Possible electrolytes for SOFC: synthesis and characterization of Bi3–xMexTaO7–δ and others fluorite-related phases

UNIVERSITÀ DEGLI STUDI DI CATANIA

FACOLTÀ DI SCIENZE MATEMATICHE, FISICHE E NATURALI

CORSO DI LAUREA IN CHIMICA INDUSTRIALE DIPARTIMENTO DI SCIENZE CHIMICHE

⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯

NICO TORINO

Possible electrolytes for SOFC: synthesis and characterization of Bi3–xMexTaO7–δ

and others fluorite-related phases

⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯ M. Sc. EXPERIMENTAL THESIS

⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯

Supervisors Prof. Salvatore Scirè

Dr. Aurélie Rolle (ENSCL)

⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯ ACADEMIC YEAR 2010 - 2011

⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯

Abstract

How to supply the world's energy needs in a safe and clean way has become

one of the most pressing task to achieve and the establishment of an electron

economy seems to be a way to succeed in it. In such economy solid oxide fuel

cells could play an important role in restoring electricity and to make it as

much efficient as possible the basic research looks for more performing

materials: this work focuses on electrolyte materials.

The weberite phase La3TaO7 shows poor oxide ions conducting properties

that can be increased by creating oxygen vacancies within the structure by

means of the substitution of La3+ ions for Sr2+ ones. Moreover it could be

stable under reducing and oxidizing conditions at once (as required for

electrolyte materials suitable for SOFC) given the relatively high standard

reduction potentials of the cations; however such a material has lower

conducting performances than those of yttrium-doped zirconia, the main

reference material for SOFC electrolytes.

The aim of this research is to investigate the existence of the solid solution

Bi3–xSrxTaO7 and the effect of the Sr2+ ion on the oxide ions conductivity of

Bi3TaO7, a phase with a structure similar to that of La3TaO7. The existence of

the solid solution Bi3–xLaxTaO7 was also investigated in order to find a way to

solve the reducibility issues of bismuth-based materials: by dispersing the

bismuth compound in the La3TaO7 matrix, it could be created a material that

can efficiently carry oxygen ions by exploiting the high polarizable 6s2 lone

pair of Bi3+ ion that is the key to the high conducting properties of bismuth-

based materials themselves, first of all δ-Bi2O3.

While the existence of the solid solution Bi3–xSrxTaO7 was confirmed, the

existence of the Bi3–xLaxTaO7 one remains questionable.

Once the Sr2+-doped phase was synthesized, its oxide ions conducting

properties were investigated and compared to those of Bi3TaO7, showing that

the doped compound is a worst conductor.

The existence of some others solid solutions with Ca2+ and Sr2+ ions was

explored for compounds with fluorite-related structures, namely Y3TaO7 and

Yb3TaO7, and preliminary investigations were carried out on La3TaO9 and the

phases Pr3TaO7 and Pr3NbO7.

