configuration

(i) Fabrication process based on EVD: This fabrication process is not currently in use. It was used to fabricate this banded configuration fuel cell until

the design was replaced by the sealless tubular design in 1980. The process used EVD for the electrolyte and interconnect layers. The key steps of this fabrication

process are described below: �9 The support tube (15% CaO-stabilized ZrO2) Was fabricated by

extrusion. Tube porosity was about 30%, and tube dimensions were 13 mm in

outer diameter and 1 to 1.5 mm in wall thickness [9.25,9.68]. �9 The anode (nickel/ZrO2) was applied by slurry dipping as 9-mm wide,

40-#m thick bands. The bands were separated by 1-mm insulating gaps that were formed by electrochemical etching of the continuous electrode coating. The

anode coating was sintered onto the support tube at 1600~

�9 The electrolyte (YSZ) was deposited on the porous support with anode

bands by EVD at 1200~ Areas which contacted to the anode were masked by thin powder layers. Electrolyte deposits over such loosely attached powder

layers were simply removed by mechanical means [9.25].


�9 The interconnect (MgO-doped LaCrO3) was also deposited by EVD (at

about 1300~ Early interconnects were made of Cr203 and deposited by CVD

[9.68]. �9 The cathode consisted of a porous zirconia skeleton impregnated with

a small amount (2 to 5 mg/cm 2) of activating oxides (e.g., Pr203). The skeleton

was also covered with a SnO-doped 111203 film. The zirconia skeleton was

formed by fixation of a loosely packed powder layer via EVD. The activating

oxides (used to minimize cathodic polarization losses) were obtained by

impregnation with a nitrate solution followed by thermal decomposition. The

111203 film was deposited by CVD.

(ii) Fabrication process based on plasma spraying: This fabrication

process uses plasma spraying to apply the electrolyte and interconnect onto the porous support. The electrode layers are made by flame spraying. In this design

of the banded configuration, the support is a porous (30 to 35 % porosity) A1203

tube 1 to 1.5 m long, 2 to 2.5 cm in diameter, and with walls 2 to 3 mm thick.

Presently, each support has 15 cells fabricated in the center third of the tube. To

date, the process has been used to fabricate fuel cell modules of this design up

to a 1-kW output. The fabrication of the components on the support tube (Figure

9.21) is as follows [9.62, 9.69-9. 72]: �9 Gastight AI203 layers are first applied on the support tube by plasma

spraying [9.69]. The starting powder has a mean particle size of about 10/zm.

The function of these thin A1203 layers is to provide gas sealing at intercon-

necting areas. Thin copper tapes are used for masking to make the required pattern on the tube.

,CATHODE !

] ,ELECTROLYTE SUPPORT TUBE / / / , A N O D E / / //INTERCONNECT ALUMINA ~ t ~ ,, / Ni-AI END CAP

Figure 9.21. Components of banded SOFCs of segmented-cell-in-series design [9.62]

262 Chapter 9

�9 The anode (80 to 110/xm thick, 25 mm wide) is coated on the support by acetylene flame spraying using NiO powder of 60-/~m mean particle size. The flame-spraying process produces an anode layer with sufficient porosity for gas transport.

�9 The YSZ electrolyte (110 to 150 ~tm thick, 22 mm wide) is plasma-

sprayed on the anode after appropriate masking. The mean particle size of YSZ powder used in the process is about 20 ~m. During spraying, the support tube

rotates around its axis while the spray gun traverses along the tube axis. A minimum thickness of about 100 #m is required to achieve gastightness without pinholes. The gastightness of the electrolyte has been improved through low-

pressure plasma spraying [9.64]. Recently, a coating technique based on a CO2 laser has been developed for making electrolyte layers [9. 73].

�9 The interconnect (200 to 250 ~m thick) is made of NiA1 (95.5 % Ni, 4.5 % A1) and fabricated by plasma spraying. Thin layers of NiA1 and CaO- stabilized ZrO2 mixtures are used to prevent the separation of the interconnect

from the electrolyte [9. 69]. The interconnects at the ends of the support tube are

covered with an A1203 coating to prevent oxidation during cell operation. �9 The cathode is a SrO- or CaO-doped LaCoO 3 layer (150 to 200 #m

thick) fabricated by acetylene flame spraying. The starting LaCoO 3 powder used in the spraying has a mean particle size of about 40/~m. To obtain high porosity for the cathode, starting powders with broad particle size distributions are commonly used [9. 74]. During the spraying process, the crystalline structure of the cobaltite cathode can be destroyed; however, the structure is recovered by

firing at 1000~ to 1100 oc for several hours in air [9. 71, 9. 74]. Figure 9.22 shows a photograph of banded SOFC tubes fabricated by the

plasma-spraying process.

Fabrication processes for the bell-and-spigot configuration

The fabrication process for early bell-and-spigot configuration SOFCs

was based on isostatic pressing and sintering to form bell-and-spigot, conical, or cylindrical electrolyte structures. Bell-and-spigot cells, machined from YSZ

tubes, are approximately 1.1 to 1.3 cm in diameter, 4 mm in wall thickness, and

1.7 cm in length [9. 75-9. 77]. Electrodes are applied by coating. Conical and cylindrical cells are about 2.2 cm in diameter, 0.5 to 0.6 mm in wall thickness,

and 1.1 cm in length [9. 78-9.83]. Electrodes are applied by plasma spraying.

The fabrication process for current bell-and-spigot configuration SOFCs is as follows [9.63,9.65,9.67,9.84-9.86]:


Figure 9.22. Photograph of banded SOFC tube (courtesy of MHI, Nagasaki)

�9 The electrolyte (YSZ) is made by axial pressing, green-state

machining, and sintering (1550~ in air). The sintered electrolyte is a hollow cylinder 1 cm long, 1.38 cm in diameter, and with a 0.3-mm wall thickness

[9.63]. The cylinder has a 1-mm-thick rim at each end. The end surfaces of the

electrolyte cylinder must be fiat and smooth for gastight joining to form a stack. The required tolerances for the ends (in terms of flatness and smoothness) are

achieved by machining and polishing. The wall surfaces of the electrolyte must be sufficiently rough to ensure adequate adhesion of electrode layers. Suitable

roughness for the surfaces is achieved by grinding. �9 The interconnect (doped LaCrO3) is a short gastight cylinder bonded

to electrolyte cylinders by a diffusion welding process (Figure 9.23). CaZrO3

layers are used to insulate the LaCrO3 from the electrolyte [9.67].

�9 The anode (nickel cermet), approximately 100/~m thick, is formed by

spraying. The electrode material (mixture of NiO, YSZ, and CeO:) is deposited

on the internal surface of the electrolyte using a suspension of ceramic powders

in an organic solvent. After sintering, the anode porosity is about 50%.

�9 The cathode (doped LaMnO3), about 300/zm thick and 50% porous, is also made by spraying.

264 Chapter 9

H 2

H20

CATHODE

ELECTROLYTE

NONCONDUCTIVE BOND MATERIAL

INTERCONNECT

ANODE

Figure 9.23. Interconnection in bell-and-spigot configuration SOFCs [9.65]

Plasma Spraying

Plasma spraying is a deposition process in which the desired material, in

powder form, is heated above its melting point while being accelerated by a gas

stream through an electric arc in a plasma spray gun (Figure 9.24). The molten

powder is directed at the substrate, and on impact, forms a thin layer on the

substrate surface.

A plasma spray gun creates its high temperature by passing a suitable gas

through an electric arc which is confined in the bore of a nozzle, thus constrict-

ing or pinching the arc to form a plasma of very hot, fast-moving ionized matter.

| ARC

GAS POWDER

COATING i

SUBSTRATE

Figure 9.24. Schematic diagram of plasma-spraying process


The arc is started inside a small chamber of the spray gun. One end of the

chamber is a front electrode, perforated at its center to provide an orifice for the plasma. The other end of the chamber contains the back electrode. To withstand

the intense heat during spraying, the electrodes are cooled by circulating water.

The gas is introduced into the gun chamber. The powder is fed into the nozzle

bore or into the plasma just beyond the end of the nozzle. Common feed

mechanisms are based on the aspirator or on mechanical metering.

Plasma gun geometry has a marked effect on spray efficiency. The

nozzle diameter is an important variable, since changes in it affect the arc gas

velocity, the current density, the powder velocity, and particle trajectory. In

general, increasing nozzle diameter increases spraying efficiency, but weakens

the adhesion of the deposit. The choice of the gas affects spraying efficiency and

deposit quality. Nitrogen, hydrogen, argon, and helium are commonly used.

Since, for a given material, the deposition efficiency has an optimal value at a

given arc enthalpy, the correct gas should be selected for a particular material.

Purity, uniformity, and stability of the powder are critical in plasma spraying.

In any case, the powder should be dry and fairly free-flowing.

Key plasma spray process parameters include power, gas flow rate,

powder size and feed rate, spray rate, surface speed, and gun-to-work distance.

Altering one or several of the spray conditions may affect the physical, chemical,

or mechanical properties of the deposits, as well as the adhesion of deposits to

substrate. Surface preparation and temperature control of the substrate are also

important factors. Some of these parameters are discussed further below:

�9 Gun-to-work distance: Work distance should be kept fairly constant for

a given application. In general, the spray gun should be held at a distance of 5

to 15 cm from the substrate.

�9 Surface speed: Traverse speed should be such that not more than 0.2

to 0.3 mm is applied in each pass.

�9 Preheating: The substrate should be preheated to 100 ~ to 150~ This

prevents surface condensate, expands the substrate, and reduces stress in the

deposit when it subsequently cools.

�9 Spray rate: In any spray application, the shape and size of the

substrate, the powder particle size, and the substrate material determine the spray

rate for a given powder.

�9 Powder size: The mesh size of the powder determines the spray rate

and efficiency, as well as deposit density and surface finish. The maximum

266 Chapter 9

spray rates and deposition efficiency are usually attained with powders of-200

+ 400 mesh sizes. �9 Surface preparation: Surface preparation of the substrate is important

to obtaining good adhesion of the deposit. All dust, oil, and other foreign matter

must be removed from the surface. Since the bond in a plasma spray deposit is

predominantly mechanical, the surface should be roughened where possible.


Bell-and-spigot SOFCs of various sizes have been constructed and tested. In early construction, up to 20 bell-and-spigot cells were connected in series to

form a tubular stack about 24 cm long. A 20-cell stack, tested with hydrogen and air, showed an open-circuit voltage of 21 V (1.05 V per cell) at 1020~ and

produced a maximum power of 10.5 W [9.76,9.87]. Twenty of these 20-cell

stacks were incorporated into a generator. At 1000~ and 60% fuel utilization, the open-circuit voltage of the generator was 200 V, and the maximum power output was 102 W [9.87]. Later bell-and-spigot fuel cells showed significant

performance improvements. A generator consisting of four 30-cell stacks (35 to 40 cm long) produced 115 W (corresponding to 0.22 W/cm 2) at 1000~ with

hydrogen and air [9.83]. A single cell of this type was operated for more than 28,000 h. A cell voltage of about 0.8 V was obtained at a constant load of 120 mA/cm 2. Most recent bell-and-spigot SOFCs are based on the design developed

for hydrogen production by electrolysis of water vapor [9.67]. A laboratory-

scale generator of this type was successfully operated, producing about 2 kW.

The generator consisted of 100 stacks, each containing 10 cells. Banded SOFCs up to 1 kW have been fabricated and operated. Early

stacks of up to 25 connected cells (produced by the EVD process) were tested

under a variety of operating conditions [9.25, 9.49, 9.88, 9.89]. An EVD seven-

cell stack demonstrated 5,000 h of life at 1000~ The stack generated a current density of 400 mA/cm 2 at 0.72 V/cell with hydrogen fuel. This stack was also

tested with simulated coal gas (400 mA/cm 2 at 0.67 V/cell). The presence of 50

ppm H2S impurity resulted in a decrease of about 5 % in operating cell voltage;

however, this effect was found to be reversible. The stack was subjected to 11

thermal cycles (1000~ to room temperature) without any significant perfor-

mance degradation. Presently, banded SOFCs are being developed using the plasma-spraying

process. Plasma-sprayed stacks containing 4 to 20 connected cells have been


tested to demonstrate cell performance and the feasibility of the fabrication technology [9.62, 9.69, 9. 71, 9.90, 9.91]. In general, early stacks showed relatively low open-circuit voltages. For example, a 12-cell stack showed an open-circuit voltage of 11.16 V (0.93 V/cell) at 1000~ with hydrogen and oxygen [9.69]. (The maximum power output of this stack was 82.4 W.) A

generator of 48 stacks (15-cell stack) showed an open-circuit voltage of 0.89 V/cell [9.91]. Significant fuel leakage was observed. (The maximum power output of this generator was 1.2 kW.) Recent modifications in material, design, and fabrication process have resulted in cell performance improvements as shown in Figure 9.25 [9.64]. A 1-kW generator of this design (48 stacks, 12-cell stack) has been operated for more than 1,000 h at 900~ with hydrogen and air. The operating voltage is 120 V, and the current is 10 A (200 mA/cm/). A voltage degradation of 3 % per 1,000 h has been observed [9.92]. Figure 9.26 shows the voltage/current characteristics and power output of this generator [9.66].

i ,,, T e m p e r a t u r e : 9 0 0 ~ .~

/ Fuel : H y d r o g e n . - - ~ . . . . O x i d a n t : Ai r . ,.J:~"T ~ "

,J::~"".~l. ~ ~ ' ' ' " IMPROVED

.Jr " ~ . (9 < 10 --.. -J x / f ..

5 0 >

0 1 2 3 4 5

CURRENT, A

4 0

30

20 i i i

o

10

0 6

Figure 9.25. Performance of conventional and improved banded SOFCs [9.64]

2oo I

1oo ",%

FUEL = HYDROGEN OXIDANT = AIR TEMPERATURE = 900~

I | 0 5 10 15 20 25

2 . 0

- 1 . 5

- 1 . 0 i i i

0

- 0 . 5

) .0 30

CURRENT, A

Figure 9.26. Voltage~current and power output of 1-kW banded SOFC [9.66]

268 Chapter 9

9.4 MONOLITHIC DESIGN

The monolithic design is the newest SOFC stack concept. The

monolithic SOFC stack consists of many cells fabricated as a single unit. The

design has the potential to achieve high power density because of its compact and

lightweight structure. The feasibility of the monolithic design has been demonstrated on a laboratory scale.


The monolithic SOFC consists of thin cell components formed into a

corrugated structure of either gas coflow or crossflow configurations (Figure

9.27) [9.1, 9.93-9.95]. The fuel cell of this design is made of two types of multilayer ceramics, each composed of three components" anode/electrolyte/

cathode and anode/interconnect/cathode. The laminates are typically 200 to 300 #m thick. In the coflow version, the fuel cell consists of alternating layers of corrugated anode/electrolyte/cathode laminate and flat anode/interconnect/cathode laminate. Fuel and oxidant flow parallel in adjacent channels formed by the laminated layers. In the crossflow version, the fuel cell consists of alternating flat layers of anode/electrolyte/cathode laminate and anode/interconnect/cathode laminate, separated by corrugated anode and cathode layers. The anode and cathode corrugations are oriented at right angles to each other. The major differences between the two versions are power density and gas manifolding. The crossflow version shows a reduced power density when compared with the

COFLO CATHODE . , , , , ~_ CROSSFLOW

INTE . .

CATHODE ~~~~~~ ELECTROLYTE \

O X I D A N T ~ J ANODE ~ ' ~ _ . ~ ~ ' ~ . . CATHODE " ELECTROLYTE OXIDANT

FUEL - "~.~"~.~..,,,,,,P"~ CATHODE

Figure 9.27. Monolithic SOFC design (coflow and crossflow configurations) [9.1]


coflow. On the other hand, the crossflow offers a simpler means of ducting

gases into and out of the fuel cell ceramic structure. Typical properties of the

components of the monolithic SOFC are summarized in Table 9.4.

TABLE 9.4

Properties of Components of Monolithic Solid Oxide Fuel Cells

Component Material Thickness

Layer Electrolyte Anode Cathode Interconnect

Y203-stabilized ZrO2 Ni/Y~O3-ZrO2

Doped LaMnO3 Doped LaCrO3

50 to 150/~m 50 to 150/zm 50 to 150/~m 50 to 150 ~m

Corrugation (1 to 2 mm high) Single cell (coflow) Anode (crossflow) Cathode (crossflow)

200 to 300/~m 200 to 300 tzm 200 to 300/zm

Electrical connection in monolithic design

A monolithic SOFC stack is an array of cells connected in electrical

series. In an operating stack, the current traverses the multicell monolithic

structure. As an example, the current path in the coflow fuel cell is shown in

Figure 9.28. In this coflow configuration, the current (coming from the

adjoining cell) flows through the plane of the interconnect, then part way around

the circumference of the cell in the plane of the cathode. The current then

traverses the electrolyte and flows part way around the circumference of the cell

in the plane of the anode. The current then travels across the plane of the

interconnect into the next cell, and the process continues. In the monolithic

SOFC, the small distance between the interconnect layers decreases the current

path length in the electrodes. This, along with thin electrolyte and interconnect

layers, reduces the voltage losses due to internal resistance in the monolithic

design. The power density for the coflow monolithic SOFC is calculated to be

about 8 kW/kg or 4 kW/L (fuel cell only) [9.96].

270 Chapter 9

/ ELECTRON/ION PATH

ANODE

ELECTROLYTE

CATHODE

INTERCONNECT

Figure 9.28. Electron~ion current path in monolithic SOFC

Gas manifolding in monolithic design

The proposed gas manifolding concept for the coflow monolithic SOFC

is shown in Figure 9.29 [9.97]. The coflow manifold design involves a transition

section. The fuel and oxidant channels are separated in the transition sections by

ducts at the end of the active region as shown in Figure 9.29. The transition

ducts, each only half the height of the corrugation in the active region, turn the

gas flow 45 degrees to the gas flow in the active area. (Turning fuel and oxidant

ducts at right angles in the transition region forms separate faces for fuel and

oxidant manifolds.)

Crossflow monolithic SOFCs reduce the complexity of the gas manifold-

ing design; in the crossflow monolithic SOFC, fuel flows in one face of the stack

and out the opposite face, as does oxidant.


The key features of the monolithic design are small cell size and high

power density. The small cell size increases the active surface area and reduces

resistive losses due to short electron/ion paths. In the monolithic SOFC, the

current conducts in the plane of the thin electrolyte and interconnect and travels

short distances in the electrodes. Thus, internal resistance is low. As a result


is-\

\ _ f s z ~ / o

j 38

-4- r '-!- 4 d-

T I

.i./ 5

S \

Figure 9.29. Gas manifolding concept in coflow monolithic SOFC [9.97]

of low resistance, the monolithic SOFC can be operated at higher current densities than other designs while achieving the same output voltage. The monolithic design offers high power output per unit mass or volume. The high power density results from the higher active surface area, higher current density, and lower weight. The lower weight in the monolithic design is primarily the

result of eliminating inactive structural supports. The development of suitable materials and fabrication processes is critical

to the monolithic design. The main disadvantage of the monolithic design is the

difficulty of fabricating the corrugated structure. Because of its intricate

structure, the monolithic fuel cell is made by cofiring. Thus, the structural

integrity of such a structure depends upon one important factor: matching thermal

expansion and firing shrinkage of the four cell components. Any significant

mismatch in thermal expansion and shrinkage can cause stress in the fired bodies

and result in cracking during processing and operation. In addition, the

monolithic design requires elaborate nondestructive evaluation (NDE) of the

fabrication process to ensure reliable manufacture.

272 Chapter 9

9.4.3 Fabrication

The current approach to the fabrication of the monolithic SOFC is to

form the fuel cell in the green state and cofire the green body at elevated

temperatures to produce the sintered structure. A fabrication process for the

monolithic fuel cell must incorporate the component materials into the ceramic

structure with the following properties: (i) dense electrolyte and interconnect,

along with porous anode and cathode, (ii) good interfacial bonding between

adjacent layers, (iii) insignificant interaction and interdiffusion between

neighboring components, and (iv) reliable and defect-free structures.


Tape casting [9.98-9.100] and tape calendering [9.101-9.103] have been

developed for the fabrication of thin ceramic layers and multilayers required for

building the monolithic fuel cell. In the tape casting process, tapes of cell

components are made by spreading a ceramic slip (or slurry) to a thin layer of

controlled thickness with a doctor blade. Examples of the slip formulations used

in casting the monolithic SOFC materials are given in Table 9.5 [9.100].

TABLE 9.5

Slip Formulations for Tape Casting of Monolithic SOFC Materials

Material Weight Ratio

Component Ceramic/ Ceramic/ Ceramic/ Plasticizer/ Binder Solvent Dispersant Binder

Anode 5.6 2.8 22.2 0.31 Cathode 5.0 2.8 25.0 0.30 Electrolyte 5.0 2.8 44.4 0.30 Interconnect 6.6 2.7 28.8 0.35

Multilayer tapes required for the monolithic SOFC are formed by

sequentially casting one layer on top of another. Tapes are corrugated by folding


them onto a warm mold. Green, flat and corrugated tapes are bonded in the

appropriate order and orientation to form the required monolithic structure. The

stack is then cofired at elevated temperatures. Figure 9.30 shows a typical

microstructure of a single cell and a photograph of a crossflow stack fabricated

by tape casting.

