livrepository.liverpool.ac.uklivrepository.liverpool.ac.uk/19353/4/SimoFra_Aug2014_19353.pdf · 2...

Novel Oxide Materials for Solid Oxide

Fuel Cells Applications Thesis submitted in accordance with the requirements of the

University of Liverpool for the degree of Doctor in

Philosophy by:

František Šimo

Supervised by

Professor M. J. Rosseinsky

Dr. J. B. Claridge

2

Abstract

The work of this thesis focuses on three perovskite-based compounds:

YSr2Cu3−xCoxO7+δ cuprates, Gd2BaCo2O5+δ related phases and Sr2SnO4 Ruddlesden-Popper

structures. Both YSr2Cu3−xCoxO7+δ and Gd2BaCo2O5+δ are cathode material candidates for

solid oxide fuel cells (SOFCs). Doping of Sr2SnO4 aims to enhance the ionic conductivity of

the parent phase and explore the phases as a potential SOFCs electrolyte material.

The cobalt content in the layered perovskite YSr2Cu3−xCoxO7+δ has been increased to a

maximum of x = 1.3. A slight excess of strontium was required for phase purity in these

phases, yielding the composition Y1−ySr2+yCu3−xCoxO7+δ (where y = 0.03 and 0.05). The

potential of Y1−ySr2+yCu3−xCoxO7+δ (where x = 1 to 1.3) as a cathode material for a solid

oxide fuel cell has been explored through optimisation of processing parameters, AC

impedance spectroscopy and DC conductivity measurements. The stability of

Y0.95Sr2.05Cu1.7Co1.3O7+δ with commercial electrolytes has been tested along with the stability

under CO2. This material exhibits a significant improvement in properties compared to the

parent member, Y0.97Sr2.03Cu2CoO7+δ, and is compatible with commercially available doped

ceria electrolytes at 900 °C.

Energetics of Ln2BaCo2O7 (Ln = Gd, Nd, Ce) materials consisting of a layer of

LnBaCo2O5+δ (Ln = Gd, Nd) and a fluorite layer (CeO2 or Ln2O3, Ln = Gd, Nd) have been

studied using DFT calculations. Various reactions including binary oxides and double

perovskites were taken into an account for the formation energy calculations. Phases

favourable in DFT calculations were observed also in PXRD patterns of the materials

prepared by a solid state synthesis.

DFT prediction has been also used in the work with Ruddlesden-Popper phases. The

structures of experimentally prepared Nb- and Ta-doped Sr2SnO4 phases were investigated

using high resolution diffraction methods. The conductivity of single phased materials was

studied by AC impedance spectroscopy. A significant improvement in conductivity was

observed in Sr2Sn1−xTaxO4 compounds with x = 0.03 and 0.04. The origin of the enhancement

has been studied using different techniques such as solid state Sn-NMR, UV-vis and NIR

spectroscopy methods and it tends to be explained by an ionic contribution.

4

Acknowledgements

I would like to offer my special thanks to my primary supervisor Prof. Matthew

Rosseinsky for the given opportunity to learn new methods and techniques and for the

guidance throughout my time in the group. I would also like to thank my secondary

supervisor Dr. John Claridge for his valuable advices.

I am grateful for the assistance of past and present members of the MJR research

group I have been closely working with. In particular, Dr. Antoine Demont and Dr. Ruth

Sayers for their help through the first synthetic attempts. I would like to thank Dr. George

Darling and Dr. Matthew Dyer for the patience and support with density functional theory

applications. Thanks also go to Dr. Phil Chater, Dr. Mike Pitcher, Dr. Julia Payne and Dr

Alex Corkett for their great help in structural refinements and valuable suggestions. Advices

given by Dr. Hripsime Gasparyan and Dr. Ming Li have been a great help in electrochemical

data analysis. I wish to acknowledge the help with UV-vis spectroscopy measurements

provided by Dr. Troy Manning and Borbala Kiss. I would like to thank Dr. Frédéric Blanc for

the help with Sn solid state NMR measurement and Natasha Flack for the SEM images. In

addition, thanks go to Dr. Hongjun Niu and Mike Chatterton for their technical support.

I would like to extend my thanks to Prof. Ken Durose, Dr. Laurie Phillips and Dr.

Robert Treharne for their kind assistance with NIR spectroscopy. X-ray and neutron

diffractometers support has been provided by instrument scientists from ISIS and Diamond,

namely Prof. C. Tang, Dr. S. Thomson, Dr. J. Parker (I11), Dr. Aziz Daoud-Aladine (HRPD)

and Dr. Winfried Kockelman (GEM).

Finally, I would like to thank my family, Eftychia and friends for their constant

support and good company outside of the workplace.

Contents List of Abbreviations...............................................................................................................................9

1 Introduction ................................................................................................................................... 11

1.1 Current energy demand ......................................................................................................... 11

1.2 SOFCs ................................................................................................................................... 12

1.2.1 Principle of operation .................................................................................................... 12

1.2.2 Cell efficiency ............................................................................................................... 13

1.2.3 SOFC component requirements .................................................................................... 15

1.3 Mass transport ....................................................................................................................... 16

1.3.1 Vacancy diffusion mechanism ...................................................................................... 17

1.3.2 Interstitial diffusion mechanism .................................................................................... 17

1.4 Charge transport .................................................................................................................... 18

1.5 Ionic and electronic conductivity .......................................................................................... 19

1.5.1 Ionic conductivity ......................................................................................................... 19

1.5.2 Electronic conductivity ................................................................................................. 19

1.6 Structures description ............................................................................................................ 21

1.6.1 Perovskite ...................................................................................................................... 21

1.6.2 Ruddlesden-Popper phases ........................................................................................... 22

1.6.3 Fluorite .......................................................................................................................... 23

1.7 Materials review .................................................................................................................... 24

1.7.1 Cathodes materials ........................................................................................................ 24

1.7.2 Electrolyte materials ..................................................................................................... 26

1.7.3 Anode materials ............................................................................................................ 29

1.8 Aims of the work .................................................................................................................. 31

2 Experimental and theoretical methods .......................................................................................... 33

2.1 Material synthesis ................................................................................................................. 33

2.2 Powder diffraction techniques .............................................................................................. 34

2.2.1 Fundamentals of diffraction .......................................................................................... 34

2.2.2 Diffraction of X-rays ..................................................................................................... 35

2.2.3 Diffraction of neutrons .................................................................................................. 38

2.2.4 Powder diffraction......................................................................................................... 39

2.2.5 The Rietveld Method .................................................................................................... 40

2.2.6 Laboratory X-ray diffraction ......................................................................................... 43

2.2.7 Synchrotron X-ray powder diffraction .......................................................................... 43

2.2.8 Neutron Sources and Time-of-Flight Diffraction ......................................................... 44

6

2.3 Scanning Electron Microscopy ............................................................................................. 47

2.4 Electrical conductivity measurements .................................................................................. 47

2.4.1 Fundamentals ................................................................................................................ 47

2.4.2 The four-probe DC method ........................................................................................... 48

2.4.3 Cold Isostatic Pressing .................................................................................................. 49

2.4.4 Density measurements .................................................................................................. 50

2.5 AC Electrochemical Impedance Spectroscopy (EIS) ........................................................... 50

2.5.1 Fundamentals ................................................................................................................ 50

2.5.2 Data analysis ................................................................................................................. 52

2.6 Ultraviolet-visible and Infrared Spectroscopy ...................................................................... 53

2.6.1 Ultraviolet and visible Spectroscopy ............................................................................ 53

2.6.2 Infrared Spectroscopy ................................................................................................... 54

2.7 Solid state NMR technique ................................................................................................... 55

2.8 Iodometric Titrations ............................................................................................................ 56

2.9 Thermogravimetric Analysis................................................................................................. 57

2.10 Dilatometry ........................................................................................................................... 58

2.11 Density Functional Theory (DFT) ........................................................................................ 59

2.11.1 The energy functional ................................................................................................... 60

2.11.2 Kohn-Sham equations ................................................................................................... 61

2.11.3 Exchange Correlation Functionals ............................................................................... 63

2.11.4 Pseudo potentials........................................................................................................... 65

2.11.5 DFT+U .......................................................................................................................... 66

3 Synthesis and characterization of Y1−ySr2+yCu3−xCoxO7+δ .............................................................. 68

3.1 Introduction ........................................................................................................................... 68

3.2 Synthesis ............................................................................................................................... 70

3.3 Structural characterization .................................................................................................... 71

3.3.1 Laboratory P-XRD data ................................................................................................ 71

3.3.2 Neutron Powder Diffraction data .................................................................................. 75

3.4 DC conductivity measurements ............................................................................................ 79

3.5 Thermal stability ................................................................................................................... 82

3.6 Chemical compatibility of Y0.95Sr2.05Cu1.7Co1.3O7+δ with electrolytes................................... 86

3.7 AC impedance spectroscopy of Y0.95Sr2.05Cu1.7Co1.3O7+δ ...................................................... 89

3.7.1 AC impedance data at 500 ‒ 800 °C ............................................................................. 89

3.7.2 AC impedance data for the dwelling at 650 °C............................................................. 94

3.7.3 SEM study of symmetrical cells ................................................................................... 96

3.8 AC impedance spectroscopy of Y0.97Sr2.03Cu2CoO7+δ ........................................................... 97

7

3.9 CO2 stability tests of Y0.95Sr2.05Cu1.7Co1.3O7+δ ....................................................................... 99

3.10 Thermal expansion studies of Y0.95Sr2.05Cu1.7Co1.3O7+δ ....................................................... 102

3.11 Discussion and conclusions ................................................................................................ 104

4 Prediction and synthesis of GBCO related phases ...................................................................... 107

4.1 Introduction ......................................................................................................................... 107

4.2 Computational methods ..................................................................................................... 109

4.3 Experimental methods......................................................................................................... 113

4.4 Computational results ......................................................................................................... 113

4.4.1 CeO2 ............................................................................................................................ 113

4.4.2 LnBaCo2O5 .................................................................................................................. 115

4.4.3 LnBaCo2O5.5 ................................................................................................................ 116

4.4.4 Double perovskite with fluorite layer ......................................................................... 116

4.4.5 Formation energies ...................................................................................................... 118

4.5 Experimental results ............................................................................................................ 122

4.6 Formation energies - different lanthanides ......................................................................... 126

4.7 Other fluorite layer .............................................................................................................. 128

4.7.1 Formation energies ...................................................................................................... 128

4.7.2 Experimental results .................................................................................................... 130


5 Ruddlesden-Popper phases ‒ stannates ....................................................................................... 136

5.1 Introduction ......................................................................................................................... 136

5.2 Computational methods ...................................................................................................... 138

5.3 Experimental methods......................................................................................................... 139

5.4 Computational results ......................................................................................................... 140

5.5 Structural characterization .................................................................................................. 142

5.5.1 Laboratory P-XRD ...................................................................................................... 142

5.5.2 Synchrotron data ......................................................................................................... 146

5.5.3 High Resolution Powder Diffraction data ................................................................... 151

5.6 AC Electrochemical Impedance Spectroscopy (EIS) ......................................................... 154

5.6.1 AC impedance data at 600 ‒ 900°C ............................................................................ 154

5.6.2 AC impedance data at 300 ‒ 600°C ............................................................................ 158

5.6.3 AC impedance data at different partial oxygen pressure ............................................ 160

5.7 Thermal stability ................................................................................................................. 164

5.8 UV-vis spectroscopy measurements ................................................................................... 166

5.8.1 As made materials ....................................................................................................... 166

5.8.2 Reduced materials ....................................................................................................... 167

8

5.9 Sn Solid-state NMR ............................................................................................................ 168

5.10 IR spectra ............................................................................................................................ 171


6 General Conclusions and Perspectives ....................................................................................... 177

References...........................................................................................................................................178

APPENDIX A: EDX data of Sr2Sn0.96Ta0.04O4....................................................................................187

APPENDIX B: Lattice parameters and cell volume of Sr2Sn1−xNbxO4...............................................188

APPENDIX C: Lattice parameters and cell volume of Sr2Sn1−xTaxO4................................................189

APPENDIX D: Joint I11 and HRPD Rietveld refinement of Sr2Sn0.97Nb0.03O4..................................190

APPENDIX E: Joint I11 and HRPD Rietveld refinement of Sr2Sn0.96Ta0.04O4...................................191

List of Abbreviations

AFM Anti-ferromagnetic

ASR Area specific resistance

BSCF Ba1−xSrxCo1−yFeyO3−δ

CIP Cold isostatic pressing

CPE Constant phase element

DFT Density functional theory

Ea Activation energy

EDX Energy dispersive X-ray

EIS Electrochemical impedance

spectroscopy

eV/FU Electron volts per formula unit

FM Ferromagnetic

GBCO GdBaCo2O5+δ

GDC Gadolinia-doped ceria

GGA Generalised-gradient

approximation

GOF Goodness-of-fit

HK Hohenberg-Kohn

HRPD High resolution neutron

powder diffraction

ICSD Inorganic crystal structure

database

IR Infrared

IT-SOFC Intermediate solid oxide fuel

cell

LAMOX La2Mo2O9

LDA Local density approximation

LNO La2NiO4+δ

LSC LaCoO3

LSCF La1−xSrxCo1−yFeyO3−δ

LSF Lanthanum strontium ferrite

LSGM La1−xSrxGa1−yMgyO3−δ

LSM Lanthanum strontium

manganite

MAC Multianalyzing crystal

MCFC Molten carbonate fuel cell

NIR Near infrared

NPD Neutron powder diffraction

PAFC Phosphoric acid fuel cell

PAW Projector augmented wave

method

PBE Perdew-Burke-Ernzerhof

PEFC Polymer electrolyte fuel cell

PM Pechini sol-gel method

ppm Parts per million

PVA Polyvinyl alcohol

PXRD Powder X-ray diffraction

RE Rare earth

RP Ruddlesden-Popper

RT Room temperature

ScSZ Scandia-stabilised zirconia

SDC Samarium-doped ceria

SEM Scanning electron microscopy

SIE Self interaction error

10

SOFC Solid oxide fuel cell

ss NMR Solid state NMR

TCO Transparent conductors

TEC Thermal expansion coefficient

TGA Thermogravimetric analysis

TOF Time of flight

UV-vis Ultraviolet-visible

VASP Vienna ab initio simulation

package

XC Exchange correlation

YDC Yttria-doped ceria

YSZ Yttria-stabilised zirconia

3ap Triple perovskite

1 Introduction

1.1 Current energy demand

Traditional power generation based on fossil fuels, oil and coal presents a huge threat

for the environment with the worldwide consequences such as the climate change.1,2

Growing

population and thus increasing demand enhances the need to find environmentally friendly

energy sources. This has led to more intense search for alternative energy sources along with

novel technologies in order to reduce carbon dioxide emissions and fossil fuel dependence.

Numerous reviews bring a discussion on energy and technology alternatives.3-5

The target of this thesis includes fuel cell technologies, specifically solid oxide fuel

cells (SOFCs). Fuel cells are electrochemical power generation devices providing higher

efficiencies than conventional power production.6 Their potential application ranges from

providing power for portable devices (mobile phones, laptops) and transport applications to

stationary power applications. Recent development in fuel cell technology has brought

different types of fuel cells, characterized by an electrolyte, such as the polymer electrolyte

fuel cell (PEFC), the phosphoric acid fuel cell (PAFC), the molten carbonate fuel cell

(MCFC) and the SOFC. Main advantages of fuel cells compared to conventional power

generation devices are the higher conversion efficiency and environmental benefit of

producing less CO2 altogether with the lower fuel requirement. The principle of fuel cells was

revealed in the middle of 19th

century, when the first fuel cell was also built.7,8

The first fuel

cell operated at room temperature, using a dilute sulphuric acid electrolyte, a hydrogen anode

and an oxygen cathode.

The main concepts and principles of SOFCs are discussed in Section 1.1 considers the

transport and conductivity mechanisms (Sections 1.3 to 1.5). Sections 1.6 and 1.7 bring an

overview of the structure and material types of each of the individual components of SOFCs.

The objectives of the work presented in this thesis are explained in Section 1.8.

Chapter 1. Introduction

1.2 SOFCs

1.2.1 Principle of operation

SOFCs are complex electrochemical conversion devices (converting fuel directly into

electrical current) composed of three main components: an electrolyte, a cathode and an

anode. Their operating temperature (500‒1000 °C), relatively high compared to other fuel

cells, is required to ensure adequate ionic conductivity in the electrolyte. Moreover, operating

temperatures above 500 °C show additional benefits, such as the possibility of avoiding

expensive platinum metal-based catalysts. High temperature also reduces the need for fuel

purity, i.e. natural gas can be used as the fuel without preliminary reforming. SOFCs can be

divided into two groups according to operating temperatures; high (800‒1000 °C) and

intermediate (IT-SOFC, 500‒700 °C). Lowering the operation temperature of SOFCs towards

IT-SOFC has become one of the main SOFC research goals,9 driven by significant

restrictions for materials at high temperatures such as chemical stability and compatibility.

A schematic picture of an SOFC is shown in Figure 1.1, displaying main three

compartments of the device. An SOFC can also contains interconnect components providing

electrical connection between the cathode of one individual cell and the anode of the

neighbouring cell.10

This component should prevent mixing of cathodic and anodic gases and

thus prevent their mutual reaction.

Figure 1.1: Schematic layout of an SOFC, showing the flow of electrons in order to provide electrical

energy to an external circuit. Reactions taking place on the cathode and the anode are shown on the

right side.


13

Oxygen reduction occurs on the cathode (Equation 1.1) while hydrogen is oxidised on

the anode (Equation 1.2). Transport of O2−

anions is allowed by the electrolyte exhibiting

high oxygen ion conductivity and low electronic conductivity. A wide range of fuels from

hydrogen to hydrocarbons, is available for SOFCs. The chemical energy of the fuel is

converted into electrical power with water and CO2 as a waste product (Equation 1.3).

Cathode reaction: O2 (g) + 4 e− → 2 O

2− (1.1)

Anode reaction: 2 H2 (g) + 2 O2−

→ 2 H2O (g) + 4 e− (1.2)

Overall reaction: O2 (g) + 2 H2 (g) → 2 H2O (g) (1.3)

1.2.2 Cell efficiency

The main advantage of fuel cells compared to traditional power sources is their higher

efficiency which can normally reach up to 65%.11

The main requirements for SOFC include

power loss not exceeding 0.1% during continuous operation for 1000 h and lifetime of at least

5 years. However, SOFCs materials often react with materials from other parts of the device

(cathode ‒ electrolyte interface) causing a rapid degradation. Another drawback is presented

by poisoning of the SOFC surface by chromium compounds. This is due to the doped

lanthanum chromite (La1−xMxCrO3, M = Ca, Sr) often used as the interconnect materials.12

The chromites can be substituted by temperature-resistant steels which inquires lowering the

SOFC temperature to 500 ‒ 700 °C.9

An increased SOFC efficiency can be generally achieved by lowering of the cell

polarisation losses (η, known also as overpotential). The cell overpotential is the difference

between the theoretical and actual cell voltage present in the cell. The difference is due to the

presence of polarisation losses. These losses are expressed by Equation 1.4 and are induced

by the following four types of polarisation: charge transfer (activation) polarisation (ηa,

associated with the electrodes), diffusion (concentration) polarisation (ηm, associated with

mass transport), reaction polarisation (ηr, similar to concentration polarisation) and resistance

or ohmic polarisation (ηΩ, associated with ionic and electronic conduction and contacts

between the cell compartments).

V = E0 − ηa − ηm − ηr − jR (1.4)


14

where V is the actual voltage output, E0 is the theoretical cell voltage and jR is equal to the

ohmic polarisation losses while R is the total cell resistance.12-15

The relationship between the

polarisation losses is displayed in the voltammogram of a working SOFC (Figure 1.2).12

Figure 1.2: Current ‒ voltage curve of a working SOFC with the polarization losses, taken from.12

Ohmic losses are related to Ohm's law and thus these losses are reduced by an

increase of temperature. Concentration losses contribute significantly at high currents when

the rates of electrochemical processes are considerable. The largest contribution of losses is

given by the activation polarisation, due to the complexity of electrochemical processes

taking place on electrodes commonly involving several steps. Lowering the operation

temperature makes a huge impact on activation polarisation. Decreasing temperature slows

the electrochemical reaction on electrodes.12

The resistance of each of the individual SOFC components can be quantified by area

specific resistance (ASR, in Ω cm2), also described in this thesis (Chapter 3). Material

properties for each part of SOFC including total performance need to be estimated.13,16

Power

density of a single cell is required to be of 490 mW cm−2

(in order to deliver 0.7 A at 0.68 V).


15

The total polarisation, made up of ohmic, anodic and cathodic overpotentials, is then ≈ 0.32

V. Neglecting contributions of other polarisation effects gives a single cell resistance of

≈ 0.45 Ω cm2. As a consequence of this power density requirement, a limitation of resistance

of ≈ 0.15 Ω cm2 for each of the SOFC component needs to be applied.

17

1.2.3 SOFC component requirements

The requirements for SOFC components are discussed below. There are some general

restrictions for all of the parts of SOFC such as: material (chemical) compatibility with other

device components, similar thermal expansion behaviour compared to other elements of

SOFC, cost efficiency for material synthesis, environment impact, low operation temperature

(for IT-SOFC).

Electrolytes are required to have high ionic conductivity (oxide ion or proton), high

ionic transport numbers, negligible (or no) electronic conductivity (any contribution causes

ohmic losses in a cell). The materials must be electrochemically and mechanically stable over

a wide range of temperature and oxygen partial pressure and they have to show good

sinterability for their fabrication as gas-tight membranes with very small thickness.6,17,18

Porous cathodes are critical components in SOFCs. They are commonly mixed ionic-

electronic conductors with high electronic conductivity under oxidizing conditions, stability

in the cathodic gas atmosphere and a high electrocatalytic activity towards oxygen

reduction.6,12,17

Materials suitable for anodes are also required to have high electronic conductivity.

They need to be stable in severely reducing environments with electrocatalytic activity

towards hydrogen oxidation.6,17

In addition, the anode must have a porous structure since a

large area for gaseous fuel is required.


16

1.3 Mass transport

SOFC electrolytes exhibit high ionic conductivity which is related to mass transport

processes controlled by diffusion. Electrode materials are mixed ionic-electronic conductors

displaying additional charge transport processes compared to electrolytes.

The SOFC's operation requires a transport of oxygen ions through materials via

diffusion. Diffusion is proportional to the negative of the concentration gradient (dC/dx;

Equation 1.5), i.e. the movement of particles (atoms or molecules) from an area with a high

concentration to an area with a low concentration. A mathematical approach is stated in

Fick's first and second law. Fick's first law in one direction (x) is given by:

dCJ D

dx

(1.5)

where J is the particle flux (equals to number of particles per unit area per unit time) and D is

the diffusion coefficient. When steady conditions cannot be reached, the diffusion is

described by Fick's second law (Equation 1.6).

2

2

C J CD

t x x

(1.6)

where ∂C/∂t is the change in concentration with time.

Diffusion in oxide materials is a complex phenomenon influenced by anion and cation

sublattices of a crystal structure. In most of the studied systems, oxygen ions diffusion is

significantly faster than the diffusion of cations.19

Diffusion in crystalline materials can be

imagined as a migration of atoms away from their equilibrium positions. The way that an

atom moves is described by a diffusion mechanism. Point defects play an important role in

enabling the move.20


17

1.3.1 Vacancy diffusion mechanism

In the vacancy mechanism of diffusion, an atom jumps to a neighbouring vacancy, i.e.

by a hopping mechanism (Figure 1.3). A series of adjacent exchanges with vacancies

provides a mechanism for atoms to move through the crystal. The presence of lattice

vacancies is required and their concentration has an impact on the kinetics.21

The vacancy

mechanism is common for oxygen self-diffusion in fluorite- and perovskite-related systems

and is of great interest in SOFC electrolytes. For instance, the consequence of adding CaCl2

into a NaCl crystal structure, means every Ca2+

ion replaces two Na+ ions and creates

vacancies within the structure. Oxide ion diffusion of a number of widely used cathode

materials for SOFCs with cubic perovskite structure (e.g. La1−xSrxMO3−δ, M = Mn, Fe, Co) is

based on the vacancy migration mechanism.21

Figure 1.3: Schematic of vacancy diffusion mechanism with a vacancy indicated by a square.

1.3.2 Interstitial diffusion mechanism

There are two different interstitial mechanisms: direct or indirect, also known as the

interstitialcy mechanism. The direct mechanism (Figure 1.4a) occurs by ions jumping from

one interstitial site to a neighbouring vacant interstitial site. There is no permanent

displacement of the other ions present after a single jump is completed. This diffusion type is

characteristic for small atoms, such as hydrogen and carbon.

The indirect (interstitialcy) mechanism, shown in Figure 1.4b, is described by two

steps. Firstly, an atom of an occupied lattice site (light grey) moves to an unoccupied


18

interstitial site becoming the new interstitial atom (dark grey) and secondly, an original

interstitial atom moves to an un-occupied lattice site.

An example of a SOFC material where oxygen ion transport is driven by an interstitial

diffusion mechanism is reported for lanthanum nickelate (Section 1.7.1).22

Figure 1.4: Schematic of interstitial diffusion mechanism; a) direct interstitial mechanism and b)

indirect (interstitialcy) mechanism. Dark grey colour represents interstitial sites while light grey lattice

site.

1.4 Charge transport

The measure of charge transport is expressed by the electrical conductivity, which is a

material's ability to transport charged particles under an applied electric field.23

The electrical

conductivity is defined as the charge flux (Ji) per unit of electric field (E) for a particle i with

a charge (Zi e) given by:

i i i i i

i i i i

J Z e C v Z eC Z e

E E (1.7)

where Ci is the particle concentration, vi indicates the particle velocity and μi refers to the

particle electrical mobility per unit field of charge. The conductivity of a material, as it is

shown in Equation 1.7, depends on the particle concentration and mobility. The mobility of

smaller electronic carriers such as electrons and holes is typically 2-3 orders of magnitudes

higher than the mobility of ionic carriers (vacancies). The total electrical conductivity of a


19

material is composed of contributions of each of the individual charged particles. Electrical

conductivity in ceramics can be either electronic, ionic or mixed ionic-electronic.

1.5 Ionic and electronic conductivity

1.5.1 Ionic conductivity

The high diffusion rate of oxygen ions species is required for the SOFC electrolytes.

The mechanism of ionic conductivity can be explained via vacancy or interstitial diffusion

types (Sections 1.3.1 and 1.3.2). Thermally dependent mobility is always a part of ionic

conductivity while the temperature dependence of carrier concentration may vary, depending

on the defect type.23

An increase of the ionic conductivity of a SOFC material is typically

achieved by increasing the defect concentration usually done by introducing electron rich or

poor elements (e.g. Sm doping in ceria).24

Another task lays on finding a minimum of the

activation energy for conductivity isotherms for a relevant doped structure.25

1.5.2 Electronic conductivity26

Electrical properties of solids, electronic conductivity including, were well described

in the middle of the twentieth century when Band Theory was completed and generally

approved. In this theory, conductive electrons are not linked to any atom; they are delocalised

enabling them to move easily through the crystal. The fundamental role is played by the outer

electrons being responsible for both chemical and electronic properties. The outer electrons

occupy bands of allowed energies while the regions between these bands cannot be occupied

and are called band gaps. According to contribution, the bands may be mainly of s, p, d or f

character. Overall, the band structure of a solid is more complex with possible overlaps in a

specific crystallographic direction whilst separated in others.26

In general, materials can be

classified into metals, insulators and semiconductors (Figure 1.5). An insulator normally does

not show electrical conductivity. Previous classifications between metals and semiconductors

were based on the magnitude of the measured electrical conductivity. The difference between


20

the electronic/band structures of materials are shown in Figure 1.5 and should be used to

define the structures. Semiconductors are of great interest in SOFC material research.

Figure 1.5: Band structure of: a) an insulator; b) a semiconductor and c) a metal. The valence band is

indicated by grey areas while the conduction band is represented by white areas.

Semiconductors have similar band picture to insulators except the narrower band gap

(Eg). The separation between the empty (conduction) and filled (valence) bands is small

enough that some electrons have sufficient energy to be transferred from the valence band to

the conduction band at room temperature. As the temperature increases, more electrons will

gain the energy to cross the band gap, which leads to enhancement of the conductivity. Each

time an electron is removed from the valence band to the conduction band, two mobile charge

carriers are created: an electron and a hole.26

The value of band gap can be estimated by the

conductivity measurement of a semiconductor at various temperatures in order to obtain the

thermal band gap. Another, more conventional determination, is based on using the energy of

photons (usually from UV-vis region) to excite an electron across the band gap to estimate

the optical band gap (such as UV-vis spectroscopy, Section 2.6).

The presence of point defects may increase the electronic conductivity. A point defect

on band diagram is represented as an energy level localised in the band gap. The effects of

point defects highly depend on which part of band gap they lie in. Those sitting close to the

edges of Eg (shallow levels) play an important role in electronic conductivity while those

lying in the middle of Eg (deep levels) influence optical properties. Shallow donor levels,

lying close to the conduction band, cause the material to become n-type semiconductor (using

electrons as carriers, e.g. Si doped by P). Shallow acceptor levels, lying close to the valence

band, take up electrons from it, thus creating holes and producing p-type semiconductors (e.g.


21

Si doped by B). The same consideration can be also applied to vacancies. An anion vacancy

may increase shallow donor levels close below the conduction band resulting in an n-type

semiconductor. A cation vacancy will act the opposite way, typical for p-type

semiconductors.

The presence of shallow levels is commonly found in transition-metal oxides (typical

also for SOFC materials) which can exist in several valence states. Their conduction band is

made up mainly from metal d orbitals (possibly mixed with s orbitals), while the valence

band is derived from oxygen 2p orbitals.

1.6 Structure descriptions

1.6.1 Perovskite

Perovskites and perovskite-related structures are common material types for SOFCs,

especially for cathodes. Their general formula is ABO3, with two different cation sites, A and

B respectively. The ideal perovskite structure has cubic symmetry with space group 3Pm m

and its example is given in Figure 1.6. The perovskite structure consists of larger 12-

coordinated A-site cations and smaller 6-fold coordinated B-site cation. BO6 octahedra are

linked forming a BO3 layer with cavities occupied by A-site cation. A-site atoms are typically

presented by alkaline and alkaline earth ions or rare earths, while B-site is typically occupied

by a transition metal.27,28

Figure 1.6: Structure of cubic perovskite SrTiO3 with Sr cations (green) on A-site and Ti cations

(blue) on B-site, oxygen anions are marked with red.


22

Many perovskite structures differ from the ideal cubic symmetry. Distortions within a

structure lower the symmetry (e.g. to hexagonal or orthorhombic). Additionally, a large

amount of oxygen or cation deficiency has been observed in many compounds. Certain types

of distortions are related to the structure's properties.

In order to understand the deviations from the ideal cubic structure, ABO3 oxides are

treated as purely ionic crystals where the relationship between the radii of A, B and O2−

is

given by:

2A O B Or r r r (1.8)

where rA, rB and rO are the ionic radii of A, B and O ions respectively. The deviation from the

ideal structure (and a measure of the geometrical driving force towards distortion) can be

expressed by the value of tolerance factor, t defined as:

2

A O

B O

r rt

r r

(1.9)

The value of t in perovskites varies between approximately 0.80 to 1.10.28

The oxides with

lower t values crystallise in the ilmenite structure.28

The values of t for ideal cubic structures

are mainly close to 1. Deviations of t are represented by systems with lower symmetry such

as orthorhombic and monoclinic.

1.6.2 Ruddlesden-Popper phases

Apart from cubic perovskite, there is a large family of perovskite-related structural

types. A few of the layered perovskite structures are mentioned in Section 1.7. Layered

perovskite structures are of much interest due to their diverse applications based on ionic

conductivity, photocatalytic activity, magnetic and dielectric properties.


23

Ordered B-sites and oxygen defects are presented within Ruddlesden-Popper phases

(RP, general formula An+1BnO3n+1). Their structure is composed of ABO3 perovskite units

intercepted by the AO rock salt unit (Figure 1.7). Based on the intergrowth of different

number of individual layers (n), various structures are accessible (such as RP1, RP2 phases

for n = 1 and 2 respectively). The list of possible compounds is expanded by structures with

two different A cations forming perovskite and rock salt layer (such as LaO∙nSrFeO3) or

phases with two different anions (i.e. SrFeO3∙SrF).28

Figure 1.7: Schematic of the structure of Sr2RuO4 and related compounds of Ruddlesden-Popper

series of Srn+1RunO3n+1, where n indicates the number of repeating RuO2 layers.29

1.6.3 Fluorite

Fluorite structures are common for SOFC electrolytes. It has a general formula given

by AO2 and an example of this structural type, CaF2, is shown in Figure 1.8. The cubic

closest-packed structure of CaF2 consists of eight-coordinated calcium cations with fluoride

ions adopting a tetrahedral arrangement. Oxide ion conductivity in AO2 fluorite oxides can be

enhanced by heterovalent substitution of A-site cation resulting in oxygen deficiency, such as

in yttria-stabilised zirconia (YSZ)30


24

Figure 1.8: Structure of fluorite CaF2 with Ca cations in blue and F anions in green.

1.7 Materials review

1.7.1 Cathodes materials

Several structural types have been examined for potential use as SOFC cathodes. This

review of cathode materials includes the three main structural families: perovskites,

Ruddlesden-Popper phases and layered perovskites.

Perovskite oxides ABO3−δ with 3d-transitional metals such as B = Mn, Fe, Co, Ni and

Cu were found to be most promising for SOFC cathodes.12

The most common material is the

doped lanthanum strontium manganite (LSM) showing high electrical conductivity

(320 S cm−1

at 800 °C for La0.6Sr0.4MnO3−δ composition),31

high electrochemical activity and

no compatibility problems.32

It has become an important cathode choice for high temperature

operations (700 ‒ 900 °C) although it exhibits no significant oxide-ion conductivity.12

Fe-

containing perovskite materials are favoured due to their low cost. Thus, A-site doped mixed


25

conductors of lanthanum strontium ferrite (LSF) have been studied for lower operation

temperatures.33-35

Cobalt-containing perovskites display high electronic and fast ionic conductivity both

with excellent electrocatalytic activity.17

Strontium-doped LaCoO3 (LSC) perovskites exhibit

p-type electronic conductivity due to the change of Co oxidation state and formation of

oxygen vacancies.36,37

Their total electrical conductivity is reported to be 1500 S cm−1

in air

at 600 °C,38

but the cathode application is limited by the high values of thermal coefficients

(TEC), usually 20 × 10−6

K−1

resulting in significant mismatch compared to common

electrolytes.17,39,40

A large amount of iron doping, such as in La0.6Sr0.4Co0.2Fe0.8O3−δ (LSCF),

reduces thermal expansion and provides good candidates for IT-SOFCs with ceria based

electrolytes.41-43

Reduced cathode performance due to strontium diffusion remains their main

limitation for long-term SOFC operation.44,45

Another widely studied perovskite is Ba1−xSrxCo1−yFeyO3−δ (BSCF), especially that of

x = 0.5 and y = 0.2 with a cubic structure.46

BSCF exhibits very low ASR values of

0.055‒0.071 Ω cm2 at 600 °C in Sm0.2Ce0.8O1.9 electrolyte based cells and high power

densities of about 1 W cm−2

at 600 °C.47,48

Incompatibility with the common electrolytes due

to the reactivity and high thermal expansion, typical for cobalt containing compounds, is the

main limitation for BSCF use as a SOCF cathode material.49

In addition, BSCF undergoes

phase decomposition under SOFC operating conditions50,51

and the presence of CO2 results in

the formation of thermodynamically stable BaCO3.52

Improving stability by increasing the Fe

content, such as observed in Ba0.5Sr0.5Co0.2Fe0.8O3−δ, leads to lower electrochemical

performance.53

Ruddlesden-Popper phases with general formula of An+1BnO3n+1 are of interest for

SOFC cathode application due to the high electronic conductivity caused by the mixed

valence of B-site element and the oxygen ion conductivity occurring via an interstitial

mechanism (e.g. lanthanum nickelates).22

Phases with n = 1 are most commonly studied,

although a few n = 2 and 3 phases, such as (La,Sr)n+1(FeCo)nO3n+1, have been examined as

cathodes.54,55

Depending on the nature of the cation, these compounds can crystallise in

different structural types. Nickel- and copper-containing oxides exhibit promising properties

for IT-SOFC.12

La2NiO4+δ (LNO) exhibits total conductivity of approximately 30 S cm−1

at

700 °C. An improvement of conductivity was observed for the A-site strontium-doped

compounds.56

LNO shows thermal expansion behaviour close to those for YSZ, GDC and


26

LSGM electrolytes.17,57

Chemical phase instability of LNO and its derivatives remains the

main limitation for its SOFC application.58,59

Double perovskites with general formula of AA'Co2O5+δ (where A is rare earth

element and A' = Ba, Sr) consist of double layers of square pyramidally coordinated Co

anions. These oxides are good mixed conductors with a rapid oxygen ion transport due to the

A-site ordering.17,60

Double perovskites with Gd and Pr are characterised by high coefficients

of diffusion and surface exchange of oxygen.61,62

The ASR value of GdBaCo2O5+δ (GBCO)

with GDC electrolyte was found to be 0.25 Ω cm2 at 625 °C.

62 Electrode performance of

GBCO, at the temperatures below 700 °C, has been examined also in other studies.63,64

Additionally, GdBaCo2O5+δ, as a most studied compound of the structural family (with

additional properties mentioned in Section 4.1), shows good chemical compatibility with

GDC and LSGM electrolytes and very good stability in CO2 atmospheres at temperatures up

to 700 °C.17

A vast number of AA'Co2O5+δ can be prepared and studied as potential SOFC

cathodes, including examples such as: YBaCo2O5+δ,65

LaBaCuFeO5+δ, LaBaCuCoO5+δ,66

and

SmBa0.5Sr0.5Co2O5+δ.67

Layered cuprates such as YSr2Cu2MO7+δ (M = Fe, Co) represents another group of

compounds studied as SOFC cathodes.68-70

The structural and electrochemical studies of

these materials are discussed in more details in Section 3.1.

1.7.2 Electrolyte materials

SOFC electrolytes are solely oxygen ion conductors. In order to achieve the

movement of the ions, the crystal structure of a material needs to contain unoccupied sites

equivalent to those of the lattice oxygen ions. Materials suitable for SOFC electrolytes belong

to several structural types, such as fluorite based (yttria and ceria derivatives), perovskite and

perovskite-related (LaGaO3 based oxides, brownmillerites), La2Mo2O9 (LAMOX) and

apatites.

The fluorite structure is shown in Figure 1.8. The zirconia-based oxides, are mainly

used as electrolytes for high temperatures (800 ‒ 1000 °C), more details about their properties

are available in selected reviews.71,72

Zirconia-based electrolytes show minimal electronic

contribution to the total conductivity and stable electrochemical performance over a broad


27

p(O2) range. The properties of zirconia-based electrolytes vary depending on the presence of

divalent and trivalent cation dopants.

Yttria-stabilised zirconia (YSZ) has become one of the most common SOFC

electrolytes. Zirconia (ZrO2) is known to exist in three different structures of monoclinic,

tetragonal and cubic phases.72

Yttria (Y2O3) is used to stabilise the high temperature cubic

phase, which is achievable even with a small amount dopant. A 8 mol% of Y2O3 is reported

to be the lowest concentration required to stabilise the cubic phase at room temperature.73

Substitution of Y3+

by Zr4+

also creates vacancies in the oxygen sublattice since the lower

valence of Y3+

compared to Zr4+

. This 8 mol% of Y2O3 in zirconia exhibits the conductivity

of 0.16 S cm−1

at 1000 °C.72

Operation at that temperature would bring significant constraints

to other parts of SOFC device. Scandia-stabilised zirconia (ScSZ) is an alternative to YSZ. It

shows two times higher ionic conductivity than those for any of the known zirconia-stabilised

electrolytes. The higher conductivity is attributed to the smaller ionic radii mismatch between

Zr4+

and Sc3+

. The use of ScSZ is limited by the presence of many phases, by the cost of Sc

and by the difficulty to obtain equilibrium in Sc2O3-ZrO2 system.72,74

Stoichiometric ceria (CeO2) is not a good oxygen ion conductor, but its conductivity

can be dramatically enhanced by low valance doping.24

It has been reported that the

conductivity and activation energy correlate with the ionic radius of the dopants ions, with

the lowest activation energy for the cations whose radius matches most closely to that of

Ce4+

.73

Ionic radii of Gd3+

and Sm3+

match most closely with that of the Ce4+

. Gadolinia-

doped ceria (GDC, CGO) and samarium-doped ceria (SDC) have become one of the leading

electrolytes for IT-SOFCs. High values of conductivity were also obtained for yttria-doped

ceria (YDC). All three ceria based electrolytes exhibit similar conductivities at 750 °C with

values of 6.5 × 10−2

S cm−1

for YDC,75

6.7 × 10−2

S cm−1

for GDC76

and 6.1 × 10−2

S cm−1

for SDC.77

At high temperatures and low oxygen partial pressure, ceria-doped electrolytes

exhibit n-type electronic conduction that is detrimental to SOFC operation. Co-doping has

also been carried out to improve the properties of ceria-based electrolytes. A decrease of the

electronic conductivity has been achieved for example in co-doping with Mg (in

Ce0.85Gd0.1Mg0.05O1.9)78

or with Pr (in Ce1−x−yGdxPrO2−z).79

Perovskite and perovskite-related electrolytes include two main structural types:

LaGaO3-based and Ba2In2O5 brownmillerites. Solid solutions of LaGaO3 oxide, especially

those with Sr and Mg (LSGM), exhibit better ionic conductivity than zirconia stabilised

electrolytes (Figure 1.9).80

The basic LaGaO3 compound can be doped with divalent ions


28

with appropriate sizes. Sr2+

and Mg2+

are favourable due to their low solution energies.73

Solid solution with Sr2+

and Mg2+

dopants was studied in extended range.81-83

The highest

values of conductivity of 0.17 and 0.08 S cm−1

were obtained for La0.8Sr0.2Ga0.83Mg0.17O2.815

at 800 and 700 °C respectively.84

A decrease of performance was observed in SOFCs

composed of LSGM, lanthanum-containing perovskite cathodes and Ni-LSGM anodes,

which was attributed to the anode-electrolyte reaction forming LaNiO3 and La2NiO4.85

Chemical stability of LSGM under reducing atmosphere at high temperatures is questionable

and causes significant changes in the surface morphology of the electrolyte.86

The presence of

several phases due to the vaporisation of Ga was detected by X-ray diffraction.

Figure 1.9: Conductivity data comparison of YSZ87

, CGO88

and LSGM84

; figure taken from17

Brownmillerite structures have been studied since the order-disorder transition in

Ba2In2O5 was observed.89

In order to suppress the phase transition, several different dopants


29

were tried.90,91

The conductivity of 0.12 S cm−1

obtained at 800 °C for (Ba0.3Sr0.2La0.5)InO2.75

was higher than that for YSZ.92,93

A new family of materials to be considered as electrolytes was reported by Lacorre's

group.94-96

These materials, based on La2Mo2O9 compound (also known as LAMOX),

represent alternative to common electrolytes, with the comparable conductivity properties

above 600 °C due to a phase transition from monoclinic α form to the cubic β form,

improving conductivity of the α form by two orders of magnitude. Several different doping

strategies have been applied to suppress the phase transition and improve the stability.97-100

The practical derivatives of La2Mo2O9 are limited due its reactivity towards some electrodes

and its high thermal expansion coefficients.

Apatites, with the general formula M10(XO4)6O2+y (where M is rare-earth element, X

is P, Si or Ge), have been proposed as potential electrolytes.101-103

Their values of

conductivity at higher temperatures are comparable with those for YSZ, but the conductivity

may be anisotropic, due to the structure, as in the case of lanthanum silicate apatites. The

conductivity in apatites involves interstitial oxide-ions, which differentiates them from

fluorite and perovskite based electrolytes.104-107

The conductivity of apatites is encouraging,

while high temperatures required for the synthesis and to material densification remains

unsolved.108

1.7.3 Anode materials

SOFC anodes provide the electrochemical oxidation of fuels (see Section 1.2.1) with

relevant requirements for the materials properties and fabrication (Section 1.2.3). Several

reviews about these anode materials are available,109-113

while here only a summary

highlighting the most common materials is outlined.

Early stage SOFC anode research was focused on the materials based on graphite,

platinum, iron, cobalt and nickel.73

Graphite was found to electrochemically corrode;

platinum did not show sufficient mechanical properties due to the water-vapour evolution.

Iron at a certain partial oxygen pressure oxidises, resulting in formation of a red iron oxide.

Cobalt was more stable but more expensive for fabrication. Nickel showed a big thermal

expansion mismatch with YSZ electrolyte and a tendency to aggregate at higher


30

temperatures. None of the metallic candidates showed ideal match for materials requirements.

Later on, a nickel-zirconia concept was reported114

and has become the leading

material for SOFC anodes. Ni provides the required electronic conductivity and catalytic

activity for both oxidation and for steam reforming. The YSZ provides ionic conductivity and

improves the thermal expansion match with YSZ electrolyte. The typical amount of Ni is at

least 30 vol% to achieve the percolation threshold for electronic conductivity.115

Few

disadvantages are connected with the use of Ni-YSZ. Since the catalytic activity of Ni is for

both steam reforming and hydrogen cracking, carbon deposition occurs when hydrocarbon is

used as the fuel and thus it prevents this material's use for operation with hydrocarbon fuels.

109,110,116,117 Impurities in the fuel stream, especially sulphur, are detrimental to the Ni-YSZ

anode, due to the low tolerance to sulphur.118

The problem of carbon deposition can be solved

by decreasing operating temperature88

or by a deposition of (Y2O3)0.15(CeO2)0.85 porous film

between YSZ and Ni-YSZ anode.119

Cu containing materials represent an option to replace Ni-YSZ cermets. For instance,

an anode composed of Cu‒CeO2‒YSZ/SDC showed improved cell performance, especially

for hydrocarbon fuels.120,121

Electrocatalytic properties of Cu-based materials are not as good

as those with Ni. Moreover, Cu is a relatively low melting point metal, having compatibility

problems with many standard SOFC fabrication procedures. An alternative is represented by

Cu-containing alloys with a second metal providing sufficient catalytic activity, with Ni as a

reasonable option.14

The bimetallic anodes (such as Cu/Ni, Cu/Co in CeO2/YSZ) exhibited

improved performance and in the case of Cu/Co also improved carbon tolerance.122

A tri-

metal anode, FexCo0.5−xNi0.5/Sm0.2Ce0.8O1.9(SDC) cermet, with GDC electrolyte and

Sm0.5Sr0.5CoO3/SDC cathode (with x = 0.25) exhibits much better electrochemical

performance than this observed for Ni/SDC.123

Alternative anode materials may be categorised into several groups depending on the

structural type they posses. Ceramics based on CeO2 having a mixed ionic-electronic

conductivity, are the most common fluorite candidates for SOFC anodes. Addition of Ni, Cu,

Pt, Ru improves their catalytic activity.124-126

The most studied perovskite materials are the

titanates and chromites due to their stability in reducing conditions,109,127

with

La0.75Sr0.25Cr0.5Mn0.5O3−δ as an examples having similar electrochemical performance to that

of Ni/YSZ cermets and good catalytic activity.128,129

Tungsten bronze materials were also

studied as SOFC anodes showing poor oxygen exchange kinetics.130,131

Pyrochlores, based on

Gd2Ti2O7, have been found to have a high mixed conductivity but their potential use as SOFC


31

anodes is reduced by the stability problems causing changes of electrochemical performance

depending on p(O2) region.132,133

1.8 Aims of the work

The work for this thesis is focused on new perovskite-based materials for solid oxide

fuel cells (SOFCs) including the development of both cathode and electrolyte materials which

can be divided into three structural types: YSr2Cu3−xCoxO7+δ cuprates (Chapter 3), GBCO

related phases (Chapter 4) and Ruddlesden-Popper structures (Chapter 5). Both experimental

and theoretical methods have been applied during the research described within this thesis.

Previously reported work on YSr2Cu2MO7+δ (M = Fe, Co) revealed the triple

perovskite materials as a potential SOFCs cathode candidate.69,70

Cobalt plays an important

role in catalysis which is necessary for the SOFC operation.17

By increasing the cobalt

content in Y1−ySr2+yCu3−xCoxO7+δ compounds (Chapter 3), and thus introducing electronic

carriers, improved conductivity properties are expected. Perovskite cuprates are found to

have thermal and compatibility problem,68,70,134-137

which may be detrimental to SOFCs

application. Therefore, the increasing of the electrochemical stability and compatibility of

Y1−ySr2+yCu3−xCoxO7+δ represents the other challenge of the work with triple perovskite

materials.

Expansions in unit cell size driven by the ordering of cations and oxide vacancies (in

order to satisfy local coordination preferences) are known for many layered perovskite

structures. Double perovskites are one of the superstructures of interest for SOFCs studies.

The ordering of Gd and Ba and accompanied oxygen vacancy ordering in GdBaCo2O5+δ

results in high oxide mobility, producing excellent SOFC cathode performance.60,62,138

By

adding a fluorite layer to a mixed double perovskite conductor, an improvement of ionic

conductivity is expected. DFT studies are used as a tool for predicting the chemical stability

of different double perovskite structures with different fluorite layers (Chapter 4). The

structural model is based on Y2SrFeCuO6.5 structure.139,140

The most stable compounds

predicted are then synthesized using a solid state chemistry synthesis.

SOFC research on Ruddlesden-Popper phases is focused mainly on the mixed ionic-

electronic conductors such as A2NiO4+x phases (A = La, Sr, Pr, Nd).56,141-143

The presented


32

work on the RP1 phases aims to find ionic conductors with interstitial oxygen. The target is

to find pure ionic conductor with a potential use as a SOFC electrolyte. The RP1 structures

can support interstitial and vacant oxygen sites (as reported for La1.5+xSr0.5−xCo0.5Ni0.5O4+δ).144

Several A- and B-site doped RP1 structures (with Sr, Ba on A-site and Sn, Zr, Hf on B-site)

are looked at using DFT calculations to investigate interstitial oxygen energetics (Chapter 5).

The theoretical work is followed by the structural and electrochemical characterization of

synthesized phases of doped Sr2SnO4 phases.

Chapter 2. Experimental and theoretical methods

2 Experimental and theoretical methods

2.1 Material synthesis

The majority of the materials discussed within this thesis were prepared via the

standard ceramic methods.145,146

The starting powdered materials (oxides and carbonates)

were treated in a way to remove any moisture absorbed in the powder when stored at room

temperatures. Most of the starting oxides were calcined at ≈ 200 °C except lanthanide oxides,

which were kept at ≈ 900 °C due to their stronger hygroscopic properties. Carbonates were

stored in a desiccator. Precise stoichiometric amounts of starting materials were weighed out

and ground together using a pestle and mortar. Mixing of the starting materials is required in

order to ensure the homogenity of the mixture, to reduce the size of the grains and thus

maximising the surface contact between the reagent grains which enhances the reaction rate.

A few synthetic routes and physical property measurements (e.g. DC conductivity, AC

impedance measurements) require improved mixing with smaller particle sizes, which is

ensured by ball milling.146

A FRITSCH Pulverizette 7 planetary mill was used for ball

milling of the materials in this thesis. The ground reaction mixture is then pressed to form a

pellet, which maximises the surface contact between reagent grains and enhances the reaction

rate. Some sample may require a calcination step at a lower temperature to begin the

chemical reaction and improve the homogeneity of the sample. A calcination step is also

required for the synthesis of volatile starting materials.

Once the mixing and/or calcination step has been completed, the powder or pellet is

sintered to high temperatures (typically over 1000 °C) for several hours or days. Sintering at

high temperatures allows the high kinetic barrier to reaction to be overcome, allowing the

thermodynamically favoured product to form.27

The sintering process needs to be commonly

repeated with re-grinding of the sample in order to produce a single phase sample. Planetary

milling used at the very beginning of the synthesis may decrease the number of sintering

cycles. The vast majority of the syntheses in this thesis were carried out in ambient air.

Sintering at different atmospheres (oxygen, nitrogen) may be used to alter the oxygen content

and to change the final product. It should also be taken into account that syntheses in

reducing atmospheres may lead to the production of materials that are not stable in certain


34

conditions. Since SOFCs materials are required to be stable at oxidising/reducing conditions,

which depends on the part of the (Section 1.2.3), reducing atmosphere was not used for the

synthesis. The final powder products were then characterised using the methods described in

the following sections.

Several Ruddlesden-Popper phase were synthesised using Pechini sol-gel method

(PM).147

Samples of outstanding homogenity can be prepared by this polymerizable complex

method. PM is based on in situ polymerization of monomers specially introduced into

solution in addition to the required metal cations. Suitable metal salts are introduced into the

ethylene glycol after dissolution of citric acid forming metal-citrate complex. The mixture is

then heated to ≈ 100 °C while stirring to enhance the formation of polyester. After the

formation of a plastic-like gel, the temperature is increased in order to remove the excess of

ethylene glycol. The obtained mixture is then heated at 500 ‒ 600 °C in order to oxidise the

organic compounds and to form precursor powders ready for further annealing at

temperatures depending on the product required.

2.2 Powder diffraction techniques27

2.2.1 Fundamentals of diffraction

The main technique used to describe structures of the solids is diffraction. A crystal

structure consists of an infinite array of constituents units (a unique unit of atoms or

molecules) forming a three-dimensional lattice. The lattice can be described by the unit cell ‒

chosen as the repeating unit. In order to define the unit cell, three translation vectors a, b, c

and the interaxial angles α, β, γ are required. Based on the values of the six lattice parameters,

seven crystal systems are defined. In the case where a ≠ b ≠ c and α ≠ β ≠ γ, the lattice has

intrinsic translation symmetry and may have inversion symmetry, but it has no rotation axes

or mirror planes. Such a lattice needs to be defined by six unique lattice parameters and is

described as a 'triclinic' system. On the other hand, the 'cubic' system has the highest number

of non-translation symmetry elements and is defined by only two unique lattice parameters a

(= b, c) and α (= β = γ = 90°).

There are fourteen independent ways of arranging points in space giving rise to 14

Bravais lattices. If the unit cell contains just a single lattice point, it is described as primitive.


35

Sometimes it is necessary to define a unit cell with more than one lattice point and such a cell

is described as centred. Based on the studies of the symmetry the crystals may have one or

more than ten basic symmetry elements (e.g. proper rotation axes, inversion or improper

axes, centres of inversion and mirror planes). A set of symmetry elements intersecting at a

common point within a crystal is called the point group. The 10 basic symmetry elements

along with their 22 possible combinations constitute the 32 crystal classes. There are two

additional symmetry elements: screw axis and glide plane, which arise in crystal but have no

counterpart in molecular symmetry. A combination of these elements with the point group

symmetry is called a space group. There are 230 possible space groups in total. The extended

translation symmetry in crystals allows diffraction methods to be used as a powerful tool for

structural characterization and determination, as described in the following sections.

2.2.2 Diffraction of X-rays148,149

The discovery of X-rays opened a vast area of scientific research in many fields. It

plays a crucial role in crystallography. For diffraction to take place, the wavelength of the

incident light has to be of the same order of magnitude as the spacings of the grating (X-ray

wavelength has a magnitude of approximately 1 × 10−10

m or 1 Å). The first theory of the X-

rays diffraction was developed by Laue at the beginning of 20th

century. The crystal was

considered as a three dimensional lattice with rows of regularly spaced atoms, but the theory

can be illustrated by a one dimensional lattice (Figure 2.1). A parallel monochromatic X-ray

beam is incident upon a row of regularly-spaced atoms. Each atom in the row behaves as a

point scatterer ‒ a secondary source of monochromatic X-rays with a spherical wavefront.

Interference then occurs between these scattered beams and maximum constructive

interference will occur when the path lengths differ by an integer (n) of wavelength (λ)

according to:

(cos cos )i i dAB CD X n (2.1)

where AB and CD are path lengths (Figure 2.1). This equation is valid for a one-dimensional

system. A real crystal is a three dimensional system with three Laue equations, corresponding


36

to each orthogonal x, y, z axis, which must be simultaneously satisfied for a constructive

interference to occur and can be written as follows:

0(cos cos )n xa n (2.2)

0(cos cos )n yb n (2.3)

0(cos cos )n zc n (2.4)

Laue equations indicate that a periodic lattice produces diffraction maxima at specific

angles, which are defined by lattice parameters. The equations give the most general

representation of a three dimensional diffraction pattern and they may be used in the form of

Equations 2.2-2.4 to describe the geometry of a single crystal diffraction. The solution of the

Laue equations involves the determination of up to 12 variables. This complexity limits their

usage in practical crystallography.

Figure 2.1: Schematic diagram illustrating Laue diffraction from a lattice in the x direction. The

labelled distances and points are used to construct the first Laue equation (Equation 2.1).


37

An alternative formulation to the Laue's model was suggested by W. L. Bragg.

Bragg's equation describes the principle of X-ray diffraction in terms of reflection of X-rays

by sets of lattice planes. Lattice planes are crystallographic planes characterized by the index

triplet hkl known as Miller indices. Parallel planes have the same indices and are equally

spaced, separated by the distance dhkl. Bragg's analysis deals with X-rays like visible light

being reflected by the surface of a mirror (Figure 2.2). In contrast to the visible light, the X-

rays penetrate inside the material. This causes additional reflections at many consecutive

parallel planes. Since all X-rays are reflected in the same direction, superposition of the

scattered rays takes place. For a constructive interference to occur, the path difference

(B'C − BC) has to be equal to an integer number of wavelengths, resulting in the following

condition:

' ; ( ' ) cos 2sin

hkldB C BC B C

(2.5)

therefore for the maximum constructive interference:

cos 2

sin sin

hkl hkld dn

(2.6)

which is simplified to the equation known as Bragg's equation:

2 sinhkld n (2.7)

Other useful equivalent variations of the Bragg's equation are:

0

2sin 1h s s

d

(2.8)

4 sin 2Q

d

(2.9)

The vector Q is the physicist's equivalent of the crystallographer's scattering vector h while s

and s0 are unit vectors along the directions of the diffracted and incident beams respectively.


38

The physical meaning of Q is the momentum transfer on scattering and differs from h by a

factor of 2π.

Figure 2.2: Schematic diagram of Bragg diffraction from planes of atoms, θ indicates the incident and

diffracted beam angles and d(dhkl) the inner spacing within the crystal structure with Miller indices

hkl.

2.2.3 Diffraction of neutrons

The diffraction of neutrons is described in the same way as X-ray diffraction

(Section 2.2.1). Neutrons can be diffracted by a crystal if they have an appropriate energy.

The scattering mechanism in the case of neutrons is slightly different since neutrons are

scattered by atomic nuclei via the nuclear force. The nuclear force acts at a different range

(10−15

m, compared to a common incident wavelength of 10−10

m). Thus the scattering objects

for the neutron diffraction are extremely small and the atomic scattering factors are

independent of the incident angle (θ) of the beam.

Neutron scattering amplitude is not related to electron count but it depends on the

scattering length of atoms. Every isotope has a unique coherent scattering length (the

effective size of the nucleus), which may take a negative value and is not related to chemical

properties. Typical scattering lengths range from −5 fm to 16 fm. This makes neutron

diffraction an outstanding tool to differentiate light elements (such as oxygen) in the presence


39

of heavy ones. Sample containers are generally made from vanadium which posses a very

small coherent scattering length (≈ −0.38 fm).

Another important feature of neutrons is their magnetic moments (spin 1/2 particle).

This means that the neutrons can also be diffracted by a lattice of ordered magnetic moments.

A neutron diffraction pattern is thus a superposition of a structural pattern and a magnetic

pattern. Magnetic structure crystallography represents another important application of

neutron diffraction.

2.2.4 Powder diffraction

Diffraction methods mentioned in previous sections were related to single crystals. In

practice, it is difficult to grow single crystals for many materials. Therefore, instead of single

crystals, polycrystalline samples have to be used. A distribution of crystallite sizes of powder

samples should be of a few microns in order to have a distribution from many crystallites.

Ideally, every crystallite of the sample should be randomly orientated to the incident beam, so

that each hkl reflection gives a uniform cone-shaped diffracted beam (Debye-Scherrer cone)

with contributions from many crystallites (Figure 2.3). Although the powder diffraction data

lack the three-dimensionality, the powder diffraction pattern represents a one-dimensional

snapshot of the three-dimensional reciprocal lattice of a crystal.148

The quality of the pattern

depends on the nature and the energy of the applied radiation, the resolution of the

instrument, and the physical and chemical conditions of the sample. In typical powder

diffraction experiments the intensity of the diffracted beam is measured as a function of 2θ.

Since all collected information is condensed into one dimension, hkl peaks with identical or

similar d-spacing overlap to some degree. Peak overlap tends to be more problematic for the

materials with lower crystal symmetry. Thus, powder methods are mostly used for high-

symmetry solids and less for complex molecular solids. Problems with overlapping most

often increase for multi-phased materials.

Phase and overlapping problems make the use of powder diffraction difficult as a tool

for a direct structure solution. After recent developments of instruments and characterization

methods, powder diffraction techniques have become more powerful and useful as a tool for

phase purity identification. High quality data (synchrotron ‒ Section 2.2.7, neutron diffraction


40

‒ Section 2.2.8) combined with Rietveld refinement149-151

can reveal structural details that in

turn can explain material properties and phenomena.

Figure 2.3: Schematic diagram of a Debye-Scherrer reflection cone from a diffraction of an incident

beam by a polycrystalline powder material.

2.2.5 The Rietveld Method

The Rietveld method was originally developed for powder neutron diffraction by H.

M. Rietveld150,151

to derive the maximum amount of information from a polycrystalline

powder. It is not a structure solution method, instead it relies on the input of a starting model

which generates a theoretical diffraction pattern. Rietveld refinements shown in this thesis

were carried out using the Topas152

software package.

Since a powder diffraction pattern is a set of peaks with overlapping and slightly

varying background, a Rietveld refinement may present a complex curve fitting problem. To

identify the true contribution to Bragg reflections to an observed peak, correct intensity to

each peak needs to be assigned. The Rietveld method does not use the integrated intensity of

each peak, but it considers the intensity of each individual data point as a single observable.

The intensity of a data point yi may contain contributions from several Bragg reflections. The

Fhkl values are generated from the structural model and are used to calculate theoretical

intensity for this point (yci).


41

The sum of other contributions to an observed intensity is as follows:

2(2 2 )ci bi hkl hkl hkl i hkl

hkl

y y s L P A F (2.10)

where s ‒ scale factor, Lhkl ‒ combined Lorentz polarisation and multiplicity, Phkl ‒ preferred

orientation and A ‒ absorption are four factors depending mainly on experimental set-up;

θ ‒ reflection profile function which describes the peak shapes and ybi ‒ the background

intensity. Fhkl refers to the structure factor for a suggested structural model.

Structural factor (Fhkl) contains information on the atomic scattering contribution at

each reflection from all of the atoms of the unit cell and is defined by:

2

1

( ) ( )j j jn

i hx ky lzj j j

hkl

j

F g t s f s e

(2.11)

where:

n denotes the number of atoms of the unit cell,

gj − fractional occupancy of atoms j, where atom j fully occupies the site g

j,

(s) − the angular dependence,

f j − the atomic scattering factor depending on the type of the diffraction,

exponential term includes the fractional coordinates (x, y, z) for atom j by the

corresponding h, k or l value.

Once a pattern of the material is calculated, corresponding values of yci are obtained

and the difference between the theoretical and observed patterns (yi − yci) is calculated. The

difference is then incorporated into a 'residual function' Sy. Adjusting the structural or set-up

parameters allows the value of Sy to be minimised by a least-squares method:

2( )i ciy

i i

y yS

y

(2.12)

The least squares optimisation is a feedback loop with a certain user-defined parameter of a

starting model and an improved model as an output model of the calculation which is used for


42

the next loop. When the refinement is finished, an adequate model closely corresponding to

the observed pattern is obtained.

The accuracy of the refinement can be expressed by several statistical outputs, known

as R-factors. The factors describing the quality of the full-pattern fitting are based on directly

observed intensities. These include the R-pattern (Rp) and R-weighted pattern (Rwp) expressed

by:

i ci

ip

i

i

y y

Ry

(2.13)

2

2

( )i i ciyi

wp

i i i i

i i

w y yS

Rw y w y

(2.14)

where wi is the weighting factor. The final important R-factor is the R-expected factor (Rexp).

Rexp can be calculated from a pattern, which can be considered as the best expected fit for a

pattern.149

varexp 2

obs

i i

i

N NR

w y

(2.15)

where Nobs and Nvar are the number of observables and refined parameters, respectively. For a

common powder experiment Nobs ≈ 2000, therefore Nobs ≫ Nvar. The only difference between

the expected R factor and Rwp value if the model was perfect is due to statistical fluctuations.

The end of the refinement is typically quoted by a goodness-of-fit (GOF), χ, related to both

Rwp and Rexp:

var exp

y wp

obs

S R

N N R

(2.16)

For an ideal powder diffraction experiment with sufficient data point number and

counting time (background intensity is less dominated in yci) with a superb structural model,

the value of GOF is expected to be in the range 1.0 ≤ χ ≤ 1.3. Practically the χ are higher,


43

mostly due to insufficiency of at least one of the set-up parameters. It needs to be noted that

the statistical evaluation should not be the only criterion taken into an account. A graphical

representation is an excellent way to judge the quality of the refinement and it should not be

omitted. A model is considered to be acceptable if it fits the observed data well, with the

reasonable statistical factors and it makes physical and chemical sense.

2.2.6 Laboratory X-ray diffraction

To generate an X-ray in laboratory based instruments, electrons are produced by

heating a metal filament, accelerated using high voltages of 40 kV through a vacuum tube

into an anode material, typically made of copper, cobalt of molybdenum. The incident

electrons have energy sufficient to initiate a number of electronic transitions within the metal.

This leads to the emission of photons with a broad spectrum of X-rays at most voltages (low

intensities) and with a sharp emission (where the exact voltage is needed) at specific

wavelengths typical for each anode metal. There are many electronic transitions available in

metals, therefore it is common to observe multiple characteristic emissions. The core

electrons of Cu atoms include 2p or 3p electrons. Kα photon is given by a 2p → 1s transition,

when Kβ photon is given by a 3p → 1s transition. Due to the spin-coupling, the transitions are

complicated yielding in two different spin states and thus resulting in two characteristic

wavelengths (Kα1 and Kα2). Since a monochromatic X-ray beam is required, one radiation is

typically selected and filtered using an appropriate monochromator. The instruments used for

the laboratory Powder X-ray Diffraction (PXRD) characterization had cobalt (a Panalytical

X'pert Pro diffractometer, λ = 1.789 Å) and copper target (a Bruker D8 Advance

diffractometer, λ (Cu Kα1= 1.541 Å).

2.2.7 Synchrotron X-ray powder diffraction

Synchrotron sources show notable advantages compared to a common laboratory

sources. The more intense beam allows faster data collection with an improved resolution of

the obtained diffraction patterns. Synchrotrons are considered, together with time-of-flight


44

neutron diffraction, as the highest available resolution tools for powder crystallography.

Another advantage of the synchrotron-based experiments is the variation of available

wavelengths. This can be useful for extreme environment studies such as high-pressure

diffraction.

The synchrotron data presented in this thesis were collected using I11 High-

Resolution Diffractometer at Diamond Light Source, U. K. The beamline is described in

more details by Thompson et al.153

The source, an in-vacuum undulation, is inserted in an

array of permanent magnets, and it produces X-ray ranging from 11 to 20 keV. The optical

hutch, consisting of slits, monochromators and harmonic rejection mirrors, focuses the beam

resulting in a small final beam size. The instrument works in Debye-Scherrer geometry. Si

and Ge are known as a suitable analyzers for synchrotron powder beamlines.154

I11 beamline

detecting system is based on multianalyzing crystal (MAC) assembly. To enhance the data

collection, five identical nine-crystal MAC-arms of total of 45 detectors are mounted at 30 °

intervals. Each individual detector corresponds to Si (111) analyser crystal. A picture of I11

experimental hutch with the closer look at the sample detector area is shown in Figure 2.4.

Samples for I11 beamline are mounted in spinning glass capillaries.

Figure 2.4: a) I11 experimental hutch showing the diffractometer (DIF), 5 arms for MAC-detectors

(MACs), robotic arm (ROB), carousel with 200 specimen positions (CAR), and heavy duty Table

(XYZ); b) sample-detector arrangement with 5 sets of MAC detectors. Both images are taken from.153

2.2.8 Neutron Sources and Time-of-Flight Diffraction

Neutrons for diffraction experiments can be generated using two different ways, by

reactor sources or by spallation sources. Neutron diffraction data presented in this thesis were


45

collected at the ISIS facility (Rutherford Appleton Laboratory, U. K.). ISIS sources

(schematic of which is shown in Figure 2.5), as other spallation sources, do not produce a

continuous beam. The neutron generation starts with the formation of hydride ions which are

accelerated in a linear accelerator. A beam of protons, generated after the collision of hydride

ions with an aluminium oxide target, are then accelerated in a synchrotron to obtain an energy

of 800 MeV. The high energy proton beam then hit a tungsten target resulting in neutron

production. A pulsed neutron beam causes energy-disperse mode of the neutron diffraction

experiments known as time of flight (TOF) diffraction.

The behaviour of neutrons is described by the de Broglie relationship:

n

h h

p m v (2.17)

where h is Planck constant, p is momentum, mn is neutron mass, and v is velocity.

Considering the fixed path length between moderator and detector L, and the time t for

neutron to reach the detector Equation 2.17 can be modified to:

n

ht

m L (2.18)

which can be substituted into Bragg's equation (Equation 2.7) giving the relationship between

the neutron TOF and d spacing:

2 sinhkl

n

htd

m L (2.19)

The values of dhkl can be determine as the variable in TOF experiments due to the detectors at

fixed 2θ positions and fixed path length L.

The dependence of the intensities of neutron Bragg reflections to 2θ is not strong.

Neutron diffraction is described by the same equations as the X-ray diffraction but the

scattering mechanism is different since neutrons are scattered by atomic nuclei via the

nuclear force. Typical wavelength of the nuclear force is at a range of 10−15

m, compared to

10−10

m incident wavelength. Thus, the scattering objects for neutron diffraction are

extremely small and are independent of the incident angle. In order to prevent any


46

overlapping of pulses, the neutron beam is not heavily moderated at ISIS. As a consequence

of this, a relatively high population with short wavelengths is observed. Therefore, reflections

with short d-spacing are common and their intensities are stronger than those observed by

synchrotron. The θ position of the detector affects the range of the observed d-spacing, which

can be seen from Equation 2.18. Thus the detector at a low scattering angle will show the

longest d-spacing and vice versa, a detector at high scattering angles the smallest d-spacing.

Having a multiple detector banks in TOF experiments is common. Low angle banks are

typically used for magnetic structure refinements, since magnetic peaks are common for a

long-range structural ordering. High-angle banks provide the highest resolution data of

reactions with small d-spacing. An intermediate bank is used to collect the data with

intermediate counting statistics and resolution for intermediate d-spacing. The High

Resolution Neutron Powder Diffraction (HRPD) instrument, used for the TOF data collection

shown in Section 5.5.3 of this thesis, has three banks of detectors, fixed at 30, 90 and 168 °.

The data can be collected over the d-spacing between 0.3 ‒ 16.5 Å. The NPD data at 740 °C

shown in Section 3.3.2 were collected using GEM instrument. The GEM diffractometer

allows collection of the data over a large range of d-spacing using six banks of detectors in a

relatively short time.155

Figure 2.5: Schematic of the ISIS neutron source, with all of the associated neutron diffraction

instrumentation highlighted, figure taken from the STFC-ISIS webpage.


47

2.3 Scanning Electron Microscopy

Electron microscopy is often used for the study of structure, morphology, crystallite

size or defects of solids. An electron beam used in Scanning Electron Microscopy (SEM) is

focused on a sample surface.27,156

The electron beam is produced by heating a tungsten

filament and accelerating to energies between 2 and 40 keV. Beam electrons interact with the

surface of the sample producing secondary electrons. These electrons are detected giving a

topographic map of the studied material. Secondary electrons can be observed only from the

near-surface region of a conducting surface. Therefore, gold or carbon thin coating is

required for insulating materials. Images of the scanned material with various range of

magnification can be achieved. The very short wavelength of the electron allows resolution to

0.1 nm. The SEM study of Y0.95Sr2.05Cu1.7Co1.3O7+δ material presented in this thesis

(Section 3.7.3) was performed using a Hitachi S-4800 scanning electron microscope and the

analysis of the sample was carried out by a low 3 kV electron beam.

2.4 Electrical conductivity measurements

2.4.1 Fundamentals

One of the key properties of solids is their electrical resistivity. Different materials

exhibit resistivity varying over 20 orders of magnitude. There is no single method available

to measure all of the materials. Electrical resistivity (ρ) of a material describes how the

material resists the flow of electricity. In a simplified microscopic model electricity is shown

as a simple movement of electrons157

. Electrons may collide with atoms of the material.

Every collision slows down the electron. A material, which produces few collisions is a low-

resistivity material, those producing lots of collisions is referred as a high-resistivity material.

Electrical resistivity is thus geometry dependent. The resistivity can significantly vary with a

change of temperature. The resistivity of metals increases as temperature increases while the

resistivity of semiconductors usually decreases. Electrical conductivity (σ) is defined as the

inverse of electrical resistivity and is given by:


48

RA

l and

1

(2.20)

where A is the cross-sectional area, l is the length and R is the electrical resistance of a

sample. The electrical resistance is defined by voltage (V) and an electric current (I)

following Ohm's law:

VR

I (2.21)

It should be noted that the resistance R depends on the size and shape of the measured

material while the resistivity ρ is independent of the shape and size of specimen.

2.4.2 The four-probe DC method

All of the DC measurements mentioned in this thesis (Section 3.4) were carried out

using the four-probe measurement technique (using a Keithley 220 Current Source and a

Keithley 2182 Nanovoltameter). The schematic of the measurement arrangement is shown in

Figure 2.6. The measurement is performed on a bar of material with four wires (contacts)

attached to it. The contacts are made of a wire and paste (typically made from gold or silver).

An ammeter measures the current I passing through the specimen. Inner contacts are

connected to voltmeter measuring the voltage V. The four-probe resistivity is then expressed

by:

1

Vwh

ll (2.22)

where w is the width of the sample bar, h ‒ height of the sample bar, l ‒ the distance between

the two outer contacts (where the ammeter is connected) and l1 is the distance between the

two inner contacts (connected with voltmeter). All the values of distance are in metres. The

quality of contacts is essential for a DC measurement. That includes the geometry of the


49

contacts where l1 distance should be much larger than the thickness of the bar. The contacts

should be thin as possible for the satisfactory accuracy and completely independent. Another

necessary condition is the density of a measured material, which is expected to be over 90%.

Figure 2.6: A four-probe DC technique of a bar of material. The voltmeter measures the voltage

between inner contacts whilst the ammeter is connected to outer contacts.

2.4.3 Cold Isostatic Pressing

Cold IsostaticPressing (CIP) is a method of applying pressure from multiple direction

(using hydrostatic pressure) to a sample in order to obtain great uniformity of compaction

over the entire surface.158

The CIP is commonly used to increase the density of a material for

physical property measurements and it was used during the work presented in this thesis;

before the DC (Section 3.4) or AC impedance measurements (Section 5.6.1). Samples were

first pelletised, then sealed in waterproof bags and lowered into the hydraulic fluid. Dense

pellets were then made by CIP using an Autoclave Engineers Cold Isostatic Press, under a

pressure of 206.85 MPa. The samples were dwelled at the pressure for 3 min, then the

pressure was slowly released and the sample bag was taken out of the autoclave.


50

2.4.4 Density measurements

The density of ceramic materials before the conductivity measurements was obtained

using an Archimedes' Principle balance. According to Archimedes' principle, the upward

force that is exerted on a body immersed in a fluid (fully or partially submerged) is equal to

the weight of the fluid that the body displaces. Sample densities (ρs) were calculated from the

following formula:

1

3 2

s l

m

m m

(2.23)

where m1 is the mass of dry sample, m2 is the mass of immersed sample, m3 is the mass of

soaked sample and ρl is the density of a liquid. All of the samples were immersed in distilled

water. The relative density (ρrel) of the prepared pellets was calculated as a fraction of the

actual density of the sample (ρs) to the theoretical crystallographic density (ρtheor, calculated

using lattice parameter data from the Inorganic Crystal Structure Database, ICSD) according

to:

(%) 100srel

theor

(2.24)

2.5 AC Electrochemical Impedance Spectroscopy (EIS)

2.5.1 Fundamentals

The use of Ohm’s law is limited to only one circuit element ‒ an ideal resistor. The

real elements (electrolytes or electrodes in our studies) are more complex and need to be

described by impedance (Z) rather than resistance. Like resistance, impedance defines the

ability of a circuit to resist a flow of electrical current. Electrochemical impedance159

is

typically carried out by applying an AC potential to an electrochemical cell and measuring

the current flowing through the cell. The current response, It, is shifted in phase (by θ, which

is related to the radial frequency ω = 2 πf) and has a different amplitude compared to a

sinusoidal signal of applied voltage (Vt) expressed by following equations:


51

sint mV V t (2.25)

sint mI I t (2.26)

where Vm and Im are magnitudes of voltage and current respectively.

An equation to calculate impedance, analogous to Ohm’s law, is given by:

sin sin

sin sin

mtm

t m

V t tVZ Z

I I t t

(2.27)

with the impedance magnitude Zm = Vm/Im.

It is common to express the impedance as a complex number composed of real (Z’, ReZ) and

imaginary (Z’’, ImZ) terms:

cos sini

m m mZ Z e Z i Z

(2.28)

where Z’ = |Zm| cos(θ) and Z’’ = |Zm| sin(θ). Figure 2.7 shows one of the common impedance

data representations ‒ the Nyquist plot. The impedance is presented as a vector of length |Z|

with θ, the angle between the vector and the x-axis.

Figure 2.7: Nyquist plot obtained from impedance data collection (taken from160

) consisting of a

single semicircle.


52

2.5.2 Data analysis

Electrochemical impedance plots often contain more than one semicircle. Each of the

individual semicircles is analysed by fitting to an equivalent electrical circuit model. Most of

the models are based on common electrical elements such as resistors, capacitors and

inductors with their impedance defined by following equations:

:resistor Z R (2.29)

:inductor Z j L (2.30)

1:capacitor Z

j C (2.31)

In the cases where the impedance semicircles are depressed, a constant phase element (CPE)

can be used for data analysis. The CPE impedance is expressed by:

( )Z A j (2.32)

where A is the inverse of the capacitance (= 1/C) and α, an exponent which is equal to 1 for a

capacitor. The capacitance of CPE is related to the depression angle by:

1

1( )n nC R Q (2.33)

where R is the element resistance, Q is the pseudo-capacitance and n is a parameter related to

the depression angle. The output of impedance data analysis contains further information

about the physical processes occurring within the measured electrochemical cell. That is

often due to the various frequencies at which the different physical phenomena relax. The

values of capacitance extracted from the equivalent circuit fits depend on different

contributions161

(examples given in Table 2.1).


53

Table 2.1: Capacitance values and their possible interpretation.161

Capacitance (F m−1

) Phenomenon responsible

10−12

Bulk of electrolyte

10−11 Minor, second phase

10−11

‒10−8

Grain boundary

10−4

Electrochemical reactions

Once all of the equivalent circuits are assigned, the value of bulk and grain boundary

conductivity (σ) can be calculated from the sample dimension: l ‒ length and SA ‒ surface

area (Equation 2.34). The cathode performance of a material is commonly expressed by the

value of area specific resistance (ASR) which is obtained using total resistance for the

cathode surface area divided by 2 (taking the cell symmetry into account, as is shown in

Equation 2.35.

l

R SA

(2.34)

2

totR SAASR

(2.35)

2.6 Ultraviolet-visible and Infrared Spectroscopy

2.6.1 Ultraviolet and visible Spectroscopy

Ultraviolet-visible (UV-vis) spectroscopy uses light in the visible and adjacent (near

UV and near infrared ‒ NIR) ranges. It is a technique for determining if a material is able to

absorb radiation in the UV or visible region. That can be found from the optical absorption of

the material determined from the optical absorption coefficient as a function of the energy of

photon corresponding to wavelengths of light. The optical absorption spectrum of a

semiconductor contains information on its band-gap. UV-vis spectroscopy is commonly

carried out in solutions, since transmission in solid materials is very low. The actual


54

absorption coefficient cannot be measured directly. The Kubelka Munk (KM) remission

function F(R) is used to represent the absorption coefficient α, and is calculated from the

reflectance of the solid sample, with s ‒ the scattering coefficient according to:

2

1( )

2

RF R

R s

(2.36)

The variation in absorption coefficient α as a function of photon energy can be fitted to a

power law relationship (Equation 2.37), where B is a constant, Eg is the band gap energy and

n takes on a value that depends on the nature of the transition; values of n = 1/2 and n = 2

result in a linear fit for direct and indirect transitions respectively.162

( ) ,n

g gE E B E E E E (2.37)

Equation 2.37 is valid only for photon energies greater or equal to the band gap energy. The

value of the band gap energy is obtained as x axis intercept from a plot of (Eα(E))n vs E.

The UV-vis study of the Sr2SnO4 related materials shown in this thesis (Section 5.8)

deals with the direct band gap determination. The values of the band gaps were determined

from the plots (F(R) × E)1/2

vs E. The linear part of the plot was fitted to straight line using

Origin while the linear equation parameter was calculated. The value of the direct band gap is

equal to the intercept of the x-axis (when y = 0). The indirect band gap values can be obtained

by a similar method by plotting (F(R) × E)2 vs E.

2.6.2 Infrared Spectroscopy

Infrared spectroscopy deals with the radiation with a longer wavelength than visible

light. The IR measurements shown in Section 5.10 used the radiation from the higher energy

near-infrared (NIR) part which is commonly measured in the region of 800 ‒ 2500 nm. IR

spectroscopy uses the fact that molecules absorb specific ‒ resonant frequencies characteristic

for their structure and functional group within the structure. IR spectroscopy of solid state


55

materials includes mainly identification of OH groups. NIR spectra application is often used

for the opto-electronic characterisation of transparent materials.163

NIR spectra of Sr2SnO4 related phases were carried out using a Shimadzu SolidSpec-

3700 UV-VIS-NIR Spectrophotometer with three detectors: photomultiplier tube detector for

UV-VIS region and InGaAs and PbS detectors for near infrared region. The InGaAs detector

can be switched to PbS detector in the region of 1600 to 1800 nm. Although the use of all of

the three detectors provides smooth spectra over a wide range, a noise represented by a bump

of the spectra is usually accompanied the detector switchover. Thus, it is advised to collect

the spectra in various wavelength ranges in order not to mis-interpret the spectra peaks.

2.7 Solid state NMR technique164,165

Solid state NMR spectroscopy is a technique providing information about the

structure of diamagnetic materials and about the dynamics of processes occurring within the

studied structures. Unlike the X-ray diffraction methods, solid state NMR can be applied for a

study both of disordered (melts and colloid gels) and ordered single crystals materials.

Solution NMR spectra consist of a series of very sharp transitions which are caused by

averaging of anisotropic NMR interactions by rapid random tumbling. Solid state NMR

spectra are on the contrary very broad due to many effects, such as anisotropic or orientation-

dependent interactions. The shape of the spectra provides additional information on

chemistry, structure and dynamics in the solid state. Determination of the relationship

between the experimental spectrum and recurring structural motifs is necessary in NMR. This

can be done by a ‘fingerprinting’ procedure when a database of NMR spectra is established

for related materials and are compared with an unknown structure. Another way of solving

the relationship is represented by ab initio theoretical calculation of the electric field gradient

at a nucleus of an atom in a known crystal structure.

Most of the NMR active nuclei in the periodic table are also available for the solid

state NMR. To produce an NMR spectrum, a nucleus must have a nuclear spin. The most

important for solid state NMR are nuclei with odd mass numbers (e.g. 29

Si, 27

Al). Another

factor, which needs to be considered, is a natural abundance of an NMR nucleus. In order to

minimize anisotropic NMR interactions several methods have been developed, e.g. Magic-

angle spinning, Dilution, Multiple-Pulse Sequences, Cross Polarization. Solid state NMR


56

technique used in this thesis includes the experiments with Sn. Three NMR isotopes of Sn

exist: 119

Sn (8.5% natural abundance), 117

Sn (7.5%) and 115

Sn (0.35%), all are spin 1/2. For

practical reasons (better sensitivity, higher natural abundance combined with higher

resonance frequency) 119

Sn is used. The Sn NMR data shown in Section 5.9 were collected

and analysed by Dr. Frédéric Blanc.

2.8 Iodometric Titrations

Iodometry is a volumetric chemical analysis, where an addition of a precise amount of

standardised compound to a sample leads to the redox reaction characterised by a colour

change of a solution. Colour change is often enhanced by the addition of a small amount of

indicator. Starch solution is used to indicate the equivalence point in iodometry. Iodometric

titrations are performed to obtain the information about the relative molecular mass or

oxygen content of an oxide material as it was performed for Y0.95Sr2.05Cu1.7Co1.3O7+δ

(Section 3.3.1). Iodometry involves indirect titration of iodine and it can be described by

following reaction:

2,2 2

2 2 3 4 62 2H H O

I S O I S O

(2.38)

An excess of I− (typically in form of KI) is required to reduce the sample and to

generate I2. Prior to iodometric titration, standardisation of thiosulphate (Na2S2O3) has to be

carried out. The most common way for titration of Na2S2O3 is by using potassium iodate

(KIO3). KIO3 is a source of I2, which is then titrated by Na2S2O3 standard in acidic solution.

The titration can be sum up as follows:

2,

3 3

H H OKIO K IO

(2.39)

2,

3 2 25 6 3 3H H O

IO I H I H O

(2.40)

In every studied oxide material, reduced by this method, it is important to identify the

reduced products. The titrations performed on Y0.95Sr2.05Cu1.7Co1.3O7+δ were based on the


57

reduction of Cu and Co from an unknown oxidation state to +1 and +2 respectively. The

presence of Cu in the titrated system requires an addition for the iodometric titrations. An

excess of KI is easily absorbed by CuI and needs to be released before the end point of the

titration.166

CuI interferes with the sharpness of the end point. To prevent that, KSCN is

added after the starch end point.167

The oxygen content ([O]) was calculated using following formula:

( )

1

w red

red

Os

ex s

MO O

Mm

n m

(2.41)

where [Ored] is the oxygen content of the reduced sample, Mw(red) is its molecular mass, ms is

the mass of the titrated sample, MO is the atomic weight of oxygen, and nex indicates the

number of moles of titrated I2, what is determined from:

2

s sex

c Vn

(2.42)

where Cs and Vs refers to the concentration and volume of Na2S2O3 solution used for titration.

2.9 Thermogravimetric Analysis

Thermogravimetric analysis (TGA) measures the weight of a sample as a function of

time as the temperature is increased at a controlled uniform rate.27

The sample is loaded into

a platinum or alumina pan and enclosed in a furnace. The sample mass is controlled by a

balance relative to the empty pan. The experiment can be performed at ambient air or under

different oxidising or reducing atmosphere. Monitored changes in mass are typically due to

oxidation or adsorption of gas (observed as a weight gain) while dehydratation, reduction and

decomposition cause a weight loss. Materials from this thesis (Y0.95Sr2.05Cu1.7Co1.3O7+δ and

Sr2SnO4 derivatives) were studied by TGA to inspect their stability at a define temperature

range in ambient air using a TA Thermogravimetric Analyzer TGA Q600.


58

TGA analysis under CO2 or Ar atmosphere were carried out using a TA Thermogravimetric

Analyzer TGA Q500.

2.10 Dilatometry

Dilatometry measurements mentioned in this thesis (Section 3.10) were done in order

to measure the thermal expansion, which is one of the important properties of SOFCs

components. The thermal expansion coefficient α at temperature T is given by:168

/T

dL dT

L (2.43)

where L is the length of the sample at room temperature. The dilatometry measurement was

carried out using a Nietzch 402 C Dilatometer. A schematic of the instrument is shown in

Figure 2.8. A dense pellet of a 6 mm diameter is placed in alumina sample carrier. The

temperature of both the sample and furnace is monitored by the thermocouple. The change in

length whilst the sample is heated is not only due to the expansion of the sample. The

changes in length of the sample support and the pushrod are also monitored and measured.

The sample length change is corrected by the expansion of the sample support as this change

is not included in the sample expansion. The expansion results in a change in voltage,

transformed by an amplifier to DC voltage, which is proportional to the displacement.


59

Figure 2.8: DIL 402 C measuring unit: 1 - tube, 2 - vacuum flange, 3 - slider knob, 4 - fan, 5 -

furnace, 6 - reversing tube with shut-off valve and clamping nut, 7 - protective tube, 8 - front panel, 9

- support, 10 - retaining nut, 11 - sample carrier with thermocouple and pushrod, 12 - sample, 13 -

stop plate. Adopted from Operating Instructions DIL 402 C.

2.11 Density Functional Theory (DFT)

Development of screening methods for materials, their properties and simulations of

processes in materials has been attracting attention in recent decades. Density Functional

Theory (DFT) based predictions represent a powerful and accurate tool to accelerate

materials discovery process by orienting experimental chemists to computationally predicted

compounds. Its advantage lies on the expression of electron-electron interactions in many

electron systems as an effective one-electron potential, which is a function of the electron

density only.169

The basic problem of describing the studied materials falls on the presence of many

electrons. The Hamiltonian operator corresponding to the total energy of the system can be

expressed:

22 22 2

, , ,

1 1ˆ2 2 2 2

I JIi I

i i I i j I I Ji I I Ji j

Z Z eZ e eH

m r R M R Rr r

(2.44)


60

An approximation from an early stage of quantum mechanics was suggested by Born-

Oppenheimer, whereby the nuclei were fixed and the kinetic terms of nuclei were neglected.

If we were able solve the Schrödinger equation for many electron (Equation 2.45), we could

predict the behaviour of any electronic system.

1 2 1 2ˆ ( , ,..., ) ( , ,..., )N NH r r r E r r r (2.45)

Many electron wave function Ψ(r1,r2,...,rN) is a function of 3N variables, where N is the

number of electron. The construction of many-electron wave function is very demanding

problem due to the complexity of the system. Three main approaches of many-body problem

can be applied. Independent electron approximation represented by Hartree-Fock, statistical

Quantum Monte Carlo method and DFT – considering the electron density instead of the

wave function. Since the theoretical work of this thesis is based on DFT calculations

(performed using Vienna Ab Initio Simulation Package – VASP) this section is concerned

about the basics of DFT, its modifications and applications.

2.11.1 The energy functional

The basic proof that the density of electron charge is a unique description of a system

and that it corresponds to the energy minimum was outlined by Hohenberg and Kohn.170

When we have a system of N electrons in the ground state of an atom, a molecule or a solid,

the ground state is characterized by the positions of the nuclei, their potential Vext(ri)

including the electrostatic potential and some other scalar components, and the number of

electrons N. The Hamiltonian is then expressed:

0ˆ ˆ ˆ( ) ( )ee ext ext i

i

H T V V H V r (2.46)

Where, the operator 0H contains the kinetic energy operator T and the electron-electron

interaction potential Vee. If the external potential is fixed, then a number of N electrons leads

to a unique wave function Ψ of the system, which results in a unique density of charge ρ(r).


61

The main statement of the first Hohenberg-Kohn (HK) theorem states the reverse of

this finding, that the potential Vext(r) and the many-electron wave function Ψ are uniquely

determined by the density of charge ρ(r). The proof is based on the opposite statement ‒ the

assumption that two wave functions leads to the same density of charge resulting in a

contradiction. Since the opposite statement is evidently wrong, the original must be correct.

The second theorem proves that the ground state density is not only unique, but also the only

one which minimizes the total energy. This can be proven in two steps. First, the total energy

is written as a functional of density (E), since the density is unique.171

ˆ ˆ ˆee ext

ee ext

E H E T V V d

T V V

r r r r

r r r

(2.47)

The property that the ground state is the state of minimum energy is applied for the second

step. Hence, if we have two different densities, one of which corresponds to the ground state

density ρ(r) when the other ρ(rʹ) does not correspond. The inequality is as follows:

ˆ ˆ' '

' ' ' 'ee ext

E H H

ETE V V

r

r r r r r

(2.48)

This gives a formula to find the ground state density in a numerical procedure: to start with a

trial density, to evaluate the total energy and to change the density until a minimum is

reached.

2.11.2 Kohn-Sham equations

An early attempt to construct the functionals before the HK theorem was the Thomas-

Fermi model.172,173

It does not bring an accurate solution since it misses essential features of a

physical system in condensed matter.174

For better understanding of HK theorem, an extended

system with non-interacting electron is considered. In this system we deal only with the

kinetic energy and the external potential (μ) and the Euler-Lagrange equation171

of the

problem takes the simply form:


62

00ext

TV

r

r (2.49)

Notice, that the kinetic energy, T0, is the kinetic energy of non-interacting electrons and it is

different from the kinetic energy of interacting electrons T. The many electrons wave

function in a non-interacting electron gas is described exactly. If every local orbital λ is

occupied by two electrons, the kinetic energy is given by a sum over the same orbital171

:

2/2

*

0

1

22

n

T dm

r r r (2.50)

The single electron orbitals are obtained by solving the single particle Schrödinger equation:

2

2V

m

r r r (2.51)

After addressing the problem of the kinetic energy functional and the functional of

electron-electron interactions, the main idea of Kohn and Sham was then to formulate the

problem of interacting electrons in exactly the same way by just changing the potential:

00eff

Tv

r

r (2.52)

The effective potential (υeff) counts all missing parts of the electron interactions in.

0,

e

eff

e TVV

Tv

r r

r r r (2.53)

The charge density is expressed by:

/2

2

1

2n

r r (2.54)


63

If the effective potential is described, a comprehensive solution for the problem can be found.

Thus the focus of finding the ground state density is changed to finding the universal

effective potential, which can only be approached by various approximations.

2.11.3 Exchange Correlation Functionals

An important feature about the Kohn-Sham theory is that fact that the theory is not

based on the density alone. The density is obtained from the density of single electron states.

The total energy is given by171

:

* *

0, , H xcE T E E V d r r r (2.55)

The exchange correlation energy (Exc[ρ] represents the difference between the energy of

interacting electrons (first bracket on the right side of Equation 2.56) and the energy of

electrons interacting only via their Coulomb interaction (second bracket):

0xc ee HE V E T T (2.56)

Finally, after expressing Hartree energy (EH[ρ]) and the energy due to electron-ion

interactions, the exchange correlation energy is given by:

*

;

' ' '

xc

xc xc

x

xc

x

xc c

c

EE

V

Ed V

V d

V

r r r rr

r r r r r r rr

(2.57)

Thus, the effective potential is expressed as follows:

eff ext H xcv V eV V r r r r (2.58)


64

where the exchange correlation (XC) potential is the functional derivative:

0

xc

xc ee H

EV V E T T

r

r r (2.59)

This potential depends only on the number of electrons and has to be calculated separately.

That means in every density functional simulation a table with the values of the potential for

a given charge density is looked up. Several different methods have been developed in order

to calculate XC potential.

Local Density Approximation (LDA)175

functional have been mostly used in early

stage of DFT calculations. LDA is based on an assumption that the exchange and correlation

potentials depend only on the value of the charge density at a specific point of the system.

The exchange-correlation energy for LDA can be given by:

3 homLDA

xc xcE n d rn n r r (2.60)

where εxchom

is the XC energy per electron in a homogeneous electron gas with the

corresponding electron density n(r). LDA tends to overestimate the bond strength in solids.

The calculated lattice parameters are too small, cohesive energies are overestimated, and

energy gaps in semiconductors and insulators are vastly underestimated.

The Generalised-Gradient Approximation (GGA)175

takes the value of the density at

each point as well as the magnitude of the gradient of the density.

3 homGGA

xc xc xcE n d r n F s r r (2.61)

with an additional term Fxcs(r) compared to LDA which indicates (with the term s(r)) how

the electron density gradient varies. The GGA makes improvement over LDA for many

cases. It corrects the overbinding tendency in LDA, with a certain overcorrection in few

situations. Other equilibrium properties that are sensitive to lattice constant such as phonon

frequencies, bulk modulus, and magnetic moments are sometimes also over-corrected by

GGA.176


65

2.11.4 Pseudo potentials177

As we have systems with increasing number of electrons (e.g. transition metals), the

computational cost is bigger. A method commonly used in DFT is to create a pseudo

potential. This involved typically two approximations. Firstly, the core electrons are treated

as a frozen core assuming that the interactions of core electrons with the surrounding

chemical environment are negligible and thus only the valence electrons interact with the

chemical environment. Core-electron wave functions are localised. For the plane wave basis

set, the number of planes would need to be very large (> 106) to describe localised core-

electrons. To make the plane wave more feasible, pseudo potential needs to be applied.

A pseudo potential replaces nucleus and core electrons by a fixed effective potential.

Only valence electrons are taken into account in the calculations. The core state is removed

from the spectrum (see schematic in Figure 2.9). In practise the core electrons wave functions

are generated from calculations on the isolated atoms, libraries of which are provided by DFT

packages. In some cases in order to improve the accuracy even some of the sub-valent

orbitals are treated as if they were valence. Pseudo potential generation can be summed into

three main steps:

1) Calculate exact all electron wave functions for a reference atom.

2) Replace the exact wave function by a node less pseudo-wave function.

3) Invert Schrödinger equation to obtain the pseudo potential.

Pseudo potentials must conserve exactly the scattering properties of the original

atomic configuration. They must be generated with the same functional that will be later used

in DFT calculations. It should also be noted that the choice of pseudo potential is not unique;

there is a lot of freedom to construct them. A pseudo potential is called soft when a few plane

waves are needed. If a pseudo potential can be used in various environment (molecule, solid,

metal) it is called transferable. Creation of a good pseudo potential must meet requirements

for both soft and transferable pseudo potentials. A variety of different methods using pseudo

potential is available.178-181

A related method used within this thesis is the Projector

Augmented Wave method (PAW)178

implemented in VASP.


66

Figure 2.9: Schematic of pseudo potential concept taken from.182

The pseudo wave function and

potential are labelled in red; the all electron wave function and potential are in blue.

2.11.5 DFT+U

DFT+U represent one of the recent and essential developments in DFT.183

An

inherent problem of the DFT methods comes from the Coulomb term in the Kohn-Sham

equations, in which an electron interacts with the total electron density. That results in the

electron interacting with itself, what is known as the self interaction error (SIE). The SIE

causes localised electron orbitals to be destabilised and so DFT often results in orbitals being

spread out spatially in order to minimise self interaction. The DFT+U includes the orbital

dependence of the self-energy operators which are missing from the Kohn-Sham potential,

neglecting the fine details of the spatial variation of the Coulomb potential. In order to obtain

correct computation of band structures and total energies of systems that should have

localised orbitals (such as in semiconductors with partially filled d and f orbitals), corrections

for SIE are required. In DFT+U for each atom we apply a numerical parameter Ueff to

specified orbitals (d orbitals in this thesis) with Ueff defined as177

:


67

effU U J (2.62)

Where U and J are parameterised. Ueff is a numerical parameter which requires fitting to a

relevant material, which is commonly done by varying the value U, with J remaining fixed.

The parameter sets used in this theses were taken from published related DFT works (given

in relevant chapter) in order to determine reaction energies to form perovskite related oxides

and oxygen gas where available.

3 Synthesis and characterization of Y1−ySr2+yCu3−xCoxO7+δ

3.1 Introduction

Previous research in the area of perovskite cuprates was focused on the

superconductor YBa2Cu3O7-δ. It has been shown as a possible cathode material for SOFCs

(high electronic conductivity and mixed conductivity at temperatures above 300 °C).68,184-186

However at temperatures above 900 °C and at low current densities materials, react with YSZ

electrolyte and decompose.68

This was indicated by the presence of significant amounts of

Y2BaCuO5 and BaZrO3 phases. Despite the insufficient stability, this material showed quite

good electronic and oxide ion conductivity.187,188

There is also electrochemical

decomposition observed at low current densities (20 mA cm−2

).68

The instability of

YBa2Cu3O7-δ was explained by electrostatic effects, observed as an increase in the Cu(1)-

O(1) distance.189

In order to improve thermal stability, copper can be partially replaced by

other metal cations with different coordination preferences (M = Fe, Co, Ga, Al, etc).134-137

Replacing Ba with Sr and partial substitution of copper with Co increased the stability of the

material.189

Figure 3.1: Structure of YSr2Cu2CoO7+δ adopted from,190

with Y (grey-blue) and Sr (green) on the A-

sites, with the square pyramids of Cu (orange) and the tetrahedral Co ordering (dark blue) on B-site:

a) view along c-axis; b) rotated view around a- and b-axis.

Chapter 3. Synthesis and characterisation of Y1−ySr2+yCu3−xCoxO7+δ

69

The orthorhombic structure of YSr2Cu2CoO7+δ (Figure 3.1) was first reported by

Huang et al.190

and later work showed that it exhibits a transformation to a tetragonal

structure at high temperature (above 800 °C).70

Its structure may be described as an ordered

perovskite, with Y, Sr, Co and Cu layers ordered along the a-axis. CuO5 units adopt a corner

sharing, square based pyramidal geometry, whilst corner sharing CoO4 tetrahedra adopt a

disordered arrangement with both oxygen atoms adopting a split site model. Strontium and

yttrium are also ordered, due to their different local coordination preferences. The disordered

CoO4 tetrahedral layers are between the two strontium layers, whilst the CuO layers sit either

side of the yttrium containing layers. This structure has two copper sites: the one (Cu2) is

associated with the superconducting and a second site (Cu1) is connecting Cu2 layers along

the c-axis. Previous structural works based on high-resolution transmission-electron

microscopy demonstrated that all the Co atoms are located in Cu1 sites.191,192

The fact that

Co2+/3+

prefers tetrahedral coordination to square planar leads to cooperative oxygen ordering

and therefore to an expanded unit cell.

Previous research on YSr2Cu3−xCoxO7+δ and related phases was focused on structural

studies.189,192-194

The most detailed neutron diffraction study193,194

confirmed previous

structural works192,195

and compared them with superconductor YBa2Cu3O7−δ material.

Reported studies on Co-containing phases included materials of x = 1. The structure of Co-

1212 phase was found to be best described in Ima2 space group.193

The overall oxygen

content 7.01(2) was determined from refined oxygen occupancies from neutron powder

diffraction (NPD) data.193

Later studies on triple copper perovskites (M = Fe, Co) were

focused on the materials as a potential SOFCs cathode,69,70

showing promising conductivity

values, for Fe σ550°C = 35 S cm−1

and for Co σ900°C = 15 S cm−1

, but also significant

compatibility problems with common electrolytes (reactivity with LSGM and ceria-based

electrolytes).70

The compatibility tests were carried out at the temperatures between 900 and

1000 °C with the reactions in all cases except for the ceria-based electrolytes at 900 °C.

Thermogravimetric analysis showed that no oxygen loss occurred in the 25 – 900 °C

temperature range under both air and nitrogen atmospheres.70

The work reported in this chapter is focused on Y1−ySr2+yCu3−xCoxO7+δ compounds

including synthesis, structural and electrochemical characterization of a series of materials

with different Co doping levels. Increasing the cobalt content improves conductivity


70

properties of studied phases by introducing electronic carriers.17

Cobalt also plays important

role in catalysis which is necessary for the operation of SOFC (see Section 1.6.1).17

The

beginning of the work includes structural characterisation of these materials by powder

diffraction, followed by thermal and compatibility studies and comparison with previously

reported works of YSr2Cu2CoO7+δ phase. Since the work is oriented towards SOFC

application, conductivity measurements of the symmetrical cells play an important role and

are presented at the end of this chapter, together with CO2 stability tests and thermal

expansion studies.

3.2 Synthesis

Y1−ySr2−yCu3−xCoxO7+δ samples with nominal values of x = 1.00, 1.10, 1.20, 1.25,

1.30, 1.40, and 1.50 and of y = 0, 0.03, and 0.05 were prepared by solid state synthesis.

Stoichiometric mixtures of pre-dried Y2O3 (99.99%), SrCO3 (99.99%), Co3O4 (99.9%) and

CuO (99.995%) – all purchased from Alfa Aesar – were ground by hand and heated in

alumina crucibles to 1050 °C in air at a rate of 5 °C min−1

and then immediately cooled down

to room temperature at a rate of 5 °C min−1

. The samples were then re-ground and pelletised

before re-firing under the same conditions. In total, the samples were re-fired four times until

phase pure materials were obtained.

The initial syntheses followed the literature synthesis conditions for the parent phase

YSr2Cu2CoO7+δ (1000 °C for 16 h).69

Attempts to synthesise YSr2Cu3−xCoxO7+δ with larger

cobalt contents yielded the yttrium containing impurity phases Y2Cu2O5 and

Y2SrCu0.6Co1.4O6.5. The presence of these impurity phases was eliminated by controlling the

Y:Sr ratio to compensate partially for the increasing positive charge on the B-site, yielding

compositions of type Y1−ySr2+yCu3−xCoxO7+δ. In the range x = 1.00 – 1.25, phase pure samples

were obtained for y = 0.03, and at x = 1.30 a single phase sample was obtained with y = 0.05.

For x > 1.30 single-phase samples could not be obtained. At x = 1.0, samples were

synthesised with y = 0.0 and 0.03.


71

3.3 Structural characterization

3.3.1 Laboratory P-XRD data

Phase purity was verified by powder XRD (PXRD) collected at room temperature

using a Panalytical X'pert Pro diffractometer using Co Kα1 radiation in Bragg-Brentano

geometry (measured over a 2θ range of 5 – 120°, step size 0.0167°, time per step 3 s).

Rietveld-quality PXRD data was collected on a Bruker D8 Advance in Debye-Scherrer

geometry using a 0.2 mm glass capillary with Mo Kα1 (λ = 0.7093 Å) radiation (d-spacing

range 0.58 – 20.25 Å). Structural parameters were refined by the Rietveld method150,151

using

the Topas academic software.152

Samples of YSr2Cu3−xCoxO7+δ with x = 1.00 and 1.10 were

prepared as single phases, with all diffraction lines indexed on the basis of a

6ap × ap√2 × ap√2 body-centred unit cell. The observed decrease in a lattice parameter

(a = 22.7691(5) Å for x = 1.00 and a = 22.7563(5) Å for x = 1.10) suggested that

compositions with different cobalt contents are accessible. Rietveld refinement of

Y0.97Sr2.03Cu2CoO7+δ is shown in Figure 3.2. A decrease of lattice parameters of

YSr2Cu2CoO7+δ compared to Y0.97Sr2.03Cu2CoO7+δ was observed (Figure 3.3). The higher Co

content and higher Sr content is required to obtain single phase samples. With increasing Co

content, there is a decrease in cell parameters, in agreement with Vegard’s law (observed for

the constant Y0.97Sr2.03 ratio for x = 1 to 1.25). Figure 3.3 and Table 3.2 show the evolution

of cell parameters with cobalt content of Y1−ySr2−yCu3−xCoxO7+δ samples (x = 1 to 1.5 and y =

0, 0.03 and 0.05). The values of the parameters were obtained from Rietveld refinements of

the PXDR data collected with lanthanum hexaboride, LaB6, used as a standard. The presence

of Y2Cu2O5 and Y2SrCu0.6Co1.4O6.5 impurity phases was the limiting factor in obtaining

single phase Y1−ySr2+yCu3−xCoxO7+δ for x >1.3 (Figure 3.4).

Rietveld refinement of the room temperature structure was carried out using

laboratory PXRD data. The ambient temperature PXRD pattern was indexed to an

orthorhombic cell of dimensions 22.744(1) × 5.4210(3) × 5.4430(3) Å, with systematic

absences consistent with Imcm or Ima2 symmetry. The structure of YSr2Cu2CoO7 published

by Huang et al.190

with orientationally disordered CoO4 tetrahedra in space group Imcm was

selected as a suitable starting point for Rietveld refinement. Cell parameters, peak shape,

zero error and 18 background parameters were refined. The Cu:Co ratio on the Cu site was

fixed at 0.85:0.15, in line with the nominal composition, due to the lack of contrast between

Co and Cu using XRD data alone. The individual Y, Sr, Co and Cu coordinates and isotropic


72

temperature parameters were refined. The oxygen coordinates were fixed at the published

values and a common isotropic temperature factor was refined. This model provided a good

fit to the data, as displayed in Figure 3.2. The refined structural parameters are listed in Table

3.1.

Figure 3.2: Rietveld refinement of Y0.95Sr2.05Cu1.7Co1.3O7+δ to laboratory PXRD data at room

temperature based on the published model of YSr2Cu2CoO7 in Imcm, Rwp = 1.51%.

Table 3.1: Structural parameters obtained for Y0.95Sr2.05Cu1.7Co1.3O7 following Rietveld refinement

using laboratory PXRD data. Space group Imcm. a = 22.744(1) Å, b = 5.4210(3) Å, c = 5.4430(3) Å.

Rwp = 1.51 %, χ = 1.64.

Site Wycoff

Site

x y z occ Beq (Å2)

Y1 4a 0 0 0 0.95 0.5(2)

Sr1 4a 0 0 0 0.05 0.5(2)

Sr2 8h 0.3452(2) 0.000(5) 0 1 1.7(1)

Cu1 8h 0.4258(3) 0.494(3) 0 0.85 1.4(1)

Co1 8h 0.4258(3) 0.494(3) 0 0.15 1.4(1)

Co2 4e 0.25 0.526(6) 0 1 2.9(3)

O1a 8i 0.25 0.614 0.386 0.39 1.0(2)

O1b 8i 0.25 0.386 0.386 0.11 1.0(2)

O2 8g 0.4327 0.75 0.25 1 1.0(2)

O3 8g 0.4352 0.25 0.75 1 1.0(2)

O4 8g 0.3255 0.473 0 1 1.0(2)


73

Figure 3.3: Lattice parameter evolution of Y1−ySr2+yCu3−xCoxO7+δ (samples with YSr2(○),

Y0.97Sr2.03(∆), and Y0.95Sr2.05(□) ratio) as a function of Co doping level (x): a) lattice parameter a; b)

lattice parameters b and c; c) cell volume; d) (b+c)/2. Errors bars are within the symbol size.


74

Figure 3.4: PXRD pattern of Y0.95Sr2.05Cu1.5Co1.5O7+δ, with the main reflections of Y2SrCu0.6Co1.4O6.5

(*) and Y2O3 (+) impurity phases.

Table 3.2: Lattice parameters and cell volume of Y1−ySr2+yCu3−xCoxO7+δ plotted against different Co

doping level x (x = 1 to 1.5, for y = 0, 0.03 and 0.05). The values of the parameters were obtained after

Rietveld refinement and compared with previously reported data on x = 1 material.193

x y a (Å) b (Å) c (Å) V (Å3)

1-reported193

0 22.7987(2) 5.45150(5) 5.40890(5) 672.26(1)

1 0 22.7618(4) 5.4525(1) 5.4088(2) 671.28(3)

1 0.03 22.7691(5) 5.4545(1) 5.4099(1) 671.88(3)

1.1 0.03 22.7563(5) 5.4567(1) 5.4135(1) 672.22(2)

1.2 0.03 22.7405(5) 5.4475(1) 5.4156(1) 670.88(2)

1.25 0.03 22.7338(4) 5.4425(1) 5.4169(1) 670.22(2)

1.3 0.05 22.7294(4) 5.4394(1) 5.4179(1) 669.83(2)

1.4 0.05 22.7362(5) 5.4380(1) 5.4174(1) 669.81(2)

1.5 0.05 22.743(1) 5.4460(3) 5.4190(2) 671.20(5)


75

Iodometric titrations were carried out to determine the oxygen content. Samples of

Y0.95Sr2.05Cu1.7Co1.3O7+δ were titrated by standard sodium thiosulfate solution (0.1 M),

standardized by potassium iodate (99.995% purchased from Sigma Aldrich). For the

titrations, an approximate 0.05 g of the Y0.95Sr2.05Cu1.7Co1.3O7+δ sample was weighed out and

dissolved in a mixture of 10 ml distilled water and 10 ml of concentrated HCl. The solution

was kept under argon atmosphere and constantly stirred while 1.0 g of KI was added and the

solution was then titrated by thiosulfate standard solution using starch indicator. A 5 ml

portion of KSCN was added after the starch end point (see Section 2.8). The same procedure

was done four times and an average (mean) titre value was obtained. Details about the sample

mass, volume of used standard solution and calculated oxygen content are shown in Table

3.3. The oxygen content was determined by iodometric titration and calculated to be

Y0.95Sr2.05Cu1.7Co1.3O7.02(3). This is in good agreement with literature values for related

phases, e.g. an oxygen content of YSr2Cu2CoO7.03(4) was determined for the parent non-

substituted phase by Sansom et al.69

using thermogravimetric methods.

Table 3.3: Iodometry determination of Y0.95Sr2.05Cu1.7Co1.3O7+δ sample.

Sample Sample weight (g) volume (Na2S2O3) (ml) O content

1 0.0497 2.2 6.93(3)

2 0.0508 2.5 7.08(3)

3 0.0496 2.4 7.06(3)

4 0.0516 2.4 7.00(3)

average 0.0504 2.4 7.02(3)

3.3.2 Neutron Powder Diffraction data

In order to determine the structure at typical SOFC operating temperatures, time-of-

flight powder neutron diffraction data were collected on the high-intensity, medium-

resolution GEM diffractometer at ISIS at 740 °C. Data from banks 2, 3, 4, 5 and 6 were used

simultaneously in the refinement. The refined room temperature structure from laboratory

PXRD was used as the starting model in the refinement. Cell parameters, atomic coordinates


76

and anisotropic temperature parameters, along with 12 background parameters, a scale factor

and peak shape were refined. The composition of the Y1 site was fixed at a nominal value of

Y0.95Sr0.05 due to poor scattering contrast between Y (7.75 fm) and Sr (7.02 fm). However

there is good contrast between Co (2.49 fm) and Cu (7.718 fm) and the occupancies of their

crystallographic sites were refined to allow Cu/Co sharing, whilst constraining to a total

occupancy of 1 at each site. In the final stages of the refinement all thermal displacement

parameters were allowed to refine anisotropically. The partially-occupied O1a and O1b sites

were constrained to a common thermal displacement ellipsoid and their site occupancies

refined with a constraint to maintain the overall composition. The refined Co2 and O4 sites

show elongated thermal displacement ellipsoids consistent with orientational disorder of the

CoO4 tetrahedra. Split-site models for both the Cu1 and O4 sites were tested (as applied to

YSr2Cu2CoO7 by Babu et al.)192

by moving both atoms from the ..m plane (8h) to a general

position (16j) with an occupancy of 0.5. However the coordinates of both atoms were found

to refine to within 1 estimated standard deviation of their original 8h positions, with no

significant change to the fit, so this site splitting was not employed in the final refinement.

No Cu was found to occupy the tetrahedral cobalt site and its partial occupancy was

subsequently fixed to zero in the final stages of the refinement. A significant fraction of Co

on the square-pyramidal Cu sites was found, with a refined Co occupancy of 0.163(4), in

agreement with the nominal cobalt content of 1.30 per formula unit. The partial occupancy of

the Cu site by Co is a key structural difference between this non-stoichiometric material and

the stoichiometric parent phase, which may play an important part in improving the

material’s properties with respect to the parent material. The final refined structural

parameters are presented in Table 3.4-3.5 and corresponding Rietveld fits are shown in

Figure 3.5. The refined structure is illustrated in Figure 3.6. Bond lengths (Table 3.6) and O-

Cu-O angles (Table 3.7) are in good agreement with those published by Huang et al.190


77

Figure 3.5: Fits to GEM neutron powder diffraction data at 740 °C using Imcm (a = 22.9740(2) Å,

b = 5.51692(6) Å, c =5.47661(5) Å). Rwp = 2.12%: a) Fit to bank 3; b) Fit to bank 4; c) Fit to bank 5;

d) Fit to bank 6.

Table 3.4: Refined parameters obtained from Rietveld refinement of Y0.95Sr2.05Cu1.7Co1.3O7+δ using

GEM 740 °C neutron powder diffraction data. Space group Imcm (a = 22.9740(2) Å,

b = 5.51692(6) Å, c = 5.47661(6) Å).

Site Wycoff Site x y z Occ

Y1 4a 0 0 0 0.95

Sr1 4a 0 0 0 0.05

Sr2 8h 0.34813(4) 0.0069(4) 0 1

Cu1 8h 0.42638(5) 0.500(5) 0 0.836(3)

Co1 8h 0.42638(5) 0.500(5) 0 0.163(3)

Co2 4e 0.25 0.558(1) 0 1

O1a 4e 0.25 0.6210(7) 0.4006(9) 0.414(3)

O1b 4b 0.25 0.347(4) 0.312(4) 0.086(3)

O2 8g 0.4351(1) 0.75 0.25 1

O3 8g 0.43631(9) 0.25 0.75 1

O4 8h 0.32453(7) 0.4639(5) 0 1


78

Table 3.5: Anisotropic temperature factor coefficients obtained from the Rietveld refinement of

Y0.95Sr2.05Cu1.7Co1.3O7 using powder neutron diffraction data collected at 740 °C.

Site U11 U22 U33 U12 U13 U23

Y11 0.0195(9) 0.0095(9) 0.021(1) 0 0 0

Sr1 0.0195(9) 0.0095(9) 0.021(1) 0 0 0

Sr2 0.0237(8) 0.034(1) 0.031(1) 0.006(2) 0.017(3) 0.0(1)

Cu1/Co1 0.0274(8) 0.0118(7) 0.0095(8) -0.002(1) 0.00(1) 0.00 (1)

Co2 0.018(3) 0.027(6) 0.072(6) 0 0 0

O1a 0.0371(2) 0.031(2) 0.012(3) 0 0 0

O1b 0.0371(2) 0.031(2) 0.012(3) 0 0 0

O2 0.057(3) 0.018(1) 0.015(2) 0.00(5) 0.00(9) -0.005(2)

O3 0.027(2) 0.016(1) 0.016(2) 0.0(3) 0.0(2) -0.008(1)

O4 0.0189(9) 0.042(2) 0.065(2) 0.010(2) 0.010(5) 0.043(3)

Table 3.6: Selected bond lengths obtained from Rietveld refinement of Y0.95Sr2.05Cu1.7Co1.3O7+δ.

Y-O (Å) Sr-O (Å) Cu-O (Å) Co-O (Å)

Y-O(2):

2.4488(15)

Sr1-O(2): 2.4488(15) Cu-O(2): 1.9536(3) Co1-O(2): 1.9536(3)

Y-O(3):

2.4325(13)

Sr1-O(3): 2.4325(13) Cu-O(3): 1.9565(3) Co1-O(3): 1.9565(3)

Sr2-O(1a): 2.4177(19) Cu-O(4): 2.350(2) Co1-O(4): 2.350(2)

Sr2-O(1c): 2.576(7) Co2-O(1a): 1.848(8)

Sr2-O(2): 2.807(2) Co2-O(1a): 2.229(4)

Sr-O(4): 2.761(4) Co2-O(1c): 2.080(17)

Sr-O(4): 2.574(3) Co2-O(1c): 2.493(15)

Sr-O(4): 3.048(3) Co2-O(4): 1.795(3)

Table 3.7: Selected bond angles obtained from Rietveld refinement of Y0.95Sr2.05Cu1.7Co1.3O7+δ.

O-Co-O (°) O-Cu-O (°)

O(1a)-Co-O(1a): 106.54(19) O(2)-Cu1-O(2): 88.990(19)

O(1c)-Co-O(1c): 80.74(18) O(2)-Cu1-O(3): 89.741(12), 167.53(12)

O(1a)-Co-O(4): 106.54(19) O(3)-Cu1-O(3): 88.822(18)

O(1c)-Co-O(4): 106.0(4) O(4)-Cu1-O(2): 93.06(8)

O(4)-Co-O(4): 145.4(4) O(4)-Cu1-O(2): 99.39(9)


79

Figure 3.6: Refined structure of Y0.95Sr2.05Cu1.7Co1.3O7+δ at 740 °C: a) orthorhombic unit cell with

atoms plotted as 99% displacement ellipsoids; b) the tetrahedral CoO4 layer viewed along the stacking

axis with atoms plotted isotropically to allow representation of oxide occupancies by pie-charts.

Grey = Y3+

, green = Sr2+

, blue = Co2+

, orange = Cu2+

and red = O2-

.

3.4 DC conductivity measurements

DC conductivity data were collected on single phase materials of

Y1−xSr2+xCu3−yCoyO7+δ (1≤ x ≤ 1.3, y = 0.03 and 0.05). In order to obtain dense pellets for the

measurements, single phase samples were ball-milled using a FRITSCH Pulverizette 7

planetary ball mill for 12 h in ethanol using zirconium oxide balls. The resulting fine powder

was mixed with a 2% polyvinyl alcohol (PVA) solution and dried at 80 °C overnight. The

samples were then pressed into pellets using cold isostatic pressing (Section 2.4.3) and

sintered in alumina crucibles under ambient air in a box furnace at 1050 °C with a heating

rate of 5 °C min−1

(with no dwell time at temperature). This resulted in pellets with densities

greater than 93% of the theoretical density, which was checked using Archimedes’ Principle

balance (Section 2.4.4).

The DC conductivity data were collected using the standard four probe technique

(Section 2.4.2) on a bar with approximate dimensions of 2 × 2 × 13 mm3. Gold paste was


80

used to bond the gold wires in a four-in-a-line contact geometry. DC conductivity data were

collected (using a Keithley 617 Programmable Electrometer) in air as a function of

temperature (300 – 800 °C) at each 50 ° with a 90 min equilibrium time both on heating and

cooling. Sample purity before and after the measurements was confirmed by using X-ray

diffraction.

The data collection of each of the samples started at 600 °C. Then the samples were

heated to 800 °C cooled down to 300 °C and re-heated to 800 °C. The plots with DC

conductivity data of Y1-ySr2+yCu3-xCoxO7+δ (1 ≤ x ≤ 1.3, y = 0.03 and 0.05) are shown in

Figure 3.7. The materials showed semiconducting behaviour, decreasing the resistivity by

increasing temperature. Increasing the Co doping level increases the measured DC

conductivity. The Arrhenius plot of conductivity for YSr2Cu2CoO7+δ (x = 1 sample) shows an

abrupt increase above 800 °C which has been attributed to an orthorhombic-tetragonal phase

transition.69

A significant increase in conductivity above 700 °C is observed for all of the

measured samples (Figure 3.7b). Y0.95Sr2.05Cu1.7Co1.3O7+δ exhibited the highest conductivity

with a value of 191 S cm−1

at 800 °C. This is one order of magnitude greater than the value

obtained for Y0.97Sr2.03Cu2CoO7+δ (20.7 S cm−1

at 800 °C) and the value reported for the

parent, YSr2Cu2CoO7+δ material in previous work (15 S cm−1

at 800 °C).69

The activation

energies (Table 3.8) were obtained from a linear fit of the Arrhenius plots in the 300 to

600 °C temperature region. The activation energy increases with increasing Co doping level,

from a value of 0.14 eV for x = 1 to 0.26 eV for x = 1.3. This results in a more modest

increase in conductivity with Co doping at lower temperatures. At 650 °C, the DC

conductivity of Y0.95Sr2.05Cu1.7Co1.3O7+δ is four times higher than in Y0.97Sr2.03Cu2CoO7+δ.

The PXRD patterns of Y0.95Sr2.05Cu1.7Co1.3O7+δ collected before and after the DC

measurement are shown in Figure 3.8.


81

Figure 3.7: DC conductivity of Y1−ySr2+yCu3−xCoxO7+δ. a) linear temperature scale of DC conductivity

as a function of Co doping level; b) logarithmic scale of DC conductivity as a function of temperature.

Table 3.8: Values of DC conductivity at 800 °C of Y1-ySr2+yCu3-xCoxO7+δ materials with the values of

activation energies (Ea) compared with reported value for YSr2Cu2CoO7+δ.69

x DC conductivity at 800 °C (S cm−1

) Ea (eV)

1 – reported69

15 –

1 20.7 0.14

1.1 26.8 0.18

1.2 29.3 0.25

1.25 36.3 0.24

1.3 191 0.26


82

Figure 3.8: PXRD of Y0.95Sr2.05Cu1.7Co1.3O7+δ – before and after DC measurement with the mark of

the most intense reflection of triple perovskite phase and gold.

3.5 Thermal stability

Since thermal instability of the related cuprate materials was a significant

problem,68,70

thermal stability behaviour of the studied ap materials was inspected. Thermal

stability tests of the most conducting material Y0.95Sr2.05Cu1.7Co1.3O7+δ were investigated for

different dwelling times (6 h and 24 h) at various temperatures (900, 950, and 1000 °C). Prior

to heating, 0.1 g of Y0.95Sr2.05Cu1.7Co1.3O7+δ material was weighed out and pressed into 6 mm

diameter pellets. The samples were sintered in alumina crucibles in a box furnace at ambient

air with heating rate 5 °C min−1

. After the tests each of the samples was re-ground manually

using a pestle and a mortar. PXRD patterns collected after the annealing (Figure 3.9 and

Figure 3.10) indicated that no impurity phases were present. For further investigation,

Rietveld refinements in Topas152

were done. The values of lattice parameters of the samples

after the tests are summarised and compared with as made powder in Table 3.9. The changes


83

of lattice parameters after the thermal stability tests were negligible compared to the values

obtained from as made Y0.95Sr2.05Cu1.7Co1.3O7+δ (Figure 3.11).

Figure 3.9: PXRD data of Y0.95Sr2.05Cu1.7Co1.3O7+δ with the most intense 3ap reflections labelled. The

data were collected after the 6 h thermal stability tests carried out at various temperatures.


84

Figure 3.10: PXRD data of Y0.95Sr2.05Cu1.7Co1.3O7+δ with the most intense 3ap reflections labelled.

The data were collected after the 24 h thermal stability tests carried out at various temperatures.

Table 3.9: Summary of the lattice parameters of Y0.95Sr2.05Cu1.7Co1.3O7+δ obtained by Rietveld

refinement from PXRD data collected after the thermal stability tests at different temperatures

compared with as made powder.

Conditions a (Å) b (Å) c (Å) V (Å3)

As made powder 22.7210(8) 5.43220(5) 5.43140(4) 670.37(2)

900 °C, 6 h 22.7181(8) 5.4298(5) 5.4303(5) 669.8(1)

900 °C, 24 h 22.7261(8) 5.4303(6) 5.4311(5) 670.2(1)

950 °C, 6 h 22.7062(5) 5.4293(5) 5.4286(5) 669.2(1)

950 °C, 24 h 22.7131(7) 5.4305(9) 5.4304(9) 669.7(2)

1000 °C, 6 h 22.7207(3) 5.4306(8) 5.4299(8) 669.9(2)

1000 °C, 24 h 22.7210(7) 5.4300(6) 5.4304(7) 670.0(1)


85

Figure 3.11: Lattice parameters of Y0.95Sr2.05Cu1.7Co1.3O7+δ obtained from Rietveld refinements using

laboratory PXRD data a) to c) show the samples annealed for 6 h and d) to f) show the samples

annealed for 24 h compared to as made material Y0.95Sr2.05Cu1.7Co1.3O7+δ. a) and d) show the a cell

parameters; b) and e) show the b and c cell parameters and c) and f) show the cell volume.

Simultaneously with thermal stability tests, a TGA experiment with

Y0.95Sr2.05Cu1.7Co1.3O7+δ material was conducted. The sample was annealed in an alumina pan

with the heating rate 5 °C min−1

in air to 900 °C and cooled down to room temperature with

the same heating rate. Figure 3.12 shows the weight loss changes during the heating and

cooling process. There is no significant weight loss, which could be assigned to any

decomposition process. The small change between the 25 – 150 °C temperatures is mostly

due to the moisture taken by the material before the measurement. The PXRD data collected

after the TGA experiment shows no impurity phases.


86

Figure 3.12: TGA measurement of Y0.95Sr2.05Cu1.7Co1.3O7+δ material at ambient air on heating from

RT to 900 °C and cooling back to room temperature.

3.6 Chemical compatibility of Y0.95Sr2.05Cu1.7Co1.3O7+δ with electrolytes

It is important to understand the reactivity of potential cathode materials with an

electrolyte, as any reaction could create an interface containing impurities with properties that

are detrimental to the operation of the SOFC. The stability of the Y0.95Sr2.05Cu1.7Co1.3O7+δ

cathode with SDC20, GDC10, and LSGM electrolytes was tested by annealing a thoroughly

ground mixture of Y0.95Sr2.05Cu1.7Co1.3O7+δ and electrolyte at temperatures of 900, 950 and

1000 °C for one week. Compatibility tests were carried out using SOFC electrolytes of

followed properties:

SDC20, Ce0.80Sm0.20O2−x, TC Grade, surface area: 6.0 m2 g

−1, purchased from

fuellcellmaterial.com


87

GDC10, Gd0.10Ce0.90O2-x, TC Grade, surface area: 5.9 m2 g

−1, purchased from

fuellcellmaterial.com

LSGM, La0.9Sr0.1Ga0.8Mg0.2Oxide, 99.9%, surface area: 0.50 m2 g

−1, purchased

from PRAXAIR Surface Technologies

Prior to annealing in ambient air, mixture of cathode and electrolyte (1:1 mass ratio,

0.05 g each) was ground using a pestle and a mortar and pressed into a 6 mm pellet and put in

an alumina crucible in a box furnace. After one week annealing the pellets of the cathode and

electrolyte were cooled down to room temperature (5 °C min−1

rate). The composition and

the lattice parameters of the materials after the tests were investigated from the collected

PXRD data.

Y0.95Sr2.05Cu1.7Co1.3O7+δ reacts with LSGM at all the studied temperatures (Table 3.10)

and is unsuitable for use with this electrolyte. The reaction gave rise to the formation of

perovskite (LaCuO3, La0.6Sr0.4CoO3, LaSrCuGaO5) and Y2Cu2O5 impurities.

After annealing the SDC20 and GDC10 electrolytes with Y0.95Sr2.05Cu1.7Co1.3O7+δ, at

900 °C, PXRD patterns showed that no impurity phases were present, indicating that no

reaction between the electrolyte and Y0.95Sr2.05Cu1.7Co1.3O7+δ had occurred (as shown in

Figure 3.13). Increasing the temperature to 950 °C resulted in observation of

Y1.5Ce0.5Sr2Cu2CoO9+δ impurities in the PXRD data, indicating reaction between the

Y0.95Sr2.05Cu1.7Co1.3O7+δ material and the electrolyte. This is in agreement with previous

electrolyte compatibility studies of YSr2Cu2CoO7+δ with GDC10 which also showed the

presence of Y1.5Ce0.5Sr2Cu2CoO9+δ impurity phases after annealing at 1000 °C for 1 week.69

The compatibility test show a possibility to use Y0.95Sr2.05Cu1.7Co1.3O7+δ material with

ceria based electrolytes at lower temperatures (900 °C and less – depends on adhesion step

required for processing). Table 3.11 shows the comparison of the lattice parameters of as

made powder with the lattice parameters obtained from the Rietveld refinements of the data

collected after the compatibility tests with ceria based electrolytes. There are small changes in

volume and lattice parameters but since there were no additional phases observed after one

week annealing, these electrolytes (SDC20, GDC10) were used for further work including

AC impedance spectroscopy on symmetrical cells (see Section 3.7).


88

Figure 3.13: PXRD data of Y0.95Sr2.05Cu1.7Co1.3O7+δ after one week of annealing with GDC10 at 900

and 1000 °C. The reflections of the different phases are identified as follows: + for GDC10, * for

Y1.5Ce0.5Sr2Cu2CoO9+δ, † for Co3O4, and ‡ for Gd0.18Y1.82O3.

Table 3.10: Summary of the reactivity and additional phases observed after compatibility tests of

Y0.95Sr2.05Cu1.7Co1.3O7+δ with SDC20, GDC10, and LSGM respectively at various temperatures for

one week.

Temp. (°C) SDC20 GDC10 LSGM

900 No reaction No reaction SrLaGa3O7

950 Y1.5Ce0.5Sr2Cu2CoO9+δ Y1.5Ce0.5Sr2Cu2CoO9+δ SrLaGa3O7

1000 Y1.5Ce0.5Sr2Cu2CoO9+δ Y1.5Ce0.5Sr2Cu2CoO9+δ Multi phased


89

Table 3.11: Comparison of the lattice parameters of Y0.95Sr2.05Cu1.7Co1.3O7+δ after compatibility test

with SDC20 and GDC10 respectively with as made powder and reported YSr2Cu2CoO7+δ phase.193

Material a (Å) b (Å) c (Å) V (Å3)

+SDC20 at 900°C 22.746(7) 5.4524(7) 5.4114(6) 671.11(7)

+GDC10 at 900°C 22.653(4) 5.4500(3) 5.4169(2) 668.76(5)

Y0.95Sr2.05Cu1.7Co1.3O7+δ 22.7294(4) 5.4394(1) 5.4179(8) 669.83(2)

YSr2Cu2CoO7+δ 193

22.7987(2) 5.45150(5) 5.40890(5) 672.26(1)

3.7 AC impedance spectroscopy of Y0.95Sr2.05Cu1.7Co1.3O7+δ

3.7.1 AC impedance data at 500 ‒ 800 °C

To further Y0.95Sr2.05Cu1.7Co1.3O7+δ investigate as a cathode material for SOFCs, AC

impedance spectroscopy measurements were performed (see Section 2.5 for more details).

Since the compatibility tests showed no reaction between the potential cathode material and

ceria based electrolytes (GDC10, SDC20) the symmetrical cells were fabricated of ceria

based electrolytes and of Y0.95Sr2.05Cu1.7Co1.3O7+δ or Y0.97Sr2.03Cu2CoO7+δ cathode. Pellets of

SDC20 and GDC10 (with densities over 95%) were prepared by annealing in air at 1400 °C

for 5 h. The surface of the electrolytes was polished with SiC paper using a polishing

machine (Struers Tegramin-30). Cathode inks were produced from powders of

Y1−ySr2+yCu3−xCoxO7+δ ball milled with ethanol for 12 h. Following evaporation of the

solvent, the dried powder was mixed with an organic binder (Paste Vehicle purchased from

Fuel Cell Materials) in a powder : binder mass ratio of 0.67 : 0.33, before milling the mixture

for a further 12 h. The resulting ink was used for screen printing to both sides of the

electrolyte and the adherence of the ink to the surface of the electrolyte was achieved after

calcinations at 800 or 900 °C for 1 h. Various processing conditions (drying temperature after

screen printing, number of screen printed cathode layers, different level of polishing of the

electrolyte surface) were applied to improve the electrochemical performance of the cells.

Gold mesh with gold paste was used as current collector for the AC impedance

spectroscopy measurements. AC impedance spectroscopy was carried out using a Solatron

1255B Frequency Response Analyzer and Solatron SI 1287 Electrochemical Interface. The

data were collected over a frequency range of 1 MHz – 0.1 Hz with a modulation potential of


90

10 mV, over the temperature of 500 – 800 °C in static air with a 90 min step every 50 °C.

Impedance measurements and corresponding equivalent circuit modelling were performed

using the ZPlot and ZView software196

. The ASR of the cathode was calculated by

normalising the measured resistance for the electrode area and dividing by two to take into

account the symmetry of the cell (Equation 2.34, Section 2.5.2).

Figure 3.14 shows the AC impedance response of the cathode at different

temperatures, from which ASR values were obtained. The ohmic impedance associated with

the GDC10 and SDC20 electrolytes has been subtracted from the spectra and the AC

impedance arcs normalised to zero on the real (x) axis for easy comparison.

Figure 3.14: Nyquist plot of measured AC impedance arcs of the cell made of

Y0.95Sr2.05Cu1.7Co1.3O7+δ (dried at 800 °C) and unpolished GDC10 at temperatures 600 ‒ 700 °C.

It is generally accepted that the cathode processing conditions can have a dramatic

effect on the physical properties (e.g. porosity, surface area, interface, adherence) of cathodes

and therefore altering the processing conditions (such as drying temperature after screen

printing, polishing of the electrolyte surface and the number of printed layers) may cause

different ASR values which are difficult to predict.47,62

To investigate this, various processing

conditions for symmetrical cells of Y0.95Sr2.05Cu1.7Co1.3O7+δ were applied, as outlined in


91

Table 3.12. Throughout this chapter, we will use a labelling system (e.g. GDC Cell 1) for the

symmetrical cells as outlined in Table 3.12. Here we will use GDC or SDC to denote the

electrolyte used, whilst cell followed by a number (1-4) will indicate the processing

conditions used (e.g. GDC Cell 1 describes a symmetrical cell made from GDC10, with a

polished electrolyte surface and a cathode drying temperature of 800 °C). Eight symmetrical

cells (4 with GDC10, 4 with SDC20 electrolyte) were prepared and measured on first heating

(600 – 800 °C), cooling (800 – 600 °C) and second heating (600 – 800 °C). Figure 3.15 and

Table 3.13 show the ASR values for all cells collected during the first heating. The lowest

ASR values of 0.08 Ω cm2 at 700 °C, were given by GDC Cell 3 (dried at 800 °C and

polished electrolyte surface; Table 3.13, Figure 3.15a and Figure 3.16).

Table 3.12: Processing conditions for symmetrical cells containing one layer of

Y0.95Sr2.05Cu1.7Co1.3O7+δ, as applied to two ceria based electrolytes (GDC10 and SDC20). The cells

will be referred to as e.g. “GDC Cell 1”, which would refer to a cell with GDC10 electrolyte

processed under condition (1).

Symmetrical cell Drying temperature

(°C)

Electrolyte surface Printed layers

(1) 800 Unpolished 1

(2) 900 Unpolished 1

(3) 800 Polished 1

(4) 900 Polished 1

Figure 3.15: AC impedance spectroscopy data of symmetrical cells made of Y0.95Sr2.05Cu1.7Co1.3O7+δ

as a cathode and using different processing conditions: (1) to (4) ‒ see Table 3.12, with a) GDC10 and

b) SDC20 electrolyte.


92

Table 3.13: Values of ASR and Ea calculated from the AC impedance data collected on the

symmetrical cells made by altering processing conditions (Table 3.12).

Processing

condition

ASR (Ω cm2) Ea

(eV)

(650 °C) (700 °C) (750 °C) (800 °C)

GDC Cell 1 0.287 0.136 0.122 0.090 0.35

GDC Cell 2 0.236 0.184 0.092 0.070 0.25

GDC Cell 3 0.207 0.081 0.065 0.057 0.27

GDC Cell 4 0.235 0.155 0.083 0.073 0.26

SDC Cell 1 0.530 0.239 0.172 0.115 0.61

SDC Cell 2 0.311 0.174 0.152 0.094 0.36

SDC Cell 3 0.243 0.117 0.083 0.063 0.32

SDC Cell 4 0.206 0.172 0.156 0.114 0.20

Although the Y0.95Sr2.05Cu1.7Co1.3O7+δ material shows promising ASR values for

symmetrical cells with both GDC10 and SDC20 (Figure 3.15, Table 3.13), there is an

increase of ASR values of one order of magnitude (in the 600 – 700 °C region) between the

first and second heating cycles as shown in Figure 3.16. The same trend was observed for all

of the symmetrical cells shown in Figure 3.15. This is presumably the reason for the

deviation observed in Arrhenius-type plots (Figure 3.15 and Figure 3.16). PXRD data of the

symmetrical cells after the AC impedance (Figure 3.17) confirmed the presence of other

phases (Y2Cu2O5, Co3O4, Sr2CeO4), and this has been attributed to in-situ electrochemical

decomposition. In-situ electrochemical decomposition has also been observed in

YBa2Cu3O7−δ at low current densities (≈ 20 mA cm−2

).68

This is different compared to the

thermal stability tests of Y0.95Sr2.05Cu1.7Co1.3O7+δ (Section 3.5), where no impurity phases

were observed. In order to improve the stability of Y0.95Sr2.05Cu1.7Co1.3O7+δ, AC impedance

measurements were carried out with a reduced maximum temperature of 650 °C.


93

Figure 3.16: The comparison of ASR values between the 1st and 2

nd heating of the cell made of

polished GDC, dried at 800 °C (GDC Cell 3).

Figure 3.17: XRD data of Y0.95Sr2.05Cu1.7Co1.3O7+δ before and after the AC impedance spectroscopy

(GDC Cell 3) with the marked reflections of GDC10 electrolyte, gold, the main reflections of

additional phases presented after the measurement: +Y2Cu2O5, †Co3O4, *Sr2CeO4 and missing (200)

reflection after the AC impedance spectroscopy.


94

3.7.2 AC impedance data for the dwelling at 650 °C

As mentioned in previous section, Y0.95Sr2.05Cu1.7Co1.3O7+δ material exhibited

promising ASR values for symmetrical cells with ceria based electrolytes, but also in-situ

electrochemical decomposition at higher temperatures (800 °C). Current research for SOFCs

cathodes is focused for the materials working in intermediate temperature range

(600 − 800 °C). Materials with a good cathode performance (processing, stability,

conductivity, etc.) at the temperature close to 600 °C are of significant interest for both

research and industry.9,17

Hence, the AC impedance of Y0.95Sr2.05Cu1.7Co1.3O7+δ material was

carried out at temperatures below 800 °C. From the previous results on symmetrical

cells (Table 3.13), the triple perovskite showed promising ASR values at 650 °C. The other

processing factor, which has not been considered is the thickness of the cathode layer of a

symmetrical cell. The thickness can be changed by the number of screen printed layers of the

cathode material.

The symmetrical cells were prepared following the procedure mentioned in

Section 4.7.1. The processing conditions (polishing, drying temperature and type of

electrolyte) were selected from previously measured cell – from the one with the best ASR

values. Thus symmetrical cells of the Y0.95Sr2.05Cu1.7Co1.3O7+δ material and polished GDC10

dried at 800 °C (GDC Cell 3) were constructed. The cells were constructed with 1 or 6 layers

(GDC Cell 3 L1 and GDC Cell 3 L6 respectively) of screen printed Y0.95Sr2.05Cu1.7Co1.3O7+δ

in order to see if there was any influence of cathode thickness on cathode performance and

electrochemical stability.

Figure 3.18 shows the ASR values obtained whilst dwelling at 650 °C for 12 h. For

both 1 and 6 layered cells, there is no significant increase of ASR observed (Table 3.14),

indicating that the number of screen printed cathode layers does not affect the

electrochemical stability. The total increase of ASR for both the 1 and 6 layered cells after

12 h dwelling was within 5% of the starting values. The Y0.95Sr2.05Cu1.7Co1.3O7+δ symmetrical

cell with 6 layers shows lower (improved) ASR values compared to the cell with 1 screen

printed cathode layer. The starting ASR value of 0.17 Ω cm2 for the 6 layered cell of the

x = 1.3 sample increased after 12 h annealing to 0.18 Ω cm2. After the measurement, PXRD


95

confirmed that the sample was phase pure (Figure 3.19), indicating that in-situ

electrochemical decomposition had been avoided.

Figure 3.18: ASR values obtained when annealing GDC Cell 3 L1 (1 cathode layer) and GDC Cell 3

L6 (6 cathode layers) at 650 °C for 12 h.

Table 3.14: Comparison of the ASR values obtained after 1 h and 12 h dwelling at 650 °C of

symmetrical cells of Y0.95Sr2.05Cu1.7Co1.3O7+δ and GDC10 (GDC Cell 3 L1 and GDC Cell 3 L6).

Symmetrical cell ASR (Ω cm2) (1 h) ASR (Ω cm

2) (12 h)

GDC Cell 3 L1 (1 layer) 0.39 0.41

GDC Cell 3 L6 (6 layers) 0.17 0.18


96

Figure 3.19: PXRD data of Y0.95Sr2.05Cu1.7Co1.3O7+δ after AC impedance at 650 °C for 12 h. The main

reflections of GDC10 and Au are indicated by + and * respectively. All other reflections are indexed

to Y0.95Sr2.05Cu1.7Co1.3O7+δ.

3.7.3 SEM study of symmetrical cells

The morphology of symmetrical cell of Y0.95Sr2.05Cu1.7Co1.3O7+δ cathode with GDC

Cell 3 L6 (6 layered) was evaluated by SEM (instruments details available in Section 2.3).

The impedance data of the cells are shown in Section 3.7.2 and suggests that alternating the

number of screen-printed layers of the cathode material can improve the ASR values.

Studying the cell by SEM provides a better insight into the real thickness of the screen-

printed cathode layer. Any other additional reaction on boundaries between the layers of the

symmetrical cells can be also observed.

Figure 3.20 shows the SEM data with different magnifications collected on the

symmetrical cell made of 6 layers of Y0.95Sr2.05Cu1.7Co1.3O7+δ (with GDC Cell 3 L6) after the

AC impedance. We can clearly see all three parts of the symmetrical cell (Au current


97

collector, cathode, and electrolyte layer). No additional interfacial layers were observed. The

total cathode thickness (for a total of 6 screen printed layers) is 40 μm.

Figure 3.20: Cross-sectional image of the symmetrical cell that exhibited best performance

measurements (with ASR650 °C = 0.18 Ω cm2) after AC impedance spectroscopy data collection. The

cell was constructed of 6 layer of screen printed Y0.95Sr2.05Cu1.7Co1.3O7+δ cathode and GDC Cell 3 L6

is shown with different magnifications: a) 50 μm; b) 30 μm.

3.8 AC impedance spectroscopy of Y0.97Sr2.03Cu2CoO7+δ

The previous section (Section 3.7) was focused on the impedance measurement of the

most conductive Y0.95Sr2.05Cu1.7Co1.3O7+δ material (Section 3.4). AC impedance data from

Y0.97Sr2.03Cu2CoO7+δ phase were collected to compare with the data obtained on x = 1.3

material. GDC Cell 1 with one screen printed cathode layer was selected for this comparative

study. The ASR values and activation energies obtained from the AC impedance data of

Y0.97Sr2.03Cu2CoO7+δ phase are shown and compared with Y0.95Sr2.05Cu1.7Co1.3O7+δ in Figure

3.21 and Table 3.15. At lower temperatures (600 and 650 °C), the ASR values of

Y0.95Sr2.05Cu1.7Co1.3O7+δ are one order of magnitude lower than those obtained for

Y0.97Sr2.03Cu2CoO7+δ. The improvement of the electrochemical performance of

Y0.95Sr2.05Cu1.7Co1.3O7+δ compared to x = 1 sample was achieved for the whole temperature

range. The activation energy value for the x = 1.3 material is almost of one half of the

activation energy obtained from the data collected on the parental x = 1 phase.


98

Figure 3.21: ASR values comparison of symmetrical cells of Y0.95Sr2.05Cu1.7Co1.3O7+δ and

Y0.97Sr2.03Cu2CoO7+δ, with GDC Cell 1 (800 °C drying temperature and unpolished electrolyte

surface). The activation energy shown was calculated from the values in the 650 to 800 °C region.

Table 3.15: Values of ASR and activation energies (calculated for 650 – 800 °C) for symmetrical

cells of Y0.95Sr2.05Cu1.7Co1.3O7+δ and Y0.97Sr2.03Cu2CoO7+δ, with one screen printed cathode layers using

GDC Cell 1 (800 °C drying temperature and unpolished electrolyte surface).

Co doping level ASR (Ω cm2) Ea (eV)

(650 °C) (700 °C) (750 °C) (800 °C)

x = 1 2.82 1.14 0.25 0.13 0.62(9)

x = 1.3 (1st heating) 0.29 0.14 0.12 0.09 0.35(8)

x = 1.3 (2nd

heating) 2.26 0.91 0.28 0.12 0.57(3)


99

3.9 CO2 stability tests of Y0.95Sr2.05Cu1.7Co1.3O7+δ

The presence of CO2 in the air feedstock gas used at the cathode means that potential

cathode materials must be tested for their stability in CO2 containing atmospheres (typical

CO2 levels in air are in the region of 400 ppm by volume). Reaction of the cathode material

with CO2 in air at SOFC operating temperatures can lead to performance loss due to the

formation of carbonate containing impurities.

The stability of Y0.95Sr2.05Cu1.7Co1.3O7+δ was tested at low and high CO2

concentrations. Approx. 0.3 g of the material was weighed out, pressed into an 8 mm pellet

and put into an alumina crucible. A mixture of 1% CO2 in Ar was passed over

Y0.95Sr2.05Cu1.7Co1.3O7+δ for 12 h at temperatures of 600, 650, and 700 °C. The resultant

PXRD data are shown in Figure 3.22. No impurity phases were observed at 600 °C,

indicating that at this temperature Y0.95Sr2.05Cu1.7Co1.3O7+δ is stable under low CO2. At the

higher temperatures of 650 and 700 °C, SrCO3 and CuO impurities were observed in the

diffraction patterns. Lattice parameters and volume of the unit cell obtained from Rietveld

refinements (Table 3.16-3.18) show expansion of the unit cell which may be due to the

reduction of Co and Cu. To investigate this annealing in pure Ar atmospheres were

performed using the same procedure and amounts of material as in CO2 stability tests. There

were no other phases presented. Table 3.16-3.18 show the same trend of expanded lattice

parameters and volumes both in the case of annealing in pure Ar and in mixture Ar + 1%

CO2, meaning that these changes are due to the reduction and are not related to the presence

of CO2.

In order to test the behaviour of Y0.95Sr2.05Cu1.7Co1.3O7+δ in a pure CO2 atmosphere,

TGA was performed on the sample whilst annealing at 650 °C for 12 h under pure CO2. TGA

showed an increase in mass over time (Figure 3.23), indicating a reaction with CO2,

compared to no change in mass when the TGA was carried out in air (Figure 3.12). This was

confirmed by PXRD data (Figure 3.22 and Figure 3.24) collected after the TGA experiment,

which showed the presence of SrCO3 and CuO in the diffraction pattern, indicating

decomposition of Y0.95Sr2.05Cu1.7Co1.3O7+δ.


100

Figure 3.22: PXRD data of Y0.95Sr2.05Cu1.7Co1.3O7+δ after exposure to 1% CO2 in Ar (at temperatures

of 600, 650 and 700 °C) and to pure CO2 at 650 °C, * denotes the CuO reflections, + indicates SrCO3

phase.

Table 3.16: Comparison of lattice parameters of the Y0.95Sr2.05Cu1.7Co1.3O7+δ sample after the 12 h

annealing at 600 °C in Ar + 1% CO2 and pure Ar atmosphere with as made powder, obtained from

Rietveld refinements using Topas.152

600°C a(Å) b(Å) c(Å) V(Å3)

As made 22.7210(8) 5.43220(5) 5.43140(4) 670.37(2)

Ar/1% CO2 22.7470(4) 5.45489(8) 5.42001(9) 672.53(2)

Ar 22.7451(3) 5.45134(7) 5.41735(7) 671.70(1)


101




650°C a(Å) b(Å) c(Å) V(Å3)

As made 22.7210(8) 5.43220(5) 5.43140(4) 670.37(2)

Ar/1% CO2 22.7420(4) 5.4594(1) 5.4204(1) 672.99(2)

Ar 22.7320(3) 5.46414(7) 5.42150(8) 673.41(2)




700°C a(Å) b(Å) c(Å) V(Å3)

As made 22.7210(8) 5.43220(5) 5.43140(4) 670.37(2)

Ar/1% CO2 22.738(1) 5.4620(2) 5.4204(2) 673.22(5)

Ar 22.7227(3) 5.46804(6) 5.42137(6) 673.60(1)

Figure 3.23: TGA experiment of Y0.95Sr2.05Cu1.7Co1.3O7+δ material, after exposure to pure CO2 at

650 °C for 12 h.


102

Figure 3.24: Rietveld refinement of Y0.95Sr2.05Cu1.7Co1.3O7+δ sample collected after the TGA

experiment carried out in pure CO2 at 650 °C for 12 h.

3.10 Thermal expansion studies of Y0.95Sr2.05Cu1.7Co1.3O7+δ

Parent YSr2Cu2CoO7+δ material showed an orthorhombic-tetragonal phase transition

at above 800 °C.69

The unit cell contraction at the phase transition was also observed using

dilatometry.70

Any phase transition of a cathode material may be detrimental for a SOFC

device causing mechanical stress of the cell. The thermal expansion studies were carried out

on the most conductive Y0.95Sr2.05Cu1.7Co1.3O7+δ material using a dilatometry measurement.

Prior to the dilatometry measurement, Y0.95Sr2.05Cu1.7Co1.3O7+δ was weighed out,

pressed into a 6 mm pellet and sintered in ambient air at 1050 °C (following the synthesis

conditions in Section 3.2). Linear thermal expansion of the sample was measured in the range

20 – 850 °C with the heating rate 10 °C min−1

. The sample was then cooled down to 20 °C

and heated again in order to see any changes in behaviour between two heating cycles. The

data obtained during the first heating cycle are shown in Figure 3.25. There is a slight


103

decrease of the thermal expansion recorded during the second heating compared to the first

one (Table 3.19). In both cycles, a linear change of expansion is observed without any sharp

change in the values especially above 800 °C, where the phase transition was reported.70

To

ensure there is no phase transition, first derivation (dL/dt) of the linear thermal expansion can

be plotted (Figure 3.25). There is no peak presented. That indicates no phase transition (apart

slow phase transitions) in Y0.95Sr2.05Cu1.7Co1.3O7+δ in the temperature range 20 – 850 °C.

Table 3.19 shows the value of the linear thermal parameter – α obtained from the

linear thermal expansion data (see Chapter 2.10) for the first and second heating.

The α values for selected temperatures (8.4-8.8 × 10−6

K−1

for the temperatures between 600

and 850 °C) are all within narrow range and they increase with the increasing temperature.

Reported dilatometry measurement of YSr2Cu2CoO7+δ material70

clearly showed phase

transition with the values of thermal coefficient (TEC): 13.0 × 10−6

K−1

above and

12.3 × 10−6

K−1

below phase transition. Measured values for Y0.95Sr2.05Cu1.7Co1.3O7+δ material

are lower than the ones for the starting phase. The difference can be assigned to different

Y/Sr ratio and higher Co-amount.

Figure 3.25: Linear thermal expansion of Y0.95Sr2.05Cu1.7Co1.3O7+δ measured in 20 ‒ 850 °C range.

Black data points indicate the obtained values for the 1st heating whilst the red points show the first

derivation of the collected data (dL/dt).


104

Table 3.19: Linear thermal expansion coefficients (α, 100 °C – selected temperature) of the

Y0.95Sr2.05Cu1.7Co1.3O7+δ material calculated from the data collected for both heating cycles.

Temperature (°C) α (K−1

) – 1st cycle α (K

−1) – 2

nd cycle

600 8.45 × 10−6

8.43 × 10−6

650 8.54 × 10−6

8.52 × 10−6

700 8.63 × 10−6

8.61 × 10−6

750 8.70 × 10−6

8.69 × 10−6

800 8.76 × 10−6

8.76 × 10−6

850 8.78 × 10-6

8.81 × 10−6

3.11 Discussion and conclusions

An increase of the cobalt content in the layered YSr2Cu2CoO7+δ perovskite was

achieved by altering the Y:Sr cationic ratio, resulting in materials of the form

Y1−ySr2+yCu3−xCoxO7+δ (1≤ x ≤ 1.3, y = 0, 0.03 and 0.05). The higher the Co content, the

higher the Sr content is required to obtain single phase samples. This most probably is a

result of charge balance and is consistent with cobalt commonly averaging higher formal

oxidation state than copper in perovskite oxides, or more generally in transition metal oxides.

An oxygen content of YSr2Cu2CoO7.03(4) was observed for the parent non substituted phase

by Slater et al.69

using thermogravimetric methods. The oxygen content is independent of the

level of Co substitution which suggest that a good stability of the peculiar oxygen vacancy

ordering pattern, associated here with the O7 stoichiometry, is responsible for the variable Y

: Sr ratios needed to obtain phase pure samples as the Co content varies. Indeed, at a fixed

oxygen content, incorporation of extra Co needs more Sr to reduce the total charge on the A-

site, in order to compensate the increased charge on the B-site. Assuming a copper oxidation

state of (+2), the Co oxidation state in Y0.95Sr2.05Cu1.7Co1.3O7+δ is lower (+2.8) than the Co3+

found in the parent YSr2Cu2CoO7+δ phase and Co substitution therefore allows the formation

of mixed valent compounds.

The retention of the O7 stoichiometry throughout the series of Co enriched materials

allows the production of mixed valent compounds with a higher content of charge carriers,


105

which may imply easier electron delocalisation. This is reflected in the DC conductivity data,

which show that we were able to produce an enhancement of the conductivity by one order of

magnitude when comparing Y0.95Sr2.05Cu1.7Co1.3O7+δ to Y0.97Sr2.03Cu2CoO7+δ. The obtained

values for Y0.95Sr2.05Cu1.7Co1.3O7+δ are lower than those for the pure Co containing

La0.7Sr0.3CoO3−δ, which has a total conductivity of 1650 S cm−1

at 800 °C,197

but they are

comparable with dominant cathode materials such as La0.6Sr0.4Fe0.8Co0.2O3−δ (LSCF) which

has a total conductivity of 280 S cm−1

at 800 °C,40

and is better than the total conductivity of

22.7 S cm−1

(at 800 °C) reported for Ba0.5Sr0.5Co0.8Fe0.2O3−δ.198

Parent YSr2Cu2CoO7+δ showed an orthorhombic-tetragonal phase transition at above

800 °C.69

The dilatometry study of Y0.95Sr2.05Cu1.7Co1.3O7+δ revealed no phase transition

between 20 – 850 °C. The obtained linear thermal expansion coefficient (8.7 × 10−6

K−1

) is

smaller than the TEC of LSCF ((La0.6Sr0.4)1−xCo0.2Fe0.8O3−δ, for x = 0.00-0.15) series,

showing a TEC of 13.8-14.2 × 10−6

K−1

at the temperature range 25 – 700 °C.199,200

TEC

values of measured triple perovskite are comparable with LSM series with TEC of 10.8 ×

10−6

K−1

for La0.79Sr0.2MnO3−x for the temperature range 30 – 800 °C.201

The thermal expansion of Y0.95Sr2.05Cu1.7Co1.3O7+δ material matches well with YSZ

electrolyte (10.5 × 10−6

K−1

for Zr0.85Y0.15O1.93 between 30 – 800 °C)202

and LSGM

electrolytes (10.4 × 10−6

K−1

for La0.8Sr0.2Ga0.9Mg0.1O3−x between 30 ‒ 800 °C).203

Higher

mismatch in TEC values is observed in comparison with Ce-based electrolytes

(12.5 × 10−6

K−1

for Ce0.8Gd0.2O1.90, 30 – 800 °C).202

The suitability of a material as a cathode can be assessed from the ASR values

obtained during cell testing; small ASRs are desirable as the total resistance of the cell is

usually dominated by the contribution from the cathode. The small increase in Co content in

Y1−ySr2+yCu3−xCoxO7+δ results in an improved (decreased) ASR compared with the parent

material. Although the material suffers electrochemical decomposition at the higher

temperatures studied here, we have also shown that as low as 15% Co substitution in the

square based pyramidal site of the initial structure enhances the electrochemical properties of

the 3ap material, again by an order of magnitude at 650 °C, where the material is stable. The

ASR value for Y0.95Sr2.05Cu1.7Co1.3O7+δ at 700 °C is 0.08 Ω cm2 and compares favourably to

the ASR reported for some of the best reported materials such as the pure cobalt containing

double perovskite – SmBa0.5Sr0.5Co2O5+δ which exhibits an ASR value of 0.092 Ω cm2

at 700 °C. 67

Y0.95Sr2.05Cu1.7Co1.3O7+δ shows similar or improved properties compared to other

Co compounds which are considered as intermediate temperature SOFC cathodes, such as the

double perovskite GdBaCo2O5+δ (ASR700 °C = 0.1 Ω cm2),

62 and the tetrahedrally coordinated


106

Co compound YBaCo4O7 (ASR600 °C = 0.40 Ω cm2).

204,205 The ASR values of

Y0.95Sr2.05Cu1.7Co1.3O7+δ at 700 °C are lower than those reported for LSCF

(ASR700 °C = 0.18 Ω cm2).

206 Y0.95Sr2.05Cu1.7Co1.3O7+δ shows a higher ASR than that obtained

for the perovskite Ba0.5Sr0.5Co0.8Fe0.2O3−δ (which has an ASR value of 0.01 Ω cm2 at

700 °C),47

which displays one of the lowest ASR values to date, although has significant

stability problems. The good electrochemical performance we observe here may be related to

some key structural features of the Y1−ySr2+yCu3−xCoxO7+δ. The cobalt enriched compound

retains a crystal structure that should be highly favourable for oxide ion mobility with several

functional sites able to enhance it, such as a fully oxygen vacant layer at the level of Y site or

a half filled layer with ordered vacancies at the level of the Co site.207

At the same time, the

incorporation of Co in the square based pyramid containing layers may provide a catalytic

boost for the oxygen reduction reaction, as this element is generally believed to provide good

catalytic activity17

while square based pyramids are favourable environments for O2 molecule

dissociation.208

Therefore the Y1−ySr2+yCu3−xCoxO7+δ series offers a combination of crystal

chemical features which may provide a high level of electrochemical performance, while

achieving the required values of total conductivity for SOFC cathode materials, owing to the

charge carriers generated by the association of Co doping and Y : Sr ratio changes.

In conclusion, although the Co substitution has substantially improved the

electrochemical properties, further work is needed to enhance the chemical and

electrochemical stability of this class of complex oxides for SOFC cathode applications ‒ this

is not uncommon for the most active electrode systems. This class of candidate cathode

materials is based on a rare example of high spin tetrahedral Co3+

and further investigation is

needed to understand how to optimise materials of this type for fuel cell applications, as this

study tends to suggest that a high level of electrochemical performances could be achievable

within this type of 3ap structured materials.

4 Prediction and synthesis of GBCO related phases

4.1 Introduction

Ordered double perovskites have attracted attention in SOFCs and oxygen storage

research in recent years. Their composition is represented by the general formula AA'Co2O5+δ

(A = RE, Y and A' = Ba, Sr).60

RE-O and A'-O layers alternate along the c-axis (Figure 4.1)

with the A' atoms located in the Co double layers. An interesting feature of double perovskite

is the variety of possible oxygen contents within the range of 0 ≤ δ ≤ 1.0. Generation of

electron-hole pairs, or Co2+

-Co4+

states determines the transport behaviour of AA'Co2O5+δ.209-

214 It has been observed that the oxide-anion vacancies are localized in the RE-O layers due to

the preference of lower coordination number of smaller rare earth element anions compared

to larger Ba2+

or Sr2+

anions. That results in the formation of CoO5 square pyramids and

CoO6 octahedra.60,215

The magnetic and transport properties of AA'Co2O5+δ are highly

dependent on the oxygen content. The oxygen content, δ, controls the mixed valence state of

Co ions216

YBaCo2O5 is an example of the oxygen deficient compound where all Co ions

have the same environment while in RBaCo2O5.5 (R = Tb, Gd) there are two different Co

sites. 217

Figure 4.1: a) Reported structure for GdBaCo2O5.5 210

b)√2×√2×1 cell used for DFT calculation

expanded in a, b lattice parameters. Atoms coloured as follows: gadolinium (purple), barium (green),

cobalt (blue), and oxygen (red).

Chapter 4. Prediction and synthesis of GBCO related phases

108

An example of a material with the double perovskite structure is GdBaCo2O5+δ

(Figure 4.1). Its good mixed conductor properties were first demonstrated by Taskin et

al.218,219

when high rates of oxygen uptake were measured. Clear indication of the importance

of vacancy layers to this process was also demonstrated. GdBaCo2O5+δ (GBCO) as a potential

SOFC cathode materials was introduced by Chang et al.63

A further study showed an

impedance spectroscopy data of GdBaCo2O5+δ and BSCF cathodes with LSGM electrolyte.138

Reported ASR values of symmetrical cells with GBCO at the temperatures above 700 °C are

lower than 0.15 Ω cm2.138

Another electrochemically studied double perovskites were

LnBaCo2O6−δ (Ln = Pr, Nd, Sm, and Gd) phases.220

An orthorhombic to tetragonal phase

transition of LnBaCo2O5+δ (Ln = Pr, Nd and Sm) was evidenced by high temperature X-ray

diffraction data.221

The chemical stability of the LnBaCo2O5+δ materials has been tested against the

common electrolyte materials such as YSZ, LSGM, and GDC. The Ln = La, Nd, and Sm are

stable against LSGM electrolyte, while the Ln = Gd and Y structures show significant side

reaction.222

Similarly to that, the Ln = La and Nd materials are compatible with GDC, but the

Ln = Gd, Y do react with GDC electrolyte.223

Thermal stability and thermal expansion is one

of the problems of Co-containing perovskite cathodes. High values of linear TEC

(αL = 16‒14 × 10−6

K−1

) were also reported for LnBaCo2O5+δ materials.223

However, materials

with medium-size cations (Gd3+

and Y3+

) show TEC values (≈ 16 × 10−6

K−1

) comparable to

some of commonly used electrolytes, especially ceria-based ones.

The work presented in this chapter is focused on double perovskite related materials.

Studied materials, Ln2BaCo2O7 (Ln = Gd, Nd, Ce), consist of a layer of LnBaCo2O5+δ

(Ln = Gd, Nd) and a fluorite layer (CeO2 or Ln2O3, Ln = Gd, Nd). Work with these new

materials is aimed towards potential SOFC cathode materials. By adding a fluorite layer to a

mixed double perovskite conductor an improvement of an ionic conductivity is expected. A

part of the work involves the study of Sr analogues. The materials are studied both by

experimental and theoretical methods. The first part is focused on the prediction of the

stability of the material based on DFT calculations. The structures for theoretical studies are

shown in Figure 4.2. The structural model for the materials with double perovskite and

fluorite layer was based on the reported structure of Y2SrCuFeO6.5 (Figure 4.3)24,25

and it is

discussed in more details in Section 4.2. Relaxations of the individual parts (fluorite, double


109

perovskite layers) were obtained beforehand and compared with the experimental data. The

synthesis of all the materials was attempted using solid state methods. The study was also to

target materials with the different fluorite layer composed of Ln2O3 oxide (Ln = Gd, Nd).

4.2 Computational methods

All the calculations were carried out using the plane wave DFT package, Vienna Ab-

initio Simulation Package (VASP) version 4.6.26178,224

with Perdew, Burke and Ernzerhof

(PBE)225

exchange correlation functional used mainly in this work. Different exchange

correlation functional: LDA226,227

and PW91228

were used at the beginning of the work

(Section 4.4.1). For Co, the first sub-valence p orbital was treated as valence, while for Ba it

was s orbital. For lanthanides there is a special GGA potential available in which f electrons

are kept frozen in the core. The number of f electrons in the core is equal to the total number

of valence electrons minus the formal valency. These potentials were used for Gd and Nd,

unlike for Ce where the standard potential with f states treated as valence states was used.

The automatic Monkhorst-pack was used for a k-point grid. The number of k-points varied

depending on the ratio of lattice parameters for a calculated structure. Thus 3 × 3 × 3 k-point

mesh was used for CeO2 and LnBaCo2O5 (Ln = Gd, Nd) calculations, 4 × 3 × 3 for the rest of

the structures. The unit cell size and atomic co-ordinates were relaxed until forces on atoms

were less than 0.01 eV/Å. The relaxation process was stopped when the energy difference

between two steps was smaller than 10−5

eV. The cut-off energy for plane wave was set to

450 eV. The computational setup was chosen after the preliminary study on CeO2 study using

different settings and the comparison of the obtained data with the experiments.

The initial atomic coordinates of LnBaCo2O5 and LnBaCo2O5.5 structures were used

from the reported data (Table 4.2 and Table 4.3). The reported unit cell was rotated in 45°,

and lattice parameters a and b were expanded by √2 (Figure 4.1). The unit cell was expanded

in order to include the structures with both ferromagnetic (FM) and anti-ferromagnetic

ordering (AFM) due to the Co mixed valence, although G-type antiferromagnetism for

LnBaCo2O5.5 (Ln = Tb, Gd) phases has been reported.217

Blocks (CeO2, GdBaCo2O5,

GdBaCo2O5.5) were relaxed separately and their relaxed atomic positions were used for the

starting structure of the double perovskites with the fluorite CeO2 layer. The unit cell of

LnBaCo2O5(5.5) with CeO2 layer was then doubled in c parameter (Figure 4.2c) in order to


110

accommodate the fluorite layer. The atomic positions of the additional cell were shifted in a

direction for the value of a/2 compared to the ones obtained by the previous DFT relaxations.

This structure (with the shifted two blocks of a double perovskite and a fluorite block) was

reported for the Y2SrCuFeO6.5139,140

and it served as a template for the DFT relaxations with

the LnBaCo2O5(5.5) + CeO2 structures.

Figure 4.2: Structures used for DFT calculations: a) GdBaCo2O5, b) CeO2, c) GdCeBaCo2O7, with

GdBaCo2O5 and fluorite CeO2 layer between double perovskite blocks. Atoms coloured as follows:

gadolinium (purple), cerium (yellow), barium (green), cobalt (blue), and oxygen (red).

The values of total energies of the relaxed structures were used to calculate the

formation energies of the structures with fluorite CeO2 layer. These values were obtained

from various equations (Equations 4.1-4.4, where Ln = Gd, Nd, and AE = Ba, Sr) depending

on the reaction enthalpies of final products from double perovskites, binary oxides or

AECeO3 (calculated because of the well known high stability of cerates). Since LnBaCo2O5

is not very stable (thus, the formation energies are higher) it is more precise to calculate the

formation energies in comparison to the more stable compound ‒ LnBaCo2O5.5. The

formation energies were calculated in electron Volts per Formula Unit, eV/FU. The more

negative values were obtained, the more stable were the structures on the right side of the


111

equations. As it was mentioned in Section 2.11.5 in order to calculate more accurate

formation energies for d electron-containing perovskites, U must be applied to the d orbitals

of the transition metals.229

For the calculations in this chapter the value of U = 4.3 eV was

used for the Co atom.230

Corrected value −8.5 eV of O2 gas was taken into account for the

formation energy calculations.230

LnAECo2O5 + CeO2 → LnCeAECo2O7 (4.1)

LnAECo2O5.5 + CeO2 → LnCeAECo2O7 +

O2 (4.2)

AEO +

Ln2O3 +

Co3O4 + CeO2 → LnCeAECo2O7 +

O2 (4.3)

AECeO3 +

Ln2O3 +

Co3O4 → LnCeAECo2O7 +

O2 (4.4)

For the Gd-containing materials with the fluorite layer built of two layers of CeO2

Equations 4.1-4.4 were adjusted (Equations 4.5-4.8).

GdBaCo2O5 + 2 CeO2 → GdCe2BaCo2O9 (4.5)

GdBaCo2O5.5 + 2 CeO2 → GdCe2BaCo2O9 +

O2 (4.6)

BaO +

Gd2O3 +

Co3O4 + 2CeO2 → GdCe2BaCo2O9 +

O2 (4.7)

BaCeO3 +

Gd2O3 +

Co3O4 → GdCe2BaCo2O9 + BaO +

O2 (4.8)

Similar to previous equations, various equations (Equations 4.9-4.11) were used to calculate

the formation energies of the materials with different fluorite layers (Ln2O3, Ln = Gd, Nd). A

Ruddlesden-Popper phase (RP2) was taken into account as another possible structure because

of the same nominal composition. The atomic coordinates of Gd2SrCo2O7 231

(Figure 4.3)

were used as a template for relaxation of Ln2BaCo2O7 (Ln = Gd, Nd) in the RP2 structural

form. The final formation energies were calculated for both structural models and compared

each other.

LnBaCo2O5 +

Ln2O3 → Ln2BaCo2O7 −

O2 (4.9)

LnBaCo2O5.5 +

Ln2O3 → Ln2BaCo2O7 (4.10)

BaO + Ln2O3 +

Co3O4 → Ln2BaCo2O7 −

O2 (4.11)


112

Figure 4.3: Reported structure for Gd2SrCo2O7 material231

used as a template for Ln2BaCo2O7

(Ln = Gd, Nd) DFT relaxation in RP2 structural model. Atoms coloured as follows: gadolinium

(purple), strontium (green), cobalt (blue), and oxygen (red).

Different atoms of lanthanides were used in relaxations of the LnBaCo2O5 structure (Ln = Ce,

Dy, Er, Eu, Gd, Ho, La, Nd, Pr, Tb, and Tm). PBE-GGA exchange correlation functional was

applied with the same configuration for the rest of the set-up, mentioned at the beginning of

this chapter. The formation energies values of the LnBaCo2O5 double perovskites were

calculated compared to binary oxide according to Equation 4.12. Since the most common

oxidation state of Ce in an oxide is (+4) (thus CeO2 is more stable than Ce2O3 whilst in non-

oxides Ce3+

is more common), Equation 4.13 was used to calculate the formation energy of

CeBaCo2O5.

Ln2O3 + BaO +

Co3O4 → LnBaCo2O5 +

O2 (4.12)

CeO2 + BaO +

Co3O4 → CeBaCo2O5 +

O2 (4.13)


113

4.3 Experimental methods

GdBaCo2O5+δ related materials with fluorite layer were prepared by solid state

synthesis. The synthesis conditions were based upon the reported ones for GdBaCo2O5+δ

materials62,138,220

were the sintering temperatures varied from 1000 to 1100 °C with the

different dwelling times (from 5 h to 20 h). The starting oxides for the synthesis with a listed

purity: Gd2O3 and Nd2O3 – 99.9%, BaCO3 – 99.99%, Co3O4 – 99.7%, and CeO2 – 99.99% all

purchased from Alfa Aesar, were measured out in the required stoichiometric amounts. The

inhomogeneous powder mixture of the starting oxides and carbonate was ground using a

pestle and a mortar. At the beginning of the synthetic work, LnCeBaCo2O7 (Ln = Gd, Nd)

materials were heated to 1000 °C for 12 h. Samples were then re-ground and pressed into

pellets and re-fired at 1000 °C for 12 h again. The samples were fired three times. The same

process was repeated for different synthesis temperatures (1050 and 1100 °C, see

Section 4.5). The same synthesis procedure was followed for the synthesis of Sr-analogues

(LnCeSrCo2O7) and materials with different fluorite layer (Ln2O3, Ln = Gd, Nd).

Phase compositions of all of the prepared samples were verified by PXRD collected at

room temperature using a Panalytical X'pert diffractometer using Co Kα1 radiation in Bragg-

Brentano geometry (Section 2.2.5). The PXRD data were measured over a 2θ range of 5-90 °

with step size 0.0334 ° and time per step 2 s. X'pert Highscore Plus software232

was used for

phase identification using the pdf-2 database.233

Pawley fits and Rietveld refinements

(Section 2.2.4) were performed by Topas academic program.152

4.4 Computational results

4.4.1 CeO2

At the first stage of the computational work each of the blocks (CeO2 and double

perovskites) were treated individually. Atomic coordinates of relaxed structures of the blocks

were then used for the main LnAECo2O7 phase. The unit cell of CeO2 used in DFT

relaxations is shown in Figure 4.2b. Different potentials (see Table 4.1 and Figure 4.4) were

applied. All the cells of CeO2 were relaxed using various energy cut-offs and different


114

k‒points. The values of the lattice parameter a of the relaxed CeO2 cells are shown in Figure

4.4 and compared with the reported value.234

There is a significant difference ‒0.1 Å in the

case of the DFT relaxations using LDA potential. Better agreement with the reported data

was obtained from the calculations using PBE and PW91 potentials. The difference against

the experimentally obtained value is for both of the potentials in the third decimal place. The

calculated values are within the 1 percent range from the reported lattice parameter of CeO2.

Increasing the energy cut-off value from 400 to 440 eV improves the accuracy of the relaxed

lattice parameter. Further increase of the cut-off energy does not enhance the accuracy.

Energy cut-off 450 eV was used as a 'medium' setting value in INCAR file for the later DFT

relaxations.

Table 4.1: Lattice parameters of relaxed CeO2 obtained by DFT calculations using different potentials

compared with experimental data.234

Lit.234

PBE LDA PW91

a lattice parameter (Å) 5.47441(1) 5.475 5.371 5.476

Figure 4.4: Comparison of the lattice parameters obtained by DFT relaxation methods with the

experimentally obtained data using different potentials and different cut-off energy values.


115

4.4.2 LnBaCo2O5

The values of lattice parameters obtained after the DFT relaxations of the double

perovskites LnBaCo2O5 (Ln = Gd, Ba) (Figure 4.2c) for both ferromagnetic and anti-

ferromagnetic structures are shown in Table 4.2. The values are compared with the reported

values for GdBaCo2O5 and NdBaCo2O5 respectively.235,236

Since the unit cell used in DFT

calculations was expanded by √2 (Section 4.2), the reported values of a and b lattice

parameters are multiplied by the same value of √2. The values of the lattice parameters of

GdBaCo2O5 obtained from DFT calculations are smaller than the reported (experimental)

ones.235

The difference for the ferromagnetic structure varies between 1-1.5% compared to

the reported structure. The values of the calculated lattice parameters of the anti-

ferromagnetic structure differ more with the 2.9% difference from the reported value235

for

the lattice parameter b. The lattice parameters obtained for NdBaCo2O5 structures are closer

to the reported values and all of them lie within sufficient 1% range. The lattice parameters a

and b are smaller than the reported ones,236

as in the case of GdBaCo2O5. The relaxed c

lattice parameter is on the contrary bigger than that obtained experimentally. The ratio

between the calculated and experimental value of a and b gives a number of 0.99, whilst for

lattice parameter c is the value of 1.01.

Table 4.2: Lattice parameters obtained after the DFT relaxation of LnBaCo2O5 materials, compared

with experimentally obtained data.235,236

Material a (Å) b (Å) c (Å)

GdBaCo2O5 − FM 5.504 5.504 7.474

GdBaCo2O5 − AFM 5.432 5.432 7.499

NdBaCo2O5 − FM 5.543 5.543 7.575

NdBaCo2O5 − AFM 5.557 5.557 7.607

GdBaCo2O5 − lit.235

3.955 (5.593) 3.934 (5.564) 7.540

NdBaCo2O5 − lit.236

3.961 (5.602) 3.930 (5.558) 7.533


116

4.4.3 LnBaCo2O5.5

The values of the relaxed lattice parameters of LnBaCo2O5.5 for both ferromagnetic

and anti-ferromagnetic structures are shown in Table 4.3. Similarly to LnBaCo2O5

relaxations, the calculated values are compared to reported ones.210,237

Lattice parameter b is

multiplied by √2 due to the expanded unit cell used for the DFT calculations (Section 4.2).

The calculated lattice parameters b, c for both Gd and Nd containing phases are smaller than

the reported values. The values of the lattice parameter a obtained after DFT relaxations are

on the contrary bigger than the reported ones. The lattice parameters of the relaxed unit cell

of GdBaCo2O5.5 are in ± 1% from the experimental values. The lattice parameters of

NdBaCo2O5.5 after the DFT relaxations show less accuracy with experimentally obtained data

(± 2%).

Table 4.3: Lattice parameters obtained after the DFT relaxations of LnBaCo2O5.5 materials, compared

with experimental data.210,237


GdBaCo2O5.5 − FM 7.954 7.645 7.543

GdBaCo2O5.5 − AFM 7.938 7.660 7.569

NdBaCo2O5.5 − FM 8.037 7.666 7.577

NdBaCo2O5.5 − AFM 7.955 7.746 7.577

GdBaCo2O5.5 − lit.210

7.867 3.862 (7.724) 7.571

NdBaCo2O5.5 − lit.237

7.802 3.901 (7.802) 7.615

4.4.4 Double perovskite with fluorite layer

Previous sections showed the lattice parameters of the relaxed structures of the double

perovskites and the fluorite structures. The total energies of the individual blocks obtained

before are necessary for the calculation of the formation energies (Section 4.4.5). The model

for the further calculations with a structure combined of the perovskite double layer and the

fluorite CeO2 layer is mentioned and showed previously (Section 4.2, Figure 4.2c). The

lattice parameters obtained after the DFT relaxations are summarized in Table 4.4. All of the


117

calculations were also done on Sr-analogues, where the Ba atom on A-site was completely

substituted by Sr (Table 4.5).

Table 4.4: Lattice parameters of calculated unit cells of LnCeBaCo2O7 materials obtained after the

DFT relaxations.


1-GdCeBaCo2O7 ‒ Gd-FM 5.523 5.523 20.688

2-GdCeBaCo2O7 ‒ Gd-AFM 5.448 5.448 20.478

3-NdCeBaCo2O7 ‒ Nd-FM 5.560 5.547 20.943

4-NdCeBaCo2O7 ‒ Nd-AFM 5.583 5.584 20.679

Table 4.5: Lattice parameters of calculated unit cells of LnCeSrCo2O7 materials obtained after the

DFT relaxations.


1-GdCeSrCo2O7 ‒ Gd-FM 5.468 5.458 20.649

2-GdCeSrCo2O7 ‒ Gd-AFM 5.492 5.494 20.475

3-NdCeSrCo2O7 ‒ Nd-FM 5.513 5.515 20.787

4-NdCeSrCo2O7 ‒ Nd-AFM 5.528 5.537 20.534

The electronic structure of GdCeBaCo2O7 material was studied by the density of

states (DOS) from GGA+U potential based on the optimized lattice structure. Figure 4.5

shows the spin-up and spin-down DOS of GdCeBaCo2O7 for both FM and AFM magnetic

ordering. A previous study238

on the electronic and magnetic structure of GdBaCo2O5.5 was

focused on several possible spin states and their relation to magnetic, electronic, and

structural properties. Our study was targeted on the role of CeO2 between the perovskite

blocks and its influence on the electronic structure of the material. Looking at the calculated

DOS (Figure 4.5a) of the FM ordered structure there is no band gap calculated. The expected

band gap region is overlapped by the Co d-orbitals. In contrast, anti-ferromagnetic

GdCeBaCo2O7 structure (Figure 4.5b) shows a band gap of 0.7 eV, with the valence band of

Co d orbitals. In general, LDA and GGA functional underestimate band gaps for

semiconductor and sometimes incorrectly predict a metal.239,240

This needs to be also

considered in the study of the density of state of GdCeBaCo2O7. It is worth mentioning very


118

low lying Ce 4f states (Figure 4.5). This could be corrected by including a +U parameter for

Ce.

Figure 4.5: Density of states of GdCeBaCo2O7 material obtained from the relaxed structure with a)

ferromagnetic; b) anti-ferromagnetic ordering.

4.4.5 Formation energies

Equations for the calculation of the formation energies (Equations 4.1-4.4) were

selected to consider the various ways of forming a double perovskite with CeO2 fluorite

layer. Equations 4.1 and 4.2 give comparison with the starting building blocks: CeO2 and

double perovskite. Comparing the stability of LnAECo2O5 with LnAECo2O5.5 phase, the

phase with oxygen amount O5 is less stable than the one with O5.5 (for Ln = Gd, Nd). Thus

the values of the formation energies obtained from Equations 4.1 are expected to be more

negative that means LnCeAECo2O7 is more stable with respect to the O5 phase than the O5.5

phase. Equation 4.3 represents the calculation of the formation energy of LnCeAECo2O7

from binary oxides since they are starting materials for the solid state synthesis (Section 4.3).

Many reported studies with Ba and Ce containing materials showed chemical and thermal

stability of BaCeO3 phase and various cerates as well,241,242

although BaCeO3 exhibits an

instability in water containing atmospheres.243,244

Therefore, Equation 4.4 takes the stability

of cerates (AECeO3) into account.


119

The values of the formation energies of LnCeBaCo2O7 materials are shown in Figure

4.6. The lowest formation energy values were calculated according to Equations 4.1 and 4.3

with the values of −2.4 and −2.8 eV/FU for anti-ferromagnetic Gd and Nd-containing

material (Equation 4.3) and −2.2 eV/FU for both ferromagnetic and anti-ferromagnetic

GdCeBaCo2O7 material (Equation 4.1). Formation energies compared with the LnBaCo2O5.5

(Equation 4.2) are more than 1 eV/FU higher than the ones compared with the LnBaCo2O5

(Equation 4.1) and are between −1 to −0.6 eV/FU for all of the calculated materials. The

highest values of formation energies were calculated using Equation 4.4 where the stability of

cerates was considered. All of the values are above 0 eV/FU and varies from +0.6 to +0.8

eV/FU. That predicts that the LnCeAECo2O7 compounds are unstable versus cerates

(AECeO3, AE = Ba, Sr), which is confirmed in Section 4.5.

Figure 4.7 displays the value of the formation energies for LnCeSrCo2O7 materials.

The formation energies calculated from Equation 4.1 were found to have the most negative

values with the most negative value −2.3 eV/FU for ferromagnetic NdCeSrCo2O7.

Comparison with the binary oxides (Equation 4.3) gives the formation energies 1 eV/FU

higher than the data obtained from Equation 4.1. Similar to the Ba-containing materials,

reaction enthalpies with LnSrCo2O5.5 are higher (approximately 1.5 eV/FU) than the ones

with the LnSrCo2O5 with the values between −0.5 and −0.6 eV/FU. The only positive values

are those obtained from the comparison with cerates (Equation 4.4) with the highest

formation energy +0.5 eV/FU for the ferromagnetic GdCeSrCo2O7.


120

Figure 4.6: Formation energies of (Gd,Nd)CeBaCo2O7 materials, (1-4) see Table 4.4 for phase setting

calculated from the total energies obtained after the DFT relaxations by the Equations 4.1-4.4 in

Section 4.2.


121

Figure 4.7: Formation energies of (Gd,Nd)CeSrCo2O7 materials, (1-4) see Table 4.5 for phase setting

calculated from the total energies obtained after the DFT the relaxations by the Equations 4.1-4.4 in

Section 4.2.

Previous formation energies included materials with one layer of the CeO2. A part of

the study of the GdBaCo2O5+δ related materials was also focused on the materials with more

layers of the fluorite CeO2 block. Table 4.6 shows the formation energies of the GdBaCo2O5

materials (with both ferromagnetic and anti-ferromagnetic ordering) with one and two layers

of the CeO2. Equations 4.5-4.8 (Section 4.2) were used to calculate the formation energies for

the materials with two layered fluorite block. The formation energies of the materials with

two layers of CeO2 are lower compared to one layered CeO2 structures for both magnetic

orderings and for most of the equations (Table 4.6). The improvement of the formation

energies varies from 1.3-1.5 eV/FU for both of the structures using Equations 4.5-4.7 (except

the GdBaCo2O5 material with AFM ordering for Equation 4.7 where the improvement was

noticed to be 0.4 eV/FU). The formation energies calculated from Equation 4.8 were the only

ones with an increase of the formation energies, +1.4 and +1.3 eV/FU for FM and AFM

ordering respectively. This is due to the stability of the BaCeO3, relevant also for the

materials with one layer of CeO2.

Summarizing all of the formation energies, all cases show that the cerate phase is

more stable than the phase with the CeO2 block within the structure. Even though the


122

structures with the fluorite layer seem to be more stable than the binary oxides or the double

perovskite + fluorite, we might not be able to prepare these materials, at least not using a

synthetic method relying on thermodynamic stability, which would expect to form the

cerates.

Table 4.6: Comparison of the formation energies (in eV/FU) of the GdBaCo2O5 double perovskites

with one and two layers of CeO2. The energy values were calculated using corresponding equations

from the data obtained after the DFT relaxations of each of the materials.

Material Equation + 1 CeO2 layer + 2 CeO2 layers

GdBaCo2O5 FM 4.1/4.5 −2.204 −3.719

GdBaCo2O5 FM 4.2/4.6 −0.998 −2.513

GdBaCo2O5 FM 4.3/4.7 −1.248 −2.762

GdBaCo2O5 FM 4.4/4.8 +0.828 +1.390

GdBaCo2O5 AFM 4.1/4.5 −2.180 −3.498

GdBaCo2O5 AFM 4.2/4.6 −0.883 −2.282

GdBaCo2O5 AFM 4.3/4.7 −2.410 −2.844

GdBaCo2O5 AFM 4.4/4.8 +0.748 +1.308

4.5 Experimental results

Figure 4.8 shows the comparison of the PXRD of the GdCeBaCo2O7 materials

sintered to various temperatures (1000, 1050, and 1100 °C) with the main peaks of the

presented phases. In order to obtain the weight percentage (w%) of the phases, Rietveld

refinements were done (Table 4.7). For the material sintered to 1000 °C almost all of the

reflection of the PXRD pattern are indexed to the GdBaCo2O5.5 and the CeO2 phase with a

small amount (1-2 w%) of BaCoO3). For the material synthesized at higher temperatures, an

additional phase BaCeO3 phase (4-5 w%) is presented. There are no other peaks presented,

only those which could be assigned to a double perovskite + fluorite structure. The presence

of the double perovskite and the CeO2 phase indicates that the CeO2 was not introduced

between the blocks of the GdBaCo2O5.5. There are no changes in the phase compositions

comparing various sintering temperatures. The values of the main phases vary in the range

± 3 %.


123

Figure 4.8: Comparison of the PXRD data collected on the GdCeBaCo2O7 materials synthesized at

various temperatures for 12 h with the main GdBaCo2O5.5 double perovskite phase. † indicates the

main reflections of the CeO2, ‡ represent the BaCeO3 reflection, and * denotes the main reflection of

the BaCoO3 phase.

Table 4.7: Phase fraction (w%) of the GdCeBaCo2O7 materials sintered three times at various

temperatures for 12 h. Phase percentages were obtained from the Rietveld refinements in Topas

against the PXRD data.

Sint. temp GdBaCo2O5.5 CeO2 BaCeO3 BaCoO3

1000 °C 71 27 0 2

1050 °C 68 28 5 < 1

1100 °C 71 24 4 1

Similar to the GdCeBaCo2O7 materials, their Nd-analogues were prepared at various

temperatures (1000, 1050 and 1100 °C) and characterized by the PXRD (Figure 4.9). The

percentages of the presented phases are shown in Table 4.8. NdBaCo2O5.7 and CeO2 are the

main phases for all of the synthesis temperatures with 5-8% of the BaCeO3 phase. There is no

extra peak which could be indexed to a perovskite + fluorite structure. NdCoO3 is the other

phase evident, especially abundant after the synthesis at 1050 °C. The percentage of the CeO2


124

phase is constant for all of the sintering temperature, the amount of NdBaCo2O5.7 varies

depending on the presence of the NdCoO3 phase. The presence of the BaCeO3 is in

agreement with the calculated formation energies for these materials (Figure 4.7). The

formation energies calculated compared to the formation of cerates (Equation 4.4 in

Section 4.2) are positive ‒ predicting the formation of cerates which is also observed on the

PXRD data collected on the LnCeBaCo2O7 materials, Ln = Gd, Nd.

Figure 4.9: Comparison of the PXRD data collected on the NdCeBaCo2O7 materials synthesized at

various temperatures for 12 h with the main the NdBaCo2O5.7 double perovskite phase. † indicates the

main reflections of the CeO2, ‡ represent the BaCeO3 reflection, and * denotes the main reflections of

the NdCoO3 phase.


125

Table 4.8: Phase fraction (w%) of the NdCeBaCo2O7 materials sintered three times at various

temperatures for 12 h. Phase percentages were obtained from the Rietveld refinements in Topas

against the PXRD data.

Sint. temp NdBaCo2O5.7 CeO2 BaCeO3 NdCoO3

1000 °C 61 25 8 6

1050 °C 58 22 6 14

1100 °C 71 24 5 0

DFT prediction of the LnCeSrCo2O7 materials showed more stability of these

compounds compared to cerates formation (Equation 4.4 in Section 4.2) than in the case of

the LnCeBaCo2O7 materials. PXRD data of the NdCeSrCo2O7 material with the presented

phases is shown in Figure 4.10. The percentages of the phases obtained after the Rietveld

refinements are displayed in Table 4.9. There are two main Co-containing phases presented:

perovskite NdCoO3 and Ruddlesden-Popper NdSrCoO4 phase with an additional cobalt

oxide. The amount of all of the Co-containing phases stays constant for all of the synthesis

temperatures. As in the previous synthesis with the Ba-analogues, CeO2 is not incorporated

into the perovskite structure and it is observed as a CeO2 phase with similar w% (23-27%) for

all of the sintering temperatures. According to the DFT calculations a formation of cerates

was expected to be a problem for the synthesis. Although the presence of the SrCeO3 was not

observed on the PXRD data after the synthesis, double perovskite material with the CeO2

fluorite layer was not synthesized due to the presence of the other Co-containing perovskite

or perovskite based phases.


126

Figure 4.10: Comparison of the PXRD data collected on the NdCeSrCo2O7 materials synthesized at

various temperatures for 12 h with the main NdCoO3 phase. † displays the main reflection of the

CeO2, ‡ indicates the main reflections of the NdSrCoO4 phase, and * denotes the Co3O4 main

reflections.

Table 4.9: Phase fraction of the NdCeSrCo2O7 materials sintered three times at various temperatures

for 12 h. Phase percentages were obtained from the Rietveld refinements in Topas against the PXRD

data.

Sint. temp NdCoO3 NdSrCoO4 CeO2 Co3O4

1000 °C 67 2 27 4

1050 °C 52 18 24 6

1100 °C 59 13 23 5

4.6 Formation energies - different lanthanides

Previously mentioned work (see Sections 4.4.5 and 4.5) was focused on the double

perovskite materials with the CeO2 fluorite layer both from the theoretical and experimental


127

point of view. Examples of other double perovskites with some of the different lanthanides

(Pr, Sm,220

La222

) were mentioned in the introduction of this chapter. The presented work

involved materials with Gd and Nd. For further study, other lanthanides were used for the

DFT relaxations of the LnBaCo2O5 structures to look at the other potential double perovskite

as a building block for materials with the fluorite layer in between. 'Medium' setting in VASP

described in Section 4.2 were used for the DFT calculations with the double perovskites.

The values of the reaction enthalpies (means also formation energies per formula unit,

in eV) of Equations 4.12 and 4.13 are displayed in Figure 4.11. As in previous plots with

formation energies, the lower value of the energy the more stable phase on the right side of

the equation. Comparing the formation energies of the selected lanthanide double

perovskites, the most stable ones are Nd and Eu materials followed by Gd, La, Pr, and Dy.

The presented values show only comparison between the stability and the presence of the

various lanthanides in LnBaCo2O5 structure. Variety of the oxygen content is typical for

double perovskite systems thus a material can be more stable in LnBaCo2O5+ δ with δ close

to 1.

Since the lowest energy LnBaCo2O5 structure is the Nd-containing material ‒ already

used for the study, another fluorite block is going to be used for the further work.

Figure 4.11: Formation energies (calculated according the Equations 4.12, 4.13 in Section 4.2) of

LnBaCo2O5 structures with the different lanthanides.


128

4.7 Other fluorite layer

4.7.1 Formation energies

Both theoretical and experimental study of the LnCeAECo2O7 materials (Ln = Gd,

Nd, and AE = Ba, Sr) showed the problem with the synthesis of these materials involving the

stability of cerates. In order to avoid this different fluorite blocks (Gd2O3 and Nd2O3

respectively) were used. The theoretical part of the study followed a similar procedure to that

applied to the double perovskites with the CeO2 fluorite layer.

Formation energies (Equations 4.9-4.11) were selected to compare the stability of the

Ln2BaCo2O7 materials versus LnBaCo2O5, LnBaCo2O5.5 and binary oxides. Table 4.10 shows

the values of the lattice parameters after the DFT relaxations. The values are compared with

the reported values of RP2 Gd2SrCo2O7 phase231

since the presence of a RP2 phase was

observed on the PXRD data collected after the synthesis of the Nd2SrCo2O7 sample

(Section 4.7.2). The values of lattice parameters of calculated structures are higher than the

reported which is due to the higher ionic radius of Ba2+

(1.61 Å) compared to Sr2+

(1.44 Å).245

The values of formation energies are shown in Figure 4.12. All of the values are

negative, that means materials on the right side of the Equations 4.9-4.11 are more

thermodynamically stable than the materials on the left side. The lowest values are calculated

for the formation energies compared to the LnBaCo2O5 double perovskites. As it was

mentioned before, LnBaCo2O5 phases are less stable than the LnBaCo2O5.5 and thus the

formation energies with the LnBaCo2O5 included are expected to be lower. Formation energy

values compared to binary oxides were approximately 0.5 eV/FU higher than these obtained

after the comparison to LnBaCo2O5 phases. The higher values of formation energies

(between −2.4 and −3 eV/FU) were obtained from the comparison of Ln2BaCo2O7 versus

LnBaCo2O5.5 (Equation 4.10).

The stability of the Ln2BaCo2O7 materials with several possible phases was

compared. Both ferromagnetic and anti-ferromagnetic ordering together with RP2 were taken

into account as a potential structural form. The formation energy values favour the formation

of the Ln2BaCo2O7 phases. The next chapter is focused on the experimental results obtained

after the solid state synthesis of the Ln2BaCo2O7 materials.


129

Table 4.10: Lattice parameters obtained after the DFT relaxations of the Ln2BaCo2O7 materials.


1- Gd2BaCo2O7 − FM 5.473 5.473 20.512

2 - Gd2BaCo2O7 − AFM 5.483 5.483 20.631

3 - Gd2BaCo2O7 − RP2 5.435 5.435 19.612

4- Nd2BaCo2O7 − FM 5.651 5.675 21.035

5 - Nd2BaCo2O7 − AFM 5.654 5.649 21.037

6 - Nd2BaCo2O7 − RP2 5.438 5.438 20.095

Gd2SrCo2O7 – lit.231

5.37506(2) 5.37506(2) 19.35807(6)

Figure 4.12: Formation energies of the Ln2BaCo2O7 materials, Ln = Gd, Nd, (1-6) see Table 4.8 for

the phase composition, calculated from the total energies obtained after the DFT relaxations from

Equations 4.9-4.11 in Section 4.2.


130

4.7.2 Experimental results

PXRD data of the Gd2BaCo2O7 material collected after the synthesis at various

temperatures is shown in Figure 4.13. There is no difference in the collected patterns. Table

4.11 shows the percentage of presented GdBaCo2O5.5 and Gd2O3 phases obtained after the

Rietveld refinements (Figure 4.14). There are no additional reflections in the PXRD data.

The phase fraction 70 : 30 of double perovskite to Gd2O3 phase remains constant for all of the

temperatures. This fact is in contradiction to the prediction of the stability based on DFT

calculations mentioned in Section 4.7.1. The values of the formation energies compared with

the GdBaCo2O5.5 were calculated to be close to −2.6 eV/FU and thus a formation of a

Gd2BaCo2O7 phase was expected.

Figure 4.13: PXRD data collected on the Gd2BaCo2O7 material sintered three times at various

temperatures for 12 h with the main double perovskite GdBaCo2O5.5 phase. The main reflections of

the Gd2O3 are labelled with †.


131

Table 4.11: Phase fraction (w%) of the Gd2BaCo2O7 material sintered three times at various

temperatures for 12 h. Phase percentages were obtained from the Rietveld refinements of the PXRD

data in Topas.

Sint. temp GdBaCo2O5.5 Gd2O3

1000 °C 70 30

1050 °C 70 30

1100 °C 71 29

Figure 4.14: Rietveld refinement of the Gd2BaCo2O7 material sintered three times at 1000 °C. First

line of reflections corresponds to the GdBaCo2O5.5 phase and the second one to the Gd2O3,

GOF = 1.445.

Nd2BaCo2O7 materials were prepared at the same time and using the same synthesis

conditions as for the Gd2BaCo2O7 synthesis. The materials were characterized by Rietveld

refinement (Figure 4.16) against the PXRD data (Figure 4.15) collected after the solid state

synthesis with different sintering temperature. The percentages of the presented phases are

displayed in Table 4.12. The Ruddlesden-Popper phase (RP2) Nd2BaCo2O7 is the main phase

for all of the sintering temperature with the abundance of 85-92 w%. No crystallographic data

for the Nd2BaCo2O7 RP2 phase have been reported. All of the reflections of the main phase

were indexed using the symmetry and the data of the reported Gd2SrCo2O7 phase.231

This

phase was also used as a starting model for the Rietveld refinements. The rest of the


132

reflections are indexed to NdBaCo2O5.7 double perovskite and the Nd2O3 phase. According to

the calculated values of formation energy (Figure 4.12) a formation of the Nd2BaCo2O7

phase is expected. The values of formation energy of ferromagnetic, anti-ferromagnetic

ordered and RP2 structures are within a close range.

Figure 4.15: PXRD data collected on the Nd2BaCo2O7 material sintered three times at various

temperatures for 12 h with the Ruddlesden-Popper (RP2) main phase. The main reflection of the

double perovskite NdBaCo2O5.7 are marked with †, the Nd2O3 reflections are labelled with ‡.

Table 4.12: Phase fraction (w%) of the Nd2BaCo2O7 material sintered three times at various

temperatures for 12 h. Phase percentages were obtained from the Rietveld refinements of the PXRD

data in Topas.

Sint. temp NdBaCo2O5.7 Nd2O3 RP2 phase

1000 °C 10 2 88

1050 °C 12 3 85

1100 °C 5 3 92


133

Figure 4.16: Rietveld refinement of the Nd2BaCo2O7 material sintered 3 times at 1000 °C with the

presented phases as follow: RP2-phase ‒ 1st line of the reflections, NdBaCo2O5.7 ‒ 2

nd line, Nd2O3 ‒

3rd

line, GOF = 1.677.


In summary, it has been shown that DFT methods can be used to calculate the

formation energies of perovskite based materials from binary and ternary oxides. The method

was used as a prediction to guide for the synthetic isolation of double perovskite

GdBaCo2O5+δ related structures with a fluorite layer. Final structures were built from the

individual relaxed blocks. The data of the starting structures obtained from DFT calculations

were compared with the reported data.210,235-237

Both FM and AFM magnetic ordering was

used for the calculations. Several possible structures were taken into account. The synthesis

of the most stable materials (according calculated formation energies) was attempted. The

selection based on DFT calculation was predictive and time saving tool for phase screening

of selected structural types.

Calculated formation energies were based on various reaction including binary oxides

and double perovskites. For the prediction of the stability of materials is important to take


134

into account all the stable oxides/materials. An unstable material (e.g. GdBaCo2O5 in our

prediction) can result in very negative values of formation energies. Since the work was

followed by the experimental work, few changes for the formation energy calculations were

done after the first steps of the solid state synthesis. All of the collected PXRD patterns

confirmed the presence of the stable cerates. Additional calculated formation energies

showed higher positive values, which was in an agreement with the experimentally obtained

data.

PXRD data obtained after the solid state synthesis of LnCeAECo2O7 (Ln = Gd, Nd;

AE = Ba, Sr) showed the presence of the double perovskite and the CeO2, with an additional

cerate phase. Hence, no CeO2 was introduced between the double perovskite layers. A charge

imbalance between the CeO2 fluorite layer and the double perovskite layer represents maybe

another problem for the synthesis of these materials.

As a minor part of the DFT study, electronic structure of the structure relaxed

materials with the fluorite layer was studied to see the influence of the CeO2 layer for the

conduction or the valence band. Since other factors (magnetic ordering, spin states) play an

important role in the electronic structure, the effect of the additional CeO2 layer cannot be

explained individually. Structures with the two fluorite layers were built and their formation

energies were calculated. The values follow the same trend as it was observed in the materials

with the one layer of CeO2.

Since the materials with the CeO2 fluorite layer could not be synthesized, another

fluorite Ln2O3 (Ln = Gd, Nd) layer was used for both DFT prediction and solid state

synthesis to avoid the formation of cerates. Calculated formation energies for Ln2BaCo2O7

show that the existence of the structures is favourable. The comparison between included

Ln2BaCo2O7 structural types (RP2 phase and double perovskite + fluorite block) showed

similar values of formation energies. There was no Gd2BaCo2O7 phase synthesized

(GdBaCo2O5.5 and Gd2O3 phases were only presented), while collected PXRD data for Nd-

containing material indicates the presence of a RP2 phase. Gd2BaCo2O7 appearance did not

follow the energetic prediction based on DFT calculations, whilst Nd2BaCo2O7 agreed with

the favourability of RP2 phase compared to double perovskite structure.

In conclusion, DFT calculation methods were used for material synthesis prediction

based on energetics comparison of selected possible phases in order to prepare GdBaCo2O5+δ

related structures with a fluorite layer. A charge misbalance between the CeO2 fluorite layer

and the double perovskite layer was overcome in materials with fluorite layers based on


135

Ln2O3 (Ln = Gd, Nd). The data of experimentally prepared phases indicate the presence of a

RP2 phase. The formation of a RP2 phase does not meet our main goal – to synthesize a

double perovskite structure with a fluorite inter-layer, but the Nd2BaCo2O7 phase could be a

candidate for further crystallographic and material properties study.

5 Ruddlesden-Popper phases ‒ stannates

5.1 Introduction

Ruddlesden-Popper phases (general formula An+1BnO3n+1, structure represented in

Figure 5.1) represent promising structures for a SOFC cathode application due to their ability

to host oxide vacancies in perovskite layer (such as observed in Sm2BaCo2O7−δ)246

and oxide

interstitial in the rock salt layer (with an example of La1.5+xSr0.5−xCo0.5Ni0.5O4+δ).144

Recent

studies involved mainly the A2NiO4+δ (A = La, Sr) phase56,141-143

and are described in details

in Section 1.7.1. This work is focused on the n = 1 (RP1) stannate (Sr2SnO4) phases doped on

A- or B-site, where the A-site Sr, Ba atoms are substituted by La and in the case of B-site

doping, Sn atom is doped by M5+

(Nb, Ta) and M6+

(Mo, W) elements, in order to introduce

mobile interstitial oxide ions into a non-metallic system. The target of this work is to find a a

pure ionic conductor with a potential use as an electrolyte for SOFCs, thus any open-shell

transitional metals are not suitable for the consideration. Those transitional metals would

increase electronic conductivity, which is detrimental to any SOFC electrolyte application.

Figure 5.1: Schematic of Ruddlesden Popper Srn+1TinO3n+1 phases with different number of

perovskite layer within the structure: a) n = 1; b) n = 2; c) n = 3. Sr atoms are in green, Ti in blue and

O atoms are in red.

Chapter 5. Ruddlesden-Popper phases − stannates

137

The structure of Sr2SnO4 (Figure 5.2) was described as a K2NiF4-type.247,248

Later,

high resolution neutron powder diffraction studies249,250

clarified the relationship between the

SnO6 octahedral tilts and the structural description of the material. The Rietveld refinements

showed that the structure Sr2SnO4 at room temperature is well described using Pccn space

group, which is a subgroup of the both Bmab and P42/ncm. The structure can be derived from

the K2NiF4 structural type by tilting of the SnO6 octahedra along the a- and b-axis and with

the different values of tilting angles: α ≠ β ≠ 0.249

According to a variable temperature

neutron diffraction study there are two phase transition observed.250

At temperatures above

423 K, the structure changes to another orthorhombic structure, Bmab, retaining both tilts but

with equal angles α = β ≠ 0. At temperatures higher than 573 K there are no tilts and the

structure is described in the tetragonal I4/mmm space group.

Figure 5.2: Structure of Sr2SnO4 described in Pccn space group,249

with Sr (green) on A-site and Sn

(grey) in octahedral ordering on B-site, oxygen ions are labelled with red: a) view along c-axis; b)

view along a-axis.

The work presented in this chapter contains both experimentally and computationally

obtained data of A- and B-site doped Sr2SnO4 phases. The computational part deals with


138

formation energies of doped RP1 phases e.g. stannates, hafnates, and zirconates, to identify

favourable doping strategies. The experimental part is focused on Nb- and Ta-doped Sr2SnO4

materials. Prepared materials are structurally characterized by PXRD techniques. The

conductivity of single phase materials are measured by AC impedance spectroscopy and

discussed later on together with the electronic structure studies by spectroscopic methods

(UV-vis, IR, and solid state Sn-NMR).

5.2 Computational methods

All of the calculations were performed using the plane wave DFT package, Vienna

Ab-initio Simulation Package (VASP) version 4.6.26178,224

with Perdew, Burke and

Ernzerhof (PBE) exchange correlation functional.225

For the A-site atoms (strontium, barium)

and hafnium, the first sub-valent s orbital was treated as valence, as were d orbital for tin and

p orbital for the rest of the atoms. A k-point 3 × 3 × 3 grid was used for the DFT calculations.

The unit cell size and atomic co-ordinates were relaxed until forces on atoms were less than

0.01 eV/Å. The relaxation process was stopped when the energy difference between two

steps was smaller than 10−5

eV. The cut-off energy for plane wave was set to 450 eV.

The coordinates of initial un-doped phases were used from the reported data (for

stannates,250

for hafnates and zirconates251

). The lattice parameters a and b of the unit cells

used for the DFT relaxations were doubled. This was necessary for a lower doping level (1/8

and 1/16 respectively) on both A- and B-site. In the case of M6+

doping (Mo, W) one atom of

tin was replaced by M5+

element (doping level 1/16). Increased positive charge was

compensated by introducing one interstitial O atom (Oint) placed in the SrO rock salt layer.

The coordinates of Oint used where those reported for the mixed conductor

La1.5+xSr0.5−xCo0.5Ni0.5O4+δ.144

When doping of M5+

was carried out, two atoms of tin were

replaced by M5+

(doping level 1/8). Structures with various positions of the doping M5+

elements were taken into account. The relaxed structure with the lowest total energy was used

for the formation energy calculation. Similarly to M6+

doping, doping level on A-site was 1/8.

Two atoms of tin were replaced by La. Several possible configurations of La atoms were used

for DFT relaxations.

The values of the total energies of the doped Ruddlesden-Popper phases were used to

calculate the formation energies from the starting un-doped phases. Equations 5.1-5.3 were


139

used to calculate the formation energies of either A- and B-site doped materials. The

formation energies were calculated in electron Volts per Formula Unit, eV/FU. The more

negative the values obtained, the more stable were the structures on the right side of the

equations.

Sr2BO4 +

M2O5 −

SnO2 → Sr2B7/8M1/8O4+1/16 (5.1)

Sr2BO4 +

MO3 −

SnO2 → Sr2Sn15/16M1/16O4+1/16 (5.2)

Sr2BO4 +

La2O3 −

SrO → Sr2−1/8La1/8BO4+1/16 (5.3)

B = Sn, Hf, Zr

Previous equations show the formation energies of the structures with Oint within the doped

structure. Doped structures of the stannates without the Oint atom (with the tin mixed valence

oxidation state of Sn2+

/Sn4+

) were also relaxed. The formation energies of these phases were

calculated using Equation 5.4-5.6 and compared with the ones with Oint included.

Sr2SnO4 +

M2O5 −

SnO2 → Sr2Sn7/8M1/8O4 +

O2 (5.4)

Sr2SnO4 +

MO3 −

SnO2 → Sr2SnM1/16O4 +

O2 (5.5)

Sr2SnO4 +

La2O3 −

SrO → Sr2La1/8SnO4 +

O2 (5.6)

5.3 Experimental methods

Parent phase of Sr2SnO4 was synthesized by a solid state method.249,250

Sr2SnO4 was

prepared from powders of SrCO3 and SnO2. The mixture was ground and heated in an

alumina crucible in air at 1177 °C. The undoped phase was heated for one week with

repeated regrinding with a pestle and a mortar every day following the literature synthesis

protocol.249,250

Synthesis conditions of Sr2SnO4 were also followed for the doped materials.

Starting doping levels x = 0.05 and 0.10 were selected for both A- (La) and B-site (Nb, Ta,

Mo, W) doped Sr2SnO4. The starting oxides for the synthesis with a listed purity: SnO2,

SrCO3, La2O3 ‒ 99.99 %; MoO3 ‒ 99.95 %; Nb2O5, Ta2O5, ZrO2 ‒ 99.9 % all purchased from

Alfa Aesar and WO3 ‒ 99.995% purchased from Sigma-Aldrich, were measured out in the


140

required stoichiometric amounts. The inhomogeneous powder mixture of the starting oxides

and carbonate was ground using a pestle and a mortar. In the case of La-, Nb-, and Ta-doped

materials the initial synthesis produced RP1 phases with significant amounts of binary oxides

(La2O3, Nb2O5, and Ta2O5) as confirmed by the PXRD data. Synthesis conditions were

improved by experimenting with temperature, mixing technique (using a planetary ball mill

and reaction atmospheres as shown in Section 5.5.1). Final synthesis conditions for the single

phased B-site doped Sr2SnO4 were as follows: heating at 1250 °C in O2 atmosphere for Nb-

doping and heating in ambient air at 1300 °C for the Ta-doped materials. La-, Mo-, and W-

doped materials could not be prepared as single phases (even after modification of the

synthesis conditions) and thus the work was focused on the Nb- and Ta-doped phases.

The phase composition of all of the prepared samples was verified by powder PXRD

collected at room temperature using a Phillips X'pert Panalytical diffractometer using Co Kα1

radiation in Bragg-Brentano geometry. X'pert Highscore Plus software232

was used for phase

identification using the pdf-2 database.233

Pawley fits and Rietveld refinements were

performed by Topas academic program.152

5.4 Computational results

All of the studied Ruddlesden-Popper phases were relaxed using the DFT methods.

Total energy values obtained after the DFT relaxations were used to calculate the formation

energies using Equations 5.1-5.3. Figure 5.3 shows the values of the formation energies for

the A- and B-site doped stannates (Sr2SnO4), zirconates (Sr2ZrO4) and hafnates (Sr2HfO4).

The same trend of binding energy values is observed for all of the studied phases. The lowest

values of the formation energies are those for the La-doped materials with the values

−1.2 eV/FU. The more negative the value calculated, the more stable is the doped material

expected to be. All of the formation energies for M5+

doped materials (except that for the

Sr2Hf7/8Nb1/8O4+1/16) were found to have negative values. The highest values of the formation

energies (and thus expected to be least stable) were calculated for the M6+

doped phases, with

the positive values from +0.02 to +0.3 eV/FU. Comparing the values of all doped structures

for all three phases, the most stable are doped zirconates phases, followed by stannates and

hafnates.


141

Figure 5.3: Formation energies of the doped Sr2(Sn,Zr,Hf)O4 materials calculated from the total

energies obtained after the DFT relaxations by the Equations 5.1-5.3 in Section 5.2.

Most of the work presented in this chapter is focused on the stannate structures.

Figure 5.4 shows a comparison of the formation energies of the Sr2SnO4 phases compared to

the Ba2SnO4. Calculated formation energies of Ba2SnO4 derivatives (except for Nb-doping)

are higher than the ones of Sr2SnO4 phases for both A- and B-site doped materials. As it was

mentioned in Section 5.2, the Sr2SnO4 phases were also relaxed without the Oint, with the

formation energies shown in Figure 5.4. The formation energy values of the phases without

Oint are almost identical to the ones with Oint for M5+

and La-doping. That means that the

stability of the phases with and without Oint with respect to the parent undoped RP phase is

very similar. In the case of M6+

doping, the formation energies of the structures without Oint

are lower (by 0.05 for W and 0.1 eV/FU for Mo doping) than those obtained for the materials

with interstitial oxygen.


142

Figure 5.4: Formation energies of the doped Sr2SnO4 materials calculated from the total energies

obtained after DFT relaxations by the Equations 5.1-5.3 in Section 5.2 compared with the doped

materials without Oint (Equations 5.4-5.6, Section 5.2) and with the Ba2SnO4 materials.

5.5 Structural characterization

5.5.1 Laboratory P-XRD

A- and B-site doped Sr2SnO4 materials were prepared by solid state synthesis

described in Section 5.3. All of the synthesized stannates were refined using the orthorhombic

Pccn space group with the reported refined atomic positions.249

Figure 5.5 shows the

Rietveld refinement of the Sr2Sn0.95Nb0.05O4 material. Changes of the intensity of the

Sr4Nb2O9 impurity main peak were found to vary depending on the synthesis conditions.

Since the materials with starting doping levels x = 0.05 and 0.1 were synthesized with an

impurity phase, materials with lower doping levels were also prepared. For Nb-doped

Sr2SnO4 materials, the highest doping level of phase pure sample was obtained for the

x = 0.03 in Sr2Sn0.97Nb0.03O4. Maximum Ta doping for single phased stannates was achieved

for x = 0.04 in Sr2Sn0.96Ta0.04O4. PXRD data of the Ta-doped materials with x > 0.4 showed a


143

presence of Sr-Ta-O oxide (found in ICSD database with structural formula of

Sr1.907Ta0.593O2.73) secondary phase.

La-doped materials were not prepared as single phases, as a presence of La2O3 impurity was

confirmed even on PXRD data for low La doping (Sr1.97La0.03SnO4 material, Figure 5.6).

Neither Mo- or W-doped Sr2SnO4 materials were prepared as single phases due to the

presence of the Sr3(Mo,W)O6 impurity phases; these compounds were not considered in the

theoretical DFT study.

Figure 5.5: Rietveld refinement of the Sr2Sn0.95Nb0.05O4 material, with a presence of the Sr4Nb2O9

impurity phase.


144

Figure 5.6: Rietveld refinement of the PXRD data obtained after the synthesis of the Sr1.97La0.03SnO4

material at 1300 °C with the La2O3 impurity phase.

Lattice parameters of the Nb-doped Sr2SnO4 materials obtained after the Rietveld

refinement against the laboratory PXRD data are shown in Figure 5.7 and in Appendix B. The

data were collected using a powder of the synthesized material mixed in 1:1 ratio with KCl as

an internal standard. There is a step-like increase of all lattice parameters observed for the

samples with x ≥ 0.03. This increment, more significant for the c lattice parameter, could be

explained due to the presence of the Sr4Nb2O9 impurity phase (Figure 5.5) for the materials

with higher Nb doping level or due to a change in charge compensation mechanism after the

doping with Nb5+

(presence of Oint or Sn2+

/Sn4+

valence state). A significant decrease of a/c

ratio is observed, showing two regions for x < 0.03 and x > 0.03 with the different values

(Figure 5.7d). The same trend of the lattice parameter values is observed for the Ta-doped

Sr2SnO4 materials (Figure 5.8 and Appendix C). The composition of the Sr2Sn0.96Ta0.04O4

material was studied by the Energy dispersive X-ray (EDX) diffraction. The incorporation of

the Ta atom within the RP1 structure was confirmed by multiple measurements. The average

ratio between the A- and B-site atoms: Sr1.934(4)Sn1.029(4)Ta0.04(2) is in good agreement with the

expected stoichiometry (Table of EDX analysis of Sr2Sn0.96Ta0.04O4 is shown in Appendix A) .


145

Figure 5.7: Lattice parameter evolution of the Sr2Sn1−xNbxO4 materials; values were obtained after

Rietveld refinements of laboratory PXRD data with KCl as an internal standard: a) lattice parameters

a, b; b) lattice parameter c; c) cell volume; d) a/c ratio.


146

Figure 5.8: Lattice parameter evolution of the Sr2Sn1−xTaxO4 materials; values were obtained after

Rietveld refinements of laboratory PXRD data with KCl as an internal standard: a) lattice parameters

a, b; b) lattice parameter c; c) cell volume; d) a/c ratio.

5.5.2 Synchrotron data

The previous section showed the laboratory PXRD data. Sr2Sn0.97Nb0.03O4 and

Sr2Sn0.96Ta0.04O4 materials were studied using high resolution synchrotron X-ray techniques

and high resolution neutron diffraction (HRPD, Section 5.5.3). Synchrotron experiments were

carried out using the I11 instrument at the Diamond Light Source (UK). The data were

collected at the room temperature with incident wavelength λ = 0.827157 Å. Small amounts

of the Sr2Sn0.97Nb0.03O4 and Sr2Sn0.96Ta0.04O4 materials were loaded into 0.2 mm diameter

borrosilicate capillary and sealed with a gas/oxygen torch. Measurements were taken from

spinning capillaries to improve the powder average. Lattice parameters and atomic positions

obtained after the Rietveld refinement of the I11 data were used as starting data for the joint

refinements (Section 5.5.3).


147

Several models for the Rietveld refinement of the doped Sr2SnO4 materials were

applied. The parent Sr2SnO4 phase is described in two orthorhombic (Pccn, Bmab) and

tetragonal (P42/ncm) space groups due to the different tilting angles around the a- and b-

axes.250

At the room temperature, the phase is described in Pccn space group, which was used

for the starting model. The refinement was also performed in orthorhombic supergroup

(Bmab). A better refinement was achieved using Pccn space group (Figure 5.9) with all peaks

of Ruddlesden-Popper Sr2SnO4 phase indexed. However, the starting model could not fit the

data correctly since the peak anisotropy was observed.

Several peak shape functions for Sr2SnO4 were applied, the best results are given by

orthorhombic Stephens Anisotropic Peak Broadening. A closer look at the refinement shows

additional peaks (inset of Figure 5.9) and broader reflections of the main phase. The

broadening of the peaks can be explained by stacking faults (observed in several Ruddlesden-

Popper phases)252-255

and by a presence of other compound of Ruddlesden-Popper series. The

position of low angle peaks was crucial for impurity phases determination.

Figure 5.9: Rietveld refinement of the Sr2Sn0.96Ta0.04O4 material in Pccn space group, Rwp = 3.554%.

The inset shows magnified area of additional reflections found to be indexed for SrSnO3 perovskite

(P) and Sr3Sn2O7 Ruddlesden-Popper n = 2 phase (RP2).


148

The previous refinement (Figure 5.9) was improved after including of SrSnO3

perovskite and Sr3Sn2O7 RP2 phases (Figure 5.10). All of the extra reflections are indexed

using these two phases and since the main peaks of these phases overlay with the main

Sr2SnO4 reflections, anisotropic peak broadening (due to stacking faults) was fitted. SrSnO3

phase is known to exist in several space groups a Pbnm space group was used in Pawley fit.

The final Rietveld refinement of Sr2Sn0.96Ta0.04O4 with SrSnO3 and Sr3Sn2O7 as secondary

phases, is shown in Figure 5.10. The refined parameters of the main RP1 phase are displayed

in Table 5.1 whilst the refined lattice parameters are compared with the parent Sr2SnO4 phase

in Table 5.3. The occupancy of Ta was fixed to nominal value of 4%.

Figure 5.10: Final refinement of I11 data of Sr2Sn0.96Ta0.04O4 in Pccn space group with secondary

phases of SrSnO3 and Sr3Sn2O7; Rwp = 2.292%.


149

Table 5.1: Refined parameters of the Sr2Sn0.96Ta0.04O4 sample from the data collected from I11

beamline.

Atom Site x y z Beq (Å2) Occ

Sn1 4a 0 0 0 0.392(9) 0.96

Ta1 4a 0 0 0 0.392(9) 0.04

Sr1 8e −0.0050(4) 0.0014(7) 0.35184(2) 0.297(9) 1

O1 4c 0.250 0.250 0.016(1) 0.74(8) 1

O2 4d 0.750 0.250 0.006(1) 0.43(7) 1

O3 8e 0.027(2) −0.025(2) 0.1656(2) 0.71(8) 1

The I11 data refinement of Sr2Sn0.97Nb0.03O4 followed the same procedure as it was

described for Sr2Sn0.96Ta0.04O4, with the refinement of the main Ruddlesden-Popper phase.

Similar to previous refinements, the occupancy of Nb was fixed to nominal value of 3 %. The

best refinement of the main phase was obtained using Pccn space group. Similarly to

Sr2Sn0.96Ta0.04O4, a broadening of the main reflection was observed in Nb-doped material,

while the extra peaks were not observed. The refinement with SrSnO3 and Sr3Sn2O7 phases

(Figure 5.11) significantly decreased the Rwp values (1.812% compared to 2.473%) and

improved the fitting of the main RP peaks. Table 5.2 shows the refined parameters of

Sr2Sn0.97Nb0.03O4.


150

Figure 5.11: Rietveld refinement of the Sr2Sn0.97Nb0.03O4 sample in Pccn space group with secondary

phases of SrSnO3 and Sr3Sn2O7; Rwp = 1.812%.

Table 5.2: Refined parameters of the Sr2Sn0.97Nb0.03O4 sample from the data collected from I11

beamline.

Atom Site x y z Beq (Å2) Occ

Sn1 4a 0 0 0 0.77(1) 0.97

Nb1 4a 0 0 0 0.77(1) 0.03

Sr1 8e −0.0021(7) 0.0024(1) 0.35213(2) 0.148(9) 1

O1 4c 0.250 0.250 0.0187(4) 0.73(8) 1

O2 4d 0.750 0.250 −0.003(1) 0.76(7) 1

O3 8e 0.0011(2) −0.0359(8) 0.1656(2) 0.87(6) 1


151

Table 5.3: Refined lattice parameters of the Sr2Sn0.97Nb0.03O4 and Sr2Sn0.96Ta0.04O4 material from I11

data after the Rietveld refinement in Topas and their comparison with the undoped Sr2SnO4 phase.

Material a (Å) b (Å) c (Å) V (Å)

Sr2SnO4 250

5.72898(5) 5.73524(5) 12.58110(6) 413.378

Sr2Sn0.97Nb0.03O4 5.73504(6) 5.74146(6) 12.65833(7) 416.807(7)

Sr2Sn0.96Ta0.04O4 5.73495(13) 5.74169(13) 12.66376(24) 416.99(1)

5.5.3 High Resolution Powder Diffraction data

The solid state synthesis of Sr2Sn0.97Nb0.03O4 and Sr2Sn0.96Ta0.04O4 materials was

scaled up in order to synthesise larger amount of the materials for the HRPD analysis.

Approximately 5 g of the Sr2Sn0.97Nb0.03O4 and Sr2Sn0.96Ta0.04O4 materials were loaded into a

vanadium 8 mm diameter can and measured at ambient temperature by time-of flight neutron

diffraction for the time 3 h. Data were collected using all three detector banks (35 °, 90 °, and

145 °). Magnetism in these compounds is irrelevant, thus the data of first two banks only

(35 °, 90 °) were used in a joint Rietveld refinement with the I11 data. Twelve background

parameters, scale factors and sample-dependent peak shapes were refined per bank.

A representative refinement of the Sr2Sn0.97Nb0.03O4 and Sr2Sn0.96Ta0.04O4 sample is

displayed in Figure 5.11 and Figure 5.13 respectively. Starting model obtained from I11 data

analysis (orthorhombic Pccn space group of Sr2SnO4 phase) was applied for both of the

refinements. All of the reflections of RP1 phase were observed and fitted, although some

additional reflections and broadening of several peaks of the main phase were noticed, similar

to I11 data analysis. Refinements with refined coordinates of RP1 phase only, obtained from

I11 fits, gave imperfect with Rwp = 2.871% for Sr2Sn0.97Nb0.03O4 and Rwp = 4.163% for

Sr2Sn0.96Ta0.04O4. Initially, models with one phase only were applied to improve the joint

refinements. Inspection of Fourier maps showed regions of negative nuclear density where

interstitial oxygen could be placed. Several positions of Oint were tried all together with the

Oint position evidenced in La1.5+xSr0.5−x Co0.5Ni0.5O4+δ,144

but no one improved the refinements

significantly. It also did not solve the problem with additional and broad reflections. Many

RP phases are known to contain stacking faults in their crystal structures.252-255

A model

including a stacking fault (with doubled oxygen sites placed in positions obtained from


152

Fourier maps) improved the joint refinement by modelling the peak anisotropy, though did

not index the additional peaks.

After trying several models of one RP phase, the model with secondary phases of

SrSnO3 and Sr3Sn2O7 was applied for the joint refinement of I11 and HRPD data. The

refinements were improved (with Rwp values of 2.083% for Sr2Sn0.97Nb0.03O4 and 2.724% for

Sr2Sn0.96Ta0.04O4 respectively) and the additional peaks were all indexed. Including of Oint

(placed in rock salt layer with coordinates: 0, 0.5, 0.25) did not affect the refinements and

showed similar fits to previous ones, with Rwp value of 2.077% for Nb-doped and 2.714% for

Ta-doped sample. Refined thermal parameters of Oint were negative, which could be related

to the small concentration of Oint within the structure, the wrong position of Oint or the

absence of Oint. Several other Oint positions, all obtained from Fourier maps, were used; all of

them provided negative values of thermal parameter for Oint. To maintain the charge balance

of the phase, occupancy of Nb or Ta were included in the refinements after addition of Oint.

Due to the lack of contrast between Sn and Nb or Ta in their neutron scattering lengths (6.225

fm for Sn, 7.054 fm for Nb and 6.91 fm for Ta), the occupancy of Nb and Ta was fixed to

nominal values of 3% and 4% respectively. The final joint refinement of Sr2Sn0.97Nb0.03O4

and Sr2Sn0.96Ta0.04O4 are shown in Figure 5.12 and Figure 5.13; refined structures with the

refined parameters of both compounds are available in Appendices D and E. Refined lattice

parameters (Table 5.4) compared to parental Sr2SnO4 phase show the same trend as observed

in the values obtained only from I11 data analysis: an increase in volume and lattice

parameter c after the doping in comparison with Sr2SnO4. The presence of Oint within the

Sr2SnO4 structure after M5+

doping (Nb, Ta) remained questionable as the dopant levels are

relatively small and the adding of Oint did not improve the refinements significantly. Other

techniques (including spectroscopic methods) were used to explore the doping chemistry of

Sr2Sn1−xMxO4 (M = Nb, Ta).


153

Figure 5.12: HRPD data refinement of Sr2Sn0.97Nb0.03O4 using Pccn space group with the presence of

secondary phases of SrSnO3 and Sr3Sn2O7, a) bank 1; b) bank 2; Rwp = 2.083%.

Figure 5.13: HRPD data refinement of the Sr2Sn0.96Ta0.04O4 material using Pccn space group with the

presence of secondary phases of SrSnO3 and Sr3Sn2O7, a) bank 1; b) bank 2; Rwp = 2.724%.

Table 5.4: Refined lattice parameters of the Sr2Sn0.97Nb0.03O4 and Sr2Sn0.96Ta0.04O4 material from

HRPD data after the Rietveld refinement in Topas and their comparison with the undoped Sr2SnO4

phase.

Material a (Å) b (Å) c (Å) V (Å)

Sr2SnO4 250

5.72898(5) 5.73524(5) 12.58110(6) 413.378(6)

Sr2Sn0.97Nb0.03O4 5.73644(6) 5.74219(6) 12.65926(6) 416.992(6)

Sr2Sn0.96Ta0.04O4 5.73478(3) 5.74230(3) 12.66306(6) 417.005(4)


154

5.6 AC Electrochemical Impedance Spectroscopy (EIS)

5.6.1 AC impedance data at 600 ‒ 900°C

The conductivity of single phased (confirmed by laboratory PXRD data analysis) Nb-

and Ta-doped Sr2SnO4 phases was measured using AC impedance spectroscopy and

compared with the values obtained for the undoped phase. A powder of materials was ball

milled in EtOH for overnight. Ball milled and dried powder was pressed in a 10 mm pellet

using a uni-axial press at a pressure of 3.5 tons. Approximately 0.6 g of the material was

weighed out in order to obtain 2 mm thick electrolyte. Cold isostatic pressing (CIP,

Section 2.4.3) was used to increase the density of the pellets applying a pressure

of ≈ 200 MPa. Samples were then heated for 24 h following the synthesis conditions for each

of the doped materials (Section 5.3). The density of all of the measured pellets was calculated

from the obtained values of mass and volume and it was over 97% of crystallographic density

for each of the pellets. Since synthesised materials showed hygroscopic properties,

Archimedes' Principle balance could not be used for density determination. The surface of the

pellets was manually polished after the sintering. Gold wire, mesh, and paste were used as

current collectors.

AC impedance data were collected in ambient air over the temperature range

500 ‒ 800 °C and at the frequencies 1 MHz to 0.01 Hz. The samples were dwelled for 90 min

each 50 °C to reach the thermal equilibrium. Each of the experiments started with 90 min

dwelling at 600 °C. The data were collected using Smart software and equivalent circuit

modelling were performed using the ZView software.196

More details about the data analysis

are described in Section 2.5.2. The spectra showed one high-frequency arc (bulk contribution)

and one intermediate frequency arc (grain boundary contribution). The high-frequency

intercepts of the impedance arcs with the Z' axis was used for the total conductivity

determination, which was calculated from the total resistance, Rtot (bulk + grain boundary).

An example of the AC impedance data collected on Sr2SnO4 material is shown in Figure

5.14. The values of total conductivities are shown in Table 5.5-5.7.


155

Figure 5.14: AC impedance spectroscopy data of the Sr2SnO4 collected at temperatures of 700, 750

and 800 °C.

Figure 5.15 shows the total conductivities for the temperatures 600 ‒ 800 °C of the

Nb-doped materials compared with the undoped Sr2SnO4 phase (with the values in Table 5.5).

All of the measured materials show an increase of conductivity with the increasing

temperature. The increase of the conductivity follows linear trend in logarithmic scale of σtot.

Sr2Sn0.97Nb0.03O4 material tends to show two linear regions (Figure 5.15). The shape of the

trend line suggests a phase transition might have occurred but more data points at the

temperatures between 650 ‒ 700 °C are needed to confirm that. Activation energy of

Sr2Sn0.97Nb0.03O4 was calculated using the data obtained at the temperatures 700 ‒ 800 °C

while for the rest of the doped Sr2SnO4 and the undoped phase data from 600 ‒ 800 °C region

were used. At the temperature of 600 °C the value of the total conductivity of the

Sr2Sn0.98Nb0.02O4 material is one order of magnitude higher than the ones collected on the

undoped phase. There is no significant improvement of conductivities of both Nb-doped

materials compared to parent Sr2SnO4 phase at the higher temperature of 700 ‒ 800 °C.


156

Figure 5.15: The total conductivity and activation energy values of the Sr2Sn1−xNbxO4 materials

(x = 0, 0.02, 0.03) collected using AC impedance spectroscopy at the temperatures 600 ‒ 800 °C.

Table 5.5: The values of the total conductivity of the Sr2Sn1−xNbxO4 materials (x = 0, 0.02, 0.03)

collected using AC impedance spectroscopy at the temperatures 600 ‒ 800 °C.

x σtot (S cm−1

)

600 °C 650 °C 700 °C 750 °C 800 °C

0 4.76 × 10−6

8.56 × 10−6

1.85 × 10−5

2.12 × 10−5

9.20 × 10−5

0.02 1.29 × 10−5

2.61 × 10−5

3.28 × 10−5

1.13 × 10−4

1.18 × 10−4

0.03 1.86 × 10−6

4.59 × 10−6

1.28 × 10−5

3.60 × 10−5

1.10 × 10−4

Total conductivity and activation energy values for the data obtained at the

temperatures 600 ‒ 800 °C of the Ta-doped Sr2SnO4 material with nominal values of x = 0,

0.2, 0.3, and 0.4 are shown in Figure 5.16 and Table 5.6. The values of the total conductivity

of Ta-doped materials are higher compared to the undoped Sr2SnO4 phase. The improvement


157

is more significant for the x = 0.03 and 0.04 samples were the σtot values are for several

temperatures over one order of magnitude higher than those collected for the undoped

material. The conductivity of Sr2Sn0.98Ta0.02O4 material remains close to the Sr2SnO4 for all

the values in 600 ‒ 800 °C. The values of activation energies are significantly increased

(0.5 eV and more compared to Ea value of Sr2SnO4) for all of the Ta-doped samples.

Figure 5.16: The total conductivity and activation energy values of the Sr2Sn1−xTaxO4 materials

(x = 0, 0.02, 0.03, 0.04) collected using AC impedance spectroscopy at the temperatures

600 ‒ 800 °C.

Table 5.6: The values of the total conductivity of the Sr2Sn1−xTaxO4 materials (x = 0, 0.02, 0.03, 0.04)

collected using AC impedance spectroscopy at the temperatures 600 ‒ 800 °C.

x σtot (S cm−1

)

600 °C 650 °C 700 °C 750 °C 800 °C

0 4.76 × 10−6

8.56 × 10−6

1.85 × 10−5

2.12 × 10−5

9.20 × 10−5

0.02 3.99 × 10−6

1.34 × 10−5

3.38 × 10−5

7.87 × 10−5

2.83 × 10−4

0.03 2.45 × 10−5

4.06 × 10−5

1.07 × 10−4

2.63 × 10−4

2.30 × 10−3

0.04 5.38 × 10−5

1.76 × 10−4

5.32 × 10−4

9.91 × 10−4

2.61 × 10−3


158

Previous impedance data show the values of the total conductivity at the temperatures

600 ‒ 800 °C. The conductivity of samples with the highest Nb and Ta doping levels

(x = 0.03 and 0.04 respectively) was also measured for the higher temperatures until 900 °C.

Table 5.7 shows the comparison of this data with parental Sr2SnO4 phase. The value of total

conductivity for Sr2Sn0.97Nb0.03O4 at 900 °C is two times lower than that for the Sr2SnO4

phase. Sr2Sn0.96Ta0.04O4 sample, similarly to lower temperatures range shows an improvement

of the total conductivity values. The highest value collected at 900 °C: 1.02 × 10−2

S cm−1

is

more than one order of magnitude higher than for the Sr2SnO4 phase. This value is

comparable (one order of magnitude lower) to YSZ256

and comparable to the reported values

of ceria based electrolytes, GDC and SDC.257,258

Table 5.7: The values of total conductivities of the Sr2SnO4 and Nb0.03 and Ta0.04 doped materials

collected using AC impedance spectroscopy at the high temperatures 800 ‒ 900 °C.

Material σtot (S cm−1

)

800 °C 850 °C 900 °C

Sr2SnO4 9.20 × 10−5

4.51 × 10−4

9.18 × 10−4

Sr2Sn0.97Nb0.03O4 1.10 × 10−4

2.87 × 10−4

4.52 × 10−4

Sr2Sn0.96Ta0.04O4 2.61 × 10−3

4.76 × 10−3

1.02 × 10−2

5.6.2 AC impedance data at 300 ‒ 600°C

Impedance data of Sr2SnO4 derivatives mentioned in previous section showed an

improvement of the conductivities for the Ta-doped materials and almost no change in the

case of Nb doping. It is important to find out the origin of the increased conductivity in

Sr2Sn1−xTaxO4 materials (i.e. due to electronic or ionic contribution). This can be due to the

few different compensation effects after the Ta doping: reduction of the Sn4+

to Sn2+

,

presence of free electron carriers (electrons) or the presence of Oint. The structural study of

the HRPD data did not give the clear answer about the presence of Oint. The presence of

interstitial oxygen would significantly increase the ionic conductivity which can be observed

at the temperature higher than 600 °C. Gradual increase of the conductivity would point

towards the electronic contribution (implying mixed-valence Sn4+

/Sn2+

).


159

The values of the impedance data at lower temperatures 300 ‒ 600 °C (collected for

90 min dwelling at every 100 °C step) are shown in Figure 5.17 and Table 5.8. The data were

collected on the undoped material and single phased materials with highest doping level of

Nb and Ta (Sr2Sn0.97Nb0.03O4 and Sr2Sn0.96Ta0.04O4 respectively). For the Nb-doped sample

the values of total conductivity are lower than the ones collected for the undoped Sr2SnO4

material. The comparison of the Sr2Sn0.96Ta0.04O4 sample with undoped material shows

higher values of the conductivity of the Ta-doped material. The difference between the

materials is decreasing with the increase of the temperature. Gradual increase of the

conductivity could be explained by the increased electronic conductivity. The trend of the

conductivity values in lower temperature region for the data collected on the undoped and

Sr2Sn0.96Ta0.04O4 material can be used as evidence tool for the conductivity contribution in

Ta-doped materials but not as a main proof. Further study of the electronic structure and its

changes after the doping is needed, including UV-vis spectroscopy and Sn solid state NMR.

These methods can provide more precise results and better evidence about the valence state

of Sn in Sr2Sn0.96Ta0.04O4.

Figure 5.17: The total conductivity values and activation energies of the Sr2SnO4 and Nb0.03 and Ta0.04

doped materials collected after the AC impedance at the temperatures 300 ‒ 600 °C.


160

Table 5.8: Total conductivity values of the Sr2SnO4 and Nb0.03 and Ta0.04 doped materials collected

after the AC impedance at the temperatures 300 ‒ 600 °C.

Material σtot (S cm−1

)

300 °C 400 °C 500 °C 600 °C

Sr2SnO4 1.75 × 10−7

2.67 × 10−7

4.12 × 10−6

1.68 × 10−5

Sr2Sn0.97Nb0.03O4 2.86 × 10−8

2.99 × 10−7

1.14 × 10−6

4.31 × 10−6

Sr2Sn0.96Ta0.04O4 6.29 × 10−7

3.31 × 10−6

9.99 × 10−6

3.52 × 10−5

5.6.3 AC impedance data at different partial oxygen pressure

One of the requirements for electrolytes for SOFCs is the stability and constant

electrochemical properties for a wide range of oxidising/reducing conditions. More details

about the electrochemical processes can be obtained from the p(O2) vs σt dependence. Thus

AC impedance data of Ta-doped materials were collected under various p(O2) pressure.

Pellets were prepared following the procedure mentioned at the beginning of Section 5.6.1.

Measurement with the Sr2Sn0.97Ta0.03O4 material was carried out in a large oxygen partial

pressure region of p(O2) from 1 to 10−4

atm for the various temperatures increasing from 600

and 700 to 800 °C. Every step at different p(O2) was dwelled for 60 min while the collected

AC impedance data were stabilised. Measurements at each individual temperature were

performed during a day. The sample was kept at 600 °C and atmospheric O2 pressure

overnight between the measurements. AC impedance data of the cell were collected at these

conditions before heating to higher temperatures. The obtained data were compared with

those collected before in order to see any changes of impedance arc, which would indicate a

presence of unwanted processes such as reduction occurring during the measurement. Pellet

of the Sr2Sn0.96Ta0.04O4 material was tested for wider p(O2) range: 1 to 10−10

atm at the

temperature of 600 °C.

Figure 5.18 shows the AC impedance data at various p(O2) collected on the

Sr2Sn0.97Ta0.03O4 material. The shape of each of the individual arcs (Figure 5.18a) remains

the same for all the measured p(O2). The total conductivity is a sum of contributions of the

bulk and grain boundary. The changes of the conductivity in studied p(O2) range were


161

negligible (Figure 5.18b) at all of the measured temperatures. The calculated slope of the data

showed in logarithmic scale is for all of the measured temperatures close to 0. That means

that the studied Sr2Sn0.97Ta0.03O4 material behaves in large oxygen partial pressure region of

p(O2) from 1 to 10−4

atm as pure ionic conductor.

Figure 5.18: EIS data of the Sr2Sn0.97Ta0.03O4 material collected at various temperature (600, 700, and

800 °C) and at different p(O2) partial pressure; a) Nyquist plot of the impedance data collected at

various p(O2) at 800 °C; b) comparison of the log σtot values against log p(O2) collected at various

temperatures.

The EIS data of Sr2Sn0.96Ta0.04O4 material collected at various p(O2) partial pressure

are shown in Figure 5.19 and Table 5.9. For the higher p(O2) pressures we can see the same

behaviour as was observed for the measurement of the Sr2Sn0.97Ta0.03O4 sample. The values

of the total conductivity are all within close range (from 7.7 to 5.8 × 10−5

S cm−1

). The main

contribution of the bulk and grain boundary remains the same for the x = 0.04 sample as for

the x = 0.03. For the values of p(O2) lower than 10−5

atm there is an abrupt increase of the

conductivity with the total conductivity 6.7 × 10−3

S cm−1

at 10−7

atm which is two orders of

magnitude higher than it is observed at p(O2) = 10−5

atm. With further decrease of p(O2)

pressure the conductivity increases more.


162

Figure 5.19: EIS data of the Sr2Sn0.96Ta0.04O4 material collected at 600 °C. Inset shows the shape of

the AC impedance arcs at various p(O2) pressure.

Table 5.9: Values of the total conductivity of the Sr2Sn0.96Ta0.04O4 material collected at 600 °C

depending on various p(O2) pressure.

p(O2) (atm) σtot (S cm−1

)

2.1 × 10−1

7.77 × 10−5

1.0 × 10−2

6.43 × 10−5

1.1 × 10−3

5.82 × 10−5

1.2 × 10−5

6.21 × 10−5

1.8 × 10−7

6.66 × 10−3

1.2 × 10−10

7.93 × 10−2

The analysis of the PXRD data of Sr2Sn0.96Ta0.04O4 after the p(O2) measurement (Figure 5.20)

confirms the presence of the SrCO3 phase. The values of the refined lattice parameters are

shown in Table 5.10. There is a small change in lattice parameter a and b (on third decimal

place). Comparison of the c lattice parameters shows a significant decrease of the parameter


163

obtained from the data after the p(O2) measurement compared to as made Sr2Sn0.96Ta0.04O4.

The decrease of the c parameter is followed by the decrease in the volume. The change of the

conductivity can be explained by the mixed Sn4+

/Sn2+

valence state. The presence of Sr-

containing impurity phase might also indicate changes on A-site perovskite layer of

Ruddlesden-Popper phase, which could be partially occupied by Sn2+

. All these changes can

lead to decrease of lattice parameter c and cell volume. More important information related to

transport properties is included in the p(O2) vs conductivity dependence. There is no

significant change of conductivity observed in large oxygen partial pressure region of p(O2)

from 1 to 10−5

atm, which means Sr2Sn0.96Ta0.04O4 behaves as a pure ionic conductor. The

same trend was also observed for the Sr2Sn0.97Ta0.03O4 material. The increase in conductivity

at p(O2) < 10−5

atm suggests a change of behaviour − transition from pure ionic conductor to

mixed ionic-electronic conductor.

Figure 5.20: Pawley fit of the data of Sr2Sn0.96Ta0.04O4 material collected after the p(O2) measurement.


164

Table 5.10: Refined lattice parameters obtained from the Rietveld refinements of the data collected

on the undoped Sr2SnO4 material with as made Sr2Sn0.96Ta0.04O4 compared with the Ta0.04 doped

sample after the p(O2) measurement.


Sr2SnO4 5.72962(8) 5.73509(8) 12.5862(1) 413.58(1)

Sr2Sn0.96Ta0.04O4 as made 5.7273(2) 5.7389(3) 12.6722(4) 416.51(3)

Sr2Sn0.96Ta0.04O4 after p(O2) 5.7215(2) 5.7326(2) 12.5895(3) 412.92(2)

5.7 Thermal stability

Thermal stability of synthesized materials was monitored by a TGA experiment

(Section 2.9). Figure 5.21 shows the TGA measurement of the Sr2Sn0.97Nb0.03O4 material.

The experiment was carried out in alumina pan. Small amount of the sample (28.1720 mg)

was annealed in air from RT to 900 °C with a dwelling for 2 h and it was cooled down back

to RT with the rate 5 °C min−1

. The small change (0.4 weight %) observed at the beginning of

the heating is due to the moisture taken by the material before the measurement. Further

heating does not decrease the mass significantly. There are no other phases observed on the

PXRD data collected after the TGA experiment. Table 5.11 shows the comparison of the

lattice parameters of as made Nb0.03 doped and the Sr2Sn0.97Nb0.03O4 material after the TGA

analysis. The differences in the values of lattice parameters obtained after the Rietveld

refinement are minimal in all cases. The thermal behaviour of Sr2SnO4 and Sr2Sn0.96Ta0.04O4

was also investigated using a TGA method. The materials exhibit the same behaviour over

the studied temperature range (RT ‒ 900 °C) as it was observed in the TGA measurement of

Sr2Sn0.97Nb0.03O4.


165

Figure 5.21: TGA of the Sr2Sn0.97Nb0.03O4 material at ambient air sintered from RT to 900 °C and

cooled back to RT.

Table 5.11: Comparison of the lattice parameters and cell volume of Sr2SnO4 and Sr2Sn0.97Nb0.03O4 as

made material and Sr2Sn0.97Nb0.03O4 after the TGA experiment. All of the parameters were obtained

using a Rietveld refinement.


Sr2SnO4 5.72962(8) 5.73509(8) 12.5862(1) 413.58(1)

Sr2Sn0.97Nb0.03O4 as made 5.74330(7) 5.73340(6) 12.6564(1) 416.759(7)

Sr2Sn0.97Nb0.03O4 after TGA 5.7410(6) 5.7334(6) 12.4654(5) 416.46(5)


166

5.8 UV-vis spectroscopy measurements

5.8.1 As made materials

Sr2SnO4 based materials have attracted attention due to various optical properties,

such as photoluminescence or machanoluminescence.259,260

Previous works on Sn containing

materials (doped SnO2, SrSnO3, Sr2SnO4) showed direct band transitions occurring.163,259,261

Direct band gaps were also calculated for Sr2SnO4 related materials of our study. Linear fits

for the direct band gap determination of corresponding material are shown in Figure 5.22c-e.

A-site Sr cation in Sr2SnO4 has a minor contribution to the electronic structure near the Fermi

level.262

The band gap depends mainly on the bonding of the octahedral network of SnO6.

Conduction band is expected to come from O 2p states whilst the valence band is

predominantly composed of Sn 5s states.262,263

UV-vis spectroscopy was used to look at the

electronic structure of parent Sr2SnO4 and the doped phases in order to find any evidence of

Sn2+

within the structure of doped materials. The diffuse reflectance of solid materials of

Sr2SnO4, Sr2Sn0.97Nb0.03O4, and Sr2Sn0.96Ta0.04O4 was measured on a Shimadzu UV-2550

UV-vis spectrometer. The powder of the measured sample was put in a solid sample holder

with a quartz glass window and placed in the UV-vis spectrometer. The reflectance was

recorded over the spectral range of 200 ‒ 800 nm (6.2 ‒ 1.55 eV). The reflectance at each

wavelength was converted to F(R) using Equation 2.35 (Section 2.6). The data obtained for

all of the samples was converted for both values of n = 1/2 for direct transition.

Parental Sr2SnO4 phase exhibited a band gap of 4.00 eV, which is smaller than the

reported 4.43 eV.259

That could be due to the different crystallinity of the material or different

synthetic route. The purpose of this study was the comparison of the band gap values and

electronic structures of the parental Sr2SnO4 phase with its derivatives. The change of the

electronic structure could be related to the presence of dopant such as in the reported works

of the doped SnO2,163,264-267

or due to the mixed Sn4+

/Sn2+

valence state. The comparison

between the parent and the doped phases is shown in Figure 5.22a, b. For both Nb- and Ta-

doped material, a blue shift towards the higher energy levels (equal to lower wavelengths)

have been observed. This trend is observed also on the values of band gap, which are higher

than that for Sr2SnO4 phase; 4.24 eV for Sr2Sn0.97Nb0.03O4 and 4.40 eV for Sr2Sn0.96Ta0.04O4.

The increase of band gap value was observed in several Sn-containing materials upon M5+

doping.267,268


167

Figure 5.22: a) Diffuse reflectance spectra of Sr2SnO4, Sr2Sn0.97Nb0.03O4, and Sr2Sn0.96Ta0.04O4

collected over the range of 200 ‒ 800 nm; b) spectra representation using F(R) function; the

extrapolation of the linear section of the diffuse reflectance of c) Sr2SnO4 d) Sr2Sn0.97Nb0.03O4, and e)

Sr2Sn0.96Ta0.04O4, plotted to determine the direct band gap, with the intercept of the x axis (direct band

gap) at c) 4.00 eV, d) 4.24 eV, and e) 4.40 eV.

5.8.2 Reduced materials

The change of the electronic structure of Sr2SnO4 related materials after the doping

and the variation (in the case of Ta-doped Sr2SnO4 phases increase, Section 5.6.1) of

conductivity could be explained due to the presence of Sn in mixed Sn4+

/Sn2+

oxidation state.

Section 5.8.1 showed the diffuse spectra of the Sr2SnO4 related materials as made. These

samples were then reduced in under N2 at 600 °C (following an established literature

procedure for indium-tin-oxide reduction).269

Direct band gaps were calculated from the

spectra obtained after the reduction of the materials (linear fits are shown in Figure 5.23c-e).

All three of the samples show a decrease of the band gap values compared to those

obtained from as made compounds and a shift in their spectra towards lower wavelengths

(Figure 5.23a, b). Reduced Sr2SnO4 sample exhibits the band gap value of 3.57 eV and a red

shift compared to the as made compound with band gap of 4.00 eV. The shift can be

explained by the presence of tin in a mixed oxidation state of Sn4+

and Sn2+

. Band gap values


168

of reduced phases of Sr2Sn0.97Nb0.03O4 and Sr2Sn0.96Ta0.04O4 were of 3.91 eV and 4.01 eV

respectively. Band gap values of all of the reduced samples are lower compared to as made

materials. One of the possible compensating effects after M5+

doping of Sr2SnO4 is the

presence of Sn4+

/Sn2+

mixed oxidation state. Reduced Sr2SnO4 (with mixed Sn4+

/Sn2+

valence

state) exhibits narrower band gap compared to as made RP1 phase. If there was a Sn4+

/Sn2+

mixed oxidation state in Sr2Sn0.97Nb0.03O4 and Sr2Sn0.96Ta0.04O4, the band gaps for as made

materials would be expected to be smaller compared to the parent Sr2SnO4 phase. The values

are on the contrary higher, which suggests that Sn is present only in Sn4+

oxidation state. This

can be confirmed by solid state 119

Sn NMR (Section 5.9).

Figure 5.23: a) Diffuse reflectance spectra of the reduced samples of Sr2SnO4, Sr2Sn0.97Nb0.03O4, and

Sr2Sn0.96Ta0.04O4 collected over the range of 200‒800 nm; b) spectra representation using F(R)

function; the extrapolation of the linear section of the diffuse reflectance of c) Sr2SnO4 d)

Sr2Sn0.97Nb0.03O4, and e) Sr2Sn0.96Ta0.04O4, plotted to determine the direct band gap, with the intercept

of the x axis (indirect band gap) at c) 3.57 eV, d) 3.91 eV, and e) 4.01 eV.

5.9 Sn Solid-state NMR

Previous work on the doped Sr2SnO4 materials showed an increase in conductivity


169

after the doping with Ta atom. Neutron diffraction did not confirm the presence of the Oint of

the doped samples. Sn solid-state NMR was used to investigate the oxidation state of the Sn

atom in the Ta-doped Sr2SnO4 materials. Solid-state NMR experiments were performed on a

9.4 T Bruker DSX NMR spectrometer equipped with a 4 mm HXY triple-resonance MAS

probe (in double resonance mode) tuned to 119

Sn at ν0(119

Sn) = 149.1 MHz. 119

Sn single pulse

experiments were performed at room temperature under magic angle spinning (MAS) at

frequency rate of νr = 12.5 kHz, using a π/2 pulse width of 3 μs (i.e. at a radio-frequency (rf)

field amplitude of ν1(119

Sn) = 83 kHz) and a recycle delay of 60 s allowing full relaxation of

the 119

Sn spins. 1024 – 4096 scans were averaged with (experimental time ranging from 17 h

to 3.5 days). 119

Sn chemical shifts were externally referenced to SnO2 at iso = −604.3 ppm.

Figure 5.24 shows the obtained 119

Sn MAS NMR spectra of Sr2Sn1−xTaxO4 (x = 0,

0.02, 0.04). The peak at −564 ppm corresponds to Sn4+

in Sr2SnO4. The peak has a multiplet

lineshape coming from spin-spin coupling between 119

Sn and 117

Sn (doublet as a spin 1/2 ‒

approximately 7 % of the signal) and between 119

Sn and 115

Sn (doublet too). This coupling

arises from crystallographically equivalent but magnetically inequivalent 119

Sn, 117

Sn and

115Sn nuclei). All three samples show residual SnO2 at −604 ppm (approximately 1% of SnO2

in Sr2SnO4 and Sr2Sn0.98Ta0.02O4 and 2% in Sr2Sn0.96Ta0.04O4). SnO2 may relax quicker than

Sr2SnO4, thus its signal will be more intense. High resolution diffraction data collected using

the I11 instrument at the Diamond Light Source (UK) did not show any presence of SnO2,

which may indicate that the amount of SnO2 is very small (< 1%). There is no Sn2+

peak

presented (expected to be at −200 ppm vs −600 for Sn4+

). The shift between Sn2+

and Sn4+

is

quite clear. The fact that there is no Sn2+

presented is important in understanding the defect

chemistry in Ta-doped Sr2SnO4. Enlarged view (Figure 5.25) of the Sn4+

region shows two

peaks at −553 and −580 ppm appearing upon Ta doping. The intensities of these peaks seem

to increase with Ta doping level. The additional peaks might be related to the secondary

phases of SrSnO3 and Sr3Sn2O7 which were used in both I11 and HRPD refinements and they

fitted low-intensity and low-angle peaks (Sections 5.5.2 and 5.5.3). Solid state NMR spectra

of the Sr2Sn1−xTaxO4 compounds show no evidence of Sn2+

presented after the Ta-doping,

implying no mixed valence of Sn4+

/Sn2+

, which was also suggested by the UV-vis

measurements. Ta-doping is compensated for another way − by the presence of Oint or free

electrons. The evidence of free electrons can be confirmed by IR spectra (Chapter 5.10).


170

Figure 5.24: 119

Sn MAS NMR spectra of Sr2Sn1−xTaxO4 obtained at 9.4 T. Asterisks denote spinning

side bands.


171

Figure 5.25: Magnified view of the 119

Sn MAS NMR spectra of Sr2Sn1−xTaxO4 obtained at 9.4 T.

5.10 IR spectra

The spectra of 10 mm pellets of Sr2SnO4, Sr2Sn0.97Nb0.03O4, and Sr2Sn0.96Ta0.04O4

materials were measured using a Shimadzu SolidSpec-3700 UV-VIS-NIR Spectrophotometer

(see Section 2.6.2 for more details).

Before the analysis of the collected spectra of Sr2SnO4 related phases, a short

description of transparent conductors must be given. Previous chapters of doped stannates


172

showed an improvement of conductivity properties after doping with Ta (Section 5.6.1). It is

important to reveal the reason of the increase in conductivity. HRPD data analysis showed no

Oint (within the experimental error limits) within the doped structure (Section 5.5.3) whilst

solid state NMR demonstrated the presence of Sn4+

only. The enhancement of conductivity

can be due to the presence of free electrons which is well known and described in transparent

conductors (TCO).163

The opto-electronic properties of conventional TCO materials can be

represented by a typical TCO representative ZnO (Figure 5.26). The gradual decrease in the

transmission starts at ≈ 1000 nm while the increase in the reflection is observed at ≈ 1500 nm.

The increase of reflectivity (R) at this region is typical for TCO. If the number of electrons in

the conduction band is increased, the wavelength shifts to shorter wavelengths, at the very

high electron concentration even in the visible region.163

This characteristic feature of TCO

was considered in the analysis of the spectra of Sr2SnO4 related materials.

Figure 5.26: Optical spectra of ZnO transparent conductor, taken from163

The spectra of Sr2SnO4, Sr2Sn0.97Nb0.03O4, and Sr2Sn0.96Ta0.04O4 materials collected

over the range of 1600 ‒ 2500 nm are shown in Figure 5.27. No increase in reflection, typical

for TCO is observed for either of the studied materials. A decrease of reflectivity at the

region of ≈ 1600 nm is found for all of the three samples. This bump lies at the lowest region

where PbS detector is used and thus a noisy spectra are commonly collected at this region.

The inset of Figure 5.27 shows the spectra of the same materials collected at wavelengths

from 1000 nm to 1700 nm using InGaAs detector especially focused on the decrease of the


173

reflectivity at ≈ 1600 nm which is found to be less than at previous spectra (1600 ‒ 2500 nm,

using PbS detector). Even at the wavelengths above 2000 nm no increase of reflectivity is

observed, this tends to indicate that no increased concentration of electrons in the conduction

band is found. It is worth noting that the presented method is mainly used on TCO materials

prepared as thin films while our spectra were collected on solid pellets. It is expected that the

presence of grain boundaries within the solid sample might reduce the sensibility of the

method and thus might not be a strong tool for pellet samples.

Figure 5.27: NIR spectra of Sr2SnO4, Sr2Sn0.97Nb0.03O4, and Sr2Sn0.96Ta0.04O4 material measured at the

wavelengths between 1600 and 2500 nm. The inset shows the spectra of the same materials collected

in 1000 ‒ 1750 nm region.


According to the theoretical study of A- and B-site doped Ruddlesden-Popper phases

e.g. A2SnO4, A2HfO4, A2ZrO4, La3+

and M5+

doping is favourable which was shown by

negative values of formation energies obtained in energetic comparison to parental phases

based on the data obtained from DFT calculations. Experimental work was focused on the


174

Sr2SnO4 related phases. Nb- and Ta- doped Sr2SnO4 were prepared by a solid state synthetic

route. Laboratory obtained PXRD data showed single phase materials with the maximum

doping level of 3% for Nb and 4% Ta. The incorporation of Ta in the RP1 structure was

confirmed by EDX and the average cation ratio was in a good agreement with the expected

stoichiometry. The lattice parameter evolution of both Nb- and Ta-doped materials shows the

same step-like increase of cell volume and c lattice parameter at the maximum doping levels

of single phased samples. There is no further increase observed for higher doping levels

(higher than 3% for Nb and 4% Ta-doped materials). The observed increase can be explained

by charge compensation effect (presence of Oint or Sn4+

/Sn2+

mixed oxidation state). La-

doped materials, the most stable (with most negative formation energies) ‒ as obtained from

DFT calculation, were not synthesised as single phases due to the presence and chemical

stability of La2O3 even after attempts to improve the synthesis conditions.

Ta-doped samples exhibit a significant increase of the conductivity obtained by AC

impedance spectroscopy whilst the Nb-doped materials show similar values as those for the

parental Sr2SnO4 phase. It is of much interest to find out the origin of the enhanced

conductivity of Sr2Sn1−xTaxO4, especially for x = 0.04 where the total conductivity in

temperature range of 700 ‒ 800 °C was increased by more than one order of magnitude

compared to Sr2SnO4. The highest total conductivity 1 × 10−2

S cm−1

of Sr2Sn0.96Ta0.04O4

compound was measured at 900 °C, which is one order of magnitude lower than that obtained

for YSZ.72

Other common electrolytes such as ceria based, exhibit similar conductivity

properties but at the lower temperatures (750 °C): 6.7 × 10−2

S cm−1

for GDC76

and

6.1 × 10−2

S cm−1

for SDC.77

La0.8Sr0.2Ga0.83Mg0.17O2.815 electrolyte shows conductivity of

0.08 S cm−1

at 700 °C.84

The conductivity of Sr2Sn0.97Ta0.03O4 and Sr2Sn0.96Ta0.04O4 were also

studied under various partial oxygen pressures. In high oxygen partial pressure region

(p(O2) ˃ 10−5

atm) both of the studied materials showed no change in conductivity, which is

common for pure ionic conductors. A significant increase of conductivity (of two orders of

magnitude) was observed at the p(O2) < 10−5

atm at 600 °C for Sr2Sn0.96Ta0.04O4. That

indicates mixed ionic-electronic conductivity due to the reduction of Sn4+

to Sn2+

. Potential

SOFC electrolyte needs to exhibit constant oxygen ion conductivity over wide range of p(O2).

Any electronic contribution to total conductivity is detrimental for any practical use as an

electrolyte for SOFCs.

Three possible mechanisms explaining the defect chemistry and improvement of

conductivity in Sr2Sn1−xTaxO4 were taken into an account: increasing of ionic conductivity

due to the presence of interstitial oxygen atoms, electronic contribution as a result of partial


175

reduction of tin and therefore mixed valence state of Sn4+

and Sn2+

and electronic

contribution due to the presence of free electrons typical for transparent conductors. The

presence of Oint was investigated using the high resolution data from synchrotron (I11) and

neutron (HRPD) sources. The analysis of the data indicates stacking defects in crystal

structure. The doping level of studied phases was rather small (3 and 4% respectively) and

even using high resolution data gives no certain answer. The interstitial oxygen

concentration, if there is any, is very low and is most likely lower than the detection limits of

HRPD. The presence of stacking faults which could be explained by the changes of synthesis

route (using a planetary ball mill to decrease the total synthesis time from seven to five days)

makes the Oint determination even more difficult.

Other methods such as Sn solid state NMR, UV-vis and infrared spectroscopy, have

been used to reveal the chemistry of M5+

doping on B-site of Sr2SnO4 materials. UV-vis

spectra of Sr2SnO4, Nb- and Ta-doped compounds show a blue shift of the doped material

compared to the parent phase expressed by the band gap values. More interesting is the

spectra comparison of the reduced samples. Reduced Sr2SnO4, with expected mixed valence

of Sn4+

and Sn2+

, shows a lower band gap value of 3.57 eV than the one of as made Sr2SnO4.

Comparison of the band gap values of the doped materials tends to show that Sn atoms are

not presented in mixed oxidation state of Sn4+

and Sn2+

. Further work based on solid state Sn-

NMR clearly showed no Sn2+

presented in either the parent or the Sr2Sn1−xTaxO4 (for x = 0.02

and 0.04) phases.

An increase of reflectivity, typical for transparent conductors with free electrons, was

not observed on the IR spectra of parental and doped materials (Sr2Sn0.97Nb0.03O4 and

Sr2Sn0.96Ta0.04O4). The increase is observed at various wavelengths, which depends on the

electrons concentration. This absence could be a result of no or very low concentration of free

electrons in the doped Sr2SnO4 phases. The method conditions (samples is in polycrystalline

pellet form instead of thin films), which reduce the sensibility, must be taken into an account

as well.

In conclusion, Ta substitution in Ruddlesden-Popper Sr2SnO4 phases has improved

the total conductivity properties. Electrochemical performance may be improved more by

altering the processing conditions. Following study for a potential use of this material as a

SOFC electrolyte, revealed severe drawbacks for a practical use such as inconstant

electrochemical behaviour under lower p(O2) atmospheres. Several different techniques have


176

been applied in order to determine the defect chemistry in Sr2Sn1−xTaxO4. The results tend to

show an ionic contribution to the enhanced conductivity.

Chapter 6. General Conclusions and Perspectives

6 General Conclusions and Perspectives

The aim of this work was to find new materials and improve the properties of already

known materials for solid oxide fuel cells. The material investigations were based on both

experimental and theoretical methods.

The first results chapter (Chapter 3) includes an examination of the cathode properties

of the triple perovskite Y1−ySr2+yCu3−xCoxO7+δ. The promising electrochemical properties

reported in the literature for YSr2Cu2CoO7+δ69,70

are devalued by significant problems with

compatibility and electrochemical stability. Further substitution of Cu for Co presented in this

work, has considerably improved the cathode performance compared to the material with the

parent material. The values of ASR are in agreement with general SOFC cathode

requirements. The limits of the nonstoichiometry on A-site, needed for the synthesis of single

phase compounds with higher dopant amount, was not studied in full and thus higher Co

doping levels might be achievable. It is worth noting that the Co enhancement retains a

crystal structure favourable for high oxide ion mobility. The crystal structure of

Y0.95Sr2.05Cu1.7Co1.3O7+δ compound was studied in more details using the neutron diffraction

data. Sufficient electrochemical stability is a common issue for state of the art cathode

material research and more stability research is, also required for 3ap structured materials.

Most of the known Co-containing cathode materials exhibit thermal expansion higher than

that of common electrolyte while the thermal expansion of Y0.95Sr2.05Cu1.7Co1.3O7+δ material

matches well with YSZ and LSGM electrolytes. The structural family of 3ap materials

represents interesting cathode candidates with further study on material optimisations

principally needed to improve the electrochemical stability.

The aim of a theoretical approach based on DFT calculations of GBCO related

structures (Chapter 4) was to find candidates with double perovskite and fluorite layer

suitable for further experimental work. Energetics of these materials were studied including

several phases as potential by-products to predict thermodynamically favourable phases,

which were then observed by PXRD data after the solid state synthesis. Synthesis of the

materials with a CeO2 fluorite layer showed the presence of the starting building blocks of

double perovskite and fluorite layer resulting in no CeO2 implementation within the double

perovskite layer. The charge misbalance between the layers was suppressed by changing the

Chapter 6. General Conclusions and Perspectives

178

fluorite layer by Gd2O3 and Nd2O3 respectively. Other structures, such as Ruddlesden-Popper

phases are thermodynamically accessible, which was also taken into an account in theoretical

study. Primary solid state synthesis showed Nd2BaCo2O7 RP2 phase as a possible candidate

for further investigation including synthesis optimisation and material properties study

although it does not possess the required structure.

Selected A- and B-site doped Ruddlesden-Popper phases and their formation energies

compared to parental phases were part of the study presented in Chapter 5. The DFT study

showed preferential B-site doping in the case of Nb and Ta, which was confirmed by the

single phased Nb- and Ta-doped Sr2SnO4 materials obtained after conventional solid state

synthesis. The conductivity properties of Sr2Sn1−xTaxO4 with x = 0.03 and 0.04 were

increased by more than one order of magnitude compared to those obtained from the parent

Sr2SnO4 phase. The following study was focused on identifying the chemical mechanism

responsible for the enhancement of the conductivity in Sr2Sn1−xTaxO4. Crystal structure

investigation based on synchrotron and neutron diffraction data revealed stacking faults and

was unable to show definitely whether oxide Oint had formed within the RP. Solid state Sn-

NMR data analysis showed no Sn2+

present, which does not support the mechanism of

increasing of the conductivity in Sr2Sn1−xTaxO4 by an electronic contribution due to the mixed

valence Sn2+

/Sn4+

state. The presence of free electrons after the Ta-doping, which was

assumed as another potential mechanism of the defect chemistry, was not confirmed by the

IR spectra collected on parent and doped materials. None of the used technique gave clear

answer about the mechanism explaining the improvement of the conductivity in

Sr2Sn1−xTaxO4, although the AC impedance measurement at various p(O2) showed

Sr2Sn1−xTaxO4 materials (x = 0.03 and 0.04) as ionic conductors in high oxygen partial

pressure region (p(O2) ˃ 10−5

atm). An increase of the dopant concentration could be helpful

to understand doping chemistry in this Ruddlesden-Popper phase. This should include an

optimisation of solid state synthesis or use of different synthetic routes, e.g. sol-gel methods.

References

References

(1) Kordesch, K. V.; Simader, G. R. Chem Rev 1995, 95, 191. (2) Dresselhaus, M. S.; Thomas, I. L. Nature 2001, 414, 332. (3) MacKay, D. Sustainable Energy - without the hot air; UIT Cambridge, 2009. (4) Craddock, D. Renewable Energy Made Easy; Atlantic Publishing Group: Florida, 2008. (5) Wengenmayr, R. a. B., T. Renewable Energy: Sustainable energy concepts for the future; Wiley-VCH: Weinheim, 2008. (6) Orera, A.; Slater, P. R. Chem Mater 2010, 22, 675. (7) Schoenbeinf, C. F. Philos Mag 1838, 14, 3. (8) Grove, W. R. Philos Mag 1839, 14, 4. (9) Steele, B. C. H.; Heinzel, A. Nature 2001, 414, 345. (10) Aguadero, A.; Alonso, J. A.; Escudero, M. J.; Daza, L. Solid State Ionics 2008, 179, 393. (11) Yamamoto, O. Electrochim Acta 2000, 45, 2423. (12) Istomin, S. Y.; Antipov, E. V. Russ Chem Rev+ 2013, 82, 686. (13) Steele, B. C. H. Solid State Ionics 2000, 134, 3. (14) Lashtabeg, A.; Skinner, S. J. J Mater Chem 2006, 16, 3161. (15) Brandon, N. P.; Skinner, S.; Steele, B. C. H. Annu Rev Mater Res 2003, 33, 183. (16) Steele, B. C. H. Solid State Ionics 1996, 86-8, 1223. (17) Jacobson, A. J. Chem Mater 2010, 22, 660. (18) Goodenough, J. B. Annu Rev Mater Res 2003, 33, 91. (19) Miyoshi, S.; Martin, M. Phys Chem Chem Phys 2009, 11, 3063. (20) Mehrer, H. Diffusion in Solids; Springer: Berlin Heidelberg, 2007. (21) Chroneos, A.; Yildiz, B.; Tarancon, A.; Parfitt, D.; Kilner, J. A. Energ Environ Sci 2011, 4, 2774. (22) Chroneos, A.; Parfitt, D.; Kilner, J. A.; Grimes, R. W. J Mater Chem 2010, 20, 266. (23) Chiang, Y.-M., Birnie, D. I. and Kingery, W. D. Physical Ceramics; John Wiley & Sons Inc, 1997. (24) Inaba, H.; Tagawa, H. Solid State Ionics 1996, 83, 1. (25) Kharton, V. V.; Wiley-VCH: Weinheim, 2009. (26) Tilley, R. J. D. Defects in Solids; John Wiley & Sons, Inc., 2008. (27) West, A. R. Solid State Chemistry and its Applications John Wiley & Sons, 1987. (28) Ishihara, T.; Springer: 2009. (29) Liu, Y. New J Phys 2010, 12. (30) Malavasi, L.; Fisher, C. A. J.; Islam, M. S. Chem Soc Rev 2010, 39, 4370. (31) Jiang, S. P. Solid State Ionics 2002, 146, 1. (32) Jiang, S. P. J Mater Sci 2008, 43, 6799. (33) Plonczak, P.; Gazda, M.; Kusz, B.; Jasinski, P. J Power Sources 2008, 181, 1. (34) Mai, A.; Haanappel, V. A. C.; Uhlenbruck, S.; Tietz, F.; Stover, D. Solid State Ionics 2005, 176, 1341. (35) Simner, S. P.; Anderson, M. D.; Pederson, L. R.; Stevenson, J. W. J Electrochem Soc 2005, 152, A1851. (36) Bak, T.; Nowotny, J.; Rekas, M.; Ringer, S.; Sorrell, C. C. Ionics 2001, 7, 380. (37) Figueiredo, F. M.; Marques, F. M. B.; Frade, J. R. Solid State Ionics 1998, 111, 273. (38) Mizusaki, J.; Tabuchi, J.; Matsuura, T.; Yamauchi, S.; Fueki, K. J Electrochem Soc 1989, 136, 2082. (39) Huang, K.; Lee, H. Y.; Goodenough, J. B. J Electrochem Soc 1998, 145, 3220.

References

180

(40) Tai, L. W.; Nasrallah, M. M.; Anderson, H. U.; Sparlin, D. M.; Sehlin, S. R. Solid State Ionics 1995, 76, 273. (41) Bae, J. M.; Steele, B. C. H. Solid State Ionics 1998, 106, 247. (42) Oishi, N.; Atkinson, A.; Brandon, N. P.; Kilner, J. A.; Steele, B. C. H. J Am Ceram Soc 2005, 88, 1394. (43) Dusastre, V.; Kilner, J. A. Solid State Ionics 1999, 126, 163. (44) Tietz, F.; Mai, A.; Stover, D. Solid State Ionics 2008, 179, 1509. (45) Tietz, F.; Fu, Q.; Haanappel, V. A. C.; Mai, A.; Menzler, N. H.; Uhlenbruck, S. Int J Appl Ceram Tec 2007, 4, 436. (46) Magnone, E. J Fuel Cell Sci Tech 2010, 7. (47) Shao, Z. P.; Haile, S. M. Nature 2004, 431, 170. (48) Shao, Z. P.; Haile, S. M.; Ahn, J.; Ronney, P. D.; Zhan, Z. L.; Barnett, S. A. Nature 2005, 435, 795. (49) McIntosh, S.; Vente, J. F.; Haije, W. G.; Blank, D. H. A.; Bouwmeester, H. J. M. Chem Mater 2006, 18, 2187. (50) Svarcova, S.; Wiik, K.; Tolchard, J.; Bouwmeester, H. J. M.; Grande, T. Solid State Ionics 2008, 178, 1787. (51) Mueller, D. N.; De Souza, R. A.; Weirich, T. E.; Roehrens, D.; Mayer, J.; Martin, M. Phys Chem Chem Phys 2010, 12, 10320. (52) Yan, A.; Cheng, M.; Dong, Y. L.; Yang, W. S.; Maragou, V.; Song, S. Q.; Tsiakaras, P. Appl Catal B-Environ 2006, 66, 64. (53) Chen, Z. H.; Ran, R.; Zhou, W.; Shao, Z. P.; Liu, S. M. Electrochim Acta 2007, 52, 7343. (54) Lee, K. T.; Bierschenk, D. M.; Manthiram, A. J Electrochem Soc 2006, 153, A1255. (55) Lee, K. T.; Manthiram, A. Chem Mater 2006, 18, 1621. (56) Aguadero, A.; Escudero, M. J.; Perez, M.; Alonso, J. A.; Pomjakushin, V.; Daza, L. Dalton T 2006, 4377. (57) Sun, C. W.; Hui, R.; Roller, J. J Solid State Electr 2010, 14, 1125. (58) Zinkevich, M.; Aldinger, F. J Alloy Compd 2004, 375, 147. (59) Amow, G.; Skinner, S. J. J Solid State Electr 2006, 10, 538. (60) Maignan, A.; Martin, C.; Pelloquin, D.; Nguyen, N.; Raveau, B. J Solid State Chem 1999, 142, 247. (61) Frison, R.; Portier, S.; Martin, M.; Conder, K. Nucl Instrum Meth B 2012, 273, 142. (62) Tarancon, A.; Skinner, S. J.; Chater, R. J.; Hernandez-Ramirez, F.; Kilner, J. A. J Mater Chem 2007, 17, 3175. (63) Chang, A. M.; Skinner, S. J.; Kilner, J. A. Solid State Ionics 2006, 177, 2009. (64) Tarancon, A.; Pena-Martinez, J.; Marrero-Lopez, D.; Morata, A.; Ruiz-Morales, J. C.; Nunez, P. Solid State Ionics 2008, 179, 2372. (65) Liu, Y. J Alloy Compd 2009, 477, 860. (66) Zhou, Q. J.; He, T. M.; He, Q.; Ji, Y. Electrochem Commun 2009, 11, 80. (67) Kim, J. H.; Cassidy, M.; Irvine, J. T. S.; Bae, J. J Electrochem Soc 2009, 156, B682. (68) Fletcher, J. G.; Irvine, J. T. S.; West, A. R.; Labrincha, J. A.; Frade, J. R.; Marques, F. M. B. Mater Res Bull 1994, 29, 1175. (69) Sansom, J. E. H.; Rudge-Pickard, H. A.; Smith, G.; Slater, P. R.; Islam, M. S. Solid State Ionics 2004, 175, 99. (70) Sansom, J. E. H.; Kendrick, E.; Rudge-Pickard, H. A.; Islam, M. S.; Wright, A. J.; Slater, P. R. J Mater Chem 2005, 15, 2321. (71) Etsell, T. H.; Flengas, S. N. Chem Rev 1970, 70, 339. (72) Badwal, S. P. S. Solid State Ionics 1992, 52, 23. (73) Fergus, J. W., Hui, R., Li, X., Wilkinson, D. P., Zhang, J. Solid Oxide Fuel Cells - Material Properties and Performance; CRC Press: Boca Raton, 2009. (74) Duwez, P.; Brown, F. H.; Odell, F. J Electrochem Soc 1951, 98, 356.

References

181

(75) Vanherle, J.; Horita, T.; Kawada, T.; Sakai, N.; Yokokawa, H.; Dokiya, M. J Eur Ceram Soc 1996, 16, 961. (76) Kudo, T.; Obayashi, H. J Electrochem Soc 1975, 122, 142. (77) Eguchi, K.; Setoguchi, T.; Inoue, T.; Arai, H. Solid State Ionics 1992, 52, 165. (78) Wang, F. Y.; Chen, S. Y.; Qin, W.; Yu, S. X.; Cheng, S. F. Catal Today 2004, 97, 189. (79) Lubke, S.; Wiemhofer, H. D. Solid State Ionics 1999, 117, 229. (80) Ishihara, T.; Matsuda, H.; Takita, Y. J Am Chem Soc 1994, 116, 3801. (81) Feng, M.; Goodenough, J. B. Eur J Sol State Inor 1994, 31, 663. (82) Huang, K. Q.; Tichy, R. S.; Goodenough, J. B. J Am Ceram Soc 1998, 81, 2565. (83) Huang, K. Q.; Tichy, R. S.; Goodenough, J. B. J Am Ceram Soc 1998, 81, 2576. (84) Huang, P. N.; Petric, A. J Electrochem Soc 1996, 143, 1644. (85) Huang, K. Q.; Feng, M.; Goodenough, J. B.; Schmerling, M. J Electrochem Soc 1996, 143, 3630. (86) Yamaji, K.; Horita, T.; Ishikawa, M.; Sakai, N.; Yokokawa, H. Solid State Ionics 1999, 121, 217. (87) Minh, N. Q., Takahashi, T. Science and Technology of Ceramic Fuel Cells; Elsevier: Amsterdam, 1995. (88) Steele, B. C. H. Solid State Ionics 2000, 129, 95. (89) Goodenough, J. B.; Ruizdiaz, J. E.; Zhen, Y. S. Solid State Ionics 1990, 44, 21. (90) Goodenough, J. B.; Manthiram, A.; Paranthaman, P.; Zhen, Y. S. Solid State Ionics 1992, 52, 105. (91) Manthiram, A.; Kuo, J. F.; Goodenough, J. B. Solid State Ionics 1993, 62, 225. (92) Kakinuma, K.; Yamamura, H.; Haneda, H.; Atake, T. Solid State Ionics 2002, 154, 571. (93) Kakinuma, K.; Arisaka, T.; Yamamura, H.; Atake, T. Solid State Ionics 2004, 175, 139. (94) Lacorre, P.; Goutenoire, F.; Bohnke, O.; Retoux, R.; Laligant, Y. Nature 2000, 404, 856. (95) Goutenoire, F.; Isnard, O.; Retoux, R.; Lacorre, P. Chem Mater 2000, 12, 2575. (96) Goutenoire, F.; Isnard, O.; Suard, E.; Bohnke, O.; Laligant, Y.; Retoux, R.; Lacorre, P. J Mater Chem 2001, 11, 119. (97) Georges, S.; Goutenoire, F.; Altorfer, F.; Sheptyakov, D.; Fauth, F.; Suard, E.; Lacorre, P. Solid State Ionics 2003, 161, 231. (98) Georges, S.; Goutenoire, F.; Laligant, Y.; Lacorre, P. J Mater Chem 2003, 13, 2317. (99) Corbel, G.; Laligant, Y.; Goutenoire, F.; Suard, E.; Lacorre, P. Chem Mater 2005, 17, 4678. (100) Georges, S.; Goutenoire, F.; Bohnke, O.; Steil, M. C.; Skinner, S. J.; Wiemhofer, H. D.; Lacorre, P. J New Mat Electr Sys 2004, 7, 51. (101) Nakayama, S.; Aono, H.; Sadaoka, Y. Chem Lett 1995, 431. (102) Nakayama, S.; Kageyama, T.; Aono, H.; Sadaoka, Y. J Mater Chem 1995, 5, 1801. (103) Nakayama, S.; Sakamoto, M.; Higuchi, M.; Kodaira, K.; Sato, M.; Kakita, S.; Suzuki, T.; Itoh, K. J Eur Ceram Soc 1999, 19, 507. (104) Islam, M. S.; Tolchard, J. R.; Slater, P. R. Chem Commun 2003, 1486. (105) Bechade, E.; Masson, O.; Iwata, T.; Julien, I.; Fukuda, K.; Thomas, P.; Champion, E. Chem Mater 2009, 21, 2508. (106) Tolchard, J. R.; Islam, M. S.; Slater, P. R. J Mater Chem 2003, 13, 1956. (107) Jones, A.; Slater, P. R.; Islam, M. S. Chem Mater 2008, 20, 5055. (108) Shaula, A. L.; Kharton, V. V.; Marques, F. M. B. Solid State Ionics 2006, 177, 1725. (109) Fergus, J. W. Solid State Ionics 2006, 177, 1529. (110) Atkinson, A.; Barnett, S.; Gorte, R. J.; Irvine, J. T. S.; Mcevoy, A. J.; Mogensen, M.; Singhal, S. C.; Vohs, J. Nat Mater 2004, 3, 17. (111) Gross, M. D.; Vohs, J. M.; Gorte, R. J. J Mater Chem 2007, 17, 3071. (112) Goodenough, J. B.; Huang, Y. H. J Power Sources 2007, 173, 1.

References

182

(113) Sun, C. W.; Stimming, U. J Power Sources 2007, 171, 247. (114) Spacil, H. S.; Patent, U., Ed. 1970; Vol. 3,558,360. (115) Dees, D. W.; Claar, T. D.; Easler, T. E.; Fee, D. C.; Mrazek, F. C. J Electrochem Soc 1987, 134, 2141. (116) Zhu, W. Z.; Deevi, S. C. Mat Sci Eng a-Struct 2003, 362, 228. (117) McIntosh, S.; Gorte, R. J. Chem Rev 2004, 104, 4845. (118) Matsuzaki, Y.; Yasuda, I. Solid State Ionics 2000, 132, 261. (119) Murray, E. P.; Tsai, T.; Barnett, S. A. Nature 1999, 400, 649. (120) Park, S. D.; Vohs, J. M.; Gorte, R. J. Nature 2000, 404, 265. (121) Lu, C.; Worrell, W. L.; Gorte, R. J.; Vohs, J. M. J Electrochem Soc 2003, 150, A354. (122) Lee, S. I.; Vohs, J. M.; Gorte, R. J. J Electrochem Soc 2004, 151, A1319. (123) Xie, Z.; Zhu, W.; Zhu, B. C.; Xia, C. R. Electrochim Acta 2006, 51, 3052. (124) Saeki, M. J.; Uchida, H.; Watanabe, M. Catal Lett 1994, 26, 149. (125) Hibino, T.; Hashimoto, A.; Yano, M.; Suzuki, M.; Sano, M. Electrochim Acta 2003, 48, 2531. (126) Hibino, T.; Hashimoto, A.; Asano, K.; Yano, M.; Suzuki, M.; Sano, M. Electrochem Solid St 2002, 5, A242. (127) Tao, S. W.; Irvine, J. T. S. Chem Rec 2004, 4, 83. (128) Tao, S. W.; Irvine, J. T. S. Nat Mater 2003, 2, 320. (129) Ruiz-Morales, J. C.; Canales-Vazquez, J.; Pena-Martinez, J.; Marrero-Lopez, D.; Nunez, P. Electrochim Acta 2006, 52, 278. (130) Slater, P. R.; Irvine, J. T. S. Solid State Ionics 1999, 124, 61. (131) Slater, P. R.; Irvine, J. T. S. Solid State Ionics 1999, 120, 125. (132) Kramer, S.; Spears, M.; Tuller, H. L. Solid State Ionics 1994, 72, 59. (133) Porat, O.; Heremans, C.; Tuller, H. L. Solid State Ionics 1997, 94, 75. (134) Awana, V. P. S.; Menon, L.; Malik, S. K. Physica C 1996, 262, 266. (135) Roth, G.; Adelmann, P.; Heger, G.; Knitter, R.; Wolf, T. J Phys I 1991, 1, 721. (136) Huang, Q.; Sunshine, S. A.; Cava, R. J.; Santoro, A. J Solid State Chem 1993, 102, 534. (137) Isobe, M.; Matsui, Y.; Takayamamuromachi, E. Physica C 1994, 222, 310. (138) Pena-Martinez, J.; Tarancon, A.; Marrero-Lopez, D.; Ruiz-Morales, J. C.; Nunez, P. Fuel Cells 2008, 8, 351. (139) Li, R. K.; Greaves, C. Phys Rev B 2000, 62, 14149. (140) Kim, J. S.; Lee, J. Y.; Swinnea, J. S.; Steinfink, H.; Reiff, W. M.; Lightfoot, P.; Pei, S.; Jorgensen, J. D. J Solid State Chem 1991, 90, 331. (141) Wan, J.; Goodenough, J. B.; Zhu, J. H. Solid State Ionics 2007, 178, 281. (142) Kharton, V. V.; Yaremchenko, A. A.; Shaula, A. L.; Patrakeev, M. V.; Naumovich, E. N.; Loginovich, D. I.; Frade, J. R.; Marques, F. M. B. J Solid State Chem 2004, 177, 26. (143) Kim, G. T.; Jacobson, A. J. Mater Res Soc Symp P 2007, 972, 175. (144) Tonus, F.; Greaves, C.; El Shinawi, H.; Hansen, T.; Hernandez, O.; Battle, P. D.; Bahout, M. J Mater Chem 2011, 21, 7111. (145) Segal, D. J Mater Chem 1997, 7, 1297. (146) Moulson, A. J. a. H., J. M. Electroceramics: materials, properties, applications; Chapman & Hall: London, 1990. (147) Pechini, M. P. 1967; Vol. U. S. Pat. No. 3 330 697. (148) Pecharsky, V. K., Zavalij, P. Y.; Springer: NY, 2009. (149) Dinnebier, R. E., Billinge, S. J. L.; RSC Publishing: Cambridge, 2008. (150) Rietveld, H. M. Acta Crystallogr 1967, 22, 151. (151) Rietveld, H. M. J Appl Crystallogr 1969, 2, 65. (152) Coelho, A. A. J Appl Crystallogr 2003, 36, 86. (153) Thompson, S. P.; Parker, J. E.; Potter, J.; Hill, T. P.; Birt, A.; Cobb, T. M.; Yuan, F.; Tang, C. C. Rev Sci Instrum 2009, 80.

References

183

(154) Hastings, J. B.; Thomlinson, W.; Cox, D. E. J Appl Crystallogr 1984, 17, 85. (155) Hannon, A. C. Nucl Instrum Meth A 2005, 551, 88. (156) Smart, E. L., Moore, E. A. Solid State Chemistry; Taylor & Francis: Boca Raton, 2005. (157) Heaney, M. B.; CRC Press LLC: 1999. (158) Groover, M. P. Fundamentals of Modern Manufacturing: Materials, Processes, and Systems; John Wiley & Sons, 2010. (159) MacDonald, J. R. Impedance Spectroscopy: Emphasising solid materials and systems; John Wiley & Sons, 1987. (160) Ulgut, B. Electrochemical Impedance Spectroscopy and Its Applications to Battery Analysis, 2014. (161) Irvine, J. T. S., Siclair, D. C., and West, A. R. Adv Mater 1990, 2, 7. (162) Smith, R. A.; Second ed. ed.; Cambridge University Press: Cambridge, 1978. (163) Ginley, D., Hosono, H., Paine, D. C. Handbook of Transparent Conductors; Sprinfer, 2011. (164) MacKenzie, K. J. D., Smith, M. E. Multinuclear Solid-State NMR of Inorganic Materials, 2002. (165) Buannic, L.; Blanc, F.; Middlemiss, D. S.; Grey, C. P. J Am Chem Soc 2012, 134, 14483. (166) Chen, W. M.; Wu, X. S.; Geng, J. F.; Chen, J.; Chen, D. B.; Jin, X.; Jiang, S. S. J Supercond 1997, 10, 41. (167) Laitinen, H. A., Harris, W. E. Chemical Analysis; McGraw-Hill: New York, 1975. (168) Hayashi, H.; Saitou, T.; Maruyama, N.; Inaba, H.; Kawamura, K.; Mori, M. Solid State Ionics 2005, 176, 613. (169) Kohn, W.; Sham, L. J. Phys Rev 1965, 140, 1133. (170) Hohenberg, P.; Kohn, W. Phys Rev B 1964, 136, B864. (171) Hanke, F., Persson, M., Hofer, W.; Liverpool, T. U. o., Ed. 2010. (172) Thomas, L. H. P Camb Philos Soc 1927, 23, 542. (173) Fermi, E. Z Phys 1928, 48, 73. (174) Teller, E. Rev Mod Phys 1962, 34, 627. (175) Burke, K.; California, U. o., Ed. Irvine, 2007. (176) Wu, Z. G.; Cohen, R. E. Phys Rev B 2006, 73. (177) Sholl, D. S. a. S., J. A.; Wiley: Hoboken, N.J, 2009. (178) Kresse, G.; Joubert, D. Phys Rev B 1999, 59, 1758. (179) Blochl, P. E. Phys Rev B 1994, 50, 17953. (180) Vanderbilt, D. Phys Rev B 1990, 41, 7892. (181) Kresse, G.; Hafner, J. J Phys-Condens Mat 1994, 6, 8245. (182) Apirath, P. DFT: basics and practice, 2007. (183) Petukhov, A. G.; Mazin, I. I.; Chioncel, L.; Lichtenstein, A. I. Phys Rev B 2003, 67. (184) Carrillocabrera, W.; Wiemhofer, H. D.; Gopel, W. Solid State Ionics 1989, 32-3, 1172. (185) Macmanus, J. L.; Fray, D. J.; Evetts, J. E. Physica C 1992, 190, 511. (186) Vischjager, D. J.; Vanzomeren, A. A.; Schoonman, J.; Kontoulis, I.; Steele, B. C. H. Solid State Ionics 1990, 40-1, 810. (187) Nowotny, J.; Rekas, M. J Am Ceram Soc 1990, 73, 1054. (188) Nowotny, J.; Rekas, M.; Weppner, W. J Am Ceram Soc 1990, 73, 1040. (189) Slater, P. R.; Greaves, C. Physica C 1991, 180, 299. (190) Huang, Q.; Cava, R. J.; Santoro, A.; Krajewski, J. J.; Peck, W. F. Physica C 1992, 193, 196. (191) Den, T.; Kobayashi, T. Physica C 1992, 196, 141. (192) Babu, T. G. N.; Kilgour, J. D.; Slater, P. R.; Greaves, C. J Solid State Chem 1993, 103, 472. (193) Awana, V. P. S.; Malik, S. K.; Yelon, W. B.; Karppinen, A.; Yamauchi, H. Physica C- Superconductivity and Its Applications 2002, 378, 155.

References

184

(194) Awana, V. P. S.; Takayama-Muromachi, E.; Malik, S. K.; Yelon, W. B.; Karppinen, M.; Yamauchi, H.; Krishnamurthy, V. V. J Appl Phys 2003, 93, 8221. (195) Sunshine, S. A.; Siegrist, T.; Schneemeyer, L. F.; Marsh, P. Mater Res Soc Symp P 1989, 156, 113. (196) Johnson, D. Copyright 1990-2005 Schribner Asociates, Inc. (197) Kakinuma, K.; Arisaka, T.; Yamamura, H. J Ceram Soc Jpn 2004, 112, 342. (198) Li, S. Y.; Lu, Z.; Huang, X. Q.; Su, W. H. Solid State Ionics 2008, 178, 1853. (199) Kostogloudis, G. C.; Ftikos, C. Solid State Ionics 1999, 126, 143. (200) Hansen, K. K.; Hansen, K. V. Solid State Ionics 2007, 178, 1379. (201) Hammouche, A.; Siebert, E.; Hammou, A. Mater Res Bull 1989, 24, 367. (202) Tietz, F.; Stochniol, G.; Naoumidis, A. Euromat 97 - Proceedings of the 5th European Conference on Advanced Materials and Processes and Applications: Materials, Functionality & Design, Vol 2 1997, 271. (203) Tietz, F. Ionics 1999, 5, 129. (204) Kim, J. H.; Kim, Y. N.; Bi, Z. H.; Manthiram, A.; Paranthaman, M. P.; Huq, A. Electrochim Acta 2011, 56, 5740. (205) Vert, V. B.; Serra, J. M.; Jorda, J. L. Electrochem Commun 2010, 12, 278. (206) Steele, B. C. H.; Bae, J. M. Solid State Ionics 1998, 106, 255. (207) Islam, M. S.; Davies, R. A.; Fisher, C. A. J.; Chadwick, A. Solid State Ionics 2001, 145, 333. (208) Huang, K., Goodenough, J. B. Solid Oxide Fuel Cell Technology; Woodhead Publishing: Cambridge, U. K., 2010. (209) Maignan, A.; Caignaert, V.; Raveau, B.; Khomskii, D.; Sawatzky, G. Phys Rev Lett 2004, 93. (210) Frontera, C.; Garcia-Munoz, J. L.; Llobet, A.; Aranda, M. A. G. Phys Rev B 2002, 65. (211) Taskin, A. A.; Lavrov, A. N.; Ando, Y. Phys Rev B 2005, 71. (212) Akahoshi, D.; Ueda, Y. J Solid State Chem 2001, 156, 355. (213) Fauth, F.; Suard, E.; Caignaert, V.; Mirebeau, I. Phys Rev B 2002, 66. (214) Martin, C.; Maignan, A.; Pelloquin, D.; Nguyen, N.; Raveau, B. Appl Phys Lett 1997, 71, 1421. (215) Moritomo, Y.; Takeo, M.; Liu, X. J.; Akimoto, T.; Nakamura, A. Phys Rev B 1998, 58, 13334. (216) Respaud, M.; Frontera, C.; Garcia-Munoz, J. L.; Aranda, M. A. G.; Raquet, B.; Broto, J. M.; Rakoto, H.; Goiran, M.; Llobet, A.; Rodriguez-Carvajal, J. Phys Rev B 2001, 64, art. no. (217) Zhang, Q. F.; Zhang, W. Y. Phys Rev B 2003, 67. (218) Taskin, A. A.; Lavrov, A. N.; Ando, Y. Appl Phys Lett 2005, 86. (219) Taskin, A. A.; Lavrov, A. N.; Ando, Y. Prog Solid State Ch 2007, 35, 481. (220) Chavez, E.; Mueller, M.; Mogni, L.; Caneiro, A. Xix Latin American Symposium on Solid State Physics (Slafes) 2009, 167. (221) Kim, J. H.; Mogni, L.; Prado, F.; Caneiro, A.; Alonso, J. A.; Manthiram, A. J Electrochem Soc 2009, 156, B1376. (222) Manthiram, A.; Prado, F.; Armstrong, T. Solid State Ionics 2002, 152, 647. (223) Kim, J. H.; Prado, F.; Manthiram, A. J Electrochem Soc 2008, 155, B1023. (224) Kresse, G.; Furthmuller, J. Phys Rev B 1996, 54, 11169. (225) Perdew, J. P.; Burke, K.; Ernzerhof, M. Phys Rev Lett 1996, 77, 3865. (226) Vosko, S. H.; Wilk, L.; Nusair, M. Can J Phys 1980, 58, 1200. (227) Perdew, J. P.; Zunger, A. Phys Rev B 1981, 23, 5048. (228) Perdew, J. P.; Wang, Y. Phys Rev B 1992, 45, 13244. (229) Wang, L.; Maxisch, T.; Ceder, G. Phys Rev B 2006, 73. (230) Lee, Y. L.; Kleis, J.; Rossmeisl, J.; Morgan, D. Phys Rev B 2009, 80.

References

185

(231) Sanchez-Andujar, M.; Senaris-Rodriguez, M. A. Z Anorg Allg Chem 2007, 633, 1890. (232) PANalytical, B. V.; edn., E., Ed. 2006. (233) International Center for Diffraction Data; edn., D., Ed. 2007. (234) Hull, S.; Norberg, S. T.; Ahmed, I.; Eriksson, S. G.; Marrocchelli, D.; Madden, P. A. J Solid State Chem 2009, 182, 2815. (235) Mevs, H.; Mullerbuschbaum, H. Z Anorg Allg Chem 1989, 573, 128. (236) Burley, J. C.; Mitchell, J. F.; Short, S.; Miller, D.; Tang, Y. J Solid State Chem 2003, 170, 339. (237) Anderson, P. S.; Kirk, C. A.; Knudsen, J.; Reaney, I. M.; West, A. R. Solid State Sci 2005, 7, 1149. (238) Pardo, V.; Baldomir, D. Phys Rev B 2006, 73. (239) Wang, C. S.; Pickett, W. E. Phys Rev Lett 1983, 51, 597. (240) Chan, M. K. Y.; Ceder, G. Phys Rev Lett 2010, 105. (241) Matskevich, N. I.; Wolf, T.; Matskevich, M. Y.; Chupakhina, T. I. Eur J Inorg Chem 2009, 1477. (242) Matsumoto, H.; Kawasaki, Y.; Ito, N.; Enoki, M.; Ishihara, T. Electrochem Solid St 2007, 10, B77. (243) Kreuer, K. D.; Paddison, S. J.; Spohr, E.; Schuster, M. Chem Rev 2004, 104, 4637. (244) Matsui, T. Thermochim Acta 1995, 253, 155. (245) Shannon, R. D. Acta Crystallogr A 1976, 32, 751. (246) Gillie, L. J.; Hadermann, J.; Hervieu, M.; Maignan, A.; Martin, C. Chem Mater 2008, 20, 6231. (247) Weiss, R.; Faivre, R. Cr Hebd Acad Sci 1959, 248, 106. (248) Kennedy, B. J. Aust J Chem 1997, 50, 917. (249) Fu, W. T.; Visser, D.; IJdo, D. J. W. J Solid State Chem 2002, 169, 208. (250) Fu, W. T.; Visser, D.; Knight, K. S.; IJdo, D. J. W. J Solid State Chem 2004, 177, 4081. (251) Scholder, R.; Rade, D.; Schwarz, H. Z Anorg Allg Chem 1968, 362, 149. (252) Detemple, E.; Ramasse, Q. M.; Sigle, W.; Cristiani, G.; Habermeier, H. U.; Keimer, B.; van Aken, P. A. J Appl Phys 2012, 112. (253) Hungria, T.; MacLaren, I.; Fuess, H.; Galy, J.; Castro, A. Mater Lett 2008, 62, 3095. (254) Suzuki, T.; Nishi, Y.; Fujimoto, M. J Am Ceram Soc 2000, 83, 3185. (255) Rodgers, J. A.; Battle, P. D.; Dupre, N.; Grey, C. P.; Sloan, J. Chem Mater 2004, 16, 4257. (256) Arachi, Y.; Sakai, H.; Yamamoto, O.; Takeda, Y.; Imanishai, N. Solid State Ionics 1999, 121, 133. (257) Hui, S. Q.; Roller, J.; Yick, S.; Zhang, X.; Deces-Petit, C.; Xie, Y. S.; Maric, R.; Ghosh, D. J Power Sources 2007, 172, 493. (258) Horita, T.; Sakai, N.; Yokokawa, H.; Dokiya, M.; Kawada, T.; Van Herle, J.; Sasaki, K. J Electroceram 1997, 1, 155. (259) Kamimura, S.; Yamada, H.; Xu, C. N. Appl Phys Lett 2012, 101. (260) Kamimura, S.; Yamada, H.; Xu, C. N. Appl Phys Lett 2013, 102. (261) Van den Eeckhout, K.; Poelman, D.; Smet, P. F. Materials 2013, 6, 2789. (262) Mizoguchi, H.; Eng, H. W.; Woodward, P. M. Inorg Chem 2004, 43, 1667. (263) Liu, Q. Z.; Li, B.; Liu, J. J.; Li, H.; Liu, Z. L.; Dai, K.; Zhu, G. P.; Zhang, P.; Chen, F.; Dai, J. M. Epl-Europhys Lett 2012, 98. (264) Zainullina, V. M. Physica B-Condensed Matter 2007, 391, 280. (265) Shanthi, E.; Dutta, V.; Banerjee, A.; Chopra, K. L. J Appl Phys 1980, 51, 6243. (266) Lee, S. W.; Kim, Y. W.; Chen, H. D. Appl Phys Lett 2001, 78, 350. (267) Brankovic, G.; Brankovic, Z.; Santos, L. P. S.; Longo, E.; Davolos, M. R.; Varela, J. A. Mater Sci Forum 2003, 416-4, 651. (268) Liu, Q. Z.; Wang, H. F.; Chen, F.; Wua, W. B. J Appl Phys 2008, 103.

References

186

(269) Shigeno, E.; Shimizu, K.; Seki, S.; Ogawa, M.; Shida, A.; Ide, A.; Sawada, Y. Thin Solid Films 2002, 411, 56.

APPENDIX A: EDX data of Sr2Sn0.96Ta0.04O4

Table A: EDX data of Sr2Sn0.96Ta0.04O4 collected on 12 different crystal areas. The data

show the percentage of each individual element and normalised (norm.) amount of Sr, Sn and

Ta in respect to nominal values for Sr2Sn0.96Ta0.04O4. The obtained average values are

compared to those expected.

Measurement

number

Sr (%) Sn (%) Ta (%) O (%) Sr

(norm.)

Sn

(norm.)

Ta

(norm.)

1 28.6 14.6 0.6 56.3 1.9589 1 0.0410

2 27.1 14.4 0.8 57.6 1.9219 1.0212 0.0567

3 28.5 14.5 0.3 56.7 1.9745 1.0046 0.0207

4 28 14.7 0.8 56.5 1.9310 1.0137 0.0551

5 27.7 14.9 0.5 57 1.9280 1.0371 0.0348

6 28.5 14.8 0.2 56.5 1.9655 1.0206 0.0137

7 27.5 14.9 0 57.6 1.9457 1.0542 0

8 26.2 15.6 0.2 58 1.8714 1.1142 0.0142

9 27.2 15.1 0.4 57.3 1.9110 1.0608 0.0281

10 27.3 14.6 0.9 57.2 1.9135 1.0233 0.0630

11 27.6 14.1 0.8 57.4 1.9482 0.9952 0.0564

12 27.6 14.4 0.6 57.3 1.9436 1.0140 0.0422

Expected 28.5714 13.7142 0.57142 57.1428 2 0.96 0.04

Average 27.65 14.7166 0.50833 57.1166 1.9344 1.0299 0.0355

188

APPENDIX B: Lattice parameters an cell volume of Sr2Sn1−xNbxO4

Table B: Lattice parameters and cell volume of the Sr2Sn1−xNbxO4 phases obtained after

Rietveld refinement in Topas (KCl used as an internal standard) and compared with

previously reported data on Sr2SnO4 material.250

Nb doping level a (Å) b (Å) c (Å) V (Å3)

0 – lit.250

5.72898(5) 5.73524(5) 12.58110(6) 413.378

0 - as made 5.72962(8) 5.73509(8) 12.5862(1) 413.58(1)

0.01 5.7283(1) 5.7352(1) 12.5874(2) 413.53(1)

0.02 5.72808(9) 5.7353(1) 12.5874(2) 413.53(1)

0.025 5.7255(3) 5.7334(3) 12.5925(3) 413.38(3)

0.03 5.7363(3) 5.7411(3) 12.6509(3) 416.62(4)

0.05 5.7352(3) 5.7418(2) 12.6493(3) 416.54(3)

0.075 5.7336(2) 5.7390(3) 12.6493(3) 416.23(3)

0.1 5.7340(2) 5.7406(2) 12.6449(4) 416.23(3)

189

APPENDIX C: Lattice parameters and cell volume of Sr2Sn1−xTaxO4

Table C: Lattice parameters and cell volume of the Sr2Sn1−xTaxO4 phases obtained after

Rietveld refinement in Topas (KCl used as an internal standard) and compared with

previously reported data on Sr2SnO4 material.250

Ta doping level a (Å) b (Å) c (Å) V (Å3)

0 – lit.250

5.72898(5) 5.73524(5) 12.58110(6) 413.378

0 - as made 5.72962(8) 5.72962(8) 12.5862(1) 413.58(1)

0.01 5.7279(1) 5.7353(1) 12.5898(2) 413.60(1)

0.02 5.7277(1) 5.7356(1) 12.5912(3) 413.65(1)

0.025 5.7264(2) 5.7356(2) 12.5910(3) 413.54(2)

0.03 5.7343(1) 5.7401(1) 12.6555(2) 416.56(1)

0.04 5.7273(2) 5.7389(3) 12.6722(4) 416.51(3)

0.05 5.7346(2) 5.7399(2) 12.6532(3) 416.50(2)

0.075 5.7339(3) 5.7390(3) 12.6507(3) 416.29(3)

0.1 5.7346(3) 5.7401(4) 12.6459(4) 416.27(4)

190

APPENDIX D: Joint I11 and HRPD Rietveld refinement of

Sr2Sn0.97Nb0.03O4

Table D: Refined parameters of the Sr2Sn0.97Nb0.03O4 sample from the joint Rietveld

refinement of the I11 and HRPD data.

Atom Site x y z Beq (Å2)

Sn1 4a 0 0 0 0.56(1)

Nb1 4a 0 0 0 0.56(1)

Sr1 8e 0.0007(1) −0.0039(5) 0.35184(2) 0.27(1)

O1 4c 0.250 0.250 0.004(1) 0.75(8)

O2 4d 0.750 0.250 −0.0158(5) 0.71(8)

O3 8e −0.0127(2) 0.026(1) 0.1638(1) 0.35(4)

Figure D: Structure of Sr2Sn0.97Nb0.03O4 obtained from joint Rietveld refinement of I11 and

HRPD data: a) view along c axis; b) view along a axis. Sn atoms are grey, Sr is green and O

atoms are red.

191

APPENDIX E: Joint I11 and HRPD Rietveld refinement of

Sr2Sn0.96Ta0.04O4

Table E: Refined parameters of the Sr2Sn0.96Ta0.04O4 material obtained from the joint

Rietveld refinement of the I11 and HRPD data.

Atom Site x y z B (Å2)

Sn1 4a 0 0 0 0.518(7)

Ta1 4a 0 0 0 0.518(7)

Sr1 8e 0.0010(6) −0.0031(4) 0.35187(2) 0.285(9)

O1 4c 0.250 0.250 0.005(1) 0.52(9)

O2 4d 0.750 0.250 −0.0153(5) 0.15(8)

O3 8e −0.011(1) 0.0335(6) 0.1629(1) 0.36(4)

Figure E: Structure of Sr2Sn0.96Ta0.04O4 obtained from joint Rietveld refinement of I11 and

HRPD data: a) view along a axis; a) view along c axis. Sn atoms are grey, Sr is green and O

atoms are red.

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