Application of Rare-Earth Doped Ceria and Natural Minerals ...1347923/FULLTEXT01.pdf · Även om...

Application of Rare-Earth Doped Ceria and

Natural Minerals for Solid Oxide Fuel Cells

Yanyan Liu

Doctoral Thesis

Stockholm 2019

Division of Heat and Power Technology

Department of Energy Technology

School of Industrial Engineering & Management

KTH Royal Institute of Technology, Sweden

TRITA-ITM-AVL 2019:24

ISBN: 978-91-7873-277-7

Akademisk avhandling

Som med tillstånd av Kungl Tekniska Högskolan framlägges till offentlig

granskning fär avläggande av teknologie doktorsexamen I fysik fredagen den

27 September 2019 kl. 10:00 I K1, Teknikringen 56, KTH, Stockholm.

© Yanyan Liu, August 2019

Tryck: Universitetsservice US-AB, Stockholm

To my husband

Doctoral thesis / Yanyan Liu

i

Abstract

Although solid oxide fuel cell (SOFC) technology exhibits considerable advantages

as compared to other energy conversion devices, e.g. high efficiency, low emission

and fuel flexibility, its high operating temperature leads to rapid component

degradation and has thus hampered commercialization. In recent years, intensive

research interests have been devoted to lowering the operating temperature from

the elevated temperature region (800-1,000 ) to intermediate or low-temperature

range (<800 ). To achieve this goal, material selection plays a dominant role,

involving improving the conductivity of existing electrolytes and developing new

exploitable materials. This dissertation is focused on enhancing the ionic

conductivity of rare-earth oxides (principally doped ceria) and exploring new

candidate materials (e.g. natural minerals) for low temperature (LT) SOFCs.

In this work, the scientific contributions can be divided into four aspects:

i) To develop desirable superionic conductors, Sm3+/Pr3+/Nd3+ triple-

doped ceria is designed to realize the desired doping for Sm3+ in bulk

and Pr3+/Nd3+ at surface domains via a two-step wet chemical co-

precipitation method. It exhibits high ionic conductivity, 0.125 S cm-1

at 600 . The SOFC device using this material as electrolyte displays

a high output power density of 710 mW cm-2 at 550 .

ii) To further clarify the individual effect of Pr3+ in the doped ceria, a

single-element (Pr3+) doped ceria is studied, exhibiting a mixed

electronic/ionic conduction property, capable of being employed as the

core component of electrolyte-layer free solid oxide fuel cells (EFFCs).

iii) To investigate various rare-earth doped-ceria materials in double- and

triple-element doping solutions for LT-SOFCs, Sm3+/Ca2+ co-doped

ceria and La3+/Pr3+/Nd3+ triple-doped ceria are synthesized and then

further incorporated with semiconductors, e.g. La0.6Sr0.4Co0.2Fe0.8O3-δ

(LSCF) or Ni0.8Co0.15Al0.05Li-oxide (NCAL), to serve as a

semiconducting-ionic conducting membrane in EFFCs.

iv) To exploit the feasibility of natural mineral cuprospinel (CuFe2O4) as

an alternative material for LT-SOFCs, three different types of fuel cell

devices are fabricated and tested. The device using CuFe2O4 as cathode

exhibits a maximum power density of 180 mW cm-2 with an open

circuit voltage of 1.07 V at 550 °C, while the device using a


ii

homogeneous mixture membrane of CuFe2O4, Li2O-ZnO-Sm0.2Ce0.8O2

(LZSDC), and LiNi0.8Co0.15Al0.05O2 (NCAL) demonstrates an

improved power output, 587 mW cm-2 under the same measurement

conditions.

Based on this work, a new triple-doping strategy is exploited to improve the ionic

conductivity of doped ceria materials by surface- and bulk-doping methodology.

Furthermore, the material developments of single-phase mixed electronic/ionic

conducting doped ceria and doped ceria/semiconductor composites are realized and

verify the feasibility of EFFC technology. Investigations on CuFe2O4 indicate the

utility of natural minerals in developing cost-effective materials for LT-SOFCs.

Keywords:

Low-temperature solid oxide fuel cells; Doped ceria; Material characterizations;

Electrochemical performances; Natural minerals.


iii

Sammanfattning

Även om fastoxid bränslecellers (SOFC) uppvisar signifikanta fördelar jämfört

med andra energiomvandlingstekniker, t.ex. hög verkningsgrad, låga emissioner

och bränsleflexibilitet, leder dess höga driftstemperatur till snabb

komponentdegradering, vilken har hindrat kommersialiseringen. Under de senaste

åren har intensiv forskning ägnats åt att sänka driftstemperaturerna från de höga

temperaturregionerna (800-1,000 °C) till mellanliggande eller låga

temperaturintervaller (<800 ). För att uppnå detta mål spelar materialvalet en

dominerande roll, vilket bland annat innebär att man förbättrar ledningsförmågan

hos befintliga elektrolyter och utvecklar nya material. Denna avhandling fokuserar

på att förbättra den jonledande förmågan hos oxider av sällsynta jordartsmetaller,

huvudsakligen dopad ceriumoxid, samt forskning på nya kandidatmaterial, t.ex.

naturliga mineraler.

I det här arbetet kan det vetenskapliga bidraget delas in i fyra aspekter:

i) Att utveckla en trippel-dopingmetodik för att syntetisera önskvärda

superjoniska ledningsförmågor i Sm3+/Pr3+/Nd3+ dopad ceriumoxid.

Detta material konstruerades med hjälp av en tvåstegs våtkemisk

samutfällningsmetod för att åstadkomma en önskad dopning för Sm3+

i bulk och Pr3+/Nd3+ vid ytdomäner. Materialet uppvisar en hög jonisk

ledningsförmåga, 0.125 S cm-1 vid 600 . En SOFC-enhet som

använder denna trippeldopade ceriumoxid som elektrolyt har uppvisat

en hög effekttäthet på 710 mW cm-2 vid 550 ;

ii) För att ytterligare klargöra den individuella effekten av Pr3+ i det

dopade ceriummaterialet studerades enfas Pr-dopade ceria, vilken

uppvisade en blandad elektronisk/jonisk ledningsegenskap som skulle

användas som kärnkomponent i avancerad elektrolytskiktsfri fastoxid

bränsleceller (EFFC).

iii) Att undersöka olika sällsynta jordartade dopade ceriummaterial i

lösningar med dubbel- och trippelelement (Sm3+/Ca2+ och

La3+/Pr3+/Nd3+) applicerade för SOFC-teknik med låg temperatur. De

dubbel- och tripeldopade ceriummaterialen var sammansatta med

halvledare, dvs Ni0.8Co0.15Al0.05Li-oxid (NCAL) och

La0.6Sr0.4Co0.2Fe0.8O3-δ (LSCF) för att fungera som en halvledande

jonisk ledande membran i EFFCs.


iv

iv) Att utnyttja naturligt kopparspinell (CuFe2O4) som ett alternativt

material för SOFC. För första gången tillverkades tre olika typer av

anordningar för att undersöka den optimala appliceringen av CuFe2O4

i SOFC. Enheten med CuFe2O4 som katodkatalysator uppvisade en

maximal effekttäthet av 180 mW cm-2 med en öppen kretsspänning

1.07 V vid 550 . En effekttäthet på 587 mW cm-2 emellertid

uppnådes från anordningen bestående av ett homogentblandat

membran med CuFe2O4, Li2O-ZnO-Sm0.2Ce0.8O2 och

LiNi0.8Co0.15Al0.05O2.

Baserat på detta arbete utnyttjades en ny strategi för att förbättra

jonledningsförmågan hos dopade ceriummaterial genom yt- och

bulkdopningsmetodik. Vidare verifierades utvecklingen av EFFC-teknikens

tillförlitlighet av enfasad, blandad elektronisk/jonledande dopade ceriumoxid samt

jonledande multidopade ceria-och halvledarkompositer. Dessa resultat visar att

naturliga mineraler kan spela en viktig roll för att utveckla kostnadseffektiva

material för bränsleceller.

Nyckelord:

Lågtemperatur fastoxidbränsleceller; Dopade ceriumoxid;

Materialkarakteriseringar; Elektrokemiska prestanda; Naturliga mineraler.


v

Acknowledgments

Doctoral studying experience is a life-enriching and life-changing journey. Over

the past four-year traveling through the ocean of knowledge, I finally got chances

to appreciate the beauty of discovering and understanding of the unknown, even if

lost and struggling in the mist often happened to me at the beginning. It would never

happen without the support from people around me. Here, I would like to express

my sincerest gratitude to all who gave me support and help during my Ph.D. study.

First, I would like to express my sincere gratitude to my principal supervisor

Professor Andrew Martin for your invaluable support, guidance, and supervision

throughout my doctoral study. You are always generous to your students with your

advice, time and encouragement whenever I encountered problems. It is my honor

to be your student and I really enjoy learning period under your supervision. And I

also would like to acknowledge the financial support from the Chinese Scholarship

Council (CSC) and Division of Heat and Power Technology (HPT), KTH Royal

Institute of Technology for my Ph.D. study and life in Stockholm, Sweden.

I would like to thank my former principal supervisor Professor Bin Zhu for giving

me the opportunity to work on this fascinating project. His inspiring speeches and

significant guidance brought me to get to know the frontier of the field from the

beginning to the end of my Ph.D. study. I also would like to thank my co-supervisor

Dr. Wujun Wang for his heart-warming encouragement and constructive advice for

my research work, especially the guidance to my dissertation writing. My gratitude

goes to Prof. Lyuba Belova, another co-supervisor, for her guidance in materials

preparation and unselfish advice for my doctoral study. In addition, special thanks

are given to my master supervisor Associate Professor Yongfu Tang, who brought

me to this fascinating road and taught me how to become a qualified researcher

when I started my research life.

Along with my study, I have received much help from my colleagues in Fuel Cell

group: Chen Xia, Muhammad Afzal, and Baoyuan Wang for discussing our

research field and endless friendship. I really enjoyed the friendly atmosphere in

our office. I am also grateful to the HPT laboratory staffs for maintaining an

efficient public lab with an open and friendly atmosphere. Furthermore, heartfelt

gratitude to Professor Muhammet Toprak for serving as the opponent at my final

seminar and advance reviewer of my dissertation. I really appreciate your valuable

time and insightful suggestion to improve my thesis. Also, many thanks to Justin

Chiu, Mahrokh Samavati and Anders Malmquist for your crucial feedback and


vi

suggestion for improving this dissertation, especially to Anders for polishing the

Swedish abstract.

Special thanks go to Professor Yongdan Li from Aalto University for your

enthusiastic advice on the direction of my future research. I really appreciate your

invaluable discussions and guidance about the solid oxide fuel cells and other

energy technologies. Besides, I would like to thank Professor Peter Lund also from

Aalto University for your unselfish help to my research work. I still remember your

warm smile and encouraging words when I gave my first oral presentation at an

international academic conference.

I also want to thank all the supportive group members, former and present. Here,

to mention few: Liangdong Fan from Shenzhen University, Yan Wu from China

University of Geosciences, Yixiao Cai from Uppsala University, Wei Zhang and

Lili Wei from Hubei University. I would also like to thank the other colleagues at

the Department of Energy Technology: Yang Gao, Simeng Tian, Tianhao Xu,

Wenshuai Miao, Yang Song, Saman Nimali Gunasekara, Göran Arntyr, Leif

Pettersson, Jens Fridh, Anneli Ylitalo-Qvarfordt, Tianrui Sun, Chang Su,

Chamindie Senaratne, Jeevan Jayasuriya, Tony Chapman and other friends in the

department for creating a friendly working environment. I also want to thank the

rest of my friends in Sweden: Xinlei Yi, Xiaolin Jiang, Ju Shu, Sheng Jiang, Filippa

Modigh, Cynthia Ho Sin Tian, Chao Zheng, Jingru Wu, Chunlei Wang, Ju Li,

Zongwei Ji, Zongfang Wu, and Yu Yang for the nice time and fun in and outside

KTH. My Ph.D. life would not have been so enjoyable without you all.

Sincere thanks to my beloved husband Shaozheng Ji for your unconditional love

and support for my life, as well as honest criticism, constructive suggestion and

inspiring discussions about my research. Thank you for surviving together with me

through the doctoral years. Your enthusiasm and kindness light up my heart

especially when I felt depressed in the difficult time. I have become a better person

and researcher because of you. Finally, I want to thank my beloved family: my

parents, my brother, my sister and sister-in-law, my aunt and uncle for their love

and support all along, as well as my little nephew for giving me fun during video-

time.

Yanyan Liu

August 2019

Stockholm, Sweden


vii

List of Appended Publications

Paper I

Liu, Y., Fan, L., Cai, Y., Zhang, W., Wang, B. and Zhu, B. Superionic conductivity

of Sm3+, Pr3+, and Nd3+ triple-doped ceria through bulk and surface two-step doping

approach. ACS applied materials & interfaces. 9.28 (2017): 23614-23623.

Paper II

Liu, Y., Wei, L., Dong, W. and Zhu, B. A single-phase mixed-conductive Pr-Doped

CeO2 membrane for advanced fuel-to-electricity technology. (Manuscript under

preparation)

Paper III

Liu, Y., Meng, Y., Zhang, W., Wang, B., Afzal, M., Xia, C. and Zhu, B. Industrial

grade rare-earth triple-doped ceria applied for advanced low-temperature

electrolyte layer-free fuel cells. International Journal of Hydrogen Energy. 42.34

(2017): 22273-22279.

Paper IV

Liu, Y., Wu, Y., Zhang, W., Zhang, J., Wang, B., Xia, C., Afzal, M., Li, J., Singh,

M. and Zhu, B. Natural CuFe2O4 mineral for solid oxide fuel cells. International

Journal of Hydrogen Energy. 42.27 (2017): 17514-17521.

Paper V

Wang, B., Wang, Y., Fan, L., Cai, Y., Xia, C., Liu, Y., Raza R., Aken P., Wang H.

and Zhu B. Preparation and characterization of Sm and Ca co-doped ceria-

La0.6Sr0.4Co0.2Fe0.8O3-δ semiconductor-ionic composites for electrolyte-layer-free

fuel cells. Journal of Materials Chemistry A, 4.40 (2016): 15426-15436.

Paper VI

Wang, B., Cai, Y., Xia, C., Liu, Y., Muhammad, A., Wang, H. and Zhu, B.

CoFeZrAl-oxide based composite for advanced solid oxide fuel cells.

Electrochemistry Communications, 73 (2016): 15-19.

Contributions of the author to the publications

In the listed papers I-IV, Y. Liu is the first author and her contributions include:

proposing the idea, performing the literature survey, designing the experiments,


viii

synthesizing the experimental samples, conducting material characterization work,

electrical property and fuel cell performance measurements, processing and

analyzing all the experimental data, as well writing and revising the manuscripts.

The other authors joined partial characterization measurements and acted as the

reviewers of the manuscripts. In paper V and VI, Y. Liu was involved in partial

contributions including synthesizing the experimental materials, performing the

electrical properties measurements, and conducting the fuel cell performance

measurements.


ix

Table of Contents

Abstract .................................................................................................................. ii

Sammanfattning ................................................................................................... iii

Acknowledgements ................................................................................................ v

List of Appended Publications ........................................................................... vii

List of Figures ...................................................................................................... xii

List of Tables ....................................................................................................... xv

Abbreviations and Nomenclature ..................................................................... xvi

Chapter 1 Introduction ......................................................................................... 1

1.1 Challenges of solid oxide fuel cell technology ............................................. 3

1.2 Materials selection for solid oxide fuel cell technology ............................... 5

1.3 Motivations and objectives ........................................................................... 7

1.4 Structure of this thesis ................................................................................... 8

Chapter 2 Solid oxide fuel cell principles and literature review ....................... 9

2.1 Electrochemistry and thermodynamics ......................................................... 9

2.1.1 Open circuit voltage ............................................................................. 10

2.1.2 Efficiency ............................................................................................. 12

2.2 Materials development ................................................................................ 13

2.2.1 Anode materials development .............................................................. 14

2.2.2 Cathode materials development ........................................................... 14

2.2.3 Electrolyte materials development ....................................................... 16

2.3 Nanocomposites for EFFC technology ....................................................... 18

Chapter 3 Experimental methods and techniques ........................................... 21

3.1 Raw materials .............................................................................................. 21

3.2 Sample preparation ..................................................................................... 21

3.2.1 SDC preparation by co-precipitation approach .................................... 21


x

3.2.2 Pr-doped ceria preparation by hydrothermal method .......................... 22

3.2.3 Sm3+, Pr3+ and Nd3+ triple-doped ceria by two-step doping method .... 22

3.2.4 Double- and triple-doped ceria preparation ......................................... 22

3.2.5 Doped ceria/semiconductor composite preparation ............................. 23

3.2.6 LZSDC nanocomposite preparation .................................................... 23

3.2.7 Natural minutral treatment ................................................................... 23

3.3 Fuel cell device fabrication ......................................................................... 23

3.3.1 Conventional SOFC device ................................................................. 24

3.3.2 EFFC device ........................................................................................ 24

3.4 Material characterizations ........................................................................... 25

3.4.1 X-ray diffraction (XRD) analysis ........................................................ 25

3.4.2 Micro-structure analysis ...................................................................... 25

3.4.3 X-ray photoelectron spectroscopy (XPS) analysis .............................. 26

3.4.4 Raman spectra analysis ........................................................................ 26

3.5 Electrical properties .................................................................................... 26

3.5.1 Electrochemical impedance spectroscopy (EIS) analysis .................... 26

3.5.2 Direct-current (DC) polarization analysis ........................................... 26

3.6 Fuel cell performance measurements ......................................................... 27

Chapter 4 Superionic conductivity of triple-doped ceria for SOFC .............. 29

4.1 Material characterizations ........................................................................... 29

4.1.1 Phase and microstructure properties .................................................... 29

4.1.2 Ultraviolet-visible (UV-vis) Raman scattering analysis ...................... 30

4.1.3 XPS analysis ........................................................................................ 33

4.2 Electrical properties .................................................................................... 35

4.3 Fuel cell performance ................................................................................. 37

4.4 Chapter summary ........................................................................................ 40

Chapter 5 Multi-component rare-earth doped ceria for EFFC ..................... 43

5.1 Single-phase Pr-doped ceria for EFFC ....................................................... 43

5.1.1 Structural and morphology characteristics .......................................... 43


xi

5.1.2 XPS analysis ........................................................................................ 44

5.1.3 Raman scattering analysis .................................................................... 47

5.1.4 Electrical conductivity ......................................................................... 48

5.1.5 Fuel cell performance analysis ............................................................. 50

5.2 Rare-earth triple-doped ceria for EFFC ...................................................... 52

5.3 Double-doped ceria for EFFC ..................................................................... 57

5.4 Fuel cell performance comparsion .............................................................. 58

5.5 Chapter summary ........................................................................................ 59

Chapter 6 Natural minerals for EFFC .............................................................. 61

6.1 Material characterization analysis ............................................................... 61

6.2 Electrochemical performance analysis ........................................................ 63

6.3 Chapter summary ........................................................................................ 67

Chapter 7 Conclusions and future outlook ....................................................... 69

References ............................................................................................................ 73


xii

List of Figures

Figure 1.1 Summary of fuel cell families [16]. ...................................................... 2

Figure 1.2 Ionic conductivity of solid electrolytes for SOFC as a function of

temperature [20]. .................................................................................................... 4

Figure 1.3 Schematic illustrations for a) conventional anode/electrolyte/cathode

three-component SOFC; and b) EFFCs. ................................................................. 5

Figure 2.1 Schematic of SOFC working principle. .............................................. 10

Figure 2.2 Ideal and actual potential vs current response in a fuel cell [11]. ....... 12

Figure 2.3 Schematic diagrams of SOFC stack, individual SOFC device and

functional SOFC components [11]. ...................................................................... 14

Figure 2.4 Crystal structures of the three-type cathodes: a) perovskite; b) layered

perovskite; and c) Ruddlesden-Popper oxides [77]. ............................................. 15

Figure 2.5 Schematic image for fluorite-structured oxide [15]. ........................... 16

Figure 3.1 Schematic illustration of the single fuel cell and measurement setup. 24

Figure 3.2 Schematic diagram of fuel cell performance testing setup. ................. 28

Figure 4.1 a) X-ray diffraction patterns of SDC and PNSDC; b) SEM and TEM

micrographs for PNSDC powder. ......................................................................... 30

Figure 4.2 SEM image of PNSDC and corresponding EDS element mapping. ... 31

Figure 4.3 a) UV-vis diffuse reflectance spectroscopy of SDC and PNSDC; UV-

visible Raman spectra of SDC and PNSDC with an excitation laser wavelength b)

λex=532 nm and c) λex=325 nm; d) Schematic description for Raman scattering by

varying the ultraviolet (UV, λex=325 nm) and visible (λex=532 nm) excitation laser

lines to collect the surface and bulk information....................................................32

Figure 4.4 a) XPS survey spectra of SDC and PNSDC samples; b) XPS O 1s

Spectra of SDC and PNSDC samples; c) XPS Ce 3d Spectra of SDC and PNSDC

samples; d) XPS Pr 3d Spectra of PNSDC sample............................................... 34

Figure 4.5 Electrochemical impedance spectra of SDC and PNSDC measured in air

at a) 400 ; b) 450 ; c) 500 ; d) 550 ; e) 600 and f) the empirical

equivalent circuit to analysis the ionic conductivity of SDC and PNSDC materials

in air. ..................................................................................................................... 36

Figure 4.6 a) Temperature dependence of total conductivity and Arrhenius plots for

SDC and PNSDC. The inset represents area specific resistance (ASR) versus

temperature of SDC and PNSDC; b) current (I) versus time (t) curves obtained at a

temperature range of 450 and 600 at an interval of 50 to determine the

electronic conductivity of PNSDC sample. .......................................................... 37

Figure 4.7 a) Typical I-V and I-P curves of fuel cells based on SDC and PNSDC

electrolyte, respectively, the mixed NCAL and electrolyte materials were used as

the electrode material. The measurements were performed at 550 with H2 and


xiii

air as the fuel and the oxidant, respectively; b) Durability test based on Cell II: Ni-

foam/PNSDC+NCAL (anode)|PNSDC (electrolyte)| PNSDC+NCAL

(cathode)/Ni- foam at a constant discharge current density 80 mA cm-2 operated at

530 .................................................................................................................... 38

Figure 4.8 Electrochemical impedance spectra of the symmetric cell using SDC and

PNSDC samples as electrolyte measured under a) air and b) H2/air conditions at

550 .................................................................................................................... 38

Figure 4.9 Schematic diagrams representing a) the overall O2‐ transfer with the

assistance of reduction conversation of Pr3+/Pr4+ ions in the fuel cell; b) the cathode

reaction process and c) brief routes for the O2‐ surface and bulk transfer and

conduction process in PNSDC electrolyte. ........................................................... 39

Figure 5.1 X-ray diffraction patterns of as-synthesized Pr-free CeO2 and Pr-CeO2

samples. ................................................................................................................. 44

Figure 5.2 Representative SEM micrographs for as-synthesized a) CeO2 and c) Pr-

CeO2 samples; TEM images of b) CeO2 and d) Pr-CeO2 samples suffering

hydrothermal process. ........................................................................................... 45

Figure 5.3 XPS spectroscopy: a) survey spectra of CeO2 and Pr-CeO2 samples; b)

Pr 3d spectra for Pr-CeO2; c) O 1s spectra of CeO2; d) O 1s spectra of Pr-CeO2; e)

Ce 3d spectra of CeO2; and f) Ce 3d spectra of Pr-CeO2. ..................................... 46

Figure 5.4 Raman Spectroscopy of CeO2 and Pr-doped CeO2 samples using an

excitation wavelength of 325 nm. ......................................................................... 47

Figure 5.5 a) Nyquist plots for Pr-CeO2 sample measured at 600 under air

atmosphere; b) Arrhenius plot and corresponding temperature dependence of total

conductivity for Pr-CeO2 sample. ......................................................................... 48

Figure 5.6 I-V-P characteristics of fuel cell devices based on a) the CeO2 and Pr-

CeO2 membrane separately measured at the temperature of 550 ; b) Pr-CeO2

operated at various temperatures of 500, 525, 550 and 600 , respectively. ...... 50

Figure 5.7 Schematic diagrams representing a brief route for the O2- mobility in

fuel cell configuration and detailed O2- transfer with the assistance of redox

Pr3+/Pr4+ couple on the surface and bulk of Pr-CeO2 electrolyte in the vicinity of

oxygen applied side. .............................................................................................. 51

Figure 5.8 SEM micrographs for a) raw material and b) as-prepared LCPN; c) TEM

microstructure and corresponding selected area electron diffraction pattern (SAED,

inset image) for LCPN sample. ............................................................................. 52

Figure 5.9 Typical SEM micrographs of a) raw material and c) LCPN; b) and d)

identified the EDS data with the representative counts of the vertical axis obtained

at the yellow rectangular areas. ............................................................................. 53

Figure 5.10 XRD patterns of LCPN and NCAL samples. .................................... 53

Figure 5.11 EIS patterns of a) the 7LCPN-3NCAL composite at different

temperatures and b) the different compositions of LCPN and NCAL (7LCPN-


xiv

3NCAL, 6LCPN-4NCAL, and 3LCPN-7NCAL) under air/H2 condition at 550 .

All the measurements were performed under air/H2 condition under open circuit

conditions. ............................................................................................................ 55

Figure 5.12 a) The I-V (current density and voltage) and I-P (current density and

peak power density) characteristics of various compositions of LCPN-NCAL

feeding with hydrogen and air at 550 ; b) OCV and maximum power density

values (Pmax) for fuel cells based on LCPN-NCAL in various compositions. ...... 56

Figure 5.13 Schematic illustration (left) and SEM micrograph (right) of EFFC

device based on the 7LCPN-3NCAL composite after the fuel cell performance

measurements. ...................................................................................................... 56

Figure 5.14 a) I-V and I-P characteristics and b) EIS of EFFC devices using the

SCDC/LSCF composites in various weight ratios; c) I-V and I-P performance for

the conventional SOFC device and the device using pure SCDC membrane. ..... 57

Figure 6.1 XRD patterns for a) raw material and b) as-treated mineral samples

sintered at 600, 700 and 900 , respectively.. ..................................................... 60

Figure 6.2 SEM images of a) raw mineral, b) the sample sintered at , and c) EDS

image of raw mineral. ........................................................................................... 62

Figure 6.3 Nyquist plots of a) the samples sintered at 600, 700 and 900 measured

in air atmosphere at 550 ; b) the sample sintered at 900 measured in H2/air at

550 . .................................................................................................................. 64

Figure 6.4 Mechanism for the ionic conductivity of the LZSDC-CuFe2O4H

composite. ............................................................................................................. 64

Figure 6.5 Schematic illustrations for a) type I, b) type II and c) type III. ........... 65

Figure 6.6 a) I-V-P curves for three types of devices measured at 550 ; b)

Durability test of Type III at a constant discharge current of 50 mA cm-2. .......... 66

Figure 6.7 Nyquist plots for a) Type I, b) Type II, and c) Type III operated at fuel

cell condition at 550 . ....................................................................................... 67


xv

List of Tables

Table 5.1 Fuel cell performance comparison based on three types of doped ceria

materials. ............................................................................................................... 58

Table 6.1 The main chemical composition of raw chalcopyrite. .......................... 63

Table 6.2 The fitting results of impedance spectra for three types of devices

measured in H2/air condition at 550 . The unit for resistance is Ω cm2. ........... 66


xvi

Abbreviations and Nomenclature

AFC Alkaline fuel cell

ASR Area specific resistance

CTE Coefficient of thermal expansion

DC Direct current

DMFC Direct methanol fuel cell

e- Electron

Ea Activation energy

EDS Energy dispersive spectrometry

EFFC Electrolyte-layer free fuel cell

EIS Electrochemical impedance spectroscopy

EMF Electromotive force

F Faraday constant

FC Fuel cell

GDC Gadolinium doped ceria

H+ Proton

HOR Hydrogen oxidation reaction

I Current

i Current density

IT Intermediate temperature

k Boltzmann constant


xvii

L Thickness, or inductance

LCPN La/Pr/Nd doped ceria

LSCF LaxSr1-xCoyFe1-yO3, La0.6Sr0.4Co0.2Fe0.8O3-δ

LT Low temperature

LZSDC Li/Zn doped SDC

MCFC Molten carbonate fuel cell

NCAL Ni0.8Co0.15Al0.05LiO2-δ

NEFC Non-electrolyte separator fuel cell

NSDC SDC composited with Na2CO3

OO Oxygen ion in an electrolyte

OCV Open-circuit voltage

ORR Oxygen reduction reaction

P Power density

PAFC Phosphoric acid fuel cell

PEMFC Proton exchange membrane fuel cell

Pi Partial pressure for hydrogen, oxygen and steam

PNSDC Sm3+, Pr3+ and Nd3+-doped ceria

R Resistance

Ro Ohmic resistance

RE Rare-earth

Rp Electrode polarization resistance

S Active area


xviii

SCDC Ca0.04Sm0.16Ce0.8O2-δ

SDC Samarium doped ceria

SEM Scanning electron microscopy

SFM Sr2Fe4/3Mo2/3O6

XPS X-ray photon spectrum

STO Strontium titanate

SIMFC Semiconductor-ion membrane fuel cell

SOFC Solid oxide fuel cell

T Temperature

TEM Transmission electron microscopy

••

OV Oxygen ion vacancy

XRD X-ray diffraction

YSZ Yttria-stabilized zirconia

Z′ Real impedance

Z″ Imaginary impedance

η Overpotential

σ Conductivity

ϕ Phase angle

ω Angular frequency


1

Chapter 1

Introduction

With continuously increasing energy demand and related global warming problems,

the development of clean and high-efficient energy utilization technologies, as

alternates to fossil fuel resources, is an urgent issue [1]. Renewable energy

resources, e.g. solar energy, wind power, hydropower, bioenergy, and wave power,

are vital towards achieving a CO2-neutral energy system and have attracted

considerable attention in recent decades. The development and implementation of

renewables-based technologies have been hindered by a number of issues,

including the intermittency of renewable energy supplies in relation to demand [2-

3]. Hydrogen as an energy carrier has the ability to bridge this gap, provided that

its supply and distribution can be achieved in a cost-effective manner [2,4,5]. A

fuel cell is an integral part of hydrogen-based technologies, enabling the supply of

electricity and potentially heat efficiently from the stored chemical energy of

hydrogen or related fuels [6,7].

Fuel cell technology has a long history since the first fuel cell device (so-called ‘gas

battery’) was demonstrated by Sir William Grove in 1839 [8]. As an

electrochemical energy conversion technology, the fuel cell is not limited by the

Carnot cycle and is capable of supplying electricity with high efficiency in a very

wide range of capacities [9,10]. However, the practical applications of fuel cells

still lag far behind other technologies, for instance, gas turbines or internal

combustion engines. Currently, fuel cell technology is making an appearance in

electrical power markets and for transportation applications [11–15].

A fuel cell device is composed of three major components: electrolyte sandwiched

between two porous electrodes (cathode and anode). In a typical oxygen ion-

conducting fuel cell, the oxygen is reduced primarily at the cathode, and the

resulting oxygen ions (O2-) are transferred through the electrolyte layer approaching


2

to the anode. Subsequently, the O2- ions react with the fuel (e.g. hydrogen or

hydrocarbon) at the anode, with the supply of electrons to an external circuit. A fuel

cell can be classified in different ways depending upon its operating conditions [1].

Based on the types of electrolyte, they are classified as follows (also see Figure

1.1):

i) Proton exchange membrane fuel cells (PEMFCs);

ii) Phosphoric acid fuel cells (PAFCs);

iii) Alkaline fuel cells (AFCs);

iv) Molten carbonate fuel cells (MCFCs);

v) Solid oxide fuel cells (SOFCs).

Figure 1.1 Summary of fuel cell families [16].

Among these fuel cell types, PEMFC, PAFC, and AFC stacks require hydrogen

with relatively high purity to supply the anode, while the elevated-temperature


3

MCFCs and SOFCs can be fed with both hydrogen and carbon monoxide [16].

Comparing to the other fuel cell types, SOFCs have at least four advantages: i) high

current and power density; ii) high fuel utilization factor; iii) available non-noble

metal catalysts; and iv) low cathode/anode polarization losses [17]. The efficiency

of SOFC converting chemical energy to electricity can reach up to 45%~60% with

potential total efficiency of >85% in combined heat and power applications

[11,18,19]. Furthermore, the fuel options for SOFCs are relatively broad, allowing

them to be employed with existing hydrocarbon fuel infrastructure [11,16].

Therefore, SOFC technology is considered to be one of the most promising

candidates currently under research and development for fuel cells.

1.1 Challenges of solid oxide fuel cell technology

The key technical issue for the development of SOFC technology is its high

operating temperature, which can cause rapid system degradation, operational

complexity and high cost [11]. In principle, a SOFC device can be designed for

operating within a wide temperature range (500-1000 ). However, the actual

operating temperature of SOFCs is limited by the ion-conducting capability of the

solid electrolytes at low temperature (see Figure 1.2). A highly ionic conducting

electrolyte is indispensable towards achieving satisfactory fuel cell performance.

Most electrolyte materials, typically (ZrO2)0.9(Y2O3)0.1, require high operating

temperature to achieve desired ionic conductivity (e.g. >0.1 S cm-1) [20]. For

instance, the yttria-stabilized zirconia (YSZ) electrolytes can exhibit good ion-

conducting properties only under high temperature of above 950 [16,21]. Over

the past several decades, considerable efforts have been paid to lowering the

operating temperature of SOFCs to an intermediate temperature range (650-800 )

and even down to a low-temperature range of <650 [11,15,16]. The distinct

benefits of intermediate- or low-temperature SOFCs are the reduced system costs

due to mitigated performance degradation and wider material choices for the

interconnects and seals [22,23]. In addition, the theoretical efficiency of SOFCs is

expected to be enhanced by reducing the operating temperature [11]. Towards this

end, Goodenough [15,24,25] proposed to design oxide-ion conductors, e.g.

La0.8Sr0.2Ga0.8Mg0.2O3 (shown in Figure 1.2), allowing rapid ion transfer at a low

enough temperature level [15,24,25]. Extensive attempts have been devoted to

exploring highly conductive ionic conductors and developing new SOFC

configurations.


4

Figure 1.2 Ionic conductivity of solid electrolytes for SOFC as a function of

temperature [20].

To achieve high ion conductivity, the key issue is to decrease polarization losses,

associated with electrode reaction kinetics and electrolyte conduction at low

temperature (e.g. [16,21,26]). A promising approach in minimizing polarization

losses of LT-SOFCs is to reduce the thickness of electrolytes, transitioning from

electrolyte-supported fuel cells to electrode-supported fuel cells [26–28]. For

example, Barriocanal et al. [12] reported a tremendous ionic conductivity

enhancement (eight orders of magnitude) in YSZ and strontium titanate (STO)

epitaxial heterostructures using ultrathin films of 1~62 nm, showing a potential

application of SOFC materials at slightly above room temperature. Thus, the


5

attainably high ion-conduction and extremely small polarization losses for ultrathin

electrolytes demonstrate a possibility for the SOFCs at low operating temperature.

Designing new fuel cell structures with desired performance is another route to

realize LT-SOFC technology. Recently, a newly developing fuel cell technology,

namely electrolyte-layer free fuel cells (EFFCs), has been given attention. The

schematic illustration of the EFFC device is presented in Figure 1.3. This

technology, with an initial name non-electrolyte separator fuel cell (NEFC), was

first presented in 2011 [30–32]. Since then, many efforts have been undertaken to

understand the underlying working principles and for materials selection [33–41].

Unlike the conventional anode/electrolyte/cathode SOFC configuration, this fuel

cell device is constructed with a homogeneous two-phase nanocomposite, i.e.

semiconductor and ionic conductor, to replace the three components in

conventional devices. Compared to traditional SOFCs, several significant

advantages have been shown from EFFCs: 1) simplified fuel cell structure, leading

to potentially streamlined manufacturing processes; 2) good thermal and chemical

compatibility; 3) and improved fuel cell performance due to the reduction of

electrode polarization losses [31,32,38,39]. Therefore, EFFC energy technology

could provide a possible path to speed up fuel cell commercialization.

Figure 1.3 Schematic illustrations for a) conventional anode/electrolyte/cathode

three-component SOFC; and b) EFFCs.

1.2 Materials selection for solid oxide fuel cell technology

As mentioned previously, a typical single SOFC device is composed of three

components: porous cathode/anode, and a dense electrolyte. For the selection of

electrode materials, three simultaneous purposes must be fulfilled: i) providing the


6

electrochemical reaction sites; ii) providing the pathways for mass transfer (ions

and electrons); iii) collecting electrical current. Hence, the requisite electrode

properties include a porous structure, high electrical conductivity, chemical

stability in reducing/oxidizing operating conditions, and good mechanical

characteristics [1]. Regarding the conventional electrolyte selection, key factors

include: i) sufficiently high ionic conductivity; ii) low electron transport; iii) good

thermodynamic and chemical stability; iv) compatibility with electrodes; v) reliable

mechanical properties [10]. YSZ is the most popular choice for solid electrolytes,

although it requires a relatively high operating temperature to realize satisfactorily

high ionic conductivity [15]. In addition, other fluorite structure oxides, e.g. doped

ceria, have been developed to explore new electrolyte materials for LT-SOFCs

[22,42].

Ceria-based oxides, e.g. typically Ce0.8Sm0.2O1.9 (SDC) and Ce0.8Gd0.2O1.9 (GDC),

have been demonstrated as excellent electrolyte candidates for intermediate- or

low-temperature SOFCs due to their low activation energy, good catalytic activity,

and high ionic conductivity [43]. As reported in the literature, the conductivity of

doped ceria can be enhanced via an appropriate doping approach using low-valence

cations, meanwhile, the materials’ stability also can be improved [43–46]. However,

the existence of redox couple Ce4+/Ce3+ could lead to redox instability and possible

electronic conducting issue, thus resulting in undesirable structural failure and fuel

cell performance degradation [47]. Several studies have focused on the long-term

durability of high-conducting doped ceria materials [22,48–50]. For instance,

Virkar et al. [48] employed an electron blocking layer, e.g. YSZ, to prevent the

electronic leakage current induced by doped ceria. Wachsman et al. [22] reported a

bismuth-oxide-based material as a favorable conductivity electrolyte applied for

LT-SOFCs. Simone et al. [50] investigated δ-Bi2O3 using alternating layers of

Er2O3-stabilized δ-Bi2O3/Gd2O3-doped CeO2 with highly coherent interfaces,

which demonstrated good chemical stability under reducing conditions at high

operating temperature. Aside from these investigations, nanocomposite and

single/double-doping approaches have also attracted numerous researchers’ interest

in improving the ceria-based oxides’ ionic conductivity [51,52]. Banerjee et al. [53]

developed a ceria-based electrolyte by Ca/Sm co-doped route displaying a

superionic conductivity of 0.122 S cm-1 at 700 . These studies indicate hints that

the performance of doped-ceria could be further improved by surface modification

and multi-doping strategies.


7

To be cost-effective, natural minerals, such as CuAlO2, CuFeO2, and CuInO2, have

been intensively investigated for the applications of solar cells and photocatalysis

owing to their prominent catalytic activity [54,55]. Natural hematite was also

explored as a promising electrolyte of SOFCs [37]. Cuprospinel, i.e. CuFe2O4, is

another example illustrating the practical applications of natural minerals in SOFCs

[56]. CuFe2O4 also offers oxidation protection for SOFC metallic interconnects

along with protection against chromium poisoning [58,59].

1.3 Motivations and objectives

As mentioned in the previous section, the main challenge for SOFC technology is

lowering its operating temperature. Clearly, materials selection is a significant

influencing factor in realizing this goal. To this end, the overall aim of this

investigation is to employ new doping strategies for developing and improving

SOFC electrolyte materials. Both ceria-based electrolytes and natural mineral

CuFe2O4 are examined for this purpose in single, double, and triple doped

arrangements. Systematic investigations on material characterizations are carried

out, and electrochemical properties and fuel cell performance are evaluated for

potential LT-SOFC applications.

The detailed objectives to be achieved in this study are specified as follows:

i) Synthesize superionic conductive rare-earth (e.g. Sm3+/Pr3+/Nd3+)

doped ceria via a two-step wet chemical co-precipitation method, and

investigation of ionic conductivity and feasibility for LT-SOFC

technology. (Paper I)

ii) Develop mixed electronic/ionic conducting single-doped ceria (e.g. Pr-

CeO2) for applications in EFFC technology. (Paper II)

iii) Further develop double- and triple-doped ceria to obtain highly ionic

conducting materials and to produce semiconductor composites for

EFFCs. (Paper III, V and VI)

iv) Exploit the natural mineral CuFe2O4 for the use in EFFC devices,

including the determination of optimal fuel cell configurations. (Paper

IV)


8

1.4 Structure of this thesis

Chapter 1 gives some general background of the hydrogen energy economy, fuel

cells, and solid oxide fuel cells. The specific research objectives of this thesis are

also described.

Chapter 2 provides an overview of recent progress on LT-SOFCs, focusing on the

development of electrolyte materials. The basic working principles of SOFCs also

are stated in terms of thermodynamics and kinetics.

Chapter 3 summarizes the materials as well as characterization methods used in this

thesis work to develop and evaluate electrolytes for SOFC applications. Sample

synthesis, material characterization, and electrochemical properties are included.

In Chapter 4, an approach is presented for improving the ionic conductivity of

doped ceria via a two-step wet-chemical co-precipitation method.

In Chapter 5, the characterization investigations of single-, double-, and triple-

doped ceria materials for LT-SOFCs are presented.

Chapter 6 covers the developments of CuFe2O4 as SOFC materials, based on the

heat-treated raw chalcopyrite. Three types of fuel cell configurations using

CuFe2O4 are] investigated to optimize the CuFe2O4 application for LT-SOFCs.

Chapter 7 presents conclusions and suggestions for future work.


9

Chapter 2

Solid oxide fuel cell principles and literature review

In this chapter fundamental characteristics of fuel cell working principles,

particularly aspects related to electrochemistry and thermodynamics, are briefly

described. This provides the framework to understand how fuel cell performance is

influenced by materials properties and operating conditions [60]. The chapter also

includes an overview of key results pertaining to materials development in SOFCs,

with emphasis placed on emerging technologies, e.g. EFFCs.

2.1 Electrochemistry and thermodynamics

A schematic of the typical SOFC operating principle is depicted in Figure 2.1,

taking the O2--conducting electrolyte case as an example. In this case dry air is

supplied to the cathode, and the molecular oxygen from the gas flow reacts with

the electrons from the external circuit, decomposing it into two oxygen ions. (N2

and other gases contained in air are assumed to be inert and are thus neglected.)

The cathode reaction is described by:

1

2𝑂2 + 2𝑒

− → 𝑂2− (2.1)

Subsequently, the generated O2- ions migrate through the O2--conducting

electrolyte to arrive at the anode. Supplied H2 fuel is oxidized and decomposed into

two H+ ions, thereby releasing two electrons. The H+ ions can combine with O2-

ions with the formation of water, and the generated electrons can be then transferred

to an external electric circuit. The anode reaction is written as:

𝐻2+𝑂2− → 𝐻2𝑂 + 2𝑒

− (2.2)

The mass transfer and electric current are realized by a complete circuit in this

SOFC device. The overall electrochemical reaction is stated as:


10

𝐻2 +1

2𝑂2 → 𝐻2𝑂 (2.3)

Figure 2.1 Schematic of SOFC working principle.

2.1.1 Open circuit voltage

When considering a steady-state energy balance of a SOFC device, the enthalpy of

the input reactants including hydrogen and oxygen is equal to the sum of product

enthalpies, the power output, and the net heat dissipated to the surroundings [60].

Normally, the Gibbs free energy concept is employed in lieu of a classical energy

balance analysis, owing to the presence of electrochemical reactions under the

assumed conditions of constant pressure and operating temperature. The total Gibbs

free energy change (𝛥𝐺𝑓) for this device is indicated by the equation:

𝛥𝐺𝑓 = (𝐺𝑓)𝐻2𝑂 − (𝐺𝑓)𝐻2 − (𝐺𝑓)𝑂2 (2.4)

Here the terms (𝐺𝑓)𝐻2𝑂, (𝐺𝑓)𝐻2 , and (𝐺𝑓)𝑂2 refer to the Gibbs free energy of H2O,

H2, and O2, respectively. This reaction process is spontaneous and

thermodynamically favored due to the higher free energy of the reactants than that

of the products [60]. In an ideal reversible system, the yielded electrical energy


11

(𝑛𝐹𝐸) equals to the released Gibbs free energy (𝛥𝐺𝑓). Therefore, the change of free

energy for this SOFC reaction can be written as:

𝛥𝐺𝑓 = −𝑛𝐹𝐸 (2.5)

in which 𝑛 is the moles of electrons derived from equation 2.3, 𝐹 refers to

Faraday’s constant (96500 C g-1) while 𝐸 represents the reversible circuit voltage.

By the transformation equation, the electromotive force (or theoretical open-circuit

voltage, abbreviated as OCV) is calculated using:

𝐸 = −Δ𝐺𝑓

2𝐹 (2.6)

The electromotive force of a fuel cell is described using the Nernst equation:

𝐸 = 𝐸0 +𝑅𝑇

2𝐹ln𝑃𝐻2 ⋅𝑃𝑂2

0.5

𝑃𝐻2𝑂 (2.7)

where 𝐸0 refers to the ‘ideal’ electromotive force under standard conditions, and

𝑃𝐻2, 𝑃𝑂2 and 𝑃𝐻2𝑂, respectively, are the partial pressures of hydrogen in the anode

chamber, oxygen in the cathode chamber along with generated steam.

The cell voltage (𝑉), by definition, is given by the difference between cathode and

anode potential:

𝑉 = 𝐸cathode − 𝐸anode (2.8)

Ideally, the theoretical voltage of a fuel cell in equilibrium is equal to the

electromotive force calculated by the Nernst equation. However, in practical

applications a fuel cell must supply an electric current, resulting in a lower cell

voltage than the theoretical value. Figure 2.2 demonstrates the actual voltage

versus current response in comparison with an ideal model of the fuel cell. The

actual cell potential is lower than the theoretical one due to various irreversible

losses. Generally, the irreversibility is linked to polarization (or over-potential)

losses, attributing primarily to activation, ohmic and concentration polarizations.

Thus, the actual voltage of a fuel cell is written as:

𝑉 = 𝐸 − ∆𝑉act − ∆𝑉ohm − ∆𝑉con (2.9)

The activation polarization (∆𝑉act) originates primarily from the electrochemical

reaction rates, i.e. hydrogen oxidation reaction occurred in anode and oxygen

reduction reaction occurred in the cathode. Both the reaction kinetics and electrode

geometric effects have an impact on activation polarization [26]. The ohmic

polarization (∆𝑉𝑜ℎ𝑚) is dominated by the ionic conductivity capacity of electrolytes,

but the electronic conductivity of electrodes contributes insignificantly to the ohmic


12

resistance. With respect to the concentration polarization (∆𝑉𝑐𝑜𝑛), the limited mass

transport of reactants and products are contributing factors.

Figure 2.2 Ideal and actual potential vs current response in a fuel cell [11].

2.1.2 Efficiency

The overall efficiency relies on the kinetics and thermodynamics of fuel cells [22].

It can be described primarily using three parameters including the

thermodynamic/Gibbs efficiency (𝜂𝑡ℎ), the voltage efficiency (𝜂𝑣) and the fraction

of fuel used or fuel utilization efficiency (𝜂𝑓). The overall efficiency (𝜂) is given

here [16]:

𝜂 = 𝜂𝑡ℎ × 𝜂𝑣 × 𝜂𝑓 (2.10)

The thermodynamic efficiency is described considering the reaction enthalpy (∆H)

and entropy (∆S). Enthalpy refers to the delivered heat in a reaction and entropy is

associated with the change of reaction order. The well-known equation relating

these terms with Gibbs free energy (𝛥𝐺𝑓) is expressed as:

𝛥𝐺𝑓 = 𝛥𝐻 − 𝑇𝛥𝑆 (2.11)

In the SOFC reaction, 𝛥𝐺𝑓 is less than 𝛥𝐻 and reaction heat is delivered from the

conversion of chemical energy stored in the fuel. Therefore, the thermodynamic

efficiency can be described as [60–62]:

𝜂𝑡ℎ =Δ𝐺𝑓

𝛥𝐻= 1 −

𝑇𝛥𝑆

𝛥𝐻 (2.12)


13

The voltage efficiency (𝜂𝑣) is calculated as the ratio between the practical potential

and the theoretical electromotive force:

𝜂𝑣 =V

E=𝐸−∆𝑉act−∆𝑉ohm−∆𝑉con

𝐸 (2.13)

The fuel utilization efficiency (𝜂𝑓) can be expressed as the ratio between the actual

fuel and the total supplied fuel. In a SOFC operation, the generated current density

is i and the input rate of fuel is 𝑣𝑓 (mol sec-1). Hence,

𝜂𝑓 =𝑖/𝑛𝐹

𝑣𝑓 (2.14)

Substituting Equation 2.9 using 2.11, 2.12 and 2.13, the overall efficiency of SOFC

is expressed as:

𝜂 = (1 −𝑇𝛥𝑆

𝛥𝐻) × (

𝐸−∆𝑉act−∆𝑉ohm−∆𝑉con

𝐸) × (

𝑖/𝑛𝐹

𝑣𝑓) (2.15)

The overall fuel cell efficiency is determined by three aspects, as described in

Equation 2.15. The thermodynamic efficiency, without the limitation of the Carnot

Cycle, suggests that low temperature is favored. The voltage efficiency is

principally affected by the properties of fuel cell materials, e.g. the ionic

conductivity of electrolytes, while the catalytic activity of the electrodes contributes

to the fuel utilization. Therefore, lowering the operating temperature and

developing improved electrode/electrolyte materials to reduce cell polarization and

enhance fuel utilization efficiency are the primary methods to improve overall

SOFC efficiency [15,22].

2.2 Materials development

In practical applications, a SOFC always functions in a system. Generally, a SOFC

stack consists of several individual fuel cells, and there are three different stack

configurations that are widely used: tubular, monolithic and planar SOFC stacks

[1]. The schematic diagrams of the SOFC stack, individual SOFC device, and

functional SOFC components are shown in Figure 2.3. As described previously,

the key technical challenge for the development of the SOFC is to reduce its

operating temperature, thus developing appropriate electrode/electrolyte materials

is the principle issue for SOFC technology. In this section, the material

developments for individual SOFC components are discussed in the context of a

literature survey.


14

Figure 2.3 Schematic diagrams of SOFC stack, individual SOFC device and

functional SOFC components [11].

2.2.1 Anode materials development

In the oxygen-conducting SOFC case, the anode functions in multiple roles: i) to

catalyze the electrochemical oxidation reaction of fuels, e.g. H2; ii) to transfer

oxygen ions from the electrolyte/anode interface into the anode structure; iii) to

transport the generated electrons; and iv) to assist the reactant and product transport

at the reaction sites [21,63,64]. Among various anode materials for SOFCs, the

most widely used ones are porous ceramic and metallic composites of nickel cermet

bonded to ion-conducting materials, most often yttria-stabilized zirconia (Ni/YSZ)

[63–67]. Ni/YSZ has several advantages: low-cost, good high-temperature stability

and matching thermal expansion to the most commonly employed electrolytes [68].

In addition, the nickel cermet serves as an excellent electron conductor and good

electrocatalyst for fuels (particularly hydrogen) at relevant operating temperature,

while YSZ acts as a porous mechanical support and provides ionic conduction

pathways [63,65,66]. However, Ni/YSZ anode has some drawbacks, such as

potential reduction/oxidation cycling instability and low-tolerance to impurities

leading to degraded fuel cell performance [63,69–71]. Therefore, current research

efforts are aimed at developing new anode materials, such as ceria-containing

anodes [72], perovskite-type ceramics [73,74], and layered materials [75].

2.2.2 Cathode materials development

In the overall SOFC process, electrochemical reactions occurring at the cathode

region have been given more attention in comparison to anode-related reactions.

This is can be interpreted by the slower reaction kinetics of oxygen reduction at the


15

cathode than that of hydrogen oxidation at another electrode. Therefore, the cathode

is the primary contributor to the overall polarization resistance of fuel cells.

For SOFC cathode materials, good catalytic activity for oxygen reduction reaction

(ORR) and electron transport functionalities are required. Furthermore, the desired

cathode materials should be able to facilitate mass transfer and be matched

thermomechanically and chemically with other components, i.e. electrolyte and

interconnections [1,10,16,76]. In order to work at low operating temperature,

appropriate cathode materials with good low-temperature catalytic activity for

ORR have been in focus for reducing cathode polarization losses and for achieving

long-term stability [77,78]. Poor ORR catalytic activity is still the key obstacle in

developing commercially viable LT-SOFCs (below 600 ) [78]. To improve

cathodic performance, Shao et. al. [76] proposed two strategies: i) developing new

cathode materials to perform good fuel cell performance in the low-temperature

regime; and ii) optimizing the microstructure of the developed cathode materials to

reduce the cathode polarization resistance.

Figure 2.4 Crystal structures of the three-type cathodes: a) perovskite; b) layered

perovskite; and c) Ruddlesden-Popper oxides [77].

As reported in the literature, three main perovskite-related types, including the

cubic-type, layered-type and Ruddlesden-Popper type perovskites, were selected as


16

promising cathode candidates for LT-SOFC applications [77]. The crystal

structures for the three-type cathode materials are displayed in Figure 2.4. As

shown, all of the three-type cathodes have an oxygen 6-fold coordinated transition

metal scaffold with the alkaline-earth or lanthanides ions-located vertices at a cube

structure [77]. Among these potential cathode materials, La1-xSrxCo1-yFeyO3

(LSCF) and La1-xSrxMnO3 (LSM) have demonstrated excellent mixed electronic-

ionic conducting characteristics and good ORR catalytic activity [76–81].

2.2.3 Electrolyte materials development

For the desired electrolyte, several requirements need to be met: i) sufficiently high

ionic conductivity; ii) negligible electronic conduction; iii) thermodynamically and

chemically stabilizable property during fuel cell operating process; iv) reliable

mechanical strength; v) chemical inertness with anode/cathode during processing;

and vi) compatible thermal expansion coefficient to the anode and cathode

materials [10].

Figure 2.5 Schematic image for fluorite-structured oxide [15].


17

In the published literature, four structural groups of O2--conducting oxides were

employed as electrolytes in SOFCs, including fluorite-type, perovskite, intergrowth

perovskite and pyrochlore-type [82–86]. Generally, oxygen ions can be transported

through the oxygen vacancies in oxygen-ion conductors and the oxygen vacancy

concentration can be enhanced by doping. In a typical fluorite-structured zirconia

(ZrO2) case, the formation of oxygen vacancy with yttria as a dopant can be

described using the defect equation in Kröger-Vink notation [10,86]:

𝑌2𝑂3 → 2𝑌𝑧𝑟′ + 3𝑂0

𝑋 + 𝑉0·· (2.16)

Figure 2.5 demonstrates the fluorite-structured oxide with the formed oxygen

vacancies. As observed, the host cations, commonly Zr4+ or Ce4+, can be replaced

by lower valent ions, e.g. Y3+ or Gd3+, to create oxygen vacancies and thus improve

their ionic conductivities.

For proton-conducting SOFCs, the operating mechanism is different from an

oxygen ion-conducting SOFC device. In this case, hydrogen is catalytically

decomposed into protons on the anode side, and thereafter protons migrate to the

cathode region across the electrolyte. Meanwhile, water vapor is formed on the

cathode side. The electrochemical reaction equations in anode, cathode and the

overall equation, respectively, are expressed as follows:

𝐻2 → 2𝐻+ + 2𝑒− (2.17)

1

2𝑂2 +𝐻

+ + 2𝑒− → 𝐻2𝑂 (2.18)

𝐻2 +1

2𝑂2 → 𝐻2𝑂 (2.19)

The migration of protons in an H+-conducting electrolyte is easier to be realized

than the migration of oxygen ions in oxygen-ion conductors, resulting in lower

activation energy. Thus, a high proton conductivity can be expected, making the

intermediate or low-temperature operation of SOFCs feasible [17,87,88]. Most of

the reported proton conductors are perovskite-type oxides, including BaCeO3- and

SrCeO3-based oxides, or Y-doped barium zirconate [17,87,89,90]. Among these

materials, BaCeO3-based perovskites are verified to demonstrate some of the

highest proton conductivities [87,91].

In the mixed proton/oxygen ion-conducting electrolyte cases, nanocomposites have

been shown to make substantial contributions, especially from the viewpoint of

improving ionic conductivity. Two categories of nanocomposites have been


18

extensively employed as electrolytes for SOFCs. The oxide/carbonate

nanocomposites and particularly ceria or doped ceria/carbonate materials, as the

first type, have attracted attention during the past decade owing to the enhancement

of ionic conductivities via carbonate coatings [92–97]. Oxide/oxide

nanocomposites, representing the second category, typically exhibit mixed

proton/oxide ion conductivities and can be classified into three groups: ionic

conductor/ionic conductor (e.g. YSZ/SrTiO3 [98]), ionic conductor/insulator (e.g.

Gd0.2Ce0.8O2/Mn-Co doped Al2O3 [99]), and semiconductor/semiconductor (e.g.

LiFeO2-LiAlO2 composite [100]).

2.3 Nanocomposites for EFFC technology

Beyond the three groups of oxide/oxide nanocomposites mentioned above, a fourth

type – semiconductor/ionic-conductor composites – has been developed

specifically for EFFC technology. This nanocomposite provides enhanced ionic

conduction with the addition of electronic conductivity. Such materials are not

suitable as conventional SOFC electrolytes owing to the internal short-circuiting

issue. In contrast, the newly developed EFFC concept is able to exploit these traits

by semiconductor/ionic-conductors as homogeneous two-phase materials for

embedding anode and cathode within the electrolyte layer. To elaborate further,

semiconductor/ionic conductor nanocomposites consist of two phases: ion-

conducting phase, and semiconductor phase. Together both are capable of

functioning as a single-component fuel cell. There are several requirements for

semiconductor/ionic nanocomposites in EFFCs [31,32,34,40,101–103]:

i) balanced ionic and electronic conduction;

ii) existence of percolating network for electrons/holes and ions;

iii) existence of P-N junctions formed inside the semiconductor-ionic

composite, to prevent the electrons through the single component layer.

As proposed in the literature, the catalytic roles for the oxidation of hydrogen and

reduction of oxygen are realized through the functional semiconductor/ionic

nanocomposite [101,102]. The electrochemical reactions in an EFFC device can be

described as follows [102]:

At H2-contacting side (hydrogen oxidation):

𝐻2 → 2𝐻+ + 2𝑒− (2.20)

At O2-contacting side (oxygen reduction):


19

1

2𝑂2 + 2𝑒

− → 𝑂2− (2.21)

Overall fuel cell reaction:

𝐻2 +1

2𝑂2 → 𝐻2𝑂 (2.22)

A key fundamental issue to be addressed for EFFC is how to realize the separation

of ions and electrons, thus avoiding the internal short-circuiting problem in a

homogenous semiconductor and ionic conductor layer. Up to now, although

considerable experimental results with high OCVs and comparable power output

have been reported, no internal short circuit occurring has been shown for EFFCs

[30–32,35,41]. Some possible principles to explain the charge separation inside

EFFCs have been proposed based on the physical bulk heterojunction [31,32],

Schottky junction [34] and the energy band alignment [36].


20


21

Chapter 3

Experimental methods and techniques

3.1 Raw materials

The main materials used in this investigation are the single, double, or multi rare-

earth elements (e.g. Sm, Pr, Nd, and La) doped-ceria materials. The used raw

materials include Ce(NO3)2·6H2O, Sm(NO3)2·6H2O, Pr(NO3)3·6H2O,

Nd(NO3)3·6H2O, La(NO3)2·6H2O and Na2CO3. The origin chemicals are analytical-

grade level and are obtained from Sinopharm Chemical Reagent Co. Ltd (China).

Several approaches, including co-precipitation method, solid-state reaction, and

hydrothermal method, are employed here to obtain the desired samples.

The natural cuprospinel comes originally from the Tonglvshan zone, Hubei

Province, China, modified by a simple heat-treatment process.

Commercial multiple oxides, i.e. Ni0.8Co0.15Al0.05Li-oxide, abbreviated as NCAL,

are acquired from Tianjin Bamo Sci. & Tech. Joint Stock Ltd (China).

3.2 Sample preparation

3.2.1 SDC preparation by co-precipitation approach

Samarium doped ceria (denoted shortly as SDC) is synthesized by a carbonate co-

precipitation solution. First of all, a designed stoichiometric ratio of 4:1 for

Ce(NO3)2·6H2O and Sm(NO3)2·6H2O are weighed and mixed with deionized water.

In parallel, a solution of sodium carbonate is prepared as the precipitant for the

above metal salt solution. The molar amounts of sodium carbonate are double that

of the total metal ions. Subsequently, the as-prepared mixed solution is stirred

vigorously for 6 h and then allowed to reset at the ambient temperature for another

6 h. Afterward, the product is washed several times in the deionized water to

remove any unreacted carbonate. The final SDC product is obtained after a 12-h

drying process followed by a 4-h calcination process at 800 .


22

3.2.2 Pr-doped ceria preparation by hydrothermal method

Pr-CeO2 sample is designed to synthesize with Pr3+ dopant as 10% in molar ratio

via hydrothermal method followed by a calcination process. The chemicals,

including Ce(NO3)3·6H2O, Pr(NO3)3·6H2O, and Na2CO3, are used in the

preparation process for Pr-CeO2 material. The precursor of Pr-CeO2 is prepared

firstly through a hydrothermal route. In this typical synthesis, the designed molar

amounts of Ce(NO3)3 and Pr(NO3)3 are dissolved in deionized water. Na2CO3

solution with a molar ratio to the total of metal ions (Ce3+ and Pr3+) of 2:1 is added

into the precursor solution as a precipitant. The precipitation-forming reaction is

carried out at 90 under continuous stirring conditions for 2 h. Subsequently, the

obtained precipitation is transferred into polytetrafluoroethylene reactor maintained

in 120 for 4 h. Then the product from hydrothermal process is sintered in a

muffle furnace at 700 for 2 h, followed by washing repeatedly with distilled

water. As a comparison, a bare CeO2 sample is synthesized in the same procedure.

3.2.3 Sm3+, Pr3+ and Nd3+ triple-doped ceria by two-step doping method

Three elements (i.e. Sm3+/Pr3+/Nd3+) combined to form a triple-doped ceria sample

are synthesized by means of a two-step co-precipitation solution for yielding the

desired elemental surface distribution map. First of all, Sm3+ is doped into the bulk

of ceria, i.e. SDC, via the carbonate co-precipitation method as mentioned above.

Praseodymium and neodymium nitrates (9:1 molar ratios) are used as the secondary

doping ions, introduced by a modified solid-state reaction method. The total molar

ratio of the praseodymium and neodymium nitrates are designed for 1.8 mol. % and

are dissolved in the deionized water to obtain transparent precursor solution.

Subsequently, SDC is added to the prepared solution and homogeneous suspension

is obtained. The final product, PNSDC (Sm3+/Pr3+/Nd3+-doped ceria), is obtained

after a 2-h calcination process at 750 followed by coarse grinding.

3.2.4 Double- and triple-doped ceria preparation

Double-doped ceria powder (Ca0.04Sm0.16Ce0.8O2-δ, SCDC) is prepared by means of

a co-precipitation route. First, the stoichiometric amounts of Ca(NO3)2 (4 mol. %),

Sm(NO3)2·6H2O (16 mol. %) and Ce(NO3)3·6H2O (80 mol. %) are dissolved in

deionized water to form 0.5 mol L-1 solution. Na2CO3 is added into the above salt-

solution as the precipitating agent with stirring vigorously. A white precipitate is

observed in the reactant solution. Subsequently, the resultant product is repeatedly

washed using deionized water to remove any residual Na2CO3. The precursor is put

into a drying oven at 120 for around 10 h. The final SCDC sample is obtained

after a 4-h calcination procedure at 800 followed by grinding.


23

The triple-doped ceria samples (i.e. doping with La3+, Pr3+, and Nd3+, denoted as

LCPN) are synthesized using a co-precipitation method. Primarily, the raw rare-

earth nitrates including La, Pr, and Nd elements are dissolved into the deionized

water and Na2CO3 is used as the precipitating agent. Then, multiple filtration and

washing with deionized water are conducted on the resulting solution followed a 2-

h calcination process at 750 to yield the final LCPN sample.

3.2.5 Doped ceria/semiconductor composite preparation

The doped-ceria/semiconductor composites are prepared via a straightforward

solid-state blending method. First of all, doped-ceria is prepared as described above.

Then four groups of the resulting doped-ceria powders are separately mixed with

NCAL or LSCF (La0.6Sr0.4Co0.2Fe0.8O3-δ) in various ratios (20 wt.%, 30 wt.%, 40

wt.%, and 50 wt.% NCAL or LSCF). These mixtures are ground and calcined at

700 for 30 minutes and then ground once again to obtain homogeneous doped-

ceria/NCAL or doped ceria/LSCF composite samples.

3.2.6 LZSDC nanocomposite preparation

LZSDC nanocomposite is synthesized as reported in the literature [104]. Briefly,

identical molar amounts of LiNO3 and Zn(NO3)2 are dissolved and citric acid is

added as the precipitating agent. (The amount of citric acid is three-quarters of the

integral molar amounts of both nitrates.) Subsequently, a continuous stirring

operation is conducted on this reactant under 120 condition to form a sol-gel

with the addition of SDC powder. The final LZSDC is obtained after a 2-h

calcination procedure.

3.2.7 Natural mineral treatment

Cuprospinel is firstly crushed and ground to 200 meshes. Subsequently, the raw

mineral is directly calcined at different temperature (i.e. 600 , 700 , and

900 ) in air atmosphere lasting for 2 h, respectively to obtain different products.

3.3 Fuel cell device fabrication

In this thesis, two kinds of fuel cell devices are fabricated in terms of the

conventional three-component SOFC configuration and the emerging EFFC

configuration. The devices are co-pressed to form a disc device under the 150 MPa

pressure. The diameter of the fuel cell device is 13 mm and the thickness is about

1 mm. The active area for fuel cell reactions is approximately 0.64 cm2.