At present, the development of the fabrication process for the monolithic

SOFC focuses on tape calendering. A schematic diagram of the fabrication

sequence based on this technique is shown in Figure 9.31 [9.103]. In this

process, ceramic powder, organic binder, and plasticizer are first mixed in a

high-shear (high-intensity) mixer. The friction resulting from the mixing action

heats the batch to form a mass with a doughy consistency that has many of the

same characteristics as a plastic. The mass is then sheeted into a thin tape using

a two-roll mill. The mill has two counterrotating rolls with independent, variable

speed and temperature controls. The spacing between the two rolls is adjustable,

and tape thickness is controlled by the spacing of the two rolls. The tape formed

by this process is flexible and can be cut to size, laminated, corrugated, or

otherwise formed before firing.

Multilayer tapes are formed by laminating individual layers in a second

rolling operation. The mechanical force during this second rolling operation

I Figure 9.30. Typical microstructure of single cell, and photograph of crossflow monolithic

SOFC fabricated by tape casting

274 Chapter 9

TAPE FORMING ~ 0 .:ii!;~ '

MOLDING~ ~~

l 1

SINTERING BINDER EXTRACTION BONDING AND ASSEMBLY

Figure 9.31. Fabrication process based on tape calendering for monolithic SOFC [7.103]

bonds the layers together. Figure 9.32 shows, as an example, the cross section of an anode/electrolyte/cathode tape made by this technique [9.104]. Corrugated layers are made by compression molding. (Compression molding involves

Figure 9.32. Cross section of typical single cell tape made by calendering [9.104]


placing a tape of plastic mix between the platens of a shaped die and applying

uniaxial pressure until the tape deforms to the shape of the die cavity.) To

produce a well-formed corrugated layer, uniform pressing pressure and good

material flow are necessary. A proper combination of temperature, pressing

pressure, pressing time, and tape thickness must be used to mold corrugations

having the desired dimensions. To form a monolithic SOFC structure,

corrugated and fiat layers are stacked in the proper sequence and bonded in the

green state to yield a stack. Figure 9.33 is a photograph of a coflow and a

crossflow stack built from calendered tapes [9.104]. Green stacks are cofired at

elevated temperatures (1300 ~ to 1400~ to remove the organic binder and sinter

the ceramic structure. Examples of sintered coflow and crossflow stacks

fabricated by this process are shown in Figure 9.34 [9.105,9.106].

Successful fabrication of the monolithic SOFC structure by cofiring

depends upon several critical factors in starting powder characteristics, green- state forming, and firing:

�9 Starting powder characteristics (especially surface area and particle

size) influence the fabrication of the monolithic SOFC in many ways, including

the amount of binder required to form the green body, sintering behavior, and

the sintered porosity/density and microstructure. For example, for high-density

layers (such as electrolyte and interconnect), the desired powder would have

small particle size and high surface area. However, these characteristics have

Figure 9.33. Green coflow and crossflow monolithic SOFCs [9.104]

276 Chapter 9

Figure 9.34. Sintered coflow and crossflow monolithic SOFCs

two important effects on the firing behavior of the layer. First, high organic binder content is required to coat each small particle. Due to the high organic content, shrinkage during the binder burn-out stage is relatively high. Second, due to the high surface area of the powder, shrinkage during the sintering stage is also high. The high firing shrinkage may result in mismatch with the other fuel cell layers. Thus, the characteristics of each material must be tailored to match the firing shrinkage profile and to produce cell components with the desired properties. An example of the tailoring of cathode powders (to attain the desired porosity and to adjust the firing shrinkage profile) is given in Figure 9.35 [9.107].

70

60

50

~ 40

0

0

2O

10

2O

I I I I

FABRICATION GOAL

I00 "/, SOLID ~ .... STATE "" 20 % LIQUID MIX tZ.Z .,z/g} '~

I I " ~ 4 0 % LIQUID MIX

I ! \

25 30 35

FIRING SHRINKAGE, %

I00 % LIQUID MIX ~ (lO.I mZ/gl

40 45

Figure 9.35. Effect of powder characteristics on cathode porosity and firing shrinkage [9.107]


�9 Green-state forming parameters must be carefully controlled to achieve

acceptable results in the fabrication of the monolithic SOFC. For example, a compromise between formability and sinterability is required when the solids loading is determined. A high solids content facilitates densification, with less shrinkage. On the other hand, high solids loading produces less flexible tapes, leading to difficulty in lamination and corrugation. Mixing and rolling temperatures are carefully controlled to prevent binder loss. Excessive organic material loss may result in a brittle tape that is difficult to form without cracking.

�9 The firing process is extremely important in the fabrication of the

monolithic fuel cell. In general, there are two stages of shrinkage during firing of a ceramic tape. At low temperatures, a small percentage of shrinkage of the

tape is observed as the organic materials are burnt out, and the ceramic particles move closer together. This amount of shrinkage is dependent on the ratio of binder to ceramic. At high temperatures, a large percentage of shrinkage is due to sintering of the ceramic particles. The shrinkage due to high-temperature sintering is dependent on the ceramic particle size, surface area, and volume loading. In firing the monolithic SOFC laminated structure, it is as critical to

match the shrinkage profiles for each of the cell components as it is to match their total firing shrinkage. Any mismatch in the firing shrinkage can cause

stress in the fired bodies, resulting in cracking and delamination. Figure 9.36 shows the firing shrinkage profiles of four monolithic SOFC tapes. As can be

seen in Figure 9.36, starting powder characteristics of the four materials must be

40

30

20 -

z n.- " r l O -

0

2OO

COMPONENT SHRINKAGES BEFORE TAILORING POWDERS

- ELECTROLYTE

............... ANODE . /

/ / INTERCONNECT / , / . CATHODE / ' , / .-"

�9 . . . ' "

/ : "

I I I | I I i _ _ _

600 1000 1400 1800

TEMPERATURE, K

40

COMPONENT SHRINKAGES AFTER TAILORING POWDERS

30

~ 20

z

0

200

--"-------- ELECTROLYTE

.............. ANODE

m _ . _ INTERCONNECT _

CATHODE / . ~ / , ~

,__)r __ . I i I i l I I

600 1000 1400 1800 TEMPERATURE�9 K

Figure 9.36. Firing shrinkage profiles of four monolithic SOFC tapes

278 Chapter 9

tailored to obtain firing shrinkage match. Another important consideration in the fabrication of the monolithic SOFC is achieving a good match of coefficients of thermal expansion among the cell materials. This matching is essential in minimizing thermal stresses generated in the fuel cell structure during cooldown after sintering [9.108, 9.109]. (This subject is discussed further in Chapter 10.)

Tape calendering

Tape calendering to form ceramic materials involves squeezing a softened thermoplastic polymer/ceramic mix between two rolls to produce a continuous sheet of material. In this process, a plastic mix is passed between two counterrotating rolls (Figure 9.37). The mix is compacted, as well as pressed to

a thickness equivalent to the spacing of the rolls. Multiple passes, at diminishing roll separation, can yield a constant-thickness tape of high uniformity. Although the basic forming operation is completed by the two-roll mill, a high-intensity mixer has to be employed for the production of the plastic mix (Figure 9.37). The tape-calendering process has been used to make thin ceramic sheets [9.110, 9.111] and to fabricate resin-bonded and rubber-bonded grinding wheels

for many years [9.15].

PLASTIC-UKE MASS

THIN TAPE

MULTILAYER TAPE

HIGH-INTENSITY MIXER TWO-ROLL MILL

Figure 9.37. Tape-calendering process


The tape calendering process involves two key steps: a mixing step to incorporate ceramic powder into an organic binder system and a rolling step to form thin ceramic tapes.

In the mixing step, ceramic powder is dispersed in an organic binder system using a high-shear mixer. The binder system for tape calendering is

composed almost exclusively of two components: binders and plasticizers. The main functions of the organic binder (a thermoplastic polymer) are to wet the

ceramic powder to aid dispersion and minimize entrapped gas, to confer

sufficient plasticity upon the powder for rolling, and to impart adequate tear

strength to the tape during forming and sufficient rigidity during initial stage of binder burnout. The plasticizer is an organic material added to the binder to

increase the plasticity (fluidity) of the ceramic/binder mix.

Mixing the ceramic with the binder polymer in a high-shear mixer is a

compounding process (where the polymer is softened, melted, and compacted, more or less to a continuum, with the dispersed ceramic powder). This type of

mixing is dispersive in nature. A typical high-shear mixer (often referred to as an internal mixer) consists of a figure-eight-shaped chamber which fits over two

signoid, counterrotating blades (Figure 9.38). The mixing process in an internal

DUST REMOVAL FAN

III DOOR

PASSAGES FOR MIXING ROTORS COOLING WATER

POWER OPERATED SLIDING DOOR

Figure 9.38. Sectional diagram of high-shear mixer

280 Chapter 9

mixer can be followed by monitoring the torque required to produce a uniform mix. Mixing torque is proportional to the viscosity of the mix and is an indication of the work required to mix the ingredients. Typically, the torque

increases rapidly as the mixing starts, reaches a peak, and then drops to a steady state. The peak indicates the melting and mixing steps. Once the torque reaches

a steady state, no additional mixing occurs, reflecting a uniformity in mix viscosity. Mixing time is the time required for a mix to reach a steady state.

The plastic mix is sheeted into a thin tape using a two-roll mill. The two counterrotating rolls of the mill, commonly made of chromium-coated stainless

steel, have independent, variable speed control. The temperature of the rolls is regulated by independently controlled circulating steam under pressure, hot water, heated oil baths, or electric resistance heaters. During rolling, the

ceramic tape being squeezed between the two rolls is elongated [9.110, 9.112]. In theory, the reduction in tape thickness (compaction) is inversely proportional

to the increase in length (elongation) (assuming the tape is incompressible). However, in practice, the actual compaction is generally higher, especially during

the first few rolling steps when the tape is still relatively thick. In general, the thickness of the tape after calendering is greater than the nip (the spacing

between the faces of the rolls), indicating a streamline pattern between the rolls

[9.110]. As a result, a sturdy calendering machine must be used to keep the expansion of the nip width within acceptable limits while the tape is being squeezed.

The important factors in tape calendering are roll speed, roll temperature, and number of passes. Good temperature control of the roll surface and a high number of passes are essential for producing tapes of uniform thickness. During calendering, the squeezing of the tape in the nip area generates heat. The tape temperature rise must be carefully controlled to prevent binder evaporation.

Evaporation of organic materials may cause gas bubbles to form in the tape and may result in a brittle tape that is susceptible to cracking during forming.


Monolithic SOFC single cells have been fabricated and operated on a

variety of fuels at 1000~ [9.113-9.116]. Single cells (anode/electrolyte/cathode

laminates) up to 10 by 10 cm in size have been cofired. The essential features of desired microstructures and strong bonding between layers have been achieved

(Figure 9.39). Figure 9.40 shows an example of performance curves obtained


ANODE ELECTROLYTE CATHODE

Figure 9.39. Micrograph of fracture surface of monolithic SOFC single cell

1 . 2 / O : 900~

1.1 ~-'~ o o : 1000 o C O~ "0~^ ZX: 1100~

1.0 FO-o v~

:>'uJ 0 " 9 / v ~ ^ O_

< 0.8 O ~ o " ~ O _ . . . . . n ~ ~o~^ ~ o > 0.7

0.6

0 . 5 Fuel: Hydrogen Oxidant: Air

0.4 . . . . . . . . . . . . . . . . . . . . . . . 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8

CURRENT DENSITY, A / c m 2

Figure 9.40. Performance curves of monolithic SOFC single cell

for a monolithic SOFC single cell. Low area-specific resistances and high

current densities have been achieved for the monolithic fuel cell; monolithic

SOFCs have been operated at current densities > 2.2 A/cm 2 with hydrogen as

fuel and air as oxidant. Operation of monolithic cells on hydrocarbons, alcohol,

natural gas, coal gas, and simulated diesel has been demonstrated [9.113, 9.117].

Fabrication of multicell monolithic SOFCs has concentrated on the

crossflow configuration. Crossflow stacks of up to 5 by 5 cm in footprint area

and 10 cells in height have been produced [9.106]. To date, testing has been

performed only on small laboratory-scale stacks. Figure 9.41 shows an example

of the polarization curve of a two-cell stack (6.25-cm 2 footprint area) obtained

282 Chapter 9

2.0

1.8

1.6 > u3 (.9 1.4 < I- " 1.2 0 >

1.0

0.8

0.6

F U E L = H Y D R O G E N

O X I D A N T = A IR

T E M P E R A T U R E = 1 0 0 0 ~

" 0

i I i ~ i I I i 0 O. 0.2 0.3 0 .4

CURRENT DENSITY, A/cm 2

Figure 9.41. Performance curve of monolithic SOFC two-cell stack

at 1000~ with hydrogen as the fuel and air as the oxidant. The area-specific

resistance of the stack is 1.0 f~.cm 2 per cell [9.117]. A two-cell crossflow stack

(9-cm 2 footprint area) has been operated under 50 mA/cm 2 for more than 700 h

[9.115]. The current technical challenge in the development of the monolithic

SOFC is the cofiring of the LaCrO3 interconnect [9.117, 9.118]. The LaCrO3 interconnect can densify if fired alone in air at 1400~ however, the material does not densify when fired in contact with the electrodes under similar firing conditions. During cofiring, the liquid phase in the LaCrO3, which is responsible for densification, tends to migrate into the electrode layers, rendering the interconnect porous. Suitable methodologies for cofiring the interconnect are being investigated.

9.5 FLAT-PLATE DESIGN

The flat-plate design, common in other types of fuel cell, has received much attention recently. The design offers simple cell geometry and multiple

fabrication options. To date, kilowatt-level multicell stacks of the fiat-plate design have been tested.


The flat-plate design consists of cell components configured as thin,

planar plates. Common plate shapes are rectangular (square) or circular.


Typical flat-plate SOFCs use the single cell and interconnect as structural

components. Individual cells are typically thicker than 200 /~m to be self-

supporting. The electrolyte layer is often the thickest component of the single

cell, thus acting as the cell structural support (although in certain cell designs,

thick electrode layers have also been used as substrates [9.119, 9.120]). The

interconnect serves as a bipolar gas separator, contacting the anode and cathode

of adjoining cells. The interconnect frequently has ribs on both sides to form gas

channels (Figure 9.42). Typical properties of the components of the fiat-plate

SOFC are summarized in Table 9.6.

TABLE 9.6

Properties of Components of Flat-Plate Solid Oxide Fuel Cells

Component Material Thickness

Electrolyte Y203-stabilized ZrO2 50 to 250/~m Anode Ni/YEO3-ZrO2 25 to 100 #m Cathode Doped LaMnO3 25 to 100/~m Interconnect Doped LaCrO3 or 200/~m to 1 mm

high-temperature alloy (2 to 6 mm including rib height)

FUEL

OXIDANT

INTERCONNECT

ANODE ELECTROLYTE

CATHODE

Figure 9.42. Flat-plate SOFC design [9.1]

284 Chapter 9

Electrical connection in flat-plate design

Flat-plate SOFC stacks are formed by stacking cell components to the

desired stack height. In an operating stack, the current (coming from the

adjacent cell) flows from the interconnect to the cathode at the contact point.

From the contact point, the current flows in the plane of the cathode, distributes

over certain electrolyte area, and travels across the plane of the electrolyte to the

anode. The current then flows in the plane of the anode to the nearest

interconnect contact point and finally across the plane of the interconnect to the

next cell, where the process continues.

Gas manifolding in fiat-plate design

The common gas manifold concept for rectangular (or square) fiat-plate

SOFCs is based on the crossflow configuration. Figure 9.43 shows a conceptual

design of a manifolded stack. A key feature of this design is the large external

manifolds which feed the gases to and collect gases from the flow channels in the

individual cells. This design requires a gas seal between the environment and the

interior of the manifold. The manifold seal must be an electrical insulator to

FUEL O U T ~

OXIDANT IN

OXIDANT OUT

FUEL IN

Figure 9.43. Gas manifolding concept in crossflow fiat-plate SOFCs


prevent cell-to-cell electrical shorts. Other variations of this manifolding concept

have been proposed [9.121-9.124]; an example is shown in Figure 9.44. In

addition, the integral manifold [9.125, 9.126] and the gas counterflow configu-

ration [9.126] have also been considered (Figure 9.45).

The common gas manifold concept for circular fiat-plate SOFCs involves circular or radial flow with the gas inlet and outlet located inside the fuel cell

stack [9.127-9.130]. Seals are required in the gas inlet and outlet ports to shield

the cathode from the flowing fuel in the fuel ports and the anode from the

flowing oxidant in the oxidant ports. Figure 9.46 shows a schematic diagram of

the gas manifold design (incorporating a heat exchanger) for a circular flat-plate

SOFC [9.1311.

OXIDANT

FUEL

FUEL

BASE PLATE

METALLIC WINDOW FOIL

ELECT RE) LYTE/ELECTRODE SEALING LEDGE


METALLIC BIPOLAR PLATE


ELECTRO LY'I'E/ELECTRO DE SEALING LEDGE


~" + . ~_L____ ~-i-- ~ BASE PLATE

Figure 9.44. Stacking and gas manifolding in fiat-plate SOFCs [9.121]

FUEL

2 8 6 Chapter 9

SEAL AREA INTERCONNECT PLATE

RIB AREA

CATHODE

ELECTROLYTE

GAS FLOW CHANNEL

GAS MANIFOLD

Figure 9.45. Counterflow flat-plate SOFC with integral manifold [9.126]

OXIDANT

t

EXHAUST

1L

FUEL DISTRIBUTION CURRENT COLLECTOR PASSAGE

\ INTERCONNECT

..! \ I i I I I ~ ~L~ < ~ OXIDANT

A ~ I I ! I I I I ~ i ~ ~ ~ ,- I kr'~,~,-~ ........... i ................. / n r . . . . . . . . . . . . . . . . . . . . . . ~ i ~ , , ~ lVl~ ~ 1 ~ - - ~

OXIDANT INLET ~ �9 ::::::::::::::::::::::::::::::::::::::::::: OXIDANT \ CELL

SUPPORT BODY =~>

FUEL

Figure 9.46. Gas manifold design for circular flat-plate SOFCs [9.131]



The flat-plate design offers improved performance and higher power

density relative to the sealless tubular and segmented-cell-in-series designs. Because of cross-plane conduction, internal resistance losses of fiat-plate SOFCs

are independent of cell area. Thus, cell components can be made very thin to minimize electrical resistance. The fiat-plate design provides more flexibility

than the other designs in terms of cell geometry and gas manifolding. For example, the fiat-plate cell components can be made into square, rectangular,

circular, or hexagonal shapes. The fiat-plate design is also simpler to fabricate. The two dense components, the electrolyte and the interconnect, can be fabricated

independently. This avoids the difficulties in cosintering LaCrO3 interconnect, minimizes chemical interactions during firing, and provides multiple fabrication options. Furthermore, the fabrication of the fiat-plate SOFC allows cell components to be assessed individually, ensuring better quality control. In addition, due to the nature of their fabrication and assembly, fiat-plate SOFCs can readily incorporate different materials such as metallic interconnects.

The fiat-plate SOFC requires high-temperature gas seals at the edges or inside the gas ports of the plates. Compressive seals, cement seals, glass seals,

and glass-ceramic seals have been proposed. However, the unforgiving nature of a compressive seal can lead to nonuniform stress distribution on the ceramic

and cracking of the layers. Cements and glasses tend to react with cell materials at the 1000~ operating temperature. Further, seals may limit the height of a cell stack. There is a higher probability of mismatches in tolerances (creating unacceptable stress levels) in taller stacks. Contact resistance can be relatively high in fiat-plate SOFCs. Some configurations incorporate limited contact area (via the interconnect) between the anode of one cell and the cathode of another. As a result, there is long path for the current in the plane of each electrode, and the resistive losses can be large. Stacking large, thin sintered ceramic layers is

expected to be difficult; this may set a limit on practical cell size for the fiat-plate design.

9.5.3 Fabrication

Fabrication and assembly appear to be simpler for the flat-plate design

as compared with the other designs. The nature of the fiat-plate design also

288 Chapter 9

permits a variety of fabrication options. by a number of fabrication methods.

Flat-plate SOFCs have been produced


The most common fabrication technique for the flat-plate SOFC is tape casting. Tape casting involves uniformly spreading a slurry of ceramic powders

dispersed in a solvent (containing various organic ingredients) onto a smooth surface, where the volatile solvent is removed. The resulting dried tape is stripped from the casting surface and fired to yield sintered plates.