Keywords: SOFC, electrolyte, oxide ion conduction, bismuth, solid state

chemistry

Contents

Preface 1

Need to look at energy production 1

Aim of the research work 3

Outline 5

Acknowledgments 6

Introduction 7

Oxide ions conducting materials 7

The fluorite structure and YSZ 7

Doped ceria 9

δ-Bi2O3 and BIMEVOX 9

Lanthanum gallate 11

Phases derived from Ba2In2O5 11

Apatites 12

Technological applications of oxide ions conductors 14

Solid oxide fuel cells 15

a. General overview on fuel cells 15

b. How they work 16

c. About the electrolyte and the others SOFC components 18

Oxygen pump 19

Oxygen sensor 20

Theory 21

Experimental 26

Sample preparation 26

XRD characterization 28

Data refining 29

Thermal analyses 31

Raman scattering spectroscopy 31

Scanning Electron Microscopy 32

Electron Probe Microanalyses 33

Electrical mesurements 34

Results and discussion 37

LaTa3O9 37

Pr3TaO7 and Pr3NbO7 40

Y3TaO7 and related phases 44

Y3TaO7 44

Y2.9Ca0.1TaO6.95 47

Y2.9Sr0.1TaO6.95 48

Yb3TaO7 and related phases 50

Yb3TaO7 50

Yb2.9Sr0.1TaO6.95 53

Yb3–xCaxTaO7–x/2 54

Bi3TaO7 and related phases 57

Bi3TaO7 58

Bi3–xLaxTaO7 66

Bi3–xSrxTaO7–x/2 70

How the strontium ion influences the conducting properties of Bi3TaO7 73

Conclusions 83

References 86

Preface

Need to look at energy production

The last decades have become the stage on which mankind is facing one of

the most challenging trial of all times: how to supply the energy needs of the

growing world population in a safe and clean way.

Such an issue has arisen in recent times most likely because on the first times

after the industrial revolution it was considered quite an oddity showing

concernments about the finite nature of resources on the planet but that kind

of attitude was drove way too far till it has become a threat to that well-being

itself, that is to say, achieved thanks to that kind of exploitation.

Lying mainly on fossil hydrocarbons and coal and minerals for nuclear fuel,

energy supply is a geopolitical and economical issue since those resources are

not widespread nor less important are the problematics connected to the

environmental impact due to the use of such commodities.

While being aware of the finiteness of resources has been a relatively easy

task, the solution of the related problem is still to be found but it's

unquestionable that it musts rely on sustainability, that is on conserving the

ecological balance by avoiding the depletion of those natural resources: it

means that the resources have to be renewable.

Preface

2

Among the various options, one of the most promising concerns the

establishment of an electron economy by means of which energy would be

harvested directly from physical processes – e.g. photovoltaic devices –

replacing the chemical carriers of today's economy. In fact, with exception of

biomass, all renewable energy is of physical nature: heat (solar, geotermal),

solar radiation (photovoltaic) and mechanical energy (wind, waves etc.) [1].

One of the technological obstacles is the storage of the energy obtained in

such ways and a practicable way could be the use of hydrogen as carrier,

made by electrolysis of water. To restore, then, the electron flow the solid

oxide fuel cells (SOFC) represent an efficient way to use that hydrogen,

attaining at up to 60% of efficiency on electricity production and up to 80%

by cogeneration of electricity and heat [2].

A SOFC is an electrochemical device that works at high temperatures,

between 800°C and 1000°C, and allows to use the electrons coming from the

oxidation of hydrogen (or hydrocarbons like methane also) as a current.

The high working temperatures result in mechanical and chemical

compatibility issues of the different components of such devices

(anode//electrolyte//cathode) thus an important task is achieving lower

working temperatures: the basic scientific research tries to answer by

exploring the possibilities to obtain more performing materials.

This is a pressing task even to lower prices, that range from ∼3500 to ∼15000

€/kWe, to make the technology affordable for a larger market [3].

The present work focuses on the electrolyte materials for SOFCs.

Preface

3

Aim of the research work

An electrolyte suitable for SOFCs must be stable under reducing and

oxidizing conditions at once, it must be an electronic insulator and it must

have an ordered framework of defects (interstitial oxide ions and/or

vacancies) through which allowing the drifting of oxygen ions.

A former work by Dr. Nicolas Preux [4] studied the compound La3TaO7,

having a weberite-type structure, constituted by chains of corner-sharing TaO6

octahedra, intercalated by La–O chains and isolated lanthanum atoms.

Fig. 1 compares the structure of La3TaO7 and of the weberite Na2MgAlF7; the

latter could be described as a defect-fluorite structure by considering only the

cationic array.

That compound was chosen because it showed the main features suitable for

developing a new electrolyte material for SOFC applications: La3+ and Ta5+

Fig. 1 Strucures of La3TaO7 and of the weberite Na2MgAlF7

Preface

4

are not so easily reduced (standard reduction potentials in water are –2.38 V

and –0.81 V respectively [5]) and they are no more oxidizable; the structure

of La3TaO7 is defective.