24

3.3.1 Conventional SOFC device

The single-phase Sm3+, Pr3+ and Nd3+ triple-doped ceria with super-high ionic

conductivity is used as an electrolyte for LT-SOFC. The NCAL combining with

PNSDC sample is employed as both cathode and anode, denoted as Cell I. A

parallel pellet (Cell II) is fabricated using SDC to replace the PNSDC in Cell I.

Commercial nickel foam is attached on both sides of the pellet as the structure

supports. The fuel cell devices can be briefly written as follow:

Cell I: Ni-foam/SDC+NCAL (anode)|SDC (electrolyte)|SDC+NCAL

(cathode)/Ni- foam

Cell II: Ni-foam/PNSDC+NCAL (anode)|PNSDC (electrolyte)| PNSDC+NCAL

(cathode)/Ni- foam

3.3.2 EFFC device

In the procedure of EFFC fabrication, the commercial NCAL is used and attached

to the devices as the supporting structure. Firstly, NCAL is mixed with terpineol to

obtain a NCAL slurry. Subsequently, the NCAL slurry is pasted on the nickel foam.

This structure is denoted as Ni-foam/NCAL.

3.3.2.1 EFFC device based on single-phase Pr-CeO2

Figure 3.1 Schematic illustration of the single fuel cell and measurement setup.

The single symmetric fuel cell is fabricated in a ‘sandwich’ configuration (Ni-

foam/NCAL | Pr-CeO2 | NCAL/Ni-foam) with Pr-CeO2 sample as core membrane

layer and semiconducting NCAL thin layer pasted onto nickel foam attached to

both sides of Pr-CeO2 layer. The same configuration is constructed using CeO2


25

sample instead of Pr-CeO2. The schematic illustration of a single fuel cell and

measurement setup are presented in Figure 3.1.

3.3.2.2 EFFC device based on triple-doped CeO2

The EFFC devices using LCPN sample are fabricated, similar to the above

configuration. The core component of this device is the LCPN/NCAL composite

with Ni-foam/NCAL employed as supporting structure on both sides of this fuel

cell pellet. In parallel, a range of weight ratios for the core LCPN/NCAL composite

are designed to fabricate various EFFC devices: LCPN:NCAL of 7:3, 6:4, and 3:7,

respectively. In addition, the device using pure LCPN as the electrolyte is also

prepared in the same procedure as the comparative case.

3.3.2.3 EFFC device based on natural mineral

To investigate the optimal fuel cell application for this natural mineral, three fuel

cell devices are fabricated in different configurations. In Type I, CuFe2O4 is used

as the cathode, while NCAL and LZSDC are employed as anode and electrolyte,

respectively. In Type II, the CuFe2O4 is mixed with LZSDC to yield a

CuFe2O4/LZSDC composite with a volume ratio of 1:1. Then the CuFe2O4/LZSDC

composite is appended to both sides of the device as a functional layer. For Type

III, CuFe2O4 material is used combing LZSDC and the semiconducting NCAL as

the core component of the EFFC device. The CuFe2O4/LZSDC composite is mixed

with NCAL in an equal volume.

3.4 Materials characterizations

3.4.1 X-ray diffraction (XRD) analysis

The physical characterizations for the materials used in this thesis are analyzed

using X-ray diffraction (XRD, D-max-2500 X-ray diffractometer, Rigaku Corp.,

Japan) with filtered Cu Kα radiation. The data are recorded in a step-scan mode of

the 2theta range of 10-80o with a step size of 0.02o.

3.4.2 Micro-structure analysis

The morphology and microscopy analysis are carried out on a field emission

scanning electron (FE-SEM, Hitachi S-4800). It is also equipped an energy-

dispersive spectrometer (EDS) to analyze the elemental information of samples.

The transmission electron microscopy (TEM) measurement is carried out on the

equipment of TECNAI20 or JEM-2100, JEOL.


26

3.4.3 X-ray photoelectron spectroscopy (XPS) analysis

X-ray photoelectron spectroscopy (XPS) analysis is performed on a Physical

Electronics Quantum 2000, Al Kα X-ray source. It is conducted to investigate the

surface chemical properties of the used samples, such as SDC, PNSDC, and Pr-

CeO2. The information related to the oxidation states of surface elements, as well

as the atom species can be identified. The IGOR software (Wavemetrics, Lake

Oswego, OR) is used here to fit the XPS original measured data. The chemical

analysis also can be performed based on the fitting results. The binding energies

are calibrated using the C 1s peak at 284.6 eV as a reference and quoted with a

precision of ± 0.2 eV.

3.4.4 Raman spectra analysis

Raman spectra analysis is carried out on a Renishaw Ramascope equipped with a

Leica LM optical microscope (a CCD camera). The excitation laser lines with the

wavelengths λex=325 and 532 nm are used in this investigation, and calibration is

conducted using a Si wafer.

3.5 Electrical properties

3.5.1 Electrochemical impedance spectroscopy (EIS) analysis

EIS technique is used to determine the resistive and capacitive characterizations of

materials. The impedance spectrum can be obtained by changing the frequency over

a small applied amplitude sinusoidal alternative current excitation signal [105]. In

this thesis, EIS measurements are carried out to study the electrical conductivities

of the as-synthesized samples and fuel cell devices. The electrical conductivities of

materials and SOFC devices are measured in air or air/H2 atmosphere by an

electrochemical workstation (CHI660B, Chen Hua Corp.) under open circuit

conditions.

3.5.2 Direct-current (DC) polarization analysis

The electronic conductivity is tested using a Hebb-Wagner ion-blocking cell by

means of the DC polarization approach. The DC measurement is performed on a

digital micro ohm meter (KD2531, Kangda Electrical Co., Ltd.) with a voltage

range of 0.01-1.0 V. The Gamry Reference 3000 instrument is employed to analyze

the electrochemical performance of materials and fuel cell pellets. The measured

frequencies range from 0.01 Hz to 1 MHz under a bias voltage of 10 mV. The


27

electrical conductivity (𝜎) for both materials and devices can be calculated using

the equation:

𝜎 =𝐿

𝑅𝐴 (3.1)

Here, 𝜎 is the total electrical conductivity. L presents the thickness of the measured

pellets. R is the total resistance and A is the active area of the pellets. At the initial

stage, the current is raised from the mixed e-/O2- movements under the Hebb-

Wagner cell polarization condition. Subsequently, the oxygen molecules are

consumed at the inner Au/Pr-CeO2 interface with decreasing chemical potential.

Finally, the constant electric current is acquired under the built-in ion-blocking

steady-state condition. In this case, the electronic conductivity can be evaluated

from the obtained current (I) versus time (t) curves.

3.6 Fuel cell performance measurements

The fuel cell performance measurements are conducted for pellet-size devices

connected to a programmable electronic load (ITECH8511, ITECH Electrical Co.,

Ltd.). Hydrogen is supplied to the anode side with flow rates of 120~130 ml min-1

at a pressure of 1 bar, while the cathode side is supplied with air at a similar flow

rate and pressure. The laboratory fuel cell performance testing system is shown in

Figure 3.2. A computerized instrument (IT7000) is employed to measure the I-V

characteristics of a single fuel cell device and corresponding power densities can

be calculated by multiplying voltage and current. The scanning rate for recording

experimental data is set to 0.05~0.1 A s-1.

In this testing system, several potential random factors might impact the accuracy

of the experimental results. One uncertainty comes from the difference between the

set fuel cell operating temperature (determined by a thermocouple) and the actual

cell pellet temperature (non-contact between the pellet and thermocouple). As

discussed in previous sections, the temperature fluctuation can influence the

thermodynamic and electrochemical properties, e.g. ionic conductivity of the solid

electrolyte, thus leading to imprecise fuel cell performance. The experimental error

from the measured furnace temperature by the thermocouple is +/- 5 .

Two different kinds of gas flow meters are employed in this work, which might

also contribute to experimental uncertainty. For the initial experiments, a rotameter-

style (Sho-RateTM 1355) is used to record the H2 flow rates (80~120 mL min-1)

and the air flow rates are not recorded. In the later experiments with an updated gas

flow controller (Bronkhorst F-201CV, for both H2 and air), however, the H2 flow

rates of 120~130 mL min-1 are required to obtain the similar results to the earlier


28

experiments. The desired air flow rate range of 120~130 mL min-1 also is monitored

for the consistent performance outcomes with the previous experimental data. An

accuracy of +/- 2 % for the F-201CV is provided by the instrument supplier. More

details on the gas flow measurements can be found in Xia [158] and Afzal [159].

Some other uncertainty factors might come from the material preparation and fuel

cell fabrication processes, such as the purity of raw materials and weighting

instrument errors. Additionally, some slight heterogeneous particle size and

compositions of composites might influence the sample properties and fuel cell

performance. Repeated experiments (at least three for each set of data) are strictly

required for the material preparation and fuel cell measurements to minimize the

uncertainty and obtain reproducible experiment data.

Figure 3.2 Schematic diagram of fuel cell performance testing setup.


29

Chapter 4

Superionic conductivity of triple-doped ceria for SOFC

This chapter presents a novel approach for modifying the surface properties of SDC

with Pr3+ and Nd3+ dopants, resulting in a triple-doped material. Pr3+ and Nd3+ are

selected due to their potential in extending the electrolytic domain boundary to

lower oxygen partial pressure as well as to enhance the catalytic activity for ORR.

For this purpose, the mechanisms of the enhanced ionic conduction and redox

capability are analyzed. This chapter is based on the published article Paper I.

4.1 Material characterizations

4.1.1 Phase and microstructure properties

Figure 4.1a displays the XRD patterns of mono-cation doping (i.e. SDC) and tri-

cation doping (i.e. PNSDC) ceria oxides. According to JCPDS 34-0394, the

diffraction peaks of the two samples could be assigned to the cubic ceria structure.

A slight shift of diffraction peaks to lower Bragg values can be observed in the

XRD patterns of PNSDC sample compared to SDC. This can be attributed to the

lattice expansion of PNSDC owing to larger ionic radius of Pr3+ ion (i.e. 1.126 Å)

and Nd3+ ion (i.e. 1.109 Å) than Ce4+ (i.e. 0.97 Å) [106]. The lattice constants are

calculated from XRD patterns, which are 5.414 and 5.415 Å, respectively, for SDC

and PNSDC samples. Both of them are higher than that of undoped ceria with a

reported lattice constant of 5.411 Å [104]. The particle sizes of SDC and PNSDC,

according to Scherrer’s equation, are estimated as 55 nm and 61 nm, respectively

[107]. The estimated particle size of PNSDC is approximately 70 nm as shown in

SEM and TEM images (Figure 4.1b). The observed results are consistent with the

calculated data according to XRD results, implying that the as-prepared PNSDC

nanoparticles by the co-precipitation method are homodisperse. The EDS element

mapping analysis (see Figure 4.2) shows that Pr3+ and Nd3+ elements are

successfully distributed onto the SDC scaffold. Therefore, a conclusion can be


30

summarized that Pr3+/Nd3+ doping-ions have been successfully introduced into the

SDC phase without the formation of a second phase.

Figure 4.1 a) X-ray diffraction patterns of SDC (black line) and PNSDC (red

line); b) SEM and TEM micrographs for PNSDC powder.

4.1.2 Ultraviolet-visible (UV-vis) Raman scattering analysis

UV-vis diffuse reflectance spectroscopy technique is employed here to analyze the

surface coordination property and the relationship of absorbance and wavelength

for the obtained samples, shown in Figure 4.3a. A strong absorption can be

observed in the UV section of < 400 nm, typically at the wavelength of 325 nm, for

SDC and PNSDC samples in the shown absorption spectra. In the visible region of

> 400 nm, a weak absorption is observed in the PNSDC line, while SDC absorption

spectrum exhibits a relatively lower absorbance, such as at the wavelength of 532

nm. It could be speculated that the PNSDC sample has an increased oxygen

vacancy concentration compared to SDC due to the replacement of Pr3+/Nd3+ ions

to Ce4+ ions in PNSDC. A related study on the relationship of oxygen vacancy and

UV absorption spectra has been conducted by Sanna et al [50]. As shown in Figure


31

4.3a, the intensity of the absorption for both SDC and PNSDC samples gradually

decreases with the increase of wavelengths. This gives a hint that the different

information from various sampling depths can be collected by means of changing

excitation laser lines using Raman spectroscopy technique. It provides a solution to

analyze the oxygen vacancy information on surface and bulk of samples by

changing the different excitation laser lines (e.g. 325 nm and 532 nm).

Figure 4.2 SEM image of PNSDC and corresponding EDS element mapping.

As reported in the literature, the UV-visible Raman spectroscopy technique is

employed to analyze the surface vibrational mode of materials, while the lattice

structural characteristics are also investigated [108,109]. Figure 4.3b and 4.3c

show the Raman spectra of the SDC and PNSDC samples using excitation laser

wavelengths (λex) of UV (λex=325 nm, Figure 4.3b) and visible (λex=532 nm,

Figure 4.3c). In a fluorite-type CeO2, the symmetrical Ce-O stretching vibration

(denoted as F2g), indexed to surrounding 460 cm-1 with an excitation laser

wavelength of λex=532 nm, plays a dominant role in Raman active modes [109–

112]. In this investigation, the F2g modes for both SDC and PNSDC with a peak

located at vicinal 461 cm-1, marked as peak α in the images, exhibit a shift to lower

energy compared to the reported results for CeO2 [113]. It is worth noting that a

new characteristic peak β, can be observed in the vicinity of 554 cm-1 for both SDC

and PNSDC in Figure 4.3b [113]. This mode indicates the formation of oxygen

vacancies owing to the substitutions of Ce4+ ions by Sm3+, Pr3+ or Nd3+ ions.


32

Compared to the SDC spectrum, the Raman feature lines of PNSDC display a

broader and more asymmetrical characteristic for both F2g mode and the one

induced by oxygen vacancies. It demonstrates a higher dopant concentration and

more oxygen vacancies in PNSDC after the secondary doping Pr3+/Nd3+ ions [114].

Figure 4.3 a) UV-vis diffuse reflectance spectroscopy of SDC and PNSDC; UV-

visible Raman spectra of SDC and PNSDC with an excitation laser wavelength b)

λex=532 nm and c) λex=325 nm; d) Schematic description for Raman scattering by

varying the ultraviolet (UV, λex=325 nm) and visible (λex=532 nm) excitation laser

lines to collect the surface and bulk information.

Figure 4.3d is a diagrammatic drawing showing the relationship of the ultraviolet

and visible excitation laser wavelengths (in this case, λex=325 and 532 nm,

respectively) with the sites of oxygen vacancy defects, e.g. inside bulk or on the

surface of the samples. Using a strong excitation laser with a wavelength of 325

nm, the surface properties of samples are detected from the escaped scattering light

while the major amount of the excitation laser is absorbed. Therefore, the surface


33

oxygen vacancy can be determined via UV Raman lines. Accordingly, the bulk

properties of PNSDC and SDC samples are obtained using a relatively weak

excitation laser, e.g. a wavelength of 532 nm [108,109,115]. Figure 4.3c exhibits

a significantly intensive peak for PNSDC (peak β) in the vicinity of 554 cm-1 using

a UV excitation laser of 325 nm, compared to that of SDC sample.

Further evidence is gathered to support the replacement of Ce4+ ions using Pr3+/Nd3+

ions, particularly on the surface regions of the PNSDC sample. As stated previously,

peaks β in Raman spectra reflects the oxygen vacancy information and peak α

represents the F2g mode information of the doped ceria with a fluorite-type structure.

The intensity ratio between the two characteristic peaks (written as Aβ/Aα) is

calculated by integration of individual peak area. The results of Aβ/Aα reflect the

oxygen vacancy concentration and indeed the defect sites combining with measured

depths of PNSDC and SDC [116]. By computation, the values of Aβ/Aα for PNSDC

and SDC are approximately 0.298 and 0.301, respectively, for the case using

excitation laser wavelength of 532 nm (Figure 4.3b). The similar results

demonstrate no apparently structural change occurring inside the bulk of PNSDC

sample after secondary Pr3+/Nd3+ doping. However, an obvious change is observed

when using the UV Raman excitation laser, i.e. 325 nm, as shown in Figure 4.3c.

The calculated Aβ/Aα value is about 1.512 for PNSDC, while 1.294 is obtained from

the calculation for the peak areas of the SDC sample. The marked difference

implies a higher concentration of the oxygen vacancy in PNSDC compared to that

of SDC. Thus, one can conclude that the Pr3+ and Nd3+ ions are doped on the surface

sites of SDC particles.

4.1.3 XPS analysis

As shown in Figure 4.4a, the survey spectra exhibit the characteristics of various

elements, i.e. Sm, Ce, Pr, Nd, C, and O in the PNSDC. The individual XPS peaks,

representing the O 1s, Pr 3d, and Ce 3d, are tested to identify the chemical states

on the surface of PNSDC. The standard Gaussian functions are employed to

analyze these relevant multiple-component XPS featuring peaks.

Figure 4.4b displays the O 1s core-level spectra related to oxygen species,

including surface chemisorbed oxygen or oxygen defects (i.e. O−, Oα) with a

binding energy of 531.0-532.6 eV and lattice oxygen with binding energy (Oβ) at

the range of 528.8-529.4 eV [117]. According to literature, the bonds associated

with Ce-O-doping cation are generated when the doping ions are introduced into

the host ceria lattice, thus resulting in more mobile oxygen species to facilitate the


34

Figure 4.4 a) XPS survey spectra of SDC and PNSDC samples; b) XPS O 1s

Spectra of SDC and PNSDC samples; c) XPS Ce 3d Spectra of SDC and PNSDC

samples; d) XPS Pr 3d Spectra of PNSDC sample.

migration of lattice oxygen species from the internal sites to the surface regions

[118]. The Ce 3d spectra of SDC and PNSDC are shown in Figure 4.4c.

Both spectra for the two samples give information that the Ce3+ and Ce4+ ions are

coexistence on the surface. Generally, the Ce 3d core-level spectrum can be

deconvolved into several different spin-split doublets aiming at analyzing the

fraction for the oxidation states of Ce3+/Ce4+ ions and thus in-depth understanding

the effects of doping ions, i.e. Sm3+/Pr3+/Nd3+, on the surface properties of PNSDC

[119,120]. According to the spectra analysis on Figure 4.4c, the percentage

composition of Ce3+ and Ce4+ species is calculated roughly using the deconvolution

peaks relevant to (v0, v′, u0, u′) and (v, v′′, v′′′, u, u′′, u′′′), respectively. The

estimated results exhibit an increasing Ce3+ ion concentration for 2% higher than

that of SDC. It is well known that the low-valance Ce3+ ions are closely related to

the oxygen vacancy, and thus a higher oxygen vacancy concentration is expected

for the PNSDC sample, especially on the surface of the sample [121,122]. As


35

reported previously [123,124], the Pr 3d spectra could be deconvoluted into the

distinguishable peaks, mirroring the information associated with the Pr3+/Pr4+

oxidation states. From Figure 4.4d, two featuring peaks for Pr4+ ions, i.e. 3d5/2 and

3d3/2, are detected respectively at surrounding 931.6 and 950.2 eV, while the

characteristic peak pairs for Pr3+ ions are observed at around 933.3/928.4 eV and

953.1/947.2 eV for 3d5/2 and 3d3/2, respectively. As observed, the PNSDC sample

has coexistence of Pr3+/Pr4+ oxidation states with much larger distinguishing peaks

of Pr3+ ions than that of Pr4+ ions, confirming a dominant role of Pr3+.

4.2 Electrical properties

EIS is carried out here to analyze the electrical property of the SDC and PNSDC

samples. In general, the intrinsic charge carriers play a dominant role in the

conduction of ionic conductor measured in air atmosphere. In a typical ion-

conducting case, the grain, grain boundary and electrode polarizations make main

contributions to the EIS including one arc at high frequency region, one arc at

middle frequency region and a following tail at low frequencies, respectively. In

some practical measurements, the high-frequency arcs might not show up due to

the not enough high frequency limited by the testing setup or operating temperature

[125].

Figure 4.5 shows the comparative Nyquist plots for SDC and PNSDC. All the plots

for both samples display a middle-frequency arc referring to grain boundary

resistance and low-frequency tail representing the electrode polarization processes.

Both high-frequency arcs fail to appear due to the limited frequency range of the

testing equipment. It is obvious that all the sizes of EIS decrease as the applied

temperature increases. ZVIEW software is employed to fit these EIS results. As

fitting data, the grain capacitors have a value level of around 10-12 F. The capacitors

for grain boundaries range from 10-10 F to 10-8 F, while the capacitors for electrode

polarization exhibit values approximately 10-4 F for both SDC and PNSDC samples.

These results are consistent with the values reported previously [126,127]. In

contrast, the PNSDC sample has a sharply decreasing grain-boundary arc at the

various temperature as compared to SDC. Figure 4.6a presents the temperature-

dependent conductivities of Arrhenius plots for PNSDC calculated by the EIS

measurements and the comparative results for SDC are also provided. As estimated,

the PNSDC sample exhibits an enhanced conductivity as 0.125 S cm-1 at 600 ,

while the conductivity of SDC is 0.0114 S cm-1 at the same temperature. The

activation energy is determined according to the Arrhenius plots: 0.41 eV for

PNSDC, lower than that of SDC, 0.65 eV. These results are comparable with the

reported data in the literature [128]. The different conduction mechanisms for


36

Figure 4.5 Electrochemical impedance spectra of SDC and PNSDC measured in

air at a) 400 ; b) 450 ; c) 500 ; d) 550 ; e) 600 and f) the empirical

equivalent circuit to analysis the ionic conductivity of SDC and PNSDC materials

in air.

PNSDC material could be speculated that the reasons are the significant

improvements in the aspects of activation energy and conductivity compared to

SDC.

To identify the possible electronic conduction of these doped ceria samples, the

Hebb-Wagner ion-blocking technique is employed here with an applied potential

range of 0.01-1.0 V. During the measuring process, the mobility of oxide ions and

electrons contribute to the current behaviors at the initial polarization stage. Due to


37

the ion transfer blockage, the chemical potential exhibits a rapidly decreasing trend

at the interfaces of Au and PNSDC sample as the testing time prolongs. The

electronic conduction property for PNSDC sample is presented in Figure 4.6b in

terms of current versus time (I-t) curves at the cooling temperature of 600, 550, 500

to 450 . At 600 , the electronic conductivity is approximately determined as

6.4×10-4 S cm-1, which is negligible compared to the value of total conductivity as

0.125 S cm-1. In conclusion, the enhanced conduction of PNSDC is attributed

primarily to ion transport, especially via conductivity of grain boundaries, and the

Pr3+/Nd3+ doping is favorable for surface ion incorporation to result in an improved

conductivity as well lower activation energy.

Figure 4.6 a) Temperature dependence of total conductivity and Arrhenius plots

for SDC and PNSDC. The inset represents area specific resistance (ASR) versus

temperature of SDC and PNSDC; b) current (I) versus time (t) curves obtained at

a temperature range of 450 and 600 at an interval of 50 to determine the

electronic conductivity of PNSDC sample.

4.3 Fuel cell performance

The fuel cell performance for cell I (SDC electrolyte) and cell II (PNSDC

electrolyte) are displayed in Figure 4.7a. The OCV for SDC cell is 0.922 V while

that of PNSDC device reaches up to 1.021 V at 550 , indicating minimal

electronic leakage for the PNSDC electrolyte. Correspondingly, the maximum

power density of cell II is 710 mW cm-2, which is much higher than the power

density of 490 mW cm-2 for SDC device. The evident enhancement can be

attributed to the reduced resistances of PNSDC electrolyte and PNSDC/NCAL

composite electrode resistances. Good durability is recorded for cell II under

operating conditions of 80 mA cm-2 discharge current density at 530 during 13

h, as shown in Figure 4.7b. Figure 4.8 presents the EIS results for analyzing the


38

Figure 4.7 a) Typical I-V and I-P curves of fuel cells based on SDC and PNSDC

electrolyte, respectively, the mixed NCAL and electrolyte materials were used as

the electrode material. The measurements were performed at 550 with H2 and

air as the fuel and the oxidant, respectively; b) Durability test based on Cell II:

Ni-foam/PNSDC+NCAL (anode)|PNSDC (electrolyte)| PNSDC+NCAL

(cathode)/Ni- foam at a constant discharge current density 80 mA cm-2 operated

at 530 .

Figure 4.8 Electrochemical impedance spectra of the symmetric cell using SDC

and PNSDC samples as electrolyte measured under a) air and b) H2/air

conditions at 550 .

fuel cell kinetic behaviors for both cells I and II measured at 550 . An empirical

equivalent mode, denoted as Ro(R1-CPE1) (R2-CPE2), is used here to represent the

O2- mobility during the fuel cell process [129]. As shown for both cases of the

applied air and fuel cell conditions, the ohmic resistance as determined from the

high-frequency intercept (Ro) is markedly reduced for the PNSDC device as

compared to the SDC device, which can be linked to enhanced oxygen ion

conduction for PNSDC [130]. Simultaneously, the oxygen ion migration from the


39

Figure 4.9 Schematic diagrams representing a) the overall O2‐ transfer with the

assistance of reduction conversation of Pr3+/Pr4+ ions in the fuel cell; b) the

cathode reaction process and c) brief routes for the O2‐ surface and bulk transfer

and conduction process in PNSDC electrolyte.

cathode through the electrolyte is improved, thus inducing a reduced electrode

polarization loss. The improved surface O2- conductivity is associated with the

reducing grain boundary of PNSDC. Figure 4.9 shows the schematic illustration

exhibiting the transfer routes of oxygen species in the PNSDC electrolyte. As

shown, the oxygen species are transferred via a redox process relying on the ion

chains [131,132]. During the redox reaction process, the O2 molecules from applied

air are initially activated and then reacted with the electrons from the connecting

external circuit. Furthermore, the electrons exchange between redox couple of


40

Pr3+/Pr4+ ions can facilitate the O2- transport in PNSDC [131,132]. This process is

described using the following formula:

Pri3++Prj

4+↔Pri4++Prj

3+ (4.1)

At the anode, the Pr4+ ions are reduced to Pr3+ by H2. Meanwhile, the electrons

escape from the reaction sites with a specific rate and their mobility is limited

largely by the intrinsic defects of PNSDC. At the cathode, the trivalent Pr ions

would be re-oxidized to tetravalent state and an entire redox couple, i.e. Pr3+/Pr4+,

is realized, where oxygen ions are involved selectively [131,132]. In this case, a

high OCV of 1.021 V at 550 is obtained arising from the characteristic redox

couple Pr3+/Pr4+ against the possible reduction of Ce4+ in this triple-doped ceria

material. To further clarify the mechanism related to the enhanced redox process

through the variable valence of single-Pr ion, a surface Pr-doping solely ceria is

developed to serve as the electrolyte of SOFCs, performing fuel cell measurement

for comparison (in Chapter 5).

As reported in the literature [133, 134], the dominant polarization loss in the low

operating SOFC system is the ORR process, and the improved cathode ORR

process for PNSDC is favorable for improved power output in the SOFC device.

The generated O2- ions at the cathode are easily transferred via both bulk and

surface routes of PNSDC in cell II. Thus, the efficiency of ion migration is

enhanced and fuel cell performance is improved. As reported in a recent study [129],

the lithiated transition metal oxide confirmed good catalytic properties for HOR as

anode and ORR as the cathode. In this study, the NCAL is employed as both the

anode and cathode to reinforce the HOR and ORR properties and the combining

PNSDC or SDC phases can extend the triple phase boundary interfaces, resulting

in the enhanced electrochemical performances [135–138]. Furthermore, the

polarization losses from both electrolyte and electrode can be alleviated owing to

the excellent electrochemical catalysis of Pr ions. Therefore, the SOFC device

using PNSDC demonstrates superior electrochemical performance.

4.4 Chapter summary

An effective approach is developed to significantly improve the oxygen ionic

conductivity by more than one order of magnitude while minimizing the redox

instability for the doped ceria-based materials. A two-step co-precipitation strategy

is employed here to realize the varied doping sites of ceria: Sm3+ in bulk and

Pr3+/Nd3+ on the surface. The grain-boundary resistance of PNSDC sample has been

verified largely decreasing with a conductivity of 0.125 S cm-1 at 600 .

Additionally, the Pr3+/Nd3+ ions are found to facilitate the surface oxygen species


41

exchange and to reinforce the electrochemical kinetic processes, thus improving

the entire fuel cell performance. In this investigation, the PNSDC-based SOFC

device exhibits excellent performance, e.g. a peak power density of 710 mW cm-2

at the operating temperature of 550 .


42


43

Chapter 5

Multi-component rare-earth doped ceria for EFFC

In this chapter, single-, double- and triple-doped approaches are proposed for the

modification of doped ceria materials aiming at applications for EFFC technology.

Single-phase Pr-doped ceria with mixed ionic and electronic conduction is

investigated as a core component of EFFC devices. Double- and triple-doped ceria

materials combined with semiconductor served as homogenous semiconductor-

ionic composites for EFFCs. Fundamental material properties and fuel cell

performance related these materials are studied systematically. Both experimental

results and data analysis are included, as reported in Paper II, III, V and VI.

5.1 Single-phase Pr-doped ceria for EFFC

5.1.1 Structural and morphology characteristics

The XRD patterns of the Pr-CeO2 crystalline powder synthesized via the

hydrothermal method in comparison with Pr-free CeO2 are shown in Figure 5.1.

Results for the two samples are related to a cubic fluorite structure with a space

group of 𝐹𝑚3𝑚 (Ref. standard CeO2 JCPDS card: NO.04-0593). A slight left shift

is observed in the Pr-CeO2. In addition, the sharp diffraction peaks indicate the

presence of nanoparticles with regular crystallinity. No other diffraction peaks of

impurities are detected within the detectability of XRD, demonstrating that a single-

phase Pr-CeO2 sample with a nominal composition of Pr0.1Ce0.9O2-δ is synthesized.

SEM and TEM images are performed to reveal overall particle morphology and

qualitatively investigate the size-distribution for the synthesized CeO2 and Pr-CeO2

particles. Figure 5.2 illustrates the representative SEM and TEM images of as-

synthesized CeO2 (a and b) and Pr-CeO2 (c and d) samples. As shown, Pr-free CeO2


44

powder is comprised of aggregated particles to form rod crystallites. However, the

Pr-doped CeO2 exhibits homogeneous particles. The small nano-crystal of the Pr-

CeO2 sample are expected to improve fuel cell performance in comparison to that

of CeO2 owing to a higher specific surface area.

Figure 5.1 X-ray diffraction patterns of as-synthesized Pr-free CeO2 and Pr-CeO2

samples.

5.1.2 XPS analysis

XPS analysis is carried out to distinguish the chemical bonding states of elements

on the surface of as-prepared CeO2 and Pr-CeO2 samples. The obtained spectra and

corresponding fitting results are presented in Figure 5.3. The survey spectra in

Figure 5.3a clearly reveal the characteristics of C, O, Ce and Pr elements. To

further identify the chemical states of the relevant elements, the O 1s, Ce 3d and Pr

3d core-level spectra are addressed using XPSPEAK 4.1 software. Figure 5.3b

illustrates the Pr 3d spectra. The characteristic peak pairs of Pr3+ 3d5/2 (located at

around 928/933 eV) and 3d3/2 (located at around 947/956 eV) are observed in the

fitting spectra. The peaks of ~931 and 952 eV can be assigned to Pr3+ 3d5/2 and 3d3/2,


45

Figure 5.2 Representative SEM micrographs for as-synthesized a) CeO2 and c)

Pr-CeO2 samples; TEM images of b) CeO2 and d) Pr-CeO2 samples suffering

hydrothermal process.

respectively. The fitting results demonstrate that both the oxidation states of Pr3+

and Pr4+ are detected in the Pr-CeO2.

As reported in the literature [139,140], the reduction of Pr4+ to Pr3+ can create

oxygen vacancies. In this case, the mixed oxidation states of Pr3+ and Pr4+ can

promote the ionic conductivity and facilitate to enhance the electrochemical

performance for Pr-CeO2. Figure 5.3 c and d show the O1s core-level spectra of

CeO2 and Pr-CeO2, respectively, with the two feature peaks relevant to the lattice

oxygen (Oα) with the binding energy range of 528-530 eV and chemically absorbed

surface oxygen species or oxygen defects (Oβ) located at around 530-533 eV[117].

Generally, the Ce-O-Pr bonds can be formed once Pr ions are incorporated into the

CeO2 lattices. Subsequently, the electro-negativity has slight differences for Pr and

Ce ions, and more changeable oxygen leads to the facilitated migrations of oxygen

from internal sites to the surface[118]. As calculated, the Oβ/Oα value of Pr-CeO2


46

Figure 5.3 XPS spectroscopy: a) survey spectra of CeO2 and Pr-CeO2 samples;

b) Pr 3d spectra for Pr-CeO2; c) O 1s spectra of CeO2; d) O 1s spectra of Pr-

CeO2; e) Ce 3d spectra of CeO2; and f) Ce 3d spectra of Pr-CeO2.

(~1.24) is higher than that of CeO2 (~1.21), indicating an increase of surface oxygen

vacancies in Pr-CeO2 sample. The Ce3+ and Ce4+ ions are detected in coexistence

on the surface of both CeO2 and Pr-CeO2, as presented in Figure 5.3 e and f. The

spectra deconvolution with several peaks in CeO2 and Pr-CeO2 are assigned to the


47

Ce 3d5/2 (vo, v and v’) and 3d3/2 (uo, u, u’ and u’’) [118,141]. In this case, the peak

groups of (uo, vo) and (u, u’ u’, u’’, v, v’) are attributed to Ce3+ and Ce4+ species,

respectively [108,118]. It can be deduced from the resultant spectra that Ce3+ and

Ce4+ ions coexist on the surface of CeO2 and Pr-CeO2.

5.1.3 Raman scattering analysis

Figure 5.4 Raman Spectroscopy of CeO2 and Pr-doped CeO2 samples using an

excitation wavelength of 325 nm.

The Raman reflectance spectra with an excitation laser wavelength of 325 nm for

CeO2 and Pr-CeO2 samples are presented in Figure 5.4. The dominant Raman

active mode is observed at around 460 cm-1 for the CeO2 sample, while both this

mode and a second one at 554 cm-1 are seen for the Pr-CeO2 sample. The peak at

around 460 cm-1 is attributed to the symmetrical oxygen lattice Ce-O stretching

vibrational mode (F2g) in a cubic fluorite-type CeO2 lattice [140,142,143], which

exhibits a decreasing energy intensity for Pr-CeO2 compared to that of CeO2.

Several factors, e.g. strain, defects, and phonon confinement, can induce the

variation of F2g peak. In this case, the slight shift in the F2g mode of Pr-CeO2 results

from the larger ionic radius of Pr3+ ions (1.013 Å) as compared to host Ce4+ ions


48

(0.92 Å) [144]. Additionally, the oxygen vacancies also contribute to the shift of

F2g peak [145]. The larger ionic radius of Pr3+ (as compared to Ce4+) contributes to

the lattice expansion [146]. The emerging feature peak, at around 554 cm-1 for the

Pr-CeO2 sample is indexed to the formation of oxygen vacancies induced by the

substitution of Pr3+ in CeO2 surface lattice, while such defect-related peaks are not

observed in Pr-free CeO2 sample. This finding indicates that the defects are

generated due to the introduction of Pr ions into the ceria crystal structure,

compared with the CeO2 spectrum without the defect mode [146]. The oxygen

vacancies are formed in the CeO2 lattice, similar to Gd-doped CeO2, by doping with

Pr3+ cations:

Pr2O32CeO2→ 2PrCe

′ + VO·· + 3OO

x (5.1)

As shown in Equation 5.1, two Pr3+ cations substitute two Ce4+ ions to generate an

oxygen vacancy to balance the charge based on the charge compensation

mechanism [147,148]. It has been reported that the asymmetry on the high-

frequency side of the Raman line in the spectrum for Pr-CeO2 is contributed from

the existence of a probable surface mode at 480 cm-1 [148]. This is ascribed to the

large defective structure and disorder in the ceria lattice.

Figure 5.5 a) Nyquist plots for Pr-CeO2 sample measured at 600 under air

atmosphere; b) Arrhenius plot and corresponding temperature dependence of

total conductivity for Pr-CeO2 sample.

5.1.4 Electrical conductivity

The electrical behaviors of the as-prepared Pr-CeO2 sample are examined by EIS

in air atmosphere operated at a temperature range between 500 and 600 with


49

an interval of 25 . Figure 5.5a demonstrates the typical Nyquist plot at 600 . It

is well known that the electrical properties are well correlated to the operating gas

atmosphere. Generally, the intrinsic charge carrier is dominant to the conductivity

in the air environment [125]. In Figure 5.5a, the Nyquist plots of Pr-CeO2 sample

demonstrate a typical medium frequency arc and low-frequency tail. In this case,

the arcs in the high-frequency range could not be detected. Prior to further EIS

analysis, a possible electronic conductivity needs to be considered since the

existence of a Pr impurity band within band gap of ceria facilitates electron hopping

in the Pr-CeO2 structure [149,150]. The estimated electronic conductivity of Pr-

CeO2 using Hebb-Wagner ion-blocking technique is approximately 0.022 S·cm-1 at

600 .

To understand the electrochemical kinetic behavior in detail, EIS simulations for

impedance plots are conducted using Zview software package. The empirical

equivalent circuit model of R0(R1-QPE1)(R2-QPE2), as shown in the inset of Figure

5.5a, is employed. In this equivalent model, R0 is assigned to ionic conduction

indexing to the intercept of the EIS curve with the Z’-axis at high-frequency range

(>104 Hz). R1 is related to the oxygen transfer at the Ag/Pr-CeO2 interface under

the frequency range of 102-104 Hz. The oxidation/reduction of oxygen processes

are associated with the resistance of R2 in the low-frequency range of 10-2-102 Hz.

QPEs are the constant phase element to express the ‘depressed’ arc in consideration

of any inhomogeneity in the measured material [150–152]. The total conductivity

values for Pr-CeO2 are estimated as 0.36, 0.16 and 0.037 S cm-1 at 600, 550 and

500 , respectively. Compared with the above-calculated results, the electronic

conductivity (~0.022 S·cm-1) is one order of magnitude lower than the ionic

conductivity (~0.338 S·cm-1) for Pr-CeO2 at 600 . The temperature-dependent

conductivities for Pr-CeO2 sample are presented in Figure 5.5b in the form of an

Arrhenius plot. The activation energy of Pr-CeO2 is estimated to be 1.3 eV,

according to the Arrhenius relationship:

σ = (σ0/T)exp(−Ea/kBT) (5.2)

where σ and σo refer to the conductivity and pre-exponential constant, respectively,

Ea is the activation energy, kB is the Boltzmann constant and T is the operating

temperature.

The unexpected high conductivity of Pr-CeO2 could be contributed to two

dominating factors: i) rapid oxygen ion transfer and surface exchange resulting

from enhanced surface oxygen vacancy concentration; generally, the oxygen

species can be quickly absorbed and transferred through oxygen vacancies, so


50

higher ionic conductivity can be expected with more vacancies; ii) promoted

electronic conduction arising from impurity Pr band, acting as an electron

conduction channel, within band gap of ceria and faster oxygen exchange [152,153].

In parallel, compared to the purely ionic conductors, the mobile electronic carriers

benefit to the catalytic activity on the surface of materials [152,154]. It can be

speculated that the electron conduction and ion transport have a strong correlation

to tremendously enhance the coordinated conductivity instead of the simple sum of

both contributions. It has verified that the ionic conductivity was significantly

enhanced by compositing with semiconductors, insulating conductors, or with

carbonates [36,155]. Analogously, the single-phase materials with electronic and

ionic conduction have similar synergistically enhanced conductivity to those two-

or three-phase composites.

5.1.5 Fuel cell performance analysis

Figure 5.6 I-V-P characteristics of fuel cell devices based on a) the CeO2 and Pr-

CeO2 membrane separately measured at the temperature of 550 ; b) Pr-CeO2

operated at various temperatures of 500, 525, 550 and 600 , respectively.

The current-voltage (I-V) and current-power (I-P) characteristics for symmetrical

fuel cells using CeO2 (as a comparison) and Pr-CeO2 as the membrane layer,

respectively, are presented in Figure 5.6a. It can be observed that both OCVs for

the fuel cells reach up to 1.0 V at 550 . The high OCV values indicate that the

fabricated devices realized the fuel cell function without electronic current leakage.

At this operating temperature, an appreciable peak power density of 744 mW cm-2

is achieved for the EFFC device using Pr-CeO2, which is superior to that of CeO2

of 497 mW cm-2. Such enhancement of the power output reflects a higher

conductivity of the Pr-CeO2 electrolyte, compared to the CeO2 electrolyte.


51

Figure 5.7 Schematic diagrams representing a brief route for the O2- mobility in

fuel cell configuration and detailed O2- transfer with the assistance of redox

Pr3+/Pr 4+ couple on the surface and bulk of Pr-CeO2 electrolyte in the vicinity of

oxygen applied side.

To further investigate the property for the Pr-CeO2 sample, the fuel cell

measurements are performed at various operational temperature, as shown in

Figure 5.6b. It demonstrated that the maximum power densities of 776 and 471

mW cm-2 are obtained at 600 and 525 , respectively. Significantly, even the

operating temperature down to 500 , a considerable peak power density of 296

mW cm-2 is still obtained.

The relatively high output power density for Pr-CeO2 is based on the following

factors: i) enhanced ionic conductivity of Pr-CeO2 membrane via Pr surface dopant;

ii) promoted ion conduction strongly correlated to electronic conductivity due to

electron hopping between neighboring Pr ions via impurity band conduction under

low oxygen partial pressure; iii) improved electron transfer under high oxygen

partial pressure owing to the redox Pr3+/Pr4+ pair. As reported by Zhu et al. [34], a

thin enriched nickel metal layer (≈ 4-6 μm) formed in the LiNi0.85Co0.15O2-δ layer

on the hydrogen-supplied side with the reduction of LiNi0.85Co0.15O2-δ. On the H2-

contacting side, a Schottky junction was verified to be established by supplying

external voltage and exhibiting a typical rectification effect in such a device. In this

study, the high OCV and good fuel cell output power density indicate that a similar

phenomenon occurs and Schottky junction is established in the H2-contacting side


52

to prevent short-circuiting. The schematic diagrams representing a brief route for

the O2- mobility in fuel cell configuration and detailed O2- transfer with the

assistance of redox Pr3+/Pr4+ couple on the surface and bulk of Pr-CeO2 in the

vicinity of oxygen applied side are presented in Figure 5.7. As shown, at the high

oxygen partial pressure region (O2-contacting side), Pr largely remains as

tetravalent Pr4+ and can be reduced by electrons from the external route during fuel

cell process. O2- ions can be transported relying on redox Pr3+/Pr4+ couple and

oxygen vacancies via bulk and surface routes. With reducing of oxygen partial

pressure, oxygen is released from the lattice with converting Pr4+ to Pr3+ and oxygen

species can be transferred flexibly through the dominative Pr-doped CeO2

membrane. As mentioned above, the fuel cell using Pr-CeO2 membrane is

significantly improved in comparison to that of Pr-free CeO2.

5.2 Rare-earth triple-doped ceria for EFFC

Figure 5.8 SEM micrographs for a) raw material and b) as-prepared LCPN; c)

TEM microstructure and corresponding selected area electron diffraction pattern

(SAED, inset image) for LCPN sample.

Microstructures of the raw material and as-obtained LCPN sample are presented in

Figure 5.8 by means of SEM and TEM images. As shown in the SEM image for

the raw rare-earth material, regular particles are shown to be agglomerated into a

range of micrometer sizes. After calcination, the LCPN sample displays the

relatively homogeneous particles with an estimated size of 40 nm. The

polycrystalline structure of the as-obtained LCPN powder is confirmed by the

SAED pattern, see inset in Figure 5.8c. To further underline the elemental

compositions, EDS experiments are conducted on the raw material and treated

LCPN sample, see Figure 5.9. During the heat-treatment process, the raw rare-

earth carbonates are reacted into rare-earth oxides with a decreasing carbon


53

composition and increasing oxygen species. The rare-earth elements (i.e. Ce, La,

Pr, Nd) and oxygen are confirmed as the dominating compositions of the obtained

LCPN material from the EDS analysis.

Figure 5.9 Typical SEM micrographs of a) raw material and c) LCPN; b) and d)

identified the EDS data with the representative counts of the vertical axis

obtained at the yellow rectangular areas.

Figure 5.10 XRD patterns of LCPN and NCAL samples.


54

Figure 5.10 exhibits the XRD results for the LCPN and NCAL. LCPN pattern can

be assigned to the cubic-fluorite ceria with a slight left-shift of the diffraction peaks

owing to the elements of La3+/Pr3+/Nd3+ with larger ionic radii compared to the Ce4+

ions. No secondary phase is detected by the XRD measurement. The commercial

NCAL is analyzed as a layer structure, indexing to the diffraction peaks of

LiNi0.8Co0.2-oxide (JCPDS, No. 87-1562) [32]. The aluminum element is not

detected due to the limitation of testing [33].

Figure 5.11a shows the electrochemical properties of the device using

LCPN/NCAL composite with a weight ratio of 7:3 measured at a temperature range

of 550-600 at intervals of 25 . All the three EIS patterns demonstrate two

components: a depressed small semicircle at the high-frequency region and another

half-arc at low-frequency range. The equivalent model (see the inset image of

Figure 5.11a) is expressed as Ro/R1(CPE1)/R2(CPE2) and is employed to analyze

the EIS results. The intercept at high frequencies represents the interfacial

resistance of charge transfer, denoted as R1, while R2 can be ascribed to the

resistance of mass transfer, indexing to the low-frequency semicircle [27]. In

general, the reaction kinetics associated with charge and mass transport can be

improved by increasing the operating temperature, which is consistent with the

measured results of Figure 5.11a. Furthermore, the EFFC devices using the

LCPN/NCAL composites with various weight ratios, i.e. 7:3, 6:4, and 3:7 for LCPN

and NCAL are investigated to select the optimum proportion of ion-conducting

phase and electron-conducting phase, as shown in Figure 5.11b. All the devices

are tested at 550 fed by air and hydrogen, exhibiting a similar characteristic EIS

curve. In contrast, the polarization resistances are reducing with the increasing

contents of the electronic phase, i.e. NCAL, reflecting a reduced semicircle in EIS

plots. The increase in electron conductor NCAL might cause possible electrical

leakage. For instance, some leakage might happen for the device using the 6LCPN-

4NCAL composite, exhibiting a slight lower OCV value compared to the pure

LCPN and 7LCPN-3NCAL devices, as shown in Figure 5.12a. According to the

fitting results of EIS, the resistance of charge transfer for the device using 7LCPN-

3NCAL is approximately 0.46 Ω cm2, while it is only 0.23 Ω cm2 for 3LCPN-

7NCAL. However, the resistances for three devices using LCPN/NCAL

composites are reduced due to the additional NCAL, indicating favorable kinetic

reactions for these devices.


55

Figure 5.11 EIS patterns of a) the 7LCPN-3NCAL composite at different

temperatures and b) the different compositions of LCPN and NCAL (7LCPN-

3NCAL, 6LCPN-4NCAL, and 3LCPN-7NCAL) under air/H2 condition at 550 .

All the measurements were performed under air/H2 condition under open circuit

conditions.

Fuel cell performance is investigated based on these fuel cell devices using various

LCPN/NCAL composites and pure LCPN, see Figure 5.12a. The device using a

pure LCPN as the electrolyte displays a high power density of 1,046 mW cm-2

measured at 550 . It indicates that the LCPN material can serve as a good

electrolyte candidate with high ionic conductivity. After composition with a

semiconductor, e.g. 7LCPN-3NCAL composite, the fuel cell device exhibits the

maximum power output (Pmax), reaching 1,187 mW cm-2 at 550 and the OCV

reaches as high as 1.07 V among others. If the electron-conducting phase is

increased to 40 wt.% in the LCPN/NCAL composite, a relatively low power density

is obtained. Further deterioration in performance is exhibited as the NCAL fraction

increases, resulting in a poor OCV (0.4 V) for the 70 wt.% NCAL sample. These

findings indicate that good EFFC performance is dependent upon proper balancing

between the ion-conducting phase, i.e. LCPN, and electron-conducting phase, i.e.

NCAL. It has been verified that a Schottky junction could be established on the

hydrogen supply side, which can block electron flow through the internal layer of

the device [1,15]. Figure 5.12b presents the various power outputs and OCV values

for the devices using different weight ratios of LCPN/NCAL composites. As shown,

a decreasing trend for the OCV of these devices is found with the increase of NCAL

composition. For example, a high OCV of 1.18 V is obtained when the pure LCPN

is used as the electrolyte of SOFC. The device using the LCPN with an addition of

30 wt.% NCAL exhibits an OCV of 1.07 V, while the one using 5LCPN-5NCAL


56

composite displays the OCV for approximately 0.92 V. Indeed, the 7LCPN-

3NCAL composition has a favorable OCV value, and it also shows the best

performance in high sustained power density across a wide range of loads. The

enlarged triple phase boundaries of LCPN/NCAL composite with concomitant

lower polarization loss contributes to the enhanced fuel cell performance.

Figure 5.12 a) The I-V (current density and voltage) and I-P (current density and

peak power density) characteristics of various compositions of LCPN-NCAL

feeding with hydrogen and air at 550 ; b) OCV and maximum power density

values (Pmax) for fuel cells based on LCPN-NCAL in various compositions.

Figure 5.13 Schematic illustration (left) and SEM micrograph (right) of EFFC

device based on the 7LCPN-3NCAL composite after the fuel cell performance

measurements.


57

Figure 5.13 displays the cross-section image for the tested EFFC device using

7LCPN-3NCAL composite and corresponding schematic sketch. As observed,

some structural changes are found in the hydrogen-contacting regions, which could

be attributed to the reduction of NCAL to yield nickel or cobalt metals. Based on

these formed metals, a Schottky junction is built inside the LCPN/NCAL composite,

via contacting with p-type NCAL semiconductor, as verified in literature [15,17,18].

From the SEM image of the EFFC device, both NCAL layers retain a dense

structure even after fuel cell testing.

5.3 Double-doped ceria for EFFC

Figure 5.14 a) I-V and I-P characteristics and b) EIS of EFFC devices using the

SCDC/LSCF composites in various weight ratios; c) I-V and I-P performance for

the conventional SOFC device and the device using pure SCDC membrane.

The typical I-V and I-P curves for EFFC devices with the configuration Ni-

NCAL/LSCF-SCDC/Ni-NCAL based on the various weight ratios of LSCF/SCDC

are presented in Figure 5.14a. All of these fuel cell devices are tested at 550 .

As observed, the best performance is obtained for a weight ratio of 1:1 (LSCF and


58

SCDC) with a maximum power density of 814 mW cm-2. Figure 5.14b shows the

EIS curves of the three fuel cells under H2/air condition in open circuit mode. As

shown, the entire polarization losses as LSCF content increased. The comparative

fuel cell performance for the conventional SOFC device and the device using pure

SCDC membrane are shown in Figure 5.14c. As shown, the conventional device

with the symmetrical configuration of NCAL/SCDC/NCAL, showed a 20% higher

power density (e.g. 368 mW cm-2) as compared to the conventional FC device using

SCDC as electrolyte and LSCF and NCAL as cathode and anode respectively (e.g.

300 mW cm-2).

5.4 Fuel cell performance comparison

Table 5.1 Fuel cell performance comparison based on three types of doped ceria

materials.

Core materials Fuel cell configuration Pmax (mW

cm-2)

Temperature

()

Paper

number

Pr-doped ceria

(single-doped

ceria)

Ni-NCAL/Pr-doped

ceria/Ni-NCAL 744 550 II

LSCF-SCDC

composite

(in a weight

ratio 1:1)

Ni-NCAL/50LSCF-

50SCDC/Ni-NCAL

814

550

V+VI

LCPN-NCAL

composite

(in a weight

ratio 7:3)

Ni-NCAL/7LCPN-

3NCAL/Ni-NCAL

1,187

550

III

In this chapter, three different types of materials are used to investigate the fuel cell

performance based on the single-, double- and triple-doped ceria materials. Table

5.1 shows a comparison of maximum power density obtained from each of these

materials for FC devices operated at 550 . Single-element Pr-doped ceria exhibits


59

mixed electronic/ionic conducting properties, enabling it to be used as the core

membrane of an EFFC device with semiconducting NCAL as anode/cathode layers.

The Sm/Ca co-doped ceria is used for EFFC by combining with semiconductor

LSCF, exhibiting a comparable peak power density (814 mW cm-2 compared to

744 mW cm-2). The device using La/Pr/Nd triple-doped ceria and NCAL

nanocomposite demonstrates the highest maximum power density 1,187 mW cm-2,

an improvement of 45% or more compared to other devices.

5.5 Chapter summary

Detailed characterization studies are conducted to investigate the superficial and

structural properties of the Pr-doped CeO2 material. Additionally, the electrical

conductivity and fuel cell performance using Pr-doped CeO2 are measured to

indicate the electrochemical properties. The results exhibit that the Pr-dopant,

primary as Pr3+ ions, are selectively distributed on the surface and thus made

positive contributions to surface oxygen ion exchange and electron conduction,

which are later found to enhance fuel cell performance.

For double- and triple-doped cases, the effect of various weight ratios of ionic

conductor LCPN and semiconductor NCAL on the EFFCs performance is

investigated. The optimal sample, i.e. 7LCPN-3NCAL, exhibits a considerable fuel

cell performance with 1,187 mW cm-2 at 550 . Therefore, the electron conduction

might balance well with the ion conductivity for this LCPN/NCAL composite with

30 wt. % NCAL. The high OCV and its fast charge transfer properties demonstrate

that the mixed semiconductor (NCAL) and ionic conductor (LCPN) composite has

the good catalytic capacity to function EFFC performance well.


60


61

Chapter 6

Natural minerals for EFFC

A natural mineral originating from the natural chalcopyrite ore (written as CuFeS2)

is considered for applications in LT-SOFCs. Three types of fuel cell devices are

fabricated aiming at exploring the optimum option for the application of CuFe2O4.

EIS is employed as the principal means to reveal possible underlying working

mechanisms for the different devices. These investigations on natural minerals have

been published in Paper IV.

6.1 Material characterization analysis

Figure 6.1 displays the XRD results of untreated natural mineral and the heat-

treated samples, respectively, at 600, 700, and 900 . The main constituent of the

natural material in Figure 6.1a relates to a chalcopyrite structure (i.e. CuFeS2)

according to JCPDS NO.37-0471. Some other phases, such as FeS2, Fe2O3, CaCO3,

and SiO2, are also found in the raw mineral (as listed in Table 6.1). After heat

treatment (calcined at 600 for 2 h), the chalcopyrite structure is destroyed, see

Figure 6.1b. When elevating the temperature to 900 , some diffraction peaks are

indexed to a face-centered cubic CuFe2O4 structure (JCPDS NO. 25-0283). In this

sample, some other phases, e.g. CaSO4 and CaO, are also detected. The structural

changes for this natural mineral could be attributed to the chemical reactions during

the heating process. As reported previously, the impurity phases of the as-obtained

sample are favorable for ion transportation to improve ionic conductivity by

combining to form composites and induce interfacial conduction [156].

Figure 6.2a and 2b display the microstructures for the raw chalcopyrite ore and the

as-sintered delafossite, respectively, by means of SEM technique. As observed, the

natural chalcopyrite consists of variable bulk agglomeration in a size scale ranging


62

from micrometer to nanometer. After heat-treatment at 900 , the obtained

CuFe2O4 sample exhibits a slight smaller bulk structure. The quantitative analysis

of elements for this mineral is carried out by EDS measurement, indicating the

distinct Fe, Cu, and S peaks and some elemental peaks of Mg, Si, Al and Ca in the

raw chalcopyrite (see Figure 6.2c).

Figure 6.1 XRD patterns for a) raw material and b) the as-treated mineral

samples sintered at 600, 700 and 900 , respectively.

Figure 6.2 SEM images of a) raw mineral, b) the sample sintered at 900 , and

c) EDS image of raw mineral.


63

6.2 Electrochemical performance analysis

Figure 6.3a demonstrates the EIS results for the three samples under air

atmosphere. As observed, all of three EIS patterns display two distinct components

including a medium-frequency arc followed by a low-frequency tail. The EIS

model comparison shows that the sample obtained at 600 has much lower

resistance compared to the other two samples. It is known that CuFeS2 has the

capacity to transfer electrons, thus the undecomposed parts inside this sample

(600 ) can contribute to the entire conductivity. The middle-frequency semicircle,

generally, is indexed to the grain boundary resistance of materials. In this case, the

sample obtained at 900 exhibits a reduced grain-boundary arc, implying that the

formation of CuFe2O4 phase might facilitate to improve the conductivity of

materials. To further evaluate the electrochemical properties of the obtained sample

at 900 (mainly CuFe2O4), the EIS measurement is conducted for this sample

under fuel cell operating conditions. Figure 6.3b demonstrates an obvious reducing

EIS curve compared to the one measured in the air for the sample (900 ). This

can be attributed to the introduction of extrinsic ions, such as protons in this case,

while the electrical conduction is contributed by the intrinsic charge carriers

operated at air [104]. It is worth noting that the interfaces of this sample created by

the impurity phases are favorable for proton transportation. Therefore, the enhanced

conductivity might be caused by the proton conduction. It has been investigated

that the conductivity of an ionic conductor is improved by combining with an

insulator, such as LiFeO2 and γ-LiAlO2 [100]. In addition, the ionic conductor

LZSDC composited with the CuFe2O4-based sample is expected to enable a dual-

ion conducting (O2-/H+) function as the SOFC material, see the schematic

illustration in Figure 6.4.

Table 6.1 The main chemical composition of raw chalcopyrite.

Element MgO Al2O3 SiO2 SO3

Wt. % 6.957 3.082 21.773 28.315

Element CaO Fe2O3 CuO

Wt. % 4.247 20.840 13.959


64

Figure 6.3 a) Nyquist plots of a) the samples sintered at 600, 700 and 900

measured in air atmosphere at 550 ; b) the sample sintered at 900 measured

in H2/air at 550 .

Figure 6.4 Mechanism for the ionic conductivity of the LZSDC-CuFe2O4

composite.

Figure 6.5 includes schematic diagrams of the three different fuel cell devices, and

the I-V and I-P characteristics for the corresponding devices are demonstrated in

Figure 6.6a. For Type I, LZSDC and CuFe2O4 are employed as the electrolyte and

ORR catalyst, respectively. This device displays a quick-response OCV (1.07 V)

and corresponding peak power density of 180 mW cm-2 at 550 . Another device

using CuFe2O4 as both anode and cathode is fabricated, exhibiting a high OCV of

1.15 V. These findings indicate that CuFe2O4 sample has good catalytic activities


65

on fuel cell reactions. In Type II, the CuFe2O4 sample, composited with LZSDC, is

used as the functional cathode/anode layers, while NCAL is pasted on nickel foam

attached on both sides as the current collectors and structural supports. This device

exhibits a good power density of 354 mW cm-2 at 550 with an OCV of 0.94 V.

In contrast, the Type III device is designed to substitute LZSDC electrolyte using

the LZSDC/CuFe2O4/NCAL composite. Remarkably, the power density of Type III

is significantly enhanced up to 587 mW cm-2 at 550 , with good obtained OCV

(0.988 V). The improved fuel cell performance for Type III is attributed to the

removal of electrode/electrolyte interfaces compared to the conventional fuel cell

structure, e.g. Type I and Type II. Figure 6.6b shows an operating stability

recording for around 5 hours to preliminarily test the practical feasibility of the

Type III device. As shown, a relatively stable potential output is shown under a

steady current condition (50 mA cm-2).

Figure 6.5 Schematic illustrations for a) Type I, b) Type II and c) Type III.

To further clarify kinetic reactions, EIS measurements are carried out on the three

devices, as shown in Figure 6.7. Type I and Type II display a similar EIS model

with a half-arc at high frequency and an upturned tail at low frequency region. They

can be fitted and analyzed by means of an empirical equivalent circuit, denoted as

Ro(R1-QPE1)W1 [157]. According to these results, the high-frequency capacitance

has a level of 10-6 F for both devices. Therefore, the resistance (R1) from the

semicircle at the high frequency region can be assigned to the charge transfer

process along with the interfaces of electrodes and electrolyte [35]. The R1 values

for Type I and Type II are 0.256 Ω cm2 and 0.112 Ω cm2, respectively. The reduced

resistance loss for Type II might be caused by the appended LZSDC/CuFe2O4

functional layers. For Type III, a new EIS feature with a depressed semicircle is

displayed in Figure 6.7c. It is fitted and analyzed using the empirical circuit of


66

Ro(R1-QPE1)(R2-QPE2), which is employed generally by the perovskite solar cells

[35]. A capacitance with a level of 10-2-10-1 F is obtained, representing the

dominant processes for oxygen adsorption or dissociation, as well as diffusion.

Simultaneously, the polarization tail, indexing to the Warburg impedance in the

conventional fuel cell devices, is not observed for Type III. All the EIS fitting

results for the three devices are summarized in Table 6.2. By comparative analysis

on the three configurations, the Type III using the LZSDC/CuFe2O4/NCAL exhibits

much less polarization loss owing to the removal the interface resistance from

electrode and electrolyte in the conventional three-component devices. Based on

the analysis above, this device using CuFe2O4 combined with LZSDC and NCAL

as core components exhibits good ORR catalytic and ion transport properties, thus

providing good fuel cell performance.

Figure 6.6 a) I-V-P curves for three types of devices measured at 550 ; b)

Durability test of Type III at a constant discharge current of 50 mA cm-2.

Table 6.2 The fitting results of impedance spectra for three types of devices measured

in H2/air condition at 550 . The unit for resistance is Ω cm2.

Parameter Ro R1 QPE1-Q QPE1-

n

W1-

R W1-T

W1-

p R2

QPE2-

Q

QPE

2-n

Type I 0.67

8 0.4 5.3E-5 1 2.42 28 0.17 - - -

Type II 0.7 0.17

6 8.8E-5 0.8 9.5 11300 0.22 - - -

Type III 0.65

6

0.07

9 0.039 0.777 - - - 3.5 0.058 0.54


67

Figure 6.7 Nyquist plots for a) Type I, b) Type II and c) Type III operated at fuel

cell condition at 550 .

6.3 Chapter summary

In this study, a natural mineral material is investigated and applied successfully for

LT-SOFCs. The one using CuFe2O4 cathode exhibits a quick-response OCV,

reaching up to 1.07 V, indicating good catalytic activity of the natural mineral for

oxygen reduction. Compared to two conventional fuel cell devices, the one using a

homogeneous LZSDC/CuFe2O4/NCAL layer demonstrates a much better power

output of 587 mW cm-2 at 550 . These findings suggest that natural minerals have

great potential to be developed as new SOFC materials.


68


69

Chapter 7

Conclusions and future outlook

In this dissertation work, rare-earth doped ceria materials are employed to explore

effective approaches for desirable solid oxide fuel cell materials. Various material

synthesis techniques are investigated including two-step wet-chemical co-

precipitation method and hydrothermal method. Furthermore, a natural mineral is

investigated as an initial exploration for SOFC applications. Several conclusions

are drawn from this thesis:

i) An effective approach is developed to significantly enhance ionic

conductivity of doped ceria oxide via a novel two-step wet-chemical

co-precipitation method realizing Sm3+ doped in bulk and Pr3+/Nd3+

doped in surface domains of ceria. The modified surface property

greatly reduced the grain boundary resistance, leading to an

exceptional electrical conductivity of 0.125 S cm-1 at 600 . The

presence of Pr3+/Nd3+ also enhances the surface oxygen exchange rate

and helps to improve the redox and electrode reaction kinetics, greatly

enhancing the voltage and power efficiency. SOFC based on PNSDC

electrolyte with a thickness of 0.5 mm obtained a maximum power

density of 710 mW cm-2 at 550 . Thus, the improved electrical and

redox properties offer a new promising approach for practical

application of ceria-based materials for high performance, low

temperature SOFCs applications, which will lead to a great impact on

SOFCs development and commercialization.

ii) A single-phase Pr-doped CeO2 is synthesized via facile template-free

hydrothermal method. The detailed characterization studies are

measured to investigate the superficial and structural properties of the

Pr-doped CeO2 material. Additionally, the electrical conductivity and

fuel cell performance are conducted to indicate the electrochemical


70

properties. The results exhibit that the Pr dopants, primary as Pr3+ ions,

prefer to dominatingly distribute on the surface and make significant

contributions to surface properties, especially surface oxygen ion

exchange and electron conduction. Furthermore, more oxygen

vacancies and emerging electronic conduction presented on the surface

of Pr-doped CeO2 material should take responsibility for the enhanced

electrical conductivity and improved fuel cell performance in

comparison to Pr-free CeO2 sample.

iii) The rare-earth doped ceria materials are successfully applied for EFFC

devices by compositing with semiconductor materials. The effects of

the weight ratios for LCPN and NCAL on the electrochemical

performance of EFFC devices are systematically investigated. The

optimal device using 30 wt. % NCAL in the LCPN/NCAL composite

exhibits a good power output with 1,187 mW cm-2 at 550 oC. The

investigations indicate that it is the indispensable prerequisite to

acquire good fuel cell performance using a well-balanced electronic

phase and ionic conductor for EFFC devices.

iv) Natural mineral is heat-treated at 900 to obtain the sample with most

CuFe2O4 and then is applied for LT-SOFCs. Three fuel cell devices are

fabricated in different configurations. Among these results, the type

with the homogenous LZSDC/CuFe2O4/NCAL composite layer

exhibits the maximum power density reaching up to 587 mW cm-2 at

550 oC. The investigations reveal the promising properties of natural

mineral for SOFC applications, including i) high cathodic ORR

catalyst functions; ii) the potential to be served as the component of

EFFC devices with good fuel cell performance; iii) the different

electrochemical processes occurred at the natural mineral-based EFFC

device. Based on these investigations, the natural minerals including

the used CuFe2O4, exhibit great potentials to apply for next-generation

mineral-based SOFCs.

Further research work is suggested to study the following aspects:

• To develop Pr-doped ceria nanocomposite with other materials, e.g.

ionic conductors, insulators, even electronic conductors.

• To explore new materials of solid oxide fuel cell to be technically

useful at relatively low operating temperature, e.g. 400 .


71

• To use the tri-doped ceria for electrolyte film of SOFC to further

enhance the ionic conductivity and lower the operating temperature.

• To develop scale-up SOFC devices using the developed materials.


72


73

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