The electrolyte for fiat-plate SOFCs is commonly made by tape casting of fine stabilized-ZrO2 powders [9.132-9.138]. A typical firing temperature of 1300 ~ to 1500~ is used to sinter the tape-cast electrolyte. Flat stabilized-ZrO2 plates 50 to 250/xm thick have been fabricated by tape casting. The formulations and process conditions for casting zirconia electrolyte plates have not been extensively published. An example is described here to illustrate the principles

of the electrolyte tape-casting process [9.99,9.137]. The formulations of this example include methylethylketone/ethanol azeotrope mixture as solvent, dibutylphthalate and polyethyleneglycol-400 as plasticizers, and KD1 (ICI) as

dispersant. In this tape casting process, YSZ powder is first dispersed using the dispersant (2.3 wt% relative to YSZ) in the solvent by ball milling with zirconia balls for 24 h. The binder and plasticizers (2:1.2:1 by weight) are dissolved in the solvent and mixed with the ceramic suspension (8 wt% binder relative to YSZ) by ball milling for 18 h. The slurry is then filtered through a l l0-/xm filter, deaired by evacuation to 2 • 104 Pa (0.2 bar), and tape cast on a polyethylene film. After 3 h of drying in a controlled environment, the casting substrate is separated from the tape. The tape is dried in air and then sintered at 1300~

The electrodes are applied on the sintered electrolyte to produce a complete single cell. Various methods have been used for electrode application:

tape casting [9.136], slurry coating [9.139, 9.140], screen priming [9.133, 9.141, 9.142], plasma spraying [9.143], and other deposition techniques [9.144-9.149]. Applied electrodes are commonly fired at temperatures up to 1300~ to form a rigid sintered structure. The electrodes have also been cofired with the

electrolyte [9.92, 9.128, 9.150]. In addition to tape casting, the electrolyte for

flat-plate SOFCs has been fabricated by slip casting [9.151,9.152]. In cases

using electrode support, the electrolyte layer is deposited by several methods

such as spray pyrolysis [9.151,9.153], slurry coating [9.151], plasma spraying


[9.154-9.158], C O 2 laser evaporation [9.159], and vacuum evaporation

[9.160,9.161]. The LaCrO3 interconnect has been fabricated by tape casting [9.162,

9.163] or hot pressing [9.162]. Ribs for gas channels are embossed or machined into the interconnect. In addition to LaCrO3 ceramics, metals can be used for

cell interconnection in flat-plate SOFCs [9.130, 9.164-9.166]. Some proposed

metallic materials are high-temperature alloys based on chromium and nickel. Alloys coated with a thin film of perovskite oxide have also been proposed

[9.167-9.170]. Flat-plate SOFC stacks are formed by stacking layers to the desired

height. Compressive loads [9.171] and conductive braze materials [9.130, 9.172] have been used to improve electrical contact between the layers.

Fabrication of thin electrolytes

Several fabrication techniques have been considered for making thin YSZ electrolytes (< 10 ~m thick) for flat-plate SOFCs. An electrolyte less than 10 ~m thick reduces ohmic losses in the fuel cell, thus allowing efficient operation

at reduced temperatures (600 ~ to 800~ The advantages of reduced-temperature operation for SOFCs include wider material choice, longer cell life, reduced

thermal stress, improved reliability, and reduced fuel cell cost [9.173]. A number of methods have been proposed and investigated for fabrication of YSZ films (< 10 ~m thick); several of these techniques are discussed below.

(i) Vapor deposition: A modified CVD process combined with an EVD has been proposed for fabrication of YSZ films [9.34,9.39]. Micron-thick electrolyte layers have been deposited on porous substrates. Other CVD-based deposition techniques have also been suggested [9.174]. For example, a plasma- assisted metal organic CVD method using a microwave discharge has been used to form YSZ thin films for flat-plate SOFC applications [9.175]. A process, termed vapor-phase electrolytic deposition (VED), uses glow-discharge plasma

as a conductive medium to deposit thin YSZ layers from gaseous phases [9,176, 9.177]. To date, work has been limited to demonstrating the feasibility of those techniques in depositing thin electrolyte layers.

(ii) Sputtering: Thin films of YSZ have been produced by rf sputtering

[9.178, 9.179] and reactive magnetron sputtering [9.180-182]. Single cells

constructed from rf-sputtered films on porous substrates have been tested.

Magnetron sputtering has been used to produce both dense electrolyte and porous

290 Chapter 9

electrode layers. Cells proposed for fabrication by this technique are multilayer

structures containing catalytically active oxide layers deposited on the electrolyte to minimize polarization losses. Single cells (consisting of 5-#m YSZ electrolyte coated with doped Bi203 and CeO2 films < 60 nm thick) have been operated at 750~

(iii) Sol-gel: Very thin YSZ films (0.5 /zm) have been deposited onto porous electrodes using the sol-gel technique [9.183]. The sol-gel chemistry is modified to permit direct film deposition on porous substrates in a single step.

A low temperature (600~ is used to form the dense film. Another proposed sol-gel method uses multiple spin coating of polymeric precursors to deposit dense films on porous support [9.184]. Heat treatment at a temperature as low as 600~ has produced fine-grain dense YSZ films of 0.1 to 2/zm in thickness.

(iv) Tape calendering: A fabrication process based on tape Calendering has been proposed for making thin electrolyte films [9.185,9.186]. The process involves progressively rolling green electrolyte and anode layers and cofiring the bilayer at elevated temperatures. Electrolyte films 1 to 10 /xm have been

fabricated by the process. Figure 9.47 shows a scanning electron microscopic photograph of a fracture surface of a thin electrolyte on an anode support fabricated by tape calendering. Single cells have been fabricated in various sizes to more than 400 cm 2 area and have been tested in the temperature range of 600 ~ to 800~

A N O D E ELECTROLYTE

Figure 9.47. Micrograph of fracture surface of thin electrolyte on anode support [9.185]


(v) Jet vapor deposition: The technique uses a sonic gas jet in a low vacuum flow to deposit ceramic films. Deposition of micron-thick YSZ electrolyte layers for SOFC applications has been demonstrated [9.187].

(v) Electrophoretic deposition: YSZ films have been produced by electrophoretic deposition [9.188]. Positively charged YSZ powders are deposited on platinum-plated anodes from YSZ/acetylacetone suspensions.

Tape casting

Tape casting has long been used as a fabrication technique for producing large-area, thin, flat ceramic plates and laminated, multilayer ceramic structures [9.189-9.192]. The tape-casting process is shown schematically in Figure 9.48. Tape casting involves spreading a slurry of ceramic powders and organic ingredients onto a flat surface where the solvents are allowed to evaporate. After drying, the resulting tape develops a leather-like consistency and can be stripped from the casting surface.

A formulated slurry (or slip) for tape casting comprises ceramic powders and a liquid system. This liquid system includes solvent, binder, plasticizer, and a deflocculant/wetting agent. The solvent, either nonaqueous or aqueous, dissolves the other organic materials and distributes them uniformly throughout the slurry. The most often used nonaqueous solvent systems are highly polar organics such as alcohols, ketones, and halogenated hydrocarbons. It is also common to use mixtures of solvents to control the drying rate of the tape. Many

i ~ SLIP DOCTOR BLADES ,~~]~:~.. CARRIER FILM

Figure 9.48. Tape-casting process

292 Chapter 9

formulations for aqueous systems have been proposed, but problems are stil associated with water-based systems such as foaming and viscosity instabilitie, of the slip. A major drawback of the aqueous system is the need to control the pH at all stages of batching, milling, and casting. The binder dissolves in the solvent and enhances the solvent viscosity. Polyvinyl butyral and various acrylic polymers are commonly used as tape-casting binders. As the solvent evaporates, the binder forms a temporary bonding medium for the ceramic particles, thus

providing green strength to the tape. Typically, 3 to 8 g of binder are added to each 100 g of ceramic powders. The plasticizer is used to modify the properties of the binder. The more common plasticizers include butyl benzyl phthalate and polyethylene glycol. Normally, as much plasticizer can be required as binder. The deflocculant/wetting agent (dispersing agent) coats the ceramic particles and keeps them in a stable suspension in the slurry because of stearic hinderance and electric repulsion. Fish oil and fatty acids (e.g., glyceral trioleate) are commonly used. As much as 3 mL of dispersing agent may be used for each 100 g of ceramic powders. In general, the organics used in tape casting remain mostly proprietary. The optimal amount and type of organics has mainly been based on empirical studies because little is known about the interaction of the different components in a ceramic slurry.

The tape-casting process consists of three key steps: milling, casting, and drying. The milling step has two stages: (i) in the first stage (a milling step), ceramic powders, solvent, and deflocculant/wetting agent are milled (especially to break down agglomerates) to produce a low-viscosity slurry; (ii) in the second stage (a mixing and homogenization step), the binder and plasticizer are dissolved into the ceramic/solvent slurry. It is critical to introduce the dispersing agent independent of other polymers to prevent competition for the ceramic particulate surfaces; this produces more uniform viscosities in the slurry. Milling is usually

performed in standard comminution equipment such as ball mills and vibrating mills. After the milling, the slurry is deaired before casting. The key

characteristic for slip control is viscosity. Since the viscosity of the slurry is very temperature-dependent, careful control of temperature of the casting slip is

required. The slip is spread to a controlled thickness with the doctor blade of a

batch or continuous casting machine. (The process is considered to be batch if the blade moves over a single rigid carrier; it is considered to be continuous if

a flexible carrier moves under a fixed doctor blade.) The casting thickness can be controlled by the blade gap. As a rule of thumb, the ratio of blade gap to


final dried green tape is approximately 2" 1. The dried green tape thickness

depends on slip viscosity, casting rate, blade gap setting, and reservoir depth behind the doctor blade. All of these conditions must be controlled to obtain

uniform tapes reproducibly. After the tape is cast, it is moved to a dryer (batch process) or through

a drying tunnel (continuous process) to remove the solvent. The drying conditions depend upon the length of the casting machine, the type of solvent

used, and the amount of heat and air flow the drying slurry can withstand related to the thickness of the tape. These conditions are carefully controlled to

minimize curling, cracking, and trapping of gas bubbles. For example, too rapid

drying of the wet slurry may seal the surface against further removal of solvent

(skinning), resulting in trapped gas bubbles in the drying tape. Use of humidity drying limits evaporation of solvent from the slurry while permitting migration

of the solvent to the surface.

Table 9.7 summarizes the effect of casting parameters on casting

behavior [9.193]. From the table, an optimized tape-casting process will have

(i) a high solids loading slip, (ii) a slip with viscosity sufficiently low to flow under blade but not to flow off the carrier, (iii) a solvent that will not promote

skinning and bubble entrapment, (iv) a fast drying slip to increase drying rate, and (v) a drying system that allows control of gas removal.


Flat-plate SOFC single cells of various sizes have been fabricated and

operated under a variety of conditions. To date, most of the tests on flat-plate cells have used hydrogen as fuel. In general, observed area-specific resistances

range from 0.25 to 1.0 f~.cm z. Performance of early flat-plate single cells tended to decrease with increasing cell area [9.119, 9.140, 9.194]. Recent material modifications and process improvements have eliminated this problem. Flat-plate

cells of large area have been tested and have shown excellent performance. For

example, single cells with an active area of 125 cm a have been operated and have

shown a cell voltage of 0.7 V at 300 mA/cm z with humidified hydrogen as fuel

and air as oxidant [9.195]. A 23-cm square cell (active area of 400 cm a) has

produced a maximum power of 97 W [9.141]. Flat-plate SOFCs have been

operated for thousands of hours [9.142, 9.196, 9.197]. Degradation rates as low

as 0.5 % per 1000 h have been achieved. Thin-electrolyte single cells have also

29

4

Chapter 9

TABLE 9.7

Effect of Casting Parameters on Casting Behavior [9.193]

Effect

Defect Generation

Parameter Solids Viscosity Skinning Drying Content Residual and New Time

Bubbles Bubbles Cracking

Increase amount Decreased Lower Reduced NA* Possibly Increased of solvent increased

Increase slip NA Lower Reduced Possibly NA Decreased temperature increased

Increase solvent NA NA NA Possibly Possibly Decreased evaporation increased increased

Increase casting NA NA NA NA NA NA rate

Increase air NA NA NA NA Possibly Increased flow increased

Dispersed NA Lower NA Lower NA NA ceramics 'not applied


operated and shown excellent performance under reduced-temperature conditions.

Figure 9.49 shows voltage/current curves for a thin-electrolyte cell fabricated by

tape calendering [9.185]. High power densities have been achieved with

humidified hydrogen fuel and air oxidant, e.g., 0.55 W/cm z at 800~

Laboratory-scale multicell stacks have been constructed. Early stacks

were small, e.g., 3 cm by 3 cm in area, 1 cm high [9.198]; recently, flat-plate

stacks of footprint area as large as 225 cm / [9.199,9.200] and height as great as

200 cells [9.196] have been fabricated. Testing has been performed on stacks of

various sizes. To date, maximum power output of tested stacks ranges from 10

to 1,300 W [9.126,9.199-9.203]. For example, a 40-cell stack of 225-cm 2

footprint area has produced more than 500 W with hydrogen and air at 1000~

[9.199]. A maximum power output of 1.3 kW has been achieved for a 30-cell

stack with a 225-cm 2 active area [9.200]. Figure 9.50 shows, as an example,

voltage/current and power/current curves of a 40-flat-plate-cell stack. Flat-plate

SOFC multicell stacks have also been tested under current load for more than

2,000 h [9.202]. Examples of a flat-plate stack with current collectors and stack

setup for laboratory-scale testing are shown in Figure 9.51.

> u./ ~D < I-- ._1 O >

1 . 2

1 . 0

0 . 8

0 . 6

' 1 F [ - - ' I '

F U E L = HYDROGEN OXIDANT = AIR �9 7 0 0 ~

" 8 0 0 ~

�9 �9 9 0 0 ~

"~'<"~It-... ~ 1 0 0 0 o C

0 1 1 0 0 ~

�9 A

, l , , I ,

0 2 0 0 4 0 0 6 0 0

CURRENT DENSITY, m A / c m 2

800

Figure 9.49. Voltage~current curve of single cell having 5-l~m electrolyte [9.185]

296 Chapter 9

501 ACTIVEAREA = 10cm 2 �9 Voltage 1 50 a Power I FUEL = HYDROGEN ~ ~ - |

40 1 OXIDANT =AIR . ~ ~ 40

> 30 30

20 20 a. >

1 10

0 0 1 2 3 4 5 CURRENT, A

Figure 9.50. Voltage~current and power output of 40-cell flat-plate SOFC [9.196]

Figure 9.M. Photographs of a fiat-plate stack and stack setup for testing (courtesy of Ceramatec)


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Easier, W.A. Ellingson, B.K. Flandermeyer. R.J. Fousek, J.J. Heiberger, T.E. Kraft, S. Majumdar, C.C. McPheeters, F.C. Mrazek, J.J. Picciolo, R.B. Poeppel,

and S.A. Zwick, in Proceedings of the 21st IECEC, August 25-29, 1986, San Diego, CA, Vol. 3, American Chemical Society, Washington, DC, 1986, p. 1634. D.C. Fee, M.C. Billone, D.E. Busch, D.W. Dees, J. Dusek, T.E. Easier, W.A. Ellingson, B.K. Flandermeyer. R.J. Fousek, J.J. Heiberger, S. Majumdar, C.C.

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9.125

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9.127 9.128

9.129

9.130

9.131 9.132

9.133

9.134

9.135

D.C. Fee, S.K. Bhattacharyya, D.E. Busch, L.W. Carlson, T.D. Claar, D.W. Dees, J. Dusek, T.E. Easier, W.A. Ellingson, B.K. Flandermeyer. R.J. Fousek, H.K. Geyer, J.J. Heiberger, K.D. Kuczen, S. Majumdar, C.C. McPheeters, F.C. Mrazek, J.J. Picciolo, and R.B. Poeppel, in Space Nuclear Power Systems 1986, M.S. E1-Genk and M.D. Hoover (eds.), Orbit Book Company, Malabar, FL, 1987,

p. 209.

N.Q. Minh, A. Amiro, T.R. Armstrong, J.R. Esopa, J.V. Guiheen, C.R. Horne,

and J.J. Van Ackeren, see Ref. 9.6, p. 607.

N.Q. Minh, T.R. Armstrong, J.R. Esopa, J.V. Guiheen, C.R. Home, F.S. Liu,

T.L. Stillwagon, and J.J. Van Ackeren, see Ref. 9.11, p. 93.

K. Koseki, N. Kusunose, K. Shimizu, H. Shundo, T. Iwata, S. Maruyama, and T. Nakanishi, see Ref. 9.40, p. 107.

H. Arai, K. Eguchi, T. Setoguchi, R. Yamaguchi, K. Hashimoto, and H.

Yoshimura, see Ref. 9.11, p. 167. E. Ivers-Tiff6e, W. Wersing, and B. Reichelt, see Ref. 9.40, p. 137. T. Hoshina, T. Yoshida, and S. Sakurada, see Ref. 9.40, p. 516.

M. Shimotsu, M. Izumi, and K. Murata, in Proceedings of the Third International

Symposium on Solid Oxide Fuel Cells, May 16-21, 1993, Honolulu, HI, S.C. Singhal and H. Iwahara (eds.), Electrochemical Society, Pennington, NJ, 1993, p. 732.

H. Takagi, H. Taira, A. Shiratori, S. Kobayashi, Y. Sugimoto, S. Sakamoto, and K. Tomono, see Ref. 9.123, p. 738. T. Hikita, M. Hishinuma, T. Kawashima, I. Yashuda, T. Koyama, and Y. Matsuzaki, see Ref. 9.123, p. 714.

Y. Akiyama, T. Yasuo, N. Ishida, S. Taniguchi, and T. Saito, see Ref. 9. 123, p. 724.

M.S. Hsu, see Ref. 9.14, p. 115. H. Takagi, S. Kobayashi, A. Shiratori, K. Nishida, and Y. Sakabe, see Ref. 9.11,

p. 99.

H. Shundo, H. Shimizu, N. Kusunose, T. Iwata, S. Maruyama, and K. Koseki, see Ref. 9.11, p. 119.

R.C. Ruhl, U.S. Patent No. 4770955, September 13, 1988.

R. Diethelm and K. Honegger, see Ref. 9.123, p. 822.

Ceramatec, Inc., Development of Planar Geometry Solid Oxide Fuel Cell Technolo-

gy, Phase I Final Report, July 1986-.luly 1987, Report No. GRI-87/0297, Gas

Research Institute, Chicago, IL, 1987.

Ceramatec, Inc., "Development of Planar Geometry SolM Oxide Fuel Cell

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F.W. Poulsen, J.J. Bentzen, and J.B. Bilde-Sorensen, see Ref. 9.64, p. 162.

M. Dokiya, N. Sakai, T. Kawada, H. Yokokawa, and I. Anzai, see Ref. 9.11, p. 127.

304 Chapter 9

9.136

9.137

9,138

9.139

9.140

9.141

9.142

9.143 9.144

9.145

9.146

9.147

9.148

9.149

9.150

9.151

9.152

9.153

9.154

9.155

9.156

9.157

9.158

9.159

9.160

9.161

9.162

9.163

9.164

J.P.P. Huijsmans, E.J. Siewers, F.H. van Heuveln, and J.P. de Jong, see Ref. 9.11, p. 113.

C. Bagger, see Ref. 9.6, p. 241.

A. Tintinelli, C. Rizzo, S. Loreti, and A. Selvaggi, see Ref. 9.11, p. 747.

T. Hoshina, T. Yoshida, and S. Sakurada, see Ref. 9.40, p. 516.

S. Mui'akami, Y. Akiyama, N. Ishida, T. Yasuo, T. Saito, and N. Furukawa, in

Addendum of Ref. 9.11, p. 59.

Y. Matsuzaki, M. Hishinuma, T. Kawashima, I. Yasuda, T. Koyama, and T.

Hikita, see Ref. 9.6, p. 237.

E. Erdle, W. D6nitz, W. Sch~ifer, and R. Schamm, see Ref. 9.6, p. 403.

L.W. Tai and P.A. Lessing, J. Am. Ceram. Soc., 74 (1991) 501.

L.G.J. de Haart, R.A. Kuipers, K.J. de Vries, and A.J. Burggraaf, see Ref. 9.32, p. 197.

H. Michibata, H. Tenmei, T. Namikawa, and Y. Yamazaki, see Ref. 9.32, p. 188.

B. Gharbage, M. Henault, T. Pagnier, and A. Hammou, Mater. Res. Bull., 26

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H. Michibata, T. Namikawa, and Y. Yamazaki, Denki Kagaku, 57 (1989) 255.

R. Okiai, S. Yoshida, I. Kaji, M. Hasegawa, H. Yamanouchi, and M. Nagata, see

Ref. 9.64, p. 191.

H. Jansen, H.-P. Buchkremer, D. St6ver, and K. Wippermann, see Ref. 9.123, p.

752. M. Dokiya, N. Sakai, T. Kawada, H. Yokokawa, T. Iwata, and M. Mori, see Ref.

9.50, p. 1547.

H. Arai, see Ref. 9.64, p. 12.

T. Setoguchi, T. Inoue, H. Takebe, K. Eguchi, K. Morinaga, and H. Arai, Solid

State Ionics, 37 (1990) 217.

T. Setoguchi, M. Sawano, K. Eguchi, and H. Arai, Solid State Ionics, 40/41 (1990)

502.

T. Nakanishi, S. Maruyama, and N. Kusunose, see Ref. 9.6, p. 407.

H. Hamatani, T. Okada, and T. Yoshida, see Ref. 9.64, p. 197.

S. Fukami, M. Kitoh, A. Bunya, H. Saitoh, T. Itoh, Y. Kaga, Y. Ohno, K. Eguchi,

and H. Arai, see Ref. 9.11, p. 205.

D. St6ver, R. Kecker, H. Jansen, and W. Mall6ner, see Ref. 9.11, p. 215.

E.A. Barringer, D.P. Heitzenrater, and M.R. Tharp, see Ref. 9.123, p. 771.

H. Nakagawa, S. Kosuge, H. Tsuneizumi, E. Matsuda, H. Mihara, and Y. Sato, see

Ref. 9.32, p. 71.

H. Michibata, T. Namikawa, and Y. Yamazaki, Denki Kagaku, 58 (1990) 1070.

Y. Yamazaki, T. Namikawa, and H. Michibata, see Ref. 9.11, p. 175.

C. Milliken and A. Khandkar, see Ref. 9.32, p. 361.

P.A. Lessing, L.W. Tai, and K.A. Klemm, see Ref. 9.32, p. 337.

M.S. Hsu, U.S. Patent No. 4629537, December 16, 1986.


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9.168

9.169

9.170

9.171

9.172

9.173

9.174

9.175 9.176 9.177

9.178

9.179

9.180

9.181

9.182

9.183

9.184

9.185

9.186

9.187

Y. Akiyama, N. Ishida, S. Murakami, and T. Saito, U.S. Patent No. 4997726, March 5, 1991. E. Ivers-Tiff6e, W. Wersing, M. SchiefSl, and H. Greiner, Ber. Bunsenges. Phys. Chem., 94 (1990) 978. T. Yoshida, T. Shima, F. Ishizaki, H. Iwasaki, I. Mukaizawa, Y. Someya, S.

Sakurada, and O. Yamamoto, U.S. Patent No. 4950562, August 21, 1990.

H. Tenmei, H. Michibata, T. Namikawa, and Y. Yamazaki, Denki Kagaku, 58

(1990) 1072.

H. Konno, M. Tokita, A. Furusaki, and R. Furuichi, Electrochim. Acta, 37 (1992)

2421.

H. Konno and R. Furuichi, in High Temperature Corrosion of Advanced Materials and Protective Coatings, Y. Saito, B. Onay, and T. Maruyama (eds.), Elsevier Science Publishers, Amsterdam, The Netherlands, 1992, p. 177.

M. Hsu and H. Tai, in Proceedings of the 26th IECEC, August 4-9, 1991, Boston, MA, Vol. 3, American Nuclear Society, La Grange Park, IL, 1991, p. 606. J. Jung, Th. Martens, H. Runge, and M. Turwitt, see Ref. 9.11, p. 144.

K. Krist and J.D. Wright, see Ref. 9.123, p. 782.

V.E.J. van Dieten, P.H.M. Walterbos, and J. Schoonman, see Ref. 9.11, p. 183.

H. Uyama, N. Oka, I. Ono, and O. Matsumoto, Denki Kagaku, 58 (1990) 564. Z. Ogumi, Y. Tsuji, Y. Uchimoto, and Z. Takehara, see Ref. 9.64, p. 203. Z. Ogumi, Y. Tsuji, Y. Uchimoto, and Z. Takehara, see Ref. 9.11, p. 201.

A. Negishi, K. Nozaki, and T. Ozawa, Solid State Ionics, 3/4 (1981) 443.

N. Nakagawa, C. Kuroda, and M. Ishida, Denki Kagaku, 57 (1989) 215.

S.A. Barnett, Energy, 15 (1990) 1.

L.S. Wang and S.A. Barnett, J. Electrochem. Soc., 139 (1992) 1134.

L.S. Wang and S.A. Barnett, J. Electrochem. Soc., 139 (1992) L89.

S.J. Visco, C. Jacobson, and L.C. De Jonghe, presented at An EPRI/GRI Workshop on Fuel Cell Technology Research and Development, April 13-14, 1993, New Orleans, LA, Electric Power Research Institute, Palo Alto, CA, 1993.

C.C. Chen, M.M. Nasrallah, and H.U. Anderson, see Ref. 9.6, p. 515.

N.Q. Minh and C.R. Horne, in Proceedings of the 14th Riso International Symposium on Materials Science, High Temperature Electrochemical Behaviour of Fast Ion and Mixed Conductors, September 6-10, 1993, Roskilde, Denmark, F.W. Poulsen, J.J. Bentzen, T. Jacobsen, E. Skou, and M.J.L. Osterg~.rd (eds.), Ris~ National Laboratory, Denmark, 1993, p. 337.

N.Q. Minh, T.R. Armstrong, J.R. Esopa, J.V. Guiheen, C.R. Horne, and J.J. Van Ackeren, see Ref. 9.123, p. 801.

B.L. Halpern, J.J. Schmitt, J.W. Golz, and Y. Di, in Proceedings of the Fourth Annual Fuel Cell Contractors Review Meeting, July 14-15, 1992, Morgantown, WV, W.J. Huber (ed.), Report No. DOE/METC-92/6127, U.S. Department of Energy, Washington, DC, 1992, p. 102.

306 Chapter 9

9.188

9.189

9.190

9.191

9.192

9.193

9.194

9.195

9.196

9.197

9.198

9.199

9.200 9.201

9.202

9.203

T. Ishihara, T. Kudo, H. Matsuda, Y. Mizuhara, and Y. Takita, see Ref. 9.123, p.

65.

R.E. Mistier, D.J. Shanefield, and R.B. Runk, in Ceramic Processing Before Firing, G.Y. Onoda, Jr. and L.L. Hench (eds.), Wiley-Interscience, New York, 1978, p.

411.

R.B. Runk and M.J. Andrejco, Am. Ceram. Soc. Bull., 54 (1975) 199.

E.P. Hyatt, Am. Ceram. Soc. Bull., 65 (1986) 637.

R.E. Mistier, Am. Ceram. Soc. Bull., 69 (1990) 1022.

E.P. Hyatt, Am. Ceram. Soc. Bull., 68 (1989) 869.

M. Dokiya, N. Sakai, T. Kawada, H. Yokokawa, T. Iwata, and M. Mori, see Ref.

9.32, p. 325. Y. Akiyama, N. Ishida, T. Yasuo, S. Taniguchi, S. Murakami, T. Saito, and N.

Furukawa, see Ref. 9.6, p. 603.

A. Khandkar, S. Elangovan, J. Hartvigsen, C. Milliken, and M. Timper, see Ref.

9.6, p. 596.

Y. Someya, T. Yoshida, and S. Sakurada, see Ref. 9.6, p. 611.

J.N. Michaels, C.G. Vayenas, and L.L. Hegedus, J. Electrochem. Soc., 133 (1986)

522. M. Hattori, A. Kato, Y. Esaki, H. Miyamoto, M. Irino, S. Naito, F. Nanjo, and

M. Funatsu, see Re f. 9.6, p. 411.

S. Sakurada and T. Yoshida, see Ref. 9.11, p.45. A. Kato, Y. Esaki, M. Hattori, S. Naito, F. Nanjo, and M. Nishiura, see Ref. 9.-

123, p. 809. T. Iwata and N. Kusunose, see Ref. 9.123, p. 792.

H. Takagi, H. Taira, A. Shiratori, S. Kobayashi, Y. Sugimoto, S. Sakamoto, and

K. Tomono, see Ref. 9.123, p. 738.

Chapter 10

MODELING AND ANALYSIS

10.1 GENERAL

A practical SOFC must have the electrical and electrochemical performance, along with the mechanical and structural integrity to meet operating requirements of specified power generation applications. In terms of electrical

and electrochemical performance, the fuel cell must be designed to achieve high voltages (at the required current density) with low ohmic and polarization losses. In terms of mechanical and structural integrity, the fuel cell design must minimize thermal stresses during fabrication and operation to avoid exceeding the

strength limits of the component materials. Mathematical modeling and analysis are common tools used in SOFC design and optimization to achieve the required electrical, electrochemical, mechanical, and structural performance.

Mathematical modeling and analysis are used to predict cell behavior under a variety of conditions and to investigate the effect and the relative importance of various processing and operating parameters. Modeling and analysis can provide information on how various processes and parameters interrelate, allowing interpretation of cell behavior with a minimum of expensive

testing. Modeling and analysis also provide a picture, for example, of stress, potential (or voltage), current density, and temperature as functions of position and time for various cell configurations and operating conditions. It may then be possible to use the information provided by the models and analysis to

optimize cell designs and operating parameters. There are several levels of modeling and analysis: the molecular type

model, electrode model, cell model, stack model, and system model [10.1]. The

discussion in this chapter is limited mainly to thermal stress analysis, electrical

analysis, and performance modeling (specifically current and temperature

distribution) of SOFC cells and stacks. (i) Thermal stress analysis: Residual thermal stresses are generated due

to differences in the coefficients of thermal expansion of various components.

308 Chapter 10

Such stresses have been analyzed in SOFC structures using mathematical models

such as finite-element techniques. (ii) Electrical analysis" Analysis of ohmic and polarization resistances is

performed to estimate cell voltage losses. Equivalent circuits have been developed to predict cell performance under a variety of operating conditions.

(iii) Performance modeling: Mathematical modeling of SOFC performance is typically based on formulating and solving simultaneous mass, energy, and potential balance equations. Such a model provides voltage/current density characteristics, current distribution, and temperature distribution of the fuel cell.

10.2 STRESS ANALYSIS

A SOFC is a layered ceramic structure composed of several different

materials. During fabrication and operation, the structure is subject to nonuniform temperature distributions and various thermal cycles. Thus, matching

thermal expansion in the cell components is critical. A small difference in the coefficients of thermal expansion of the layers can result in large thermal stresses in the fuel cell and cause cracking in the ceramic structure. In general, the problem of thermal expansion mismatch in present SOFCs mainly centers on the

anode, since nickel has a higher coefficient of thermal expansion than Y203- stabilized ZrO2 (YSZ). The anode typically must contain more than 30 vol% nickel to have sufficient conductivity. The coefficient of thermal expansion of this composition is higher than that of the YSZ electrolyte and other components,

causing thermal expansion mismatch. Stress analysis has been performed to evaluate the effects of design,

fabrication, and operation parameters on the fracture characteristics of SOFCs. Analysis of thermal stresses in the fuel cell structure typically uses simplified

stress models or the finite-element method. For example, a simple model based

on small-displacement linear elastic theory has been employed to analyze single-

cell circular plates [10.2]. Simplified stress models generally involve establishing

stress equations and deriving strain energy release rates and stress intensity

factors. The finite-element method, a numerical technique of piecewise

approximation, is conventionally used in stress analysis, especially for intricate structures.

The finite-element method models a SOFC structure as an assemblage of elements with nodes indicating where elements are connected to one another. A

finite element stress analysis typically involves the following steps"

Modeling and Analysis 309

�9 Divide the structure into finite elements.

�9 Formulate the properties of each element. (Determine nodal loads asso-

ciated with all element deformation states that are allowed.)

�9 Assemble elements to obtain the finite-element model of the structure.

�9 Apply the known loads (nodal forces and/or moments).

�9 Specify how the structure is supported.

�9 Solve simultaneous equations to determine nodal displacements.

�9 Calculate element strains from the displacements and the displacement

field interpolation, and finally calculate stresses from strains.

A simple double-layered model has been developed for the sealless

tubular cell [10.3]. The cell is approximated as two layers - - an anode layer of

thickness di~ and length 2e, and the rest of the cell (the support, cathode, and

electrolyte) as a layer of thickness 62 and length 2f. The following equations are

derived for the tensile stress, tri, in the anode and the shear stress, r, at its

interface with the rest of the cell (~ is the coefficient of thermal expansion, T is

the temperature, E is the elastic modulus, v is the Poisson's ratio, and x is the

distance from the center of the double layer, x = 0 to ___ e)"

o" 1 = (r

,111 + 11 El61 E2~ 2

(coshc coshcx)

( l -u )

coshc (Eq. 10.1)

7 =

where c and/Xl are given as

r =

#I(r r )AT sinhcx

/ilC coshc

I~1 1 1

81 E181 E282

(Eq. 10.2)

(Eq. 10.3)

~1,1 = 2(l+v)

(Eq. 10.4)

310 Chapter 10

Using this simplified model, approximate equations for strain energy release rates

for cracks in the anode have been derived.

A finite-element stress analysis of the sealless tubular cell has also been performed [10.3]. The stress model consists of 166 four-node quadrilateral and three-node triangular isoparametric elements. Stress distribution in the individual components of the cell has been obtained from the model and has been compared with that from the simple double-layered model [10.3].

A stress and fracture analysis has been developed for SOFC single cells (anode/electrolyte/cathode trilayers) during cooldown after cofiring. This model considers a circular single-cell disk of radius r containing an annular crack of

depth C c (Figure 10.1) [10.2]. The model assumes that the temperature, thermal

expansion coefficient, and elastic modulus of the disk vary only in the thickness direction, if at all.

The radial stress, r and tangential stress, frO(i), at section AiA i a r e given

as

E(zi) [gx, Or( 0 = Oo( 0 - ( l_v) l -~i-ot(z i )T(z i )

E(zi)zi My.. Di(l_v2) '

(Eq. 10.5)

where for each i = 1 to 3, E(z) is Young's modulus at zi, 1~ is Poisson's ratio, ~(z~ is the coefficient of thermal expansion at zi, T(z~ is the temperature range

at zi, zi is the distance from centroidal axis.

I< 6', ~l A, A3

z~, A3

A, A2 r _

Figure 10,1. Geometry of circular single cell containing an annular crack [10.2]


The other parameters are defined by the following equations"

fa Ezi dz i = 0 (Eq. 10.6)

xf . K~ - 1-v ,E dz~ (Eq. 10.7)

D i - 1 fa 1 - v 2 Ez/2. d z i (Eq. 10.8)

_ 1 fA,E a T dz~ (Eq. 10.9) N~, l-v

i = a Tz~ dz~ (Eq. 10.10) M~, 1-v

Thus, the strain energy release rate per unit of extension of the crack area, G, is given as

G = G n + G I =

2 2 2 2 6)2 ~3 M2 M3

~ + + +

r 2 r 3 D2(I+v) Da(l+v) (Eq. 10.11)

where, for i = 2,3

G I = 4-

/92(1+v) D 3 ( l + v ) (Eq. 10.12)

G H =

2 2 6)2 ~3 (Eq. 10.13)

i fAj %(0 dzi (Eq. 10.14)

312 Chapter 10

and

= -fa ~ dzi (Eq. 10.15) M, l

In this model, the opening mode component G~ is assumed to be more important than the shear component G,. The stress intensity factor, K~, is related to the mode I strain energy release rate by the following equation:

[EGt] '/' (Eq. 10.16) K~ = 1-v 2

The maximum K~ has been calculated as a function of the coefficient of thermal expansion of the anode for a symmetrical anode/electrolyte/anode trilayer (Figure 10.2). Correlation between calculated stress intensity factors and observed cracking behavior of trilayers indicates a fracture toughness of the YSZ electrolyte of about 1.25 MPa.n~. Figure 10.3 shows the boundaries (computed

, , ANODE

~22 N

x <

N 1

O SURVIVED COOLDOWN

�9 FRACTURED DURING COOLDOWN

0 I I 11 12 13 14

THERMAL EXPANSION COEFFICIENT OF ANODE, 10 .6 cm/cm.K

Figure 10.2. Variation of maximum mode I stress intensity factor with coefficient of thermal expansion of anode [10.2]


E 1 0 0 -

u3 a o -1- I"- < (J tL O

z (...) 10 "r" I--

A N O D E

vol % Ni o , 1 0 " e c m l c m K

(A) 40 13 (B) 30 12

CRACK AT CRACK AT CATHODE- ANODE-

( A ) I (B) ~ ELECTROLYTE ELECTROLYTE I INTERFACE INTE_~F.a~E

/ELECTROLYTE I ANODE- ~, I / M IDPLANE I E LTE;TFAR~)cL;TE ~'~

. . . . . ' i �9 ' ' ' �9 - ' ' ] ' 10 1 0 0

T H I C K N E S S OF A N O D E , ,urn

Figure 10.3. Effect of coefficient of thermal expansion of anode on cracking of asymmetrical single cell with electrolyte 38 I~m thick [10.2]

from the model) between the domains of cathode and anode thickness combina- tions that are predicted to survive a cooldown after cofiring (lower left) and those that are predicted to fail (upper right) for an electrolyte thickness of 38/zm and two different values of coefficient of thermal expansion of the anode. The finite- element method has also been applied to analyzing the residual stresses developed

in a single cell bonded to anode and cathode corrugations [10.4]. An eight-node

layered-shell element is used to model the anode/electrolyte/cathode trilayer (single cell). A quadrilateral-shell element is used to model the anode and

cathode corrugations. Based on the analysis, failure-zone and safe-zone maps have been developed to provide guidelines in selecting the proper thicknesses of

electrodes and electrolyte to prevent failure of the fuel cell during fabrication. An example of such a map is shown in Figure 10.4.

314 Chapter 10

175

/ E 150 ...........

r 125 LU z v

-r- lOO

~ / o ~ 7 5

_ J

5O

2 5 0

CORRUGATIOh

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1 25

............................... ........

50 75

THICKNESS = 2 5 0 p m

.. ............................................................

I

............................

............................ ~ m !natle s .....................

l .............................

�9 . . . . . . . . . . . . . . . . . . . . . . . . .

100 125 150 175

ELECTROLYTE THICKNESS, pm

Figure 10.4. Failure- and safe-zone map for single cell bonded to anode and cathode corrugations [10.4]

10.3 ELECTRICAL ANALYSIS

One of the objectives in stack design is to maximize electrochemical performance (minimize ohmic and polarization losses) in the fuel cell. Electrical

analysis is often used to estimate cell internal resistances and to predict cell performance potentials under specified conditions. Electrical analysis of a particular SOFC design considers electron/ion paths in the stack to model ohmic resistances of cell components and contact resistances at interfaces. Several

examples of electrical analysis are discussed in this section to illustrate the

analysis principles and techniques.

In banded SOFCs, one approach in electrical analysis is to approximate

cell resistance [10.5]. The resistance of an unit cell of the banded configuration

(Figure 10.5) is considered to be the sum of the resistances of the active cell, the

interconnect, the cathode across the anode gap, and the anode under the cathode

gap. Figure 10.6 shows plots of power, unit cell resistance, and current density

as functions of electrolyte and interconnect lengths calculated by the approximation method. In this case, at maximum power per unit of volume, the value for

electrolyte and interconnect length is 0.21 cm. Under those conditions, the resistance and the current density of a unit cell are 0.12 fl and 0.66 A/cm 2,

respectively.


ANODE GAP

ACTIVE C E k k ~ i N T E R C O N N E C T I O N |

. . ~ CATHODE GAP ELECTROLYTE - I

~ ~ : - : , : " ~ : : ~.~s :.~-... : ~"..-~.2,.".q :~;-" ;..5 . . . . . . . . . . . ' - - , . . " . , , ' . -~ . -~

UNIT C E L L INTERCONNECT

POROUS SUPPORT ~(('~t, r,-,,,,'z ~I~(( I

Figure 10.5. Unit cell of banded configuration [10.5]

~ :~176 i " ' ~ 1 ~o:~ ,,0F, i i ~ -~ ,,= 50 l i .~ o i , ;

0 . 4 - ~ |

I l ,,,~ o 3 - \ : : z i i r ~ 02- t t

~ 0.1 -

I

0 0.8

>: t-- 0.6 z ,,, I L l

-~ oE 0 . 4

z LU rr 0.2 re ::)

U 0

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 ACTIVE CELL OR INTERCONNECT LENGTH, cm

Figure 10.6. Power, unit cell resistance, and current density as functions of electrolyte and interconnect lengths for banded SOFCs [10.5]

316 Chapter 10

A more exact approach to assessing cell resistance in banded SOFCs is to use a model of distributed constants [10.5-10. 7]. In this approach, the model considers an element, Ax, of the cell and its equivalent circuit (Figure 10.7). The following equations can be derived from the equivalent circuit:

(Eq 10.17)

(Eq. 10.18)

where subscripts o, f, and e denote cathode, anode, and electrolyte. Each current (]) value in these equations can be replaced by voltage and resistance. As Ax --,0, the following equations can be obtained (d is the diameter)"

dj~,~ _ -r id v'c~) (Eq. 10.19) dx p~8~

dJ~ - ~ d v~x) (Eq. 10.20) dx p~8~

where v is the voltage, p the resistivity, and c5 the thickness. On the other hand, the relationship between the voltages is given as

p~ PlAXjg~+,~) : 0 (Eq. 10.21) ve(x) + 8 o = d j~ + Ax) - v ~x § ax) ~~ d

p,,Ax

I

%o,) f po~, v ~,.~) t

t

Figure 10.7. Equivalent circuit of banded SOFC cell [10.5]


As Ax --, 0, Eq. 10.21 becomes

dYe(x) _ 1 r PO �9 Pf. dx ; d t-~o J ~ + Ax) -5--/J1~x + A x)]

Combining and solving Eqs. 10.19, 10.20, and 10.22 give

V~x ) = I [Amsinh(Ox)+Bmcosh(Ox)]

(Eq. 10.22)

(Eq. 10.23)

I J e - [Amsinh(Ox) + BmC~ (Eq. 10.24)

p~8~

where je is the current density in the electrolyte, I is the current, and 0, Am, and Bm are given as (le is the electrolyte length)

02 = ~ 1 (Po+ P_~_f) (Eq. 10.25) PeSe 8 o 6f

An, = 1 P/ (Eq. 10.26) 0~d 81

Bm -

P o + P/cosh(0l,) (Eq. 10.27) 1 8 0 8f

sinh(0/)

The following equation has also been derived for the resistance R c (in fl):

8 6/ o/ . R~ = ~ (Eq. 10 281

6 0 8 /

This equation illustrates how the various parameters affect unit cell resistance, thus cell power output per unit of volume.

For banded SOFCs, the most desirable number of single cells on a support tube (radius r) can be obtained at the condition of minimum resistance and can be calculated from the thicknesses of the electrolyte and electrodes [10.8].

318 Chapter 10

The number of cells is given by

ll ~ 1 1 --. .[. ~ )

The minimum resistance is thus given as

2 ~ 5 e 1 1 - _ _ ( + ) R~. ~-r o, o15/ Oo6 o

(Eq. 10.29)

(Eq. 10.30)

The maximum power of the banded SOFC tube, Pwr can be estimated from the following equation (E~ is the reversible voltage):

P ~max) 4R~. | 6e (Eq. 10.31)

8~ ~ ( 1 1 ~ . 4 - ~ ) O e O71 008 0

An analytical model has been developed to evaluate performance potentials of various SOFC configurations [10.9,10.10]. In this model, the current flow through a single cell is governed by a second order differential equation of the form

A2V_ V (Eq. 10.32) ~2

where V is the voltage between the two interconnect contact points, and

= ~ P e6e (Eq. 10.33)

P/5/+ P 0 8 o

Eq. 10.32 can be solved for one-dimensional single-cell strips and disks. general solution for a single cell strip is as follows"

The

V(x ) = A vsinh(x/f~ ) + Bvcosh(x/f~ ) (Eq. 10.34)

where x is the strip coordinate in the direction of current flow, and Av and By are integration constants determined by the boundary conditions. Two contact


geometries are possible for the single cell resistance (Figure 10.8). For directly opposed line terminals, the effective specific ohmic resistance is given as

R, , , : (Eq. 10.35)

with J v - X (Eq. 10.36) ff

For diagonally opposed line terminals, the effective specific ohmic resistance is

R,~ = (p~;+p~6 ~+po 8 o)Jv(coth(Jv) + Bv[Jv-2tanh(Jvl2)]} (Eq. 10.37)

where

Po 8/

8 o P f By = (Eq. 10 38) PoSf2 (1+ ~ ) 8oPf

From these equations and information on comact and polarization resistances, potential cell voltage, power density, and cell efficiency can be estimated for various cell configurations. Some examples of predicted voltage/current curves are given in Figures 10.9 and 10.10.

ELECTR

CA1

0 x

7 / F-=-~~--ax 0 x

Figure 10. 8. Current collection contact configurations [10.9]

3 2 0 Chapter 10

- - = - - 700oc

800oc

- - , - - 900oC

1000oC

- - , - - 1100o0

> 0.80

(,.9 0 .60 < I.-.- ,,_,1 0 > 0.40

-SISTANCE -I, lOSS-PLANE ~ 0 2 0 .ANE

ONNECT o.00 J \ 0 2 4 6 8

1100 CURRENT DENSITY, kA/m 2

I + I

\ 10

Figure 10.9. Projected contributions to ohmic resistance, and voltage~current characteristics of sealless tubular SOFCs [10.9]

i - - u - - 700oc

800oc

i " - - 900~

----o---- 1000~

1100~

,oo , !

�9 ~ _ _ / > o6o

< o 60 ~ \ > o,o \

o oo +oo , (..) 0.20 , .~

Ix: 700 "'-:~..<.._/2/~--__.~"-..:..~//~ ~ IN-PLANE \ I < 800 , . , ^ ~ ~ . . . ~ ~ - - - - - - ~ ~ INTERCONN=.,.,T 0.oo m

~uu ~ < - ~ 7 / L. ~ ~ 0 2 4 6 8 10 TEMPERATURE, o C 1000 ~ CONTACT

1100 CURRENT DENSITY, kA/m 2

Figure 10.10. Projected contributions to ohmic resistance, and voltage~current characteristics of bell-and-spigot SOFCs [10.9]


10.4 MODELING OF CURRENT AND TEMPERATURE DISTRIBUTION

At each point in a SOFC, a potential balance equation relates the cell

voltage E to the reversible potential Er and to the voltage terms that depend on

the local current I. Thus, cell voltage (assuming E is constant over the cell face)

is given by

E = E r - IRoj = - (ha+nO (Eq. 10.39)

where Rohm is the local ohmic impedance of the cell, and r/a and r/~ are the anodic

and cathodic overpotentials, respectively. To model the current distribution of

a SOFC, it is necessary to solve Eq. 10.39. Knowledge of Er, Rohm, ~/a, and r/c

is required to solve Eq. 10.39.

�9 Reversible voltage: The reversible voltage Er is given by the reversible

potentials at the cathode and anode. These potentials can be calculated

using the Nernst equation, along with the local gas composition. The

local gas composition can be calculated from mass balances on all re-

acting species, usually in terms of a local degree-of-conversion vari-

able for each electrode.

�9 Ohmic impedance: The ohmic impedance can be obtained from mea-

sured cell resistance or can be estimated from the effective conductivity

of the cell components and the effective distance between the components.

�9 Overpotentials: In general, the overall polarization behavior of a SOFC

system is approximately linear over the current density range of prac-

tical interest. The quasilinear local polarization characteristics may be

determined from cell measurements or electrode modeling.

In a fuel cell stack, current and temperature distributions are strongly

coupled. Practically every chemical, material, and transport property involved

in current generation is significantly temperature dependent. Thus, the local

temperature plays an important role in determining local current density, which,

in turn, determines the local heat flux. To model temperature distribution,

knowledge of a heat balance is required. Energy balance equations can be

formulated to describe the heat transfer processes occurring in the local heat

balance. The predominant heat transfer processes in a SOFC include:

�9 Heat release or absorption arising from the electrochemical reactions,

electrical resistances, and anode fuel chemistry (for example, internal

reforming)

322 Chapter 10

�9 Convective heat transfer between the cell components and the anode

and cathode gas streams

�9 In-plane heat conduction through the cell components

�9 Heat exchange between the cell and either gas stream due to the com-

bined effect of conductive heat transfer and a mass transfer contribu-

tion associated with exchange of reacting species at the cell tempera-

ture.

Many factors affect the relative importance of these heat transfer

processes, and hence, the thermal characteristics of the fuel cell. Some of the

factors include gas flow channel shape and size, gas flow configuration (e.g.,

coflow, crossflow, counterflow), gas inlet temperature, gas flow rate, and cell

area, voltage, and pressure.

Current and temperature distributions have been calculated for the various

cell designs [10.11-10.36]. For example, a mathematical model has been

developed to simulate a SOFC of crossflow configuration (Figure 10.11). The

model considers a unit cell placed deep within the stack (such that end effects

may be neglected). In this model, the horizontal plane of the cell is divided into

small rectangular areas or elements, each defined as a node. The elements are

sufficiently small such that the nodal values of temperature, pressure, concentra-

tion, etc., can be used over the element for the calculation of local current

densities and heat fluxes.

ELECTROLYTE

NTERCONNECT

Figure 10.11. SOFC of crossflow configuration [10.20]


The Nernst potential at any node (i,j) is given by the relationship

4. / eI~ o ,av

(Eq. 10.40)

For a specific cell operating voltage, E, the current generated at the node is

Er, q - E Itj = (Eq. 10.41)

Ri/

where the total resistance at any node R U can be calculated from the resistivities

of the cell components and the electron and ion path length.

Heat is generated in the node from the resistance to current flow and

from the entropy change of the electrochemical reaction. The total heat released

is assumed in the electrolyte layer and is given by

Q~j = -~-ff(-A , - (Eq. 10.42)

The heat is removed as sensible heat with the flowing fuel and oxidant gases.

The heat transfer between the SOFC component layers and the flowing gases is

assumed to occur only by convection. For heat removal from the electrolyte, the

fuel- and oxidant-side heat transfer coefficients (h/and ho) are calculated based on a limiting Nusselt number of 3.0. For heat transfer between fuel and oxidant

across the interconnect, an overall heat transfer coefficient, U, is used. The

following energy balances are solved to determine the unknown electrolyte

temperature, T e, at each node, fuel temperature at the outlet of each node, T/, u,

and oxidant temperature at the outlet of each node, To, u:

Cathode (oxidant) energy balance

ro, l_ijG,o(L,l_ij- Td) - ro,ijG, o(To, ij- Td) + Saho(L, ij-

Anode (fuel) energy balance

ro,a) + S a e ( ro,a) - o

(Eq. 10.43)

- + SahtT,.,j- G ) * SaVG.av- G ) - 0

(Eq. 10.44)

324 Chapter 10

Electrolyte energy balance

Qij - Saho(T,,ij - To,a) - Sahy(Te,ij- T /a) = 0 (Eq. 10.45)

where Sa is the heat transfer area, Td the datum temperature, r the gas flow rate,

Cp the specific heat capacity, and

_ To,i,,z a - To,~j (Eq. 10.46) 2

Tl, inta- Tf, iJ (Eq. 10.47) 2

The model solution starts at the node where the fuel and oxidant inlet edges meet, because both inlet streams are completely defined. The average

partial pressures and exit temperatures of the node are assumed, and a first

estimate of the Nernst potential is obtained. The resistance and current are

calculated; then the energy balance equations are solved to determine the average electrolyte temperature in the node and the temperatures of the gas streams

leaving the node. The amount of fuel and oxidant consumed is calculated from the current to determine the new compositions of the gases. The compositions are used to calculate the partial pressures of the gas components at the node exit, which are, in turn, used to update the estimates of the average partial pressures. The iteration scheme is repeated until convergence of the partial pressures has been reached.

The model has been used to simulate a 9-cm square cell in a SOFC stack. An example of input geometric and process data is given in Table 10.1. The

results of the model calculations based on the input data in Table 10.1 are

summarized in Table 10.2. Figures 10.12 and 10.13 show the current density

distribution and electrolyte temperatures across the cell, respectively. In Figure

10.12, the current density is highest near the intersection of the fuel inlet and

oxidant outlet edges because (i) the hydrogen concentration is high in the fuel

stream, leading to higher Nernst potential, and (ii) the temperatures in this region

are higher, resulting in better conductivity in the electrolyte, thus lower resistances. The temperature is also highest in this region because (i) the oxidant

entering these nodes has picked up sensible heat from previous nodes, and (ii) the

high current densities generate large amounts of heat, as seen in Figure 10.13.


TABLE 10.1

Model Input Data [10.20]

Geometric data

Electrolyte thickness, mm

Electrode thickness on electrolyte, mm

Interconnect thickness, mm

Electrode thickness on interconnect, mm

Channel wall thickness, mm

Channel height, mm

Channel width,mm

Fuel edge length, mm

Oxidant edge length,mm

Process data Hydrogen content in fuel inlet, volume fraction

Water content in fuel inlet, volume fraction

Inlet temperature, ~

Fuel inlet flow, standard L/h

Oxygen content in oxidant inlet, volume fraction

Nitrogen content in oxidant inlet, volume fraction

Oxidant inlet flow, standard L/h

Cell operating voltage, V

0.180

0.025

0.120

0.025

0.15

1.00 1.00

90

90

0.97

0.03

900

11.5

0.21

0.79

434

0.75

TABLE 10.2

Model Output Data [10.20]

Fuel utilization, %

Oxygen utilization, %

Fuel efficiency, %

Gross power, W

Net power, W

Specific power, kW/kg

Average Nernst potential, V

Average current density, mA/cm 2

Average cell temperature, ~

Maximum cell temperature, ~

Maximum temperature gradient, ~

85

5.2

49

17.0

16.3

0.60

0.88

280

934

1050

187

326 Chapter 10

TO0"

~E 8o u < E

o~ z w o

z oll oc oc

" '~'~tVT EDG'E~'*t.E

0

0

8

0.00 u.

Figure 10.12. Current distribution across cell layer (Geometric and process data given in Table 10.1.) [10.20]

.. : : . . : : . : : .:::,. : . ~ : : ~ -~ . . . . . . . .

' i

/

CT"lo,,w ,,,

Figure 10.13. Temperature distribution across cell layer (Geometric and process data given in Table 10.1.) [10.20]

As mentioned earlier, gas flow configuration plays a major role in

determining current and temperature distributions in a fuel cell stack. For example, a model has been developed to simulate current and temperature

distributions in SOFCs having crossflow, coflow, and counterflow configurations

(Figure 10.14) [10.35]. This model considers a stack of square geometry with metallic interconnect. The stack has 20 channels, each 4 mm wide, 1 mm high,

and 1 dm long, for fuel and oxidant gases. The anode, electrolyte, and cathode are 120, 200, and 50/zm thick, respectively. The gas composition is 54% H2,

Modeling and Analysis 3 2 7

29% H20, 12% CO, and 5 % CO2; the air stoichiometric ratio is 2; and the inlet temperature is 800~ As can be seen from Figure 10.14, the counterflow configuration shows the highest temperature; the coflow has a more uniform current density profile; and the crossflow presents a larger thermal gradient. Internal reforming of hydrocarbons significantly affects current and temperature distributions in SOFC stacks [10.21,10.33,10.35]. Figure 10.15 shows, as an example, the differences in temperature distribution in a SOFC operated with hydrogen (no internal reforming) (90 % H2, 10 % H20) and with methane (internal reforming) (25 % CH 4, 75 % H20 ) [10.21]. In this case, the inlet temperature is

900~ and the air stoichiometric ratio is 12.

CROSSFLOW

TEMPERATURE, K

I/lll/lllIl[[/!1'0 4

2

0 0 2 4 6 6 10

DISTANCE FROM OXIDANT INLET, c.m

TEMPERATURE DISTRIBUTION

COFLOW COUNTERFLOW

1 ~ 0 " 1 ~ 0

. 1370 v

, 1 .1o ~" 2 12~o =~ I

i ~ OXIDANT : 1 1 3 0

! s o u p i lO7O _ , , r / 1

0 2 4 ' 6 " " 0

DISTANCE FROM FUEL INLET, c m

1370

1 3 1 0

1250 \ i 1190

1070 o " i " ; " ; " ; o

DISTANCE FROM FUEL INLET, cm

CURRENT DISTRIBUTION

CURRENT DENSITY, A/cm" , o

~, . _ ~

~ 6

~: 4

0 0 2 4 6 8 10

DISTANCE FROM OXIDANT INLET, c m

1.0

"E o 0.8

>-

~ 0.6 t~ Q

~" 0 .4 II:

"N,, COFLOW �9 - ....... COUNTERFLOW \\,

"~,, "\,

0 2 4 6 O 10

DISTANCE FROM OXIDANT INLET, c m

Figure 10.14. Simulated current and temperature distributions in SOFCs having crossflow, coflow, and counterflow configurations [10.35]

328 Chapter lO

WITHOUT INTERNAL REFORMING (90% H 2, 10% H20)

91o ~ i

WITH INTERNAL REFORMING (25% CH4, 75% H20)

~Jo. 9671 LU t1"1._:3 94~]

IX.l ' irt" '~" 977]

,,,:~ 907] I,-- .................

887J

Figure 10.15. Temperature distributions in SOFC operated with hydrogen (no reforming) and with methane (internal reforming) [10.211

SOFC systems can be designed to influence current and temperature

distributions within the stack, thus maximum current output. For example, a

sealless tubular SOFC system with conventional air cooling would produce an

average current density of 370 mA/cm 2 at 0.606 V with 80 % fuel utilization and

25 % air utilization [10.36]. Fuel recycling modifies the current and temperature

distributions of the fuel cell. Thus, a fuel recycling system would produce 0.926

mA/cm z at 0.811 V with 10 % fuel utilization per cycle and 100 % air utilization.

As a result, power generation density is improved.

References

10.1

10.2 10.3

10.4

10.5

J.R. Selman, in Proceedings of the International Symposium on Solid Oxide Fuel Cells, November 13 - 14, 1989, Nagoya, Japan, O. Yamamoto, M. Dokiya, and H. Tagawa (eds.), Science House, Tokyo, Japan, 1989, p. 212. S. Majumdar, T. Claar, and B. Flandermeyer, J. Am. Ceram. Soc., 69 (1986) 628. Westinghouse Electric Corporation, High-Temperature Solid Oxide Electrolyte Fuel Cell Power Generation System, Quarterly Summary Report, January 1, 1984- March 31, 1984, Report No. DOE/ET/17089--2217, U.S. Department of Energy, Washington, DC, 1984. A. Saigal and S. Majumdar, presented at 1992 ANSYS Technology Conference, May 4-7, 1992, Pittsburgh, PA, CONF-9205129--1, U.S. Department of Energy, Washington, DC, 1992. E.F. Sverdrup. C.J. Warde, and R.L. Eback, Energy Convers., 13 (1973) 129.


10.6 10.7

10.8

10.9

10.10

10.11

10.12

10.13

10.14

10.15

10.16

10.17

10.18

10.19

10.20

10.21

10.22

10.23

10.24

10.25

10.26

S. Nagata, K. Shirai, and H. Sato, Bull. Electrotech. Lab., 40 (1976) 974. S. Nagata, Y. Ohno, Y. Kasuga, Y. Kaga, and H. Sato, in Proceedings of the 19th IECEC, August 19-24, 1984, San Francisco, CA, Vol. 2, American Nuclear Society,

La Grange Park, IL, 1984, p. 827. Y. Ohno, Y. Kaga, and S. Nagata, Electr. Eng. Jpn., 107 (1987) 59 (translated

from Denki Gakkai Ronbunshi, 106B (1986) 693). U.G. Bossel, Performance Potentials of Solid Oxide Fuel Cell Configurations, Report No. EPRI TR-101109, Electric Power Research Institute, Palo Alto, CA,

1992. U.G. Bossel, in Proceedings of the Third International Symposium on Solid Oxide Fuel Cells, May 16-21, 1993, Honolulu, HI, S.C. Singhal and H. Iwahara (eds.),

Electrochemical Society, Pennington, NJ, 1993, p. 833.

T. Kudo and H. Obayashi, Energy Convers., 15 (1976) 121.

P.G. Debenedetti and C.G. Vayenas, Chem. Eng. Sci., 38 (1983) 1817.

P.G. Debenedetti, C.G. Vayenas, I. Yentekakis, and L.L. Hegedus, in American Chemical Society Symposium Series, 237 (1984) 171.

W.J. Wepfer and M.H. Woolsey, Energy Convers., 24 (1985) 477.

W.R. Dunbar and R.A. Gaggioli, in Proceedings of the 23rd IECEC, July 31-August 5, 1988, Denver, CO, American Society of Mechanical Engineers, New York, 1988,

p.257. W.J. Marner, J.W. Suitor, and C.R. Glazer, see Ref. 10.15, p. 265.

C.Y. Lu and T.M. Maloney, in 1988 Fuel Cell Seminar Abstracts, October 23-26, 1988, Long Beach, CA, Courtesy Associates, Washington, DC, 1988, p. 78.

T.M. Maloney and G.A. Coulman, in 1990 Fuel Cell Seminar Abstracts, November 25-28, 1990, Phoenix, AZ, Courtesy Associates, Washington, DC, 1990, p. 239. S. Ahmed, R. Kumar, H. Walden, and K.P. Barr, in Proceedings of the 26th IECEC, August 4-9, 1991, Boston, MA, Vol. 3, American Nuclear Society, La

Grange Park, IL, 1991, p. 594. S. Ahmed, C.C. McPheeters, and R. Kumar, J. Electrochem. Soc., 138 (1991)

2712. E. Erdle, J. GroB, H.G. Miiller, W.J.C. Miiller, H.-J. Reusch, and R. Sonnens-

chein, in Proceedings of the Second International Symposium on Solid Oxide Fuel Cells, July 2-5, 1991, Athens, Greece, F. Grosz, P. Zegers, S.C. Singhal, and O.

Yamamoto (eds.), Commission of The European Communities, Luxembourg, 1991,

p. 265.

E. Arato and P. Costa, see Ref. 10.21, p. 273.

I.V. Yentekakis, S. Neophytides, S. Seimanides, and C.G. Vayenas, see Ref. 10.21,

p. 281.

D.G. Walhood and J.R. Ferguson, see Ref. 10.21, p. 289.

A. Solheim, R. Tunold, and R. OdegArd, see Ref. 10.21, p. 297.

J.R. Ferguson, see Ref. 10.21, p. 305.

330 Chapter 10

10.27

10.28

10.29

10.30

10.31

10.32

10.33

10.34

10.35

10.36

Z. Takehara, K. Kanamura, A. Hirano, and M. Ippommatsu, see Ref. 10.21, p.

313.

N.F. Bessette II and W.J. Wepfer, in 1992 Fuel Cell Seminar Abstracts, November

29-December 2, 1992, Tucson, AZ, Courtesy Associates, Washington, DC, 1992,

p. 519.

J. Hartvigsen, S. Elangovan, and A. Khandkar, see Ref. 10.28, p. 532.

A. Solheim, see Ref. 10.10, p. 841.

T. Sira and M. Ostenstad, see Ref. 10.10, p. 851.

C. Bleise, J. Divisek, B. Steffen, U. K6nig, and J.W. Schultze, see Ref. 10.10, p.

861.

H. Karoliussen, K. Nisancioglu, A. Solheim, and R. Odeghrd, see Ref. 10.10, p.

868. J. Hartvigsen, S. Elangovan, and A. Khandkar, see Ref. 10.10, p. 878.

A. Malandrino and M. Chindemi, see Ref. 10.10, p. 885.

A. Hirano, M. Suzuki, and M. Ippommatsu, J. Electrochem. Soc., 139 (1992)2744.

Chapter 11

SYSTEM AND APPLICATION

11.1 GENERAL

SOFCs and other types of fuel cell are an attractive option for future electric power generation [11.1]. Compared with conventional generation methods, fuel cell systems offer several advantages: substantially higher energy conversion efficiency, modular construction, higher efficiency at partial loads, minimal siting restrictions, potential for cogeneration, and much lower production of pollutants. Figure 11.1 presents, as an example, the required coal usage and emission outputs for a given constant power production of 21 MW for four coal-based systems: (i) pulverized coal (PC) boiler and advanced PC boiler with flue gas scrubbing, (ii) advanced pressurized fluidized bed combustion (APFBC), (iii) integrated gasification combined cycle (IGCC), and (iv) integrated gasification fuel cell (IGFC) systems [11.2]. As seen from Figure 11.1, the fuel cell system has the most efficient coal usage, the lowest CO2 and other emissions, and the lowest amount of waste material. Because of the many desirable characteristics, fuel cells have been considered for a variety of applications, ranging from large multimegawatt electric utility power plants to compact kilowatt space power systems. For commercial terrestrial applications, economic studies have indicated that fuel cells could be produced and operated at costs that are competitive with other power generation systems [11.3].

As compared with the other types of fuel cell, the SOFC has the unique characteristics of all-solid state (mainly ceramic) construction, along with high

operating temperature (highest among the present generation of fuel cells). These characteristics offer a number of attractive features for the SOFC. For example, because all the cell components are solid, the SOFC can be configured into compact and lightweight structures unachievable in fuel cell systems having a liquid electrolyte. The high operating temperature, along with a tolerance to fuel impurities, makes the SOFC suitable for combination with coal gasification plants. The heat released from the cell can be efficiently transferred and used for

332 Chapter 11

Circa 1920 Circa 1990 A P F B C I G C C I G F C PC PC *

~ a , ~ ~ -O 0 ~ . . . . ~ ~ ~ ~s~ge \ " ' " / "~= " "

O O

r~ 1.4 1.5 1.0 0.8

To,. 0.02 < 0.001 < 0.001

SO2 ~ c~ c:> => r r.,. ~ 0.06 0.05 0.0003 0.00003

Note: Constant power output of 21.2 MW APFBC - Advanced Pressurized Fluidized Bed Combustion IGCC - Integrated Gasification Combined Cycle IGFC - Integrated Gasification Fuel Cell "with flue gas desulfurization scrubbing

Figure 11.1. Required coal usage and emissions f or four coal-based systems [11.2]

coal gasification or hydrocarbon reforming. The high operating temperature also provides high-quality byproduct heat suitable for use in cogeneration or a bottoming cycle. The main potential applications of the SOFC are the electric

utility and cogeneration sectors. SOFCs have also been proposed for use in transportation, space, and other applications.

11.2 ELECTRIC UTILITY

The electric utility application of the SOFC includes central station electric power generation and dispersed plants located near load centers. In addition, the fuel cell can be used for repowering existing electric utility units.

The preferred electric utility applications are baseload central stations (coal) in the range of 250 to 500 MW, repowering of existing power units (natural gas,

oil, coal) in the range of 150 to 300 MW, and dispersed SOFC plants (natural gas and oil) in the range of 5 to 50 MW [11.4,11.5].

System and Application 333

Comparative assessment of fuel cell electric utility power plants

integrated with coal gasification plants has established that of the three fuel cell systems (phosphoric acid, molten carbonate, and solid oxide), the SOFC is

expected to have the highest overall efficiency for conversion of coal to

electricity [11.6]. Figure 11.2 shows a flow diagram of a conceptual design of a 250-MW integrated coal gasification/SOFC-steam turbine system (IG/SOFC- ST) [11.7]. The power plant consists of five major subsystems: coal gasification,

fuel-gas cleanup and rotating equipment, SOFC, heat recovery, and steam turbine. In this design, an air-blown, fluidized bed gasifier is applied in

combination with hot gas filtration for particulate removal and zinc ferrite catalyst beds for desulfurization. The bottoming system is a high-performance, regenerative, reheat steam cycle. For this power plant, natural gas is used during startup of the coal gasification system. Thus, natural gas is available and could be readily substituted for the coal-derived fuel gas without affecting the specification of the SOFC and steam systems. Performance estimates for the

power plant are summarized in Table 11.1.

FUEL GAS C L E A N U P HEAT RECOVERY C O A L G A S I F I C A T I O N S Y S T E M A N D R O T A T I N G E Q U I P M E N T S Y S T E M SOFC S Y S T E M S Y S T E M

I I ~ Air I I ' 1~-~-~. ,

1 I I a

Rllurn ASh I I [ I F!{I~zld [ 1 "1~ urn $::: h 1~7r r -I I ~----~ I I Turtl~ UrlIL :

, - - - , L . . . . . .

, : . . . . . J L j I I ,.,_~ 1 :

L " ;? ' ' ~ " J

L _j S T E A M TURBINE S Y S T E M

Figure 11.2. Process flow diagram of an IG/SOFC-ST power plant [11.7]

334 Chapter I 1

TABLE 11.1

250 MW IG/SOFC-ST Power Plant Performance Summary [11.7]

Coal flow rate, tons/day Plant power output, MW

Fuel gas expander SOFC generator Steam turbine

Plant gross AC power, MW Plant net AC power, MW Net plant heat rate (HHV*), BTU (1.054 kJ)/kWh Bottoming cycle steam conditions

Pressure, atm (1.01 x 105 Pa) Superheat temperature, ~ Reheat temperature, ~

Final SOFC exhaust temperature, ~

2300

41 117 104 262 250 7640

100 538 538 144

"HHV = high heating value

Various options and configurations of gasifiers, SOFCs, and heat engines

can be used in the design of the power plant. For example, gasification systems

can be oxygen-blown or air-blown, and SOFCs can be atmospheric or pressur-

ized. Table 11.2 summarizes the results of a study of four conceptual 675-MW

SOFC system configurations [11.8]. The four cases are (i) an air-blown

gasification system with an atmospheric pressure (1.1 atm) SOFC and a Rankine

bottoming cycle (denoted A); (ii) an oxygen-blown gasification system with an

atmospheric-pressure SOFC and a Rankine bottoming cycle (denoted B); (iii) an

air-blown gasification system with a pressurized (11.2 atm) SOFC and a

combined Brayton/Rankine bottoming cycle (denoted C); and (iv) an oxygen-

blown gasification system with a pressurized SOFC and a combined Brayton/

Rankine bottoming cycle (denoted D). As seen from Table 11.2, the combina-

tion of a pressurized SOFC with an air-blown fluidized bed coal gasifier and a

combined Brayton/Rankine bottoming cycle has the potential to achieve

conversion efficiencies of 58%. Other studies have also shown that high

efficiencies (up to 70%) can be achieved for SOFC systems integrated with

bottoming cycles [11.9-11.11].


TABLE 11.2

Summary of Performance for Four 675-MW SOFC Systems [11.8]

CASE A B C D

Power output, MW Fuel cell 387 423 279 307 Combustion turbine 0 0 226 226 Steam turbine 212 184 153 127 Fuel gas expander 88.8 81.0 26.6 24.9 Inverter -7.8 -8.3 -6.2 -6.6 Pump -5.0 -4.7 -3.4 -3.4

Coal flow rate, 1000 lb (454 kg)/h 359 383 322 351 Heat rate, BTU (1.054 kJ)/kWh 6513 6951 5841 6366 Efficiency, % 52.4 49.1 58.4 53.4

11.3 COGENERATION

SOFCs are potentially suited to commercial and industrial cogeneration

applications (kW to MW sizes). For commercial applications, the fuel cell

system supplies the electrical and thermal requirements of residential, commer-

cial, and light industrial buildings. In this case, SOFCs offer a better match of

thermal-to-electric ratio for existing buildings than conventional cogeneration

options. The fuel cell also provides several important features for these

applications, such as low noise and low vibration, high power density, and small

footprint. Economic assessment of SOFCs for the commercial building sector

reveals that SOFCs are strong competitors against reciprocating engines in low

thermal-to-electric and low load factor sites [11.12]. For industrial cogeneration applications (e.g. chemical processing and

metallurgical industries), the thermal output of the SOFC would be used to

supply industrial process heat (most often in the form of steam), and the fuel

generally would be supplied from a byproduct stream from some chemical

process in the plant. For industrial cogeneration, the high-quality heat from the

fuel cell provides many opportunities for thermal integration. Industrial SOFC-

based cogeneration systems would also possess the benefits associated with fuel

cells, such as resource saving, low emissions, and modularity.

336 Chapter 11

Figure 11.3 shows a conceptual design of a commercial natural gas-

fueled SOFC cogeneration system rated at 200 kW net AC power [11.7,11.13]. In this particular design, the incoming fuel is preheated using heat from the

SOFC exhaust and desulfurized in a reactor containing ZnO. The heat recovery

system consists of two process air heaters and a heat recovery steam generator.

The heat recovery steam generator is positioned between the two air heaters, and

the heaters are sized to generate steam at the desired temperature and required

pressure at the generator. Figure 11.4 provides a pictorial view of the installed

200-kW cogeneration system. The estimated weight of this system is 9,300 kg,

and the projected footprint is 16 m 2. Table 11.3 summarizes system performance

[11.7].

TABLE 11.3

200-kW SOFC Cogeneration System Performance [11.7]

Natural gas flow rate, kg/h Fuel HHV input rate, kW DC power output, kW Gross AC power output, kW Net AC power output, kW Terminal DC efficiency (HHV), % System net AC efficiency (HHV), % Exhaust heat recovered for site use, kW System fuel efficiency (HHV), % Steam generation rate, kg/h Steam conditions

Pressure, atm (1.01 x 105 Pa) Temperature, ~

Potential steam chiller cooling effect, tons Final SOFC exhaust temperature, ~

34.0 479.3 223.9 210.4 200.0 46.7 41.7 172.3 77.7 262.2

8.0 170.0 60.0" 66.0

"Chiller coefficient of performance = 1.23

At design conditions, this SOFC cogeneration system produces net

electrical power of 200 kW with a system efficiency of 41.7 % based on the fuel

high heating value. The system fuel effectiveness (the percentage of the

incoming fuel energy that is converted to electric power and heat for site use) is


NATURAL GAS

PROCESS AIR

AIR PREHEATER

ZnO ~ I L ~--1--: DESULFURIZER FEEDWATER -'--A ~ STEAM

POWER CONDITIONING

HRSG"

AIR PREHEATER

"HRSG - HEAT RECOVERY STEAM GENERATOR

Figure 11.3. Conceptual design of a 200-kW SOFC cogeneration system [11.7]

CONTROLS ~ fBOILER

~ o ~ ~- - / / ~ ~ ~_ _ . ~ .,~";~,~.~7"~'2 o~

2.2~ 4.0 m ~ ~ . J ~ FUEL COMPRESSOR

Figure 11.4. Pictorial view of a 200-kW SOFC cogeneration system [11.13]

338 Chapter 11

77.7%. Recently, a 20-kW SOFC cogeneration system has been designed,

constructed, and tested [11.14].

Figure 11.5 shows a schematic diagram of a SOFC cogeneration system for an aluminum production plant [11.15]. In this preliminary design, exhaust heat is used to produce some steam and some electrical power. The design is

based on the electrical and thermal demand of the aluminum plant and its correlation to electrical and thermal output of a 220-MW SOFC generator.

11.4 TRANSPORTATION

Fuel cells are presently being considered for possible use as a power source in transportation applications [11.16-11.20]. The key motivating factors are the potential energy saving and emission reduction offered by fuel cell systems. Fuel cell powered vehicles are projected to attain efficiencies of 35 to 55 % after thermal and parasitic power losses are taken into account, whereas efficiencies for current internal combustion engine vehicles are only about 18 %

[11.16]. Fuel cell powered vehicles have negligible levels of CO, NOx, and

other pollutant emissions as compared to those of internal combustion engines

(Figure 11.6) [11.21]. Fuel cell'and internal combustion engine systems are

compared schematically in Figure 11.7 [11.22]. (It should be noted that the fuel processor may not be required for SOFC-based systems).

FUEL PREHEATER ___ ~ CLEAN COAL GAS .~j,~' ~] GAS PPELINE "-~ t -L . I-

AIR ~

STARTUP BURNER

STACK " T

AIR Z z ~

REGENERATIVE FAN AIR HEATER

AFTERBURNER

L I-

. I ]= I ECONOMIZER WASTE ....... HEAT BOILER

J 1: t i- '

FEED WATER [ ~ HEATER

FEED PUMP ' ~

SOFC

TURWNE GENERATOR

_ ~ PRETREATED MAKEUP WATER

Figure 11.5. Schematic diagram of a 220-MW industrial SOFC cogeneration system [11.15]


3 _o

..... o

O 2

tel

1

0

CO (,3.4)

NOx(2.0)

!

INTERNAL FUEL CELL COMBUSTION ENGINE

"NOT DETECTABLE

Figure 11.6. Comparison of emissions of internal combustion engine and fuel cell vehicles [11.21]

INTERNAL COMBUSTION (IC) VEHICLES

--H H H FUEL IC DRIVE I i TANK ENGINE TRANSMISSION SHAFT/ H WHEELS

AXlE ~11 FUEL CELL VEHICLES

H -i FUEL FUEL FUEL ELECTRIC WHEELS TANK PROCESSOR CELL MOTOR

....

SYSTEM CONTROLLER

Figure 11.7. Comparison of internal combustion engine and fuel cell systems [11.22]

The operating conditions of, and design constraints on, fuel cells when

used in transportation applications differ considerably from those of stationary applications. In transportation, the fuel cell must have sufficient power density

to meet the performance specifications of the vehicle while fitting within the

available space. It must be sufficiently inexpensive to compete with internal

combustion engines on an economic basis. It must be safe and reliable. The fuel

cell also must be able to start rapidly and to respond quickly to changes in power

demand. Fuel cell systems can be designed to meet these requirements. Among the various proposed fuel cell propulsion systems, the SOFC is potentially the

340 Chapter 11

simplest, due to simplified onboard fuel processing. The SOFC can internally

reform methanol or other hydrocarbon fuels and does not require external

reformers. The SOFC can also operate at high current and power densities, thus

offering compactness. Due to its all-solid state, the SOFC can be designed in a variety of orientations and configurations, an important consideration in

packaging the system for transportation applications. Figures 11.8 and 11.9

show layouts for a fuel cell passenger car [11.18] and a fuel cell bus [11.23], respectively. These layouts include external reforming for the fuel cell system;

however, the external reformer may be eliminated for SOFC-based systems.

FUEL CELL STACI 20 kW NOMINAL 60 kW PEAK

METHANOL REFORMER

INSTRUMENTATION AND AUTOMATIC CONTROL SYSTEM

ELECTRIC MOTOR MOTOR CONTROLLER

TRANSMISSION

Figure 11.8. Layout for a conceptual fuel cell passenger car [11.18]

Air Tank

/ Air Compressor Drive Motor

Fuel Cell Stacks Radiators

Starting Hydrogen Battenes DC-DC \Traction Motor Water Tank Storage Tanks (6) Converters \

Air Compressol

Power Steering

Aulomahc Transmission

Figure 11.9. Layout for a conceptual fuel cell bus [11.23]


A schematic diagram of a conceptual 60-kW SOFC propulsion system

fueled with methanol is given in Figure 11.10 [11.24]. In this system, the

methanol fuel is vaporized using the hot fuel-side exhaust gas from the SOFC as

a heat source and fed directly to the fuel cell. Spent fuel is burned to preheat the

incoming air. A fraction of the SOFC air and fuel streams is recirculated to

minimize thermal stresses in the fuel cell. In addition to the fuel cell generator,

the major components of the system are an air heater, a spent fuel burner, and

fuel and air recirculators. Other system components include a feed air blower,

a fuel injector, power conditioner, and diagnostic and control instrumentation

[11.25]. In this design, the fuel cell generator consists of 16 stacks; the operating and design parameters of each stack are summarized in Table 11.4.

Table 11.5 compares the performance of the system at 60-kW and 25-kW power

levels.

[~ TEMP. ~

AIR EJECTOR ,=

F

6:kw soft 1

L i!" ! [3ooi ' !

,

~ FUEL I [ B U R N E R l

FUEL EJECTOR

J

AIR I METHANOL FUEL 104 scfm" EXHAUST 23.9 kg/h

118 scfm" "sofm = STANDARD CUBIC FOOT PER MINUTE

Figure 11.10. Schematic diagram of a conceptual 60-kW SOFC propulsion system [11.24]

342 Chapter 11

TABLE 11.4

Operating and Design Parameters of SOFC Stack for Vehicular Propulsion System [11.24]

Stack power, kW 3.75 Current density, A/cm 2 0.735

Cell voltage, V 0.6 Gas passage height, mm 0.7

Area-specific resistance, fl.cm 2

Materials 0.122 Interfacial 0.150

Air flow (x theoretical) 4

Fuel utilization, % 85 Stack dimension, cm

Length 13.2 Width 8.9

Height 13.3 Stack power density, kW/L 2.32

TABLE 11.5

Summary of SOFC Propulsion System Performance [11.24]

60 kW

Power

25 kW

Methanol feed, kg/h

Energy input (HHV), kW

Cell energy losses, kW

Heat duty, kW

Fuel vaporizer

Air preheater

Spent fuel burner Air supply blower power, kW Net power, kW

Efficiency, %

23.9

150.8

67.0

15.3

15.5 22.6 1.17

60

39.9

8.2

51.6

18.5

5.2

9.7 7.7 0.42

25

48.6


Preliminary thermodynamic analysis indicates that SOFC propulsion systems should be operated at the highest practical electrochemical fuel utilization rates for best efficiency and dynamic response, as well as minimum component volume and cost [11.26]. In general, because of its high operating temperature,

a SOFC system must be thermally insulated and should be kept reasonably warm, even when no net power is drawn, to minimize the power-up time. Therefore, the SOFC is most attractive for applications that typically have long duty cycles, i.e., infrequent startup and shutdown with long periods of operation in between. Examples of such applications are locomotives, trucks, barges [11.25].

11.5 SPACE AND OTHER APPLICATIONS

Fuel cells have been used for electric power generation in space missions

because fuel cell systems have been shown to have the capability to satisfy onboard power requirements along with the requirements of efficiency, weight,

life, reliability, safety, and mission flexibility. Recently, regenerative fuel cells (combined fuel cells and water electrolyzers) have been identified as an enabling

technology for space missions and extraterrestrial surface energy storage applications. For example, regenerative fuel cell systems have been evaluated for use in spacecraft missions operating in the 10- to 50-kW range for many years in mid-altitude to geosynchronous orbits [11.27]. Plans for future exploration of the solar system have considered regenerative fuel cells as power

sources for permanent nuclear-powered bases at manned outposts to support base operations and human expeditions [11.28].

A regenerative fuel cell system for space applications is commonly based on solar energy (Figure 11.11) [11.29]. During periods of sunlight, photovoltaic solar arrays generate power to the fuel cell, which acts as a water electrolysis unit to produce hydrogen and oxygen. During periods of darkness, the fuel cell discharges electricity through consumption of hydrogen and oxygen. The design of a regenerative fuel cell system requires the specification of a number of

parameters including (but not limited to) fuel cell power level, type of solar array and radiator, peak and emergency power requirements, desired operating life,

mass and volume limitations, degree of voltage regulation, and level of integration with other systems.

Figure 11.12 shows the schematic diagram of a regenerative SOFC system designed to provide darkside energy storage for solar photovoltaic power

[11.30]. The major elements of this system are the SOFC converter unit, storage

344 Chapter I 1

ENERGY STORAGE SYSTEM

............................................

............................................

, E L E C T R O L Y Z E R I

IMIl~m

k W

k W

k W

M A I N

P O W E R

B U S

Figure 11.11. Regenerative fuel cell system for space applications [11.29]

SOLAR PHOTOVOLTAIC PANEL

FUEL GAS 90% H~ 10% H20

PRESSURE - 6 - 100 =tin

INTAKE/ OUTLET

THERMAL CONTROL

T,P SENSOR

INSULA_T_ED PRESSURE VESSEL , WITH SUMP J

, ELECTRICA �9 POWER

RADIATOR

SOFC CONVERTER UNIT

"INTERNAL RADIATOR LOOP OUT" - SENSOR WIRING HARNESS ~ I L ~ ,

== I FUEL INLET/OUTLET IT = ~ "

�9 INTERNAL RADIATOR LOOP IN ~ SENSOR WIRING HARNESS==

INTERNAL RADIATOR LOOP OUT --------

OXYGEN

PRESSURE - 5 - 100 mtm

~ INTAKE/ OUTLET

�9 INSULATED PRESSURE VESSEL

Figure 11.12. Regenerative SOFC system [11.30]


tanks (for oxygen, hydrogen, and water), and a radiator. The SOFC converter

is a sealed unit shown schematically in Figure 11.13. The converter comprises an insulated high-temperature section (SOFC stack and reactant preheat heat

exchangers) and a low-temperature section (gas flow and fluid transport). Within

the high-temperature section, the SOFC is bordered on four sides by two separate

sets of reactant preheat heat exchangers for the oxygen and hydrogen gas streams. The low-temperature section contains the waste heat and water removal

portion of the recirculating loop. The components of this section include a hydrogen pump/separator, a condenser, pumps, and valves. This system does

not take the advantage of the SOFC high-temperature heat rejection capability to minimize radiator size, but condenses the product water to minimize reactant

storage instead. Table 11.6 summarizes the component weights for a 10-kW

solar photovoltaic/regenerative SOFC power system.

WATER

REACTANT PREHEAT ]--~t

POWER

LIQUID WATER DRAIN

DISCHARGE VALVE INSERTION VALVE===~ Ilill I ....

HYDROGEN OXYGEN

RADIATOR LOOP IN

RA D IATO R

LOOP OUT

<I==OR ==r

CONDENSER I I

~ FLOW m,.~ PROPO RTIO NING

VALVES

COOLANT CIRCULATION PUMP

,": i "'" -.-=,

i

I 1 ..--1.

",,',,',, I

'r

. r

H I G H - T E M P E R A T U R E S E C T I O N

Figure 11.13. Schematic diagram of SOFC converter unit for a regenerative system [11.30]

346 Chapter 11

TABLE 11.6

Component Weights for 10-kW Solar Photovoltaic/Regenerative SOFC Power System [11.30]

Power System Summary

Power level (delivered) Shade time assumed Sun time assumed Required solar photovoltaic daylight power

10 kW 13h 11h 29.9 kW

Component Weight, kg

Reactant and tankage Hydrogen Tank (1.32 m diameter sphere) Oxygen Tank (1.00 m diameter sphere)

Converter unit SOFC stack (10 by 10 by 46 cm) Reactant preheat heat exchanger Insulation package Condenser heat exchanger Low-temperature section ancillary Structure and containment

Radiator and solar photovoltaic panel External radiator Solar panel and mount

11 182 42 84

20 5.3 6 1.5 38 36

91 300

Total weight 817

A regenerative fuel cell based on the monolithic SOFC has also been

proposed for coupling with a nuclear reactor for space-based pulse power systems

[11.31,11.32]. This power system concept has been shown to be capable of

producing very high power bursts for long durations with significant short

recharge times. Other system configurations have been suggested to integrate the


SOFC in an electrochemical loop or in a regenerative mode with chemical

storage for space applications [11.33].

In addition to space, fuel cell systems in general and SOFCs in particular

are suitable for marine (underwater vehicles [11.34-11.36] and surface ships

[11.34,11.35,11.37,11.38]) and military applications [11.39]. For example,

SOFCs have been found to be advantageous as replacements for diesel electric

generators for marine auxiliary power and low- to medium- (1-MW) power

propulsion systems using reformed diesel fuel [11.40]. Some possible military

applications for SOFCs include mobile power generators, transportation vehicles,

airfield lighting, communication and transmitter devices, satellite power systems,

prime power remote site generators, and other applications.

References

11.1

11.2

11.3

11.4

11.5

11.6

11.7

11.8

11.9

Notice of Market Opportunity for Fuel Cells, American Public Power Association, Washington, DC, 1988. W.T. Langan, in Proceedings of the Fourth Annual Fuel Cells Contractors Review Meeting, July, 14-15, 1992, Morgantown, WV, W.J. Huber (ed.), Report No. DOE/METC-92/6127, U.S. Department of Energy, Washington, DC, 1992, p. 3. P.D. Lilley, E. Erdle, and F. Gross, Market Potential of Solid Oxide Fuel Cell (SOFC), Report No. EUR 12249 EN, Commission of the European Communities, Luxembourg, 1989. W.G. Parker and E.R. Ray, in Proceedings of the First Annual Fuel Cells Contractors Review Meeting, May 2-4, 1989, Morgantown, WV, W.J. Huber (ed.), Report No. DOE/METC-89/6105, U.S. Department of Energy, Washington, DC, 1989, p. 26. W.G. Parker, S.M. Knable, and C.M. Zeh, in 1988 Fuel Cell Seminar Abstracts, October 23-26, 1988, Long Beach, CA, Courtesy Associates, Washington, DC, 1988, p. 248. C.J. Warde, R.J. Ruka, and A.O. Isenberg, Energy Conversion Alternate Study (ECAS), Westinghouse Phase I, Vol. 2-Fuel Cells, Report No. N76-23703, National Aeronautics and Space Administration, Washington, DC, 1976. W.L. Lundberg, in Proceedings of the 25th IECEC, August 12-17, 1990, Reno, NV, Vol. 3, American Institute of Chemical Engineers, New York, 1990, p. 218. D.Q. Hoover, E.V. Somers, E.J. Vidt, and R.E. Grimble, Evaluation of Solid Oxide Fuel Cell Systems for Electricity Generation, Report No. N83-23698 (Contract No. JPL-956252), National Aeronautics and Space Administration, Washington, DC, 1982. W. Drenckhahn and W. Schramm, in 1990 Fuel Cell Seminar Abstracts, November 25-28, 1990, Phoenix, AZ, Courtesy Associates, Washington, DC, 1990, p. 495.

348 Chapter 11

11.10

11.11

11.12

11.13

11.14

11.15

11.16

11.17

11.18

11.19 11.20 11.21 11.22 11.23 11.24 11.25

11.26

11.27

11.28

11.29

11.30 11.31

M. Hsu and D. Nathanson, in 1992 Fuel Cell Seminar Abstracts, November 29- December 2, 1992, Tucson, AZ, Courtesy Associates, Washington, DC, 1992, p. 510. A.L. Dicks, R.J. Carpenter, E. Erdle, D.F. Lander, P.D. Lilley, A.G. Melman,

and N. Woudstra, Solid Oxide Fuel Cell Systems Study-Volume 1, A.G. Melman

and N. Woudstra (eds.), Report No. EUR 13103 EN, Commission of the European Communities, Luxembourg, 1991. S.M. Knable and G.P. Merten, in Proceedings of the 23rd IECEC, July 31-August 5, 1988, Denver, CO, Vol. 2, American Society of Mechanical Engineers, New York, p. 213. W.L. Lundberg, Solid Oxide Fuel Cell Cogeneration System Conceptual Design, Report No. GRI-89/0162, Gas Research Institute, Chicago, IL, 1989. K. Shinozaki and T. Satomi, see Ref. 11.10, p. 358.

W. Feduska and A.O. Isenberg, J. Power Sources, 10 (1983) 89.

National Program Plan Fuel Cells in Transportation, Report No. DOE/CH-930 l a, U.S. Department of Energy, Washington, DC, 1993.

M. Krumpelt and C.C. Christianson, An Assessment and Comparison of Fuel Cells for Transportation Applications, Report No. ANL-89/28, Argonne National

Laboratory, Argonne, IL, 1989. R.A. Lemons, J. Power Sources, 29 (1990) 251.

P.G. Patil, J. Power Sources, 37 (1992) 171. P.G. Patil, see Ref. 11.10, p. 8. P.G. Patil, C.C. Christianson, and S. Romano, see Ref. 11.12, p. 227. M. Krumpelt, R. Kumar, J. Miller, and C. Christianson, see Ref. 11.10, p. 35. K.B. Prater, J. Power Sources, 37 (1992) 181. D.W. Dees and R. Kumar, see Ref. 11.9, p. 71. R. Kumar, M. Krumpelt, and K.M. Myles, in Proceedings ofthe Third International Symposium on Solid Oxide Fuel Cells, May 16-21, 1993, Honolulu, HI, S.C. Singhal and H. Iwahara (eds.), Electrochemical Society, Pennington, NJ, 1993, p. 948.

R. Kumar, M. Krumpelt, and B. Misra, in Proceedings of the 24th IECEC, August 6-11, 1989, Washington, DC, Vol. 3, Institute of Electrical and Electronics Engineers, New York, 1989, p. 1601.

R.K. Taenaka, E. Adler, E.J. Stifel, and K.B. Clark, in Proceedings of the 22nd IECEC, August 10-14, 1987, Philadelphia, PA, Vol. IV, American Institute of Aeronautics and Astronautics, New York, 1987, p. 2016.

P.R. Prokopius, A. Antoine, and R.B. King, in 1990 Fuel Cell Seminar Abstracts, Addendum, November 25-28, 1990, Phoenix, AZ, Courtesy Associates, Washington,

DC, 1990.

M. Warshay and P.R. Prokopius, J. Power Sources, 29 (1990) 193. D.J. Bents, see Ref. 11.27, Vol. II, p. 808. D.C. Fee, M.C. Billone, D.E. Busch, D.W. Dees, J. Dusek, T.E. Easier, W.A. Ellingson, B.K. Flandermeyer, R.J. Fousek, J.J. Heiberger, S. Majumdar, C.C.


11.32

11.33

11.34

11.35 11.36

11.37 11.38

11.39

11.40

McPheeters, F.C. Mrazek, J.J. Picciolo, and R.B. Poeppel, see Ref. 11.27, Vol. II, p. 803. D.C. Fee, S.K. Bhattacharyya, D.E. Busch, L.W. Carlson, T.D. Claar, D.W. Dees, J. Dusek, T.E. Easier, W.A. Ellingson, B.K. Flandermeyer, R.J. Fousek,

H.K. Geyer, J.J. Heiberger, K.D. Kuczen, S. Majumdar, C.C. McPheeters, F.C. Mrazek, J.J. Picciolo, and R.B. Poeppel, in Space Nuclear Power Systems 1986, M.S. EI-Genk and M.D. Hoover (eds.), Orbit Book, Malabar, FL, 1987, p. 209.

M. Hsu, in Proceedings of the 26th IECEC, August 4-9, 1991, Boston, MA, Vol. 3, American Nuclear Society, La Range Park, IL, 1991, p. 498.

Marine Applications for Fuel Cell Technology-A Technical Memorandum, OTA- TM-O-37, U.S. Congress, Office of Technology Assessment, Washington, DC, 1986. V.W. Adams, J. Power Sources, 29 (1990) 181. T.M. Maloney, see Ref. 11.10, p. 395.

W.H. Kumm, J. Power Sources, 29 (1990) 169. Y. Fukui, H. Komaki, N. Ambo, H. Morita, H. Nakamura, and R. Ono, see Ref. 11.10, p. 374.

R.R. Barthelemy, in Proceedings of the Symposium on Fuel Cells: Technology Status and Applications, November 16-18, 1981, Chicago, IL, Institute of Gas Technology,

Chicago, IL, 1982, p. 109. V.W. Adams, in Proceedings of the Second International Symposium on Solid Oxide Fuel Cells, July 2-5, 1991, Athens, Greece, F. Grosz, P. Zegers, S.C. Singhal, and

O. Yamamoto (eds.), Commission of the European Communities, Luxembourg, 1991, p. 247.


APPENDIX

SELECTED REFERENCES RELEVANT TO SOLID OXIDE FUEL CELL TECHNOLOGY

Book Chapters and Review Papers

.

~

o

,

.

,

A.J. Appleby and F.R. Foulkes, Fuel Cell Handbook, Van Nostrand Reinhold, New York, 1989. J.T. Brown, "High-Temperature Solid-Oxide Fuel Cells," Energy, 11 (1986) 209; also in Assessment of Research Needs for Advanced Fuel Cells, S.S. Penner (ed.), Report No. DOE/ER/30060-T1, U.S. Department of Energy, Washington, DC, 1985, p. 209. A. Hammou, "Solid Oxide Fuel Cells," in Advances in Electrochemical Science and Engineering, Vol. 2, H. Gerischer and C.W. Tobias (eds.), VCH, New York, 1992, p. 88. J.H. Hirschenhofer, D.B. Stauffer, and R.R. Engleman, Fuel Cells, A Handbook (Revision 3), Report No. DOE/METC-94/1006, U.S. Department of Energy, Morgantown Energy Technology Center, Morgantown, WV, 1994. H.S. Isaacs, "Solid Electrolyte Fuel Cells," in Proceedings of the Symposium on Fuel Cells: Technology Status and Applications, November 16-18, 1981, Chicago, IL, Institute of Gas Technology, Chicago, IL, 1982, p. 83. H.S. Isaacs, "Zirconia Fuel Cells and Electrolyzers," in Science and Technology of Zirconia, A.H. Heuer and L.W. Hobbs (eds.), American Ceramic Society, Columbus, OH, 1981, p. 406. K. Kinoshita, F.R. McLarnon, and E.J. Cairns, Fuel Cells, A Handbook, Report No. DOE/METC-88/6096, U.S. Department of Energy, Morgantown Energy Technology Center, Morgantown, WV, 1988. N.Q. Minh, "Ceramic Fuel Cells," J. Am. Ceram. Soc., 76 (1993) 563. (Japanese translation available from the Solid Oxide Fuel Cell Society of Japan, Japan) N.Q. Minh, "High-Temperature Fuel Cells, Part 2: The Solid Oxide Cell," CHEMTE CH, 21 ( 1991 ) 120.

352 Appendix

10.

11.

12.

13.

K.A. Murugesamoorthi, S. Srinivasan, and A.J. Appleby, "Research, Development, and Demonstration of Solid Oxide Fuel Cell Systems," in Fuel Cell Systems, L.J.M.J. Blomen and M.N. Mugerwa (eds.), Plenum Press, New York, 1993, p. 465. F.J. Rohr, "High-Temperature Fuel Cells," in Solid Electrolytes, P. Hagenmuller and W. van Gool (eds.), Academic Press, New York, 1978, p. 431. B.C.H. Steele, "High Temperature Fuel Cells and Electrolysers," in Electrode Processes in Solid State Ionics: Theory and Application to Energy Conversion and Storage, M. Kleitz and J. Dupuy (eds.), D. Reidel Publishing Company, Boston, MA, 1976, p. 367. T. Takahashi, "Solid Electrolyte Fuel Cells (Theoretics and Experiments)," in Physics of Electrolytes, Vol. 2, J. Hladik (ed.), Academic Press, New York, 1972, p. 989.

Conference Proceedings and Extended Abstracts

.

.

,

o

.

.

Proceedings of the Workshop on High Temperature Solid Oxide Fuel Cells, May 5-6, 1977, Brookhaven National Laboratory, Upton, NY, H.S. Isaacs, S. Srinivasan, and I.L. Harry (eds.), Report No. BNL 50756, BrookhavenNational Laboratory, Upton, NY, 1978. Proceedings of the Conference on High Temperature Solid Oxide Electrolytes, August 16-17, 1983, Brookhaven National Laboratory, Upton, NY, F.J. Salzano (ed.), Report No. 51728, Brookhaven National Laboratory, Upton, NY, 1983. Proceedings of the International Symposium on Solid Oxide Fuel Cells, November 13-14, 1989, Nagoya, Japan, O. Yamamoto, M. Dokiya, and H. Tagawa (eds.), Science House, Tokyo, Japan, 1989. Proceedings of the First International Symposium on Solid Oxide Fuel Cells, October 16-18, 1989, Hollywood, FL, S.C. Singhal (ed.), Electrochemical Society, Pennington, NJ, 1989. Proceedings of the Second International Symposium on Solid Oxide Fuel Cells, July 2-5, 1991, Athens, Greece, F. Grosz, P. Zegers, S.C. Singhal, and O. Yamamoto (eds.), Commission of the European Communities, Luxembourg, 1991. Proceedings of the Third International Symposium on Solid Oxide Fuel Cells, May 16-21, 1993, Honolulu, HI, S.C. Singhal and H. Iwahara (eds.), Electrochemical Society, Pennington, NJ, 1993. Extended Abstracts, The First Symposium on Solid Oxide Fuel Cells in Japan, December 15-16, 1992, Tokyo, Japan, The Solid Oxide Fuel Cell Society of Japan,

Japan, 1992. Extended Abstracts, The Second Symposium on Solid Oxide Fuel Cells in Japan, December 15-16, 1993, Tokyo, Japan, The Solid Oxide Fuel Cell Society of Japan, Japan, 1993.

Appendix 353

.

10.

11.

13.

14.

15.

16.

Proceedings of the First European Solid Oxide Fuel Cell Forum, October 3-7, 1994, Lucerne, Switzerland, U. Bossel (ed.), European SOFC Forum Secretariat, Baden,

Switzerland, 1994. Proceedings of the 1st lEA Workshop on SOFC, Charmey, Switzerland, Swiss Federal Office of Energy, Berne, Switzerland, 1989. Proceedings of the 2nd lEA Workshop on SOFC, Hertenstein, Switzerland, Swiss Federal Office of Energy, Berne, Switzerland, 1990. Proceedings of the 3rd lEA Workshop on SOFC, Holmenkollen, Norway, Senter for Industriforskning, Oslo, Norway, 1991. Proceedings of the 4th lEA Workshop on SOFC, Lausanne, Switzerland, KFA, Jiilich, Germany, 1992. Proceedings of the 5th lEA Workshop on SOFC, Jiilich, Germany, KFA, Jiilich, Germany, 1993. Proceedings of the 6th lEA Workshop on SOFC, Rome, Italy, KFA, Jtilich, Germany, 1994.


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INDEX

A

ABO 3, 56, 101, 120, 168 AC impedance, see Complex impedance Activation energy of conduction

doped BaCeO3, 103 doped CeO2, 94 doped LaMnO3, 128 doped LaCrO3, 173 doped SrCeO3, 106 nickel/YSZ cermet, 157 stabilized ZrO2, 80

Alcohol, 3, 37, 199,340, 341 Alloy/Al203 cermet, 190 Alloy interconnect, see Metallic intercon-

nect Ambipolar diffusion theory, 64 Amorphous citrate process, see Pechini

process Anode. See a lso Nickel/YSZ cermet

CO oxidation, 208 hydrogen oxidation, 200 interfacial resistance, 204,205,207,

208 overpotential, 150, 154, 155,201,204 reforming, 210 requirement, 7, 147

Arrhenius equation, 82, 128 Arrhenius conductivity plot, 82, 93, 97,

105, 106, 128, 157, 173

B

BaCeO3 activation energy of conduction, 105 ionic conductivity, 102, 103, 105 oxygen-ion conduction, 102 protonic conduction, 103 structure, 103

Band gap energy, 51 Banded configuration, see Segmented-cell-

in-series design BaZrO3, 102, 107 Bell-and-spigot configuration, see

Segmented-cell-in-series design Bi203

Bi203-MO system, 97 Bi203-M203 system, 98 Bi2Oa-MO 2 system, 101 Bi203-MO3 system, 101 BizO3-M205 system, 101 Bi203-M6Oll system, 101 Bi4V2.xMxOll.~, 101 crystallographic polymorph, 96 electronic conductivity, 98 face-centered cubic phase, 97 ionic conductivity, 96, 97, 100 rhombohedral phase, 97 stability, 97

Binder, 242,272,276,277,279,288, 291

Bipolar plate, 6, 283. See also Intercon- nect

Bottoming cycle, 1, 2, 20, 25, 28,332, 333,334

Brayton/Rankine bottoming cycle, 334

C

CaFeO3, 139 Calendering, see Tape calendering Casting, see Slip casting, Tape casting CaZrO3, 107, 132, 180, 263 Carbon formation, 37, 39, 161, 211 Carnot efficiency, 16 Cathode. See a lso LaMnO3

interfacial resistance, 221,222,223 overpotential, 214,219, 223 oxygen reduction, 212 requirement, 7, 117

CeO 2 association energy, 54

358 I n d e x

CeOz buffer layer, 223 CeO2 anode material, 161,202, 209,

212, 263 defect, 49, 92 dopant, 93 electrolytic domain, 95 electronic conductivity, 93, 95, 96 ionic conductivity, 94 ionic domain, 95 oxygen deficiency, 92 oxygen-ion transference number, 93 oxygen reduction, 217

Ceramic fuel cell, s e e Solid oxide fuel cell

Chemical potential diagram La-Cr-O, 183 La-Cr-Zr-O, 180 La-Mn-Zr-O, 133

Chemical vapor deposition, see CVD Chromate, 158, 172, 180, 185, 186 CO oxidation

effect of CO2, 208 reaction mechanism, 208 reversible voltage, 17, 19 shift reaction, 200, 208

CO tolerance, 11, 28 COE

effect on CO oxidation, 208 effect on mechanical strength of

LaCrO3, 172 COE tolerance, 11 Coal gas, 3, 28, 36, 37, 39, 199,200,

209, 252, 266, 331,332 Coal gasification, 39, 331,333,334 Cobalt/YSZ cermet, 160 CoCr204, 12, 188 Cofiring, 7, 158, 180, 187,271,272,

282,287,310 Cogeneration, 1, 2, 20, 25, 28,331,332,

335 Complex impedance

equivalent circuit, 78 hydrogen oxidation, 202 LaMnO3, 217,220 measurement, 78 nickel/YSZ cermet, 202 oxygen reduction, 217,220 YSZ, 78

Compression molding, 274 Coprecipitation, 168

Current density electronic current density, 30, 61 exchange current density, 22,204,220 ionic current density, 30, 61 limiting current density, 23, 26

Current distribution, 321,324 CVD, 72, 149, 152, 167,248,261,289

D

Defect defect association, 54 defect cluster, 54, 126 defect complex, 82 defect domain, 52 Frenkel defect, 42, 44, 49 Koch-Wagner defect, 42 lattice defect, 41 line defect, 41 nonstoichiometric defect, 41 oxygen-ion vacancy defect, 43 plane defect, 41 point defect, 41 Schottky defect, 42, 44, 45, 46, 56 stoichiometric defect, 41

Desulfurization, 37, 38, 333,336 Diffusion coefficient

ambipolar diffusion coefficient, 64 measurement, 44, 66 Nernst-Einstein equation, 60 oxygen diffusion coefficient in LaCrO3,

188 oxygen diffusion coefficient in

LaMnO 3, 137 Diffusion welding, 263 Drip pyrolysis, 168

E

Efficiency current efficiency, 20, 21, 24 electrochemical efficiency, 20, 21 energy efficiency, 31, 33 Faradaic efficiency, s e e current

efficiency fuel-to-electricity conversion efficiency,

1, 15,28 heating value efficiency, 20, 24

Index 359

maximum energy efficiency, 32, 33 system efficiency, 20, 25. See also

System, system efficiency thermal efficiency, see thermodynamic

efficiency thermodynamic efficiency, 18, 19, 20,

21 voltage efficiency, 20, 21, 27

Electric utility, 332 Electrical analysis

analytical model for performance evaluation, 318

electrical analysis of banded configuration, 314

model of distributed constants, 316 performance potential, 318

Electrical connection flat-plate design, 284 monolithic design, 269 sealless tubular design, 237 segmented-cell-in-series design, 257

Electrochemical deposition, 72 Electrochemical vapor deposition, see

EVD Electrode. See also Anode and Cathode

capacitance, 221 conductance, 208,217,221,224,225 interfacial resistance, 204,205,207,

208,221,222,223 sintering, 153, 161

Electromotive force, 17, 32, 63 Electrolyte. See also YSZ

electronic conduction, 28, 29, 44, 50, 60

mixed-conducting electrolyte, 29, 32, 36

requirements, 2, 4, 7, 69 Electrolytic domain boundary, 52, 53, 95 Electron

electron concentration in MOE-AO system, 45, 46, 47, 48

electron concentrationin perovskite, 57, 58, 59

electron conductivity, 50, 52, 66 electron migration energy, 52 electron mobility, 50, 52

Electron hole, see Hole Electronic conduction

electronic conduction in electrolyte, 28, 29, 60

electronic conductivity, 29, 50, 52 electronic current density, 30, 61 electronic transference number, 62, 64,

65, 66 Electrophoretic deposition, 72, 291 Emission, 1,331,335,338 Enthalpy change, 18, 19, 21, 25 Entropy change, 18, 20, 323 Equivalent circuit

mixed conducting electrolyte, 29 YSZ, 78

EVD EVD process, 247, 251 growth mechanism, 249 kinetics, 249 LaCrO3, 167,245,250, 261 LaMnO3, 222 parabolic rate constant, 242 YSZ, 71, 152, 149, 246,247,250,

260, 261,289 Extrusion process, 242

F

Faraday's law, 24 Fick's law, 23 Film coating, 119 Firing shrinkage, 271,277 Flame spraying, 119, 149, 262 Flat-plate design

advantage and disadvantage, 287 anode fabrication, 288 cathode fabrication, 288 electrical connection, 284 electrolyte fabrication, 288,289 fabrication, 287 feature, 282 gas flow configuration, 284 gas manifolding, 284 interconnect fabrication, 289 performance, 293 property, 283 schematic diagram, 9, 282 tape casting, 288,291 technological status, 293

Fluorite conductivity of fluorite-type oxide, 50 doped MO2, 44 electrolytic domain of fluorite-type

360 I nde x

oxide, 52 fluorite-type oxide, 43, 49 fluorite structure, 43 ionic domain of fluorite-type oxide, 52

Freeze drying, 120 Fuel, see Hydrogen Fuel processing, 16, 20, 25, 38,338,340 Fuel utilization, 21, 24 Fuel cell

characteristics, 1 comparison, 11 definition, 1 regenerative fuel cell, 343 type, 10

G

Gas manifold flat-plate design, 284 monolithic design, 270 requirement, 234 sealless tubular design, 238 segmented-cell-in-series design, 258

Gd2(ZrxTil_x)207, see GZT Gibbs free energy change, 18, 21 Glycine/nitrate process, 120, 167, 184 GZT, 101

I-I

Heat generation, 20, 323 Heat transfer coefficient, 323 Heating value

heating value efficiency, 20 high heating value, 18, 20, 25 low heating value, 18, 25

HHV, see Heating value, high heating value

Hole hole concentration in MO2-AO system,

45, 46, 47, 48, 50 hole concentration in perovskite, 57,

58, 59 hole conductivity, 50, 52, 66 migration energy, 52 hole mobility, 50, 52

H2S, 209 Hydrocarbon, see Hydrocarbon reforming

Hydrocarbon reforming, 2, 3, 28, 37, 38,200, 210, 253,327, 332, 340

Hydrogen hydrogen concentration cell, 4 hydrogen oxidation, 15,200 hydrogen/oxygencell, 18, 19

Hydrogen oxidation reaction AC impedance spectra, 202 effect of electrode material, 201 effect of H20 content, 203,204 effect of microstructure, 204 effect of YSZ electrolyte, 201 electrochemical reactive site, 202,203 triple point, 201,203,204 reaction mechanism, 202

In203 based oxide, 139, 261 Interconnect. See a lso LaCrO3

connection in flat-plate design, 284 connection in monolithic design, 267 connection in sealless tubular design,

237 connection in segmented-cell-in-series

design, 257 electrical connection, 6 metal, see Metallic interconnect requirements, 7, 165

Inverter, 40 Ionic conduction

ionic conduction in electrolyte, 28, 29 ionic conductivity, 50, 52 ionic current density, 30, 61 ionic transference number, 29, 62, 64

Ionic domain boundary, 52, 53, 95

Jet vapor deposition, 72,291 Joule heat, 30

K

Kr6ger-Vink notation, 45, 78, 125, 176, 200

KTaO3, 102, 107

I n d e x 361

L

LaCoO 3 chemical interaction with YSZ, 139,

223 critical oxygen partial pressure, 138 electrical conductivity, 138 oxygen stoichiometry, 138 phase transformation, 138 stability, 138,262 thermal expansion, 139

LaCrO3 chemical interaction with glass sealant,

180 chemical interaction with nickel/YSZ,

158 chromate formation, 172, 180, 185 chromium deficiency, 185 coating, 190, 289 defect structure, 175 dimensional change, 172 dopant, 172, 186 effect of oxygen partial pressure on

conductivity, 176 electrical conductivity, 172 gas permeability, 188 hydroxide formation, 170 lanthanum stoichiometry, 169, 170 liquid phase sintering, 172, 180, 185,

186, 187 liquid phase migration, 180, 187,282 microwave sintering, 188 oxygen stoichiometry, 170, 177 oxygen vacancy formation, 179 phase transformation, 169 powder synthesis, 167 preparation, 166 property, 168 sinterability, 183 sintering aid, 172, 187 thermal expansion, 169, 181 volatility, 171, 183

LaCrO3-LaCoO3 solid solution, 173 LaCrO3-LaMnO3 solid solution, s e e

LaMnO3-LaCrO3 solid solution LaFeO3-LaCoO3 solid solution, 139, 223 LaMnO3

chemical interaction with YSZ, 87, 132,223

critical oxygen partial pressure, 126,

131 decomposition, 123, 126, 131 defect structure, 125, 130 dopant, 127 effect of oxygen partial pressure on

defect structure, 125, 130 effect of oxygen partial pressure on

conductivity, 130 electrical conductivity, 127 hydroxide formation, 123, 126 lanthanum stoichiometry, 123, 126,

129, 133, 135, 137 orthorhombic/rhombohedral phase

transformation, 122 oxygen diffusion, 137 oxygen stoichiometry, 122, 125, 135 phase diagram, 120 powder synthesis, 119 preparation, 118 property, 122 sinterability, 137 stability, 126 structure, 120 thermal expansion, 135

LaMnO3-LaCoO3 solid solution, 136 LaMnO3-LaCrO 3 solid solution, 135,136,

173, 180 La.EZraO7, 87, 132, 133, 139, 223,246 Laser

CO2 laser evaporation, 72,262, 289 laser spraying, 72 pulse laser deposition, 72

LHV, s e e Heating value, low heating value

Liquid mix process, s e e Pechini process

M

Manganese migration, 132, 180 Manifold, s e e Gas manifold Metal vacancy concentration in MOE-AO

system, 46 Metallic interconnect, 12, 189,287,289 Microwave sintering, 71, 188 Migration energy, 52 Mixed conduction

mixed conduction in electrolyte, 28, 29, 44, 50, 60

mixed-conductingelectrolyte, 29, 32, 36

362 Index

MOE-AO system conductivity, 50 electron concentration, 45, 46, 47, 48 hole concentration, 45, 46, 47, 48 metal vacancy concentration, 45, 46,

47, 48 oxygen-on vacancy concentration, 45,

46, 47, 48 MOE-B203 system, 48, 50, 53 Mobility

electron, 50, 51, 52 hole, 50, 51, 52 oxygen-ion vacancy, 50, 51, 52

Modeling, 307 Modularity, 1, 28, 331,335 Monolithic design

advantage and disadvantage, 270 anode fabrication, 272 cathode fabrication, 272 coflow, 268,269, 270, 275 crossflow, 268,270,275,281,322 electrical connection, 269 electrolyte fabrication, 272 fabrication, 272 feature, 268 gas manifolding, 270 interconnect fabrication, 272 performance, 280 power density, 269 property, 269 schematic diagram, 9, 268 tape calendering, 272, 278 tape casting, 272 technological status, 280

N

Natural gas, 28, 36, 37, 38,200, 209, 252, 332, 333,336

Nernst equation, 17, 323 Nernst potential, 17, 323,324 Nernst-Einstein equation, 60 Nickel-MgO/YSZ cermet, 152,207 Nickel/YSZ cermet

chemical interaction with LaCrO3, 158 dimensional change, 153 effect of microstructure on anode over-

potential, 204 electrical conductivity, 156

microstructure, 150 nickel sintering (coarsening), 152, 153,

207 preparation, 149 property, 149 reduction, 150, 158 thermal expansion, 149, 159

Nitrate pyrolysis, 184 Nusselt number, 323

O

Ohmic losses, 24, 25,199,233,321. See

also Polarization, ohmic polarization Overpotential, see Polarization Overvoltage, see Polarization Oxidant, see Oxygen Oxidation plating, 119 Oxygen

oxygen oxidation, 16 oxygen reduction, 16 oxygen concentration cell, 17, 63

Oxygen reduction reaction AC impedance spectra, 217,220 effect of chemical interaction, 223 effect of electrode morphology, 216,

221 electrochemical reactive site, 216 metal electrode, 214 oxide electrode, 216 rate-determining step, 213, 214, 217 reaction at metal electrode, 214 reaction at oxide electrode, 216 reaction mechanism, 214, 217 reaction mechanism at LaMnO3 elec-

trode, 219 three-phase boundary, see triple point triple point, 216,221,224

Oxygen-ion vacancy oxygen-ion vacancy concentration in

MO2-AO system, 45, 46, 47, 48, 50 oxygen-ion vacancy concentration in

perovskite, 57, 58, 59 formation enthalpy, 53 migration enthalpy, 52 mobility, 50, 52 oxygen-ion vacancy in LaMnO3, 125 thermodynamic data for oxygen-ion

vacancy formation in LaCrO3, 179

I n d e x 363

P

Patterson map, 53 Pechini process, 119, 167 Percolation theory, 156 Performance modeling, 308 Permeation technique, 64 Perovskite

defect structure, 56 effect of oxygen partial pressure on

defect structure, 57 electron concentration, 57, 58, 59 hole concentration, 57, 58, 59 oxygen-ion vacancy concentration, 57,

58, 59 perovskite-type oxide, 56, 101, 102,

120, 138, 168,217 perovskite structure, 56, 120, 168

Plasma enhanced metal organic CVD, 72 Plasma spraying

A1203, 261 hybrid plasma spraying, 72 LaCrO3, 167 LaMnO3, 119,262,288 nickel/YSZ, 149,262, 288 plasma spraying process, 264 YSZ, 71,262,288

Plasticizer, 272,288,291 Platinum electrode, 214 Polarization

activation polarization, 21, 22, 25 anode polarization, 199, 321 cathode polarization, 199, 321 charge transfer polarization, s e e

activation polarization concentrationpolarization, s e e diffusion

polarization diffusion polarization, 21, 23,233 ohmic polarization, 21, 23, 25, 199,

233, 321 reaction polarization, 21, 23 resistance polarization, s e e ohmic

polarization Polarized cell technique, 65 Polaron conduction, 128, 129, 173, 179 Potential

chemical potential, 60, 65 electrochemical potential, 60 Nernst potential, 17,. 323,324

Power

power conditioning, 16, 20, 25, 39 power generation, 25 power output, 26, 30 maximum power, 26

PrMnO 3, 139 Pressing, 167, 188,263,289 Protonic conduction

activation energy, 105, 106 proton conductivity, 103, 105, 106 protonic conductor, 102 proton-conductor SOFC, 4, 16

Pyrochlore oxide, 101 Pyrolysis of metallic soap slurry, 152

QR

Rankine bottoming cycle, 334 Reforming, s e e Hydrocarbon reforming Regenerative fuel cell, 343,346 Reversible voltage

CO/oxygen cell, 17, 19 definition, 17 effect of pressure, 19 effect of temperature, 19 hydrogen/oxygencell, 18, 19

Rf sputtering, 72,289 Rolling, s e e Tape calendering RuOa/YSZ, 140 Ruthenium/YSZ cermet, 161

Screen printing, 119, 149, 154, 288 Sealing, 180, 235,239,255,260,284,

285,287 Sealless tubular design

advantage and disadvantage, 239 anode fabrication, 247 cathode fabrication, 244 electrical connection, 237 electrolyte fabrication, 246 EVD, 247 fabrication, 240 feature, 235 gas manifolding, 238 interconnect fabrication, 245 performance, 252 property, 236

364 I nde x

schematic diagram, 9, 235 support tube fabrication, 242 technological status, 252

Seebeck coefficient, 179 Segmented-cell-in-series design

advantage and disadvantage, 259 anode fabrication, 260, 262,263 cathode fabrication, 261,262, 263 electrical connection, 257 electrolyte fabrication, 261,262,263 EVD, 260 fabrication, 260 feature, 255 gas manifolding, 258 interconnect fabrication, 260, 261,263 performance, 266 plasma spraying, 264 property, 257 schematic diagram, 9, 255 support tube fabrication, 260 technological status, 266

Shift reaction, 200,208, 211 Sintering aid, 85, 172, 187,242 Siting, 2, 28,331 Slip casting, 288 Slurry coating

LaMnO3, 118,244, 288 nickel, 149,152,205,260,288 YSZ, 72,288

SOFC, see Solid oxide fuel cell SOFC system, see System SOFC system efficiency, see System,

system efficiency Solar photovoltaic power, 343 Sol-gel process, 72, 120, 290 Solid oxide fuel cell

characteristics, 2, 28,331 comparison with other fuel cell types,

10 component, 6 definition, 3, 15 design, 7,233,234 historical background, 10 operating principle, 3, 15 oxygen-ion-conductor solid oxide fuel

cell, 4, 16 proton-conductor solid oxide fuel cell,

4,16 reaction in solid oxide fuel cell, 4 component requirement, 7

type, 4 Space application, 341 Spray pyrolysis, 72, 120, 168,288 SrCeO 3

activation energy for conduction, 106 dopant, 106 protonic conductivity, 106

SrZrO3, 107, 132, 133 Stabilized ZrO2. See a lso YSZ

blackening, 76,203,216 defect structure, 43, 44 dopant, 74, 80 effect of dopant concentration on con-

ductivity, 80 effect of oxygen partial pressure on

defect structure, 45 electrical conductivity, 80 oxygen-ion vacancy, 43, 44 thermal expansion, 89

Stack fiat-plate design, 282 monolithic design, 268 sealless tubular design, 235 segmented-cell-in-series design, 255 stack design configuration, 7,233,234 stack design requirement, 233

Steam, see Water Stress analysis

finite-element method, 308, 310, 313 simplified stress model, 308 strain energy release rate, 311 stress intensity factor, 312

Sulfur poisoning, 39,200,209, 212 System

cogeneration system, 336 integratedcoal gasification/SOFC-steam

turbine system, 333 propulsion system, 341 regenerative system, 343 system efficiency, 333,334, 335,336,

338, 342

T

Tafel equation, 22,204,215 Tape calendering

LaCrO3, 167,272 LaMnO3, 119, 272 NiO/YSZ, 149,272

Index 365

tape calendering process, 272,290 ZrO2, 71,272,290

Tape casting LaCrO3, 167,289 LaMnO 3, 119,288 NiO/YSZ, 149,288 tape casting process, 272,278,291 ZrO2, 71,288

Temperature distribution, 321,324 Tetragonal ZrO2, see TZP Thermal expansion

alloy, 190 LaCrO3, 169, 181 LaMnO3, 122, 135 Nickel/YSZ cermet, 149, 159 ZrOz, 75, 87

Thiele modulus, 249 ThO2, 54 Transfer coefficient, 22 Transference number

electronic transference number, 62, 64, 65, 66

ionic transference number, 29, 62, 64 oxygen-ion transference number, 29,

32, 35, 44 transference number measurement, 62

Transport processes, 60 Transportation application, 338 TZP, 91

UV

Vapor-phase electrolytic deposition, 72, 289

Vacuum evaporation, 72, 119,289 Voltage

equilibrium voltage, 21, 199 open-circuit voltage, 21, 199 reversible voltage, 17, 21, 199, 321.

See also Reversible voltage thermodynamic voltage, see reversible

voltage

W

Wagner oxidation process, 247,249,250 Wagner scaling equation, 65 Water

effect of water on hydrogen oxidation, 203,204

reaction of water in EVD process, 248 steam ratio, 211, 212 steam reforming, 37, 39,211 water in protonic conduction, 103,104,

106, 107 water gas shift reaction, 200, 208, 211

XY

YCrO3, 188,246 YFeO 3, 139 YMnO3, 139 YSZ

bend strength, 75, 90 blackening, 76, 203,216 chemical interaction with LaMnO3, 87,

132,223 complex impedance, 79 conductivity aging, 86, 91 cubic solid solution, 76 defect cluster, 55, 82 deposition method, 72 electrical conductivity, 75, 78 electronic conductivity, 83 fracture toughness, 75, 80, 312 grain-boundary conductivity, 79, 84 manufacturing process, 72 microwave sintering, 71 monoclinic solid solution, 76 phase diagram, 77 powder synthesis, 72 preparation, 71 property, 74 tetragonal solid solution, 76, 91 thermal expansion coefficient, 75, 87 toughening, 90

Y2(ZrxTil_x)207, s e e YZT YZT, 101

Z

ZrO 2. See also Stabilized ZrO2 and YSZ blackening, 76,203,216 cubic phase, 44, 74 dopant, 74, 80 monoclinic phase, 44, 74

366 Index

phase transformation, 44, 74 property, 74 tetragonal phase, 44, 74

ZrO2-CaO cubic solid solution, 74 defect cluster, 55 ionic transference number, 62 Koch-Wagner defect, 43, 44 monoclinic solid solution, 74 phase diagram, 75 tetragonal solid solution, 74

ZrO2-MgO cubic solid solution, 76 monoclinic solid solution, 76 tetragonal solid solution, 76

ZrO2-M203, 76 ZrO2-Y203, see YSZ ZrO2-Y203-TiO2, 161

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Science and Technology of Ceramic Fuel Cells