Dr. Preux's work showed the existence of the solid solution La3–xSrxTaO7–x/2

for 0 < x ≤ 0.15 and compared the ionic conductivity of the phase with

x = 0.1 to that of the non-doped La3TaO7 phase. The ionic conductivity was

greatly enhanced by the substitution of lanthanum ion for the strontium one:

moving from that result, the nucleus of the present research concerns the

possibility of doping the phase Bi3TaO7, having a similar structure to that of

La3TaO7 [6]. Because of the high polarizable 6s2 lone pair of Bi3+ ion,

Bi3TaO7 presents better conducting properties than the Sr-doped La3TaO7

phase, so the purpose was to synthesize a Sr-doped Bi3TaO7 phase and to

explore its conducting properties and compare them with those of the Bi3TaO7

phase.

Moreover the substitution of bismuth for lanthanum was tried in order to

confirm or refute the existence of the solid solution Bi3–xLaxTaO7. Such

investigation was useful because the reducibility of bismuth in its compounds

poses limitation to their application as electrolyte for SOFC: dispersing the

bismuth based compound in a lanthanum based matrix could be a way to

solve such an issue.

The Sr2+-doped Bi3TaO7 phase was characterized by electrochemical

impedance spectroscopy and also microstructurally by Scanning Electron

Microscopy.

Preface

5

- Yb3TaO7 to form Yb3–xSrxTaO7–x/2 and Yb3–xCaxTaO7–x/2 phases;

- Y3TaO7 to form Y3–xSrxTaO7–x/2 and Y3–xCaxTaO7–x/2 phases;

- LaTa3O9 by adding Zr4+ ions.

The investigation looked even at the phases Pr3TaO7 and Pr3NbO7. They were

chosen because of the nature of the praseodymium atom that can be oxidized

from PrIII to PrIV: electronic conduction might so be added in a solid solution

with other tantalates, thus introducing domains of mixed electronic/ionic

conduction in a pure ionic conductor.

Outline

An overview of the main oxide ions conducting materials will be given in the

introduction along with the principle of SOFC functioning and some other

applications. Then, a theoretical chapter will briefly survey the ionic

conduction. The core of this research work will follow with the experimental

section and the presentation of the results.

Acknowledgments

6

Acknowledgments

As part of the ERASMUS programme, the present research work was carried

out with the “Oxide Materials for Energy” team from the Solid State

Chemistry group at the UCCS (Unité de Catalyse et de Chimie du Solide) of

the ENSCL (École Nationale Supérieure de Chimie de Lille), in collaboration

with Dr. Aurélie Rolle and under the supervision of Pr. Rose-Noëlle Vannier

and with the precious help of Dr. Nicolas Preux, assistant engineers Laurence

Burylo and Nora Djelal for the analysis work (X-ray diffractions, thermal and

SEM analyses), CNRS engineer Edouard Capoen during the set-up for

electrical measurements and laboratory technician Maxence Vanderwalle

providing and explaining everything was necessary for a happy lablife.

Special thanks go to Pr. Annick Rubbens who provided Raman scattering

spectra and many useful advices.

Introduction

7

Introduction

Oxide ions conducting materials

The fluorite structure and YSZ

Fluorite, CaF2, exhibits a fcc lattice of Ca2+ ions where F– ions occupy all

tetrahedral sites (fig. 2):

This kind of structure is also shown by the main compound used as electrolyte

in SOFC technology, the yttria stabilized zirconia, YSZ, of composition

(ZrO2)1–x(Y2O3)x with 0.08 ≤ x ≤ 0.1.

ZrO2 has several polymorphs and it is a poor ionic conductor. Its cubic

Fig. 2 The fluorite structure

Introduction

8

fluorite phase appears only beyond 2370°C: Zr4+ ions form a fcc lattice in

which oxide ions occupy the tetrahedral interstices. As a means to stabilize

the fluorite phase at room temperature, ZrO2 can be added of oxides of

divalent or trivalent cations such as CaO, MgO, Y2O3, Sc2O3, Yb2O3, Sm2O3

etc. Those dopant ions exactly replace part of Zr4+ ions and in turn they

introduce oxygen ion vacancies for maintaining charge neutrality: through

those vacancies, oxide ions can migrate and allow the material to behave as an

oxide ions conductor (fig. 3).

It is well understood that for zirconia, the conductivity is maximum when the

dopant is just adequate to fully stabilize the cubic fluorite phase. The best

conductivity results are obtained by adding Y2O3 for an 8% molar substitution

of zirconium ions (σ = 0.02 S/cm at 800°C and 0.1 S/cm at 1000°C).

Zr4+ and Y3+ ions form an fcc lattice and most of the tetrahedral sites of Zr4+

ions are occupied by O2– ions. At high temperatures, oxide anions tend to

vibrate in the region of their regular tetrahedral lattice sites and a few

(probably equal to the number of vacancies created) migrate from one regular

lattice site to the other while octahedral sites have no role to play in the

Fig. 3 Scheme of ionic migration in YSZ

Introduction

9

oxygen diffusion [7, 8].

High temperatures, besides, during prolonged operations may produce cation

diffusion and, as a result, the segregation of vacancies into vacancy-rich

clusters that trap the mobile vacancies. This phenomenon causes a

degradation of the performance of the electrolyte with time, so-called aging

[9].

Doped ceria

CeO2 has a stable fluorite structure from room temperature up to its melting

point (2400°C) and shows poor ionic conductivity, due mainly to

nonstoichiometry, that can be increased by means of doping: the substitution

of Ce4+ with cations of lower valence, such as Ca2+ or Y3+, creates oxygen

vacancies [8]. The best results are obtained by a 10% molar substitution in

Gd2O3: Ce0.9Gd0.1O1.95 has a ionic conductivity of 0.01 S/cm at 500°C.

Unfortunately problems in using ceria as SOFC electrolyte arise from the

partial reduction of Ce4+ to Ce3+ under the reducing conditions of the anode

[10].

δ-Bi2O3 and BIMEVOX

Bi2O3 shows a monoclinic to cubic phase transition (α → δ) that occurs at

approximately 730°C. The δ-phase is stable until the melting point of Bi2O3 at

approximately 825°C and it shows an incredibly high ionic conductivity (ca.

1 S/cm at 650°C, transition temperature to the metastable β-phase when

cooling down from the high temperature δ-phase) only if pure: doping Bi2O3

by the substitution of bismuth for rare-earth dopants (such as Y, Dy or Er) or

Introduction

10

high valency cations (Nb5+, Ta5+) stabilizes the fluorite phase at room

temperature but results in lowering its conducting properties. Moreover

bismuth based compounds are unstable at high temperature under reducing

conditions and undergo volatilization of bismuth oxide [9, 11].

BIMEVOX are solid solutions exhibiting high ionic conductivity at moderate

temperature, 400°-600°C . They are based on γ-Bi4V2O11, stabilized by partial

substitution of vanadium with transition metal cations such as Co2+, Ni2+ or

Cu2+, and possess a crystal lattice belonging to the Aurivillius series that

consists of alternating (Bi2O2)2+ and perovskite-like (VO3.5)2– layers (fig. 4)

with oxygen vacancies allowing ion migration only in the latter ones. Solid

solutions of the type Bi2V1–xMexO5.5–δ (Me = Cu, Ni and 0.07 ≤ x ≤ 0.12)

show the highest conductivities but still suffer the disadvantages of the

bismuth based compounds [10, 12].

Fig. 4 Bi4V2O11 structure

Introduction

11

Lanthanum gallate

Perovskite-type phases derived from LaGaO3 (fig. 5) show an ionic

conductivity higher than YSZ in the temperature range 500°-800°C that can

be even enhanced by substituting lanthanum with alkaline earth elements

and/or incorporating divalent metal cations, such as Mg2+, into the gallium

sublattice in order to increase oxygen vacancy concentration. The best results

are obtained for the La1–xSrxGa1–yMgyO3–δ series with 0.10 ≤ x ≤ 0.20 and

0.15 ≤ y ≤ 0.20. Disadvantages of this kind of materials include possible

reduction and volatilization of gallium oxide, reactivity with perovskite

electrodes under oxidizing conditions and with metal anodes in reducing

conditions and the relatively high cost of gallium [10].

Phases derived from Ba2In2O5

Other materials with substantially high conductivity can be derived by partial

substitution of brownmillerite-type Ba2In2O5. The structure of brownmillerite,

A2B2O5, consists of alternating perovskite layers of cornersharing BO6

octahedra and layers of BO4 tetrahedra (fig. 6) and can be considered as an

Fig. 5 Lanthanum gallate structure

Introduction

12

oxygen-deficient perovskite where the oxygen vacancies are ordered along

(010) planes forming one dimensional diffusion pathways for oxygen ion

migration. Ba2In2O5 has mixed conductivity with dominant oxygen ionic

transport in dry air (the ionic transference number at 500°C is ~ 0.93) but

becomes a pure ionic conductor above 870°-950°C when it undergoes a

transition to a disordered perovskite phase, stabilized at lower temperatures

by substituting indium with other cations such as Zn2+, Ce4+, Sn4+ or Hf4+.

Disadvantages are represented by the instability shown in humid atmospheres,

the high reactivity with CO2 and the easy reducibility [10].

Apatites

The apatite-type phases A10–x(SiO4)6O2±δ, where A corresponds to rare earth

and alkaline earth cations, possesses a significant level of oxygen ion

conductivity. A-site cations are located in the cavities created by SiO4

tetrahedra with four distinct oxygen positions and additional oxygen sites

forming channel through the structure (fig. 7): the ionic conduction is made

by means of those isolated oxygen atoms.

Fig. 6 Ba2In2O5 structure

Introduction

13

The oxygen ionic transport in Ln10Si6O27 (with Ln = La, Pr, Nd, Sm, Gd, Dy)

increases with increasing radius of Ln3+ cations, with maximum conductivity

for the La-containing phase. Due to relatively poor sintering, different

processing techniques and substantial anisotropy of ionic transport in the

apatite lattice, the conductivity values reported in the literature vary in a very

broad range. The highest ionic transport is observed when apatite contains

more than 26 oxygen ions per unit formula, suggesting a significant role of

the interstitial migration mechanism. To increase ionic conductivity many

other ions were tried as substituent of La3+: best results follow the scheme

Ba ≥ Sr ≥ Ca ≥ Y [10, 13, 14].

In the next page figure 8 shows a chart in which the oxygen ionic

conductivities for the mentioned materials are plotted versus the temperature.

Fig. 7 Apatite structure

Introduction

14

Technological applications of oxide ions conductors

Oxide ions conducting electrolytes are used mainly in electrochemical cells

where the chemical reactants at the two electrodes are either gaseous or

liquid. They are used in the form of a thin sheet of large surface area having

porous electrodes on opposite surfaces for two classes of applications: power

applications (solid oxide fuel cells and oxygen pump) and open-circuit

applications (oxygen sensor).

Fig. 8 Oxygen ionic conductivities of the mentioned solid electrolyte materials [10, 15, 16, 17, 18]

Introduction

15

Solid oxide fuel cells

a. General overview on fuel cells

A fuel cell is an electrochemical cell.

There are many types of fuel cells, all of them constituted by a stack of three

elements: an anode, an electrolyte and a cathode (fig. 9).

At the anode an oxidation reaction provides the electron flow and at the

cathode a reduction one uses the same electrons that pass through the circuit

to occur; the electrolyte is responsible for the ionic current that must balance

the electronic current in the external circuit.

Being the principle of functioning the same for all types of fuel cell devices,

they differ mainly for the working temperatures range, the electrolyte and the

nature of the ionic species involved. Table 1 summarizes the different kinds

of fuel cells: Alkaline Fuel Cell (AFC), Proton Exchange Membrane Fuel Cell

(PEMFC) and Phosphoric Acid Fuel Cell (PAFC) belong to the low

temperature devices whereas Molten Carbonate Fuel Cell (MCFC) and Solid

Oxide Fuel Cell (SOFC) belong to the high temperature ones.

Fig. 9 Schematic representation of a fuel cell

Introduction

16

low temperature fuel cells high temperature fuel cells

type of fuel cell AFC PEMFC PAFC MCFC SOFC

electrolyte KOH solution water H3PO4 Li2CO3/K2CO3 solid oxide

working temperature 25° - 90°C 70° - 90°C 200°C 600° - 650°C 800° - 1000°C

charge carrier OH– H+ H+ CO3

2– O2–

application domains

spacial transportation

transportation stationary stationary stationary transportation

stationary

Today's research is oriented mostly towards PEMFC technology, where water

through a proton exchange polymeric membrane (a sulfonated poly-

tetrafluoroethylene based copolymer) can deliver hydrogen ions from the

anode to the cathode [7], and towards SOFC where a ceramic oxide (mainly

yttria-doped zirconia) conducts oxyde ions from the cathode to the anode.

b. How they work

A schematic drawing of a SOFC is shown in fig. 10.

While air is supplied at the cathode, the oxide ions resulting from the

reduction of oxygen molecules diffuse through the conducting ceramic

membrane to react with hydrogen, at the anode, forming water as byproduct

along with heat since the reaction is exothermic:

H2 + 1/2 O2 → H2O (g) with ∆H < 0

Tab. 1 The different kinds of fuel cells [19, 20]

Introduction

17

The electrons provided by the oxidation reaction occurring at the anode can so

be used to power electrical devices.

The Nernst potential for the reaction, for a fuel cell operating at atmospheric

pressure, is given by:

The migration of oxide ions through the electrolyte membrane is a thermally

activated process thus high temperatures are needed (800° - 1000°C), limiting

indeed the applications only to big stationary and transportation systems.

Such high temperatures even imply aging problems to overcome which many

efforts are addressed to the development of technologies working at lower

temperatures (500° - 600°C), the so-called Intermediate Temperature-SOFC,

namely by researching materials that can conduct oxide ions at milder

conditions without significant dropping in conductivity.

!

Eth = EOx /R.ed0 +

RT2F

lnPH2

PO2

1/2

PH2O

Fig. 10 Functioning scheme of a solid oxide fuel cell

Introduction

18

There this work looks, focusing on phases with a defect-fluorite structure and

trying to enhance their conduction properties by introducing oxygen

vacancies within the structure.

c. About the electrolyte and the others SOFC components

The electrolyte in a SOFC performs three critical functions:

a. separation of reactants;

b. blockage of internal electronic current between the cell electrodes, thus forcing any electronic current to flow in an external circuit;

c. conduction of ions responsible for the internal ionic current that must balance the electronic current in the external circuit.

To be suitable for SOFC applications it must satisfy several requirements:

- it must be an electronic insulator but a good ionic conductor;

- it must be as denser and thinner as possible to limit the ohmic resistance;

- it must be stable under oxidant and reductive atmosphere;

- it must display good mechanical properties;

- it must be chemically and mechanically compatible with the other components of the cell (cathode and anode).

Nowadays, the main material used in SOFC application is the yttria stabilized

zirconia (YSZ).

At the cathode oxygen molecules from air dissociate and are reduced to oxide

ions. The cathode is a porous structure (20 - 40% of porosity) that must allow

rapid mass transport of reactant and product gases. Strontium-doped

lanthanum manganite (LSM), a p-type semiconductor, is the most commonly

used cathode material.

Introduction

19

At the anode hydrogen is oxidized to water. Similarly to the cathode, the

anode has a high porosity so that mass transport of reactant and product gases

is not inhibited. The anode of state-of-art SOFCs is a cermet made of metallic

nickel and a YSZ skeleton [7].

Oxygen pump

At the negative electrode of an oxygen pump, the reaction

O2 + 4 e– → 2 O2–

provides oxide ions that flow through the electrolyte membrane driven by the

application of a voltage Vapp across the electrodes. Once on the opposite side,

at the positive electrode, the reverse reaction gives back oxygen molecules

(fig. 11) [21].

The technique is used for the separation of oxygen from air or for industrial

scale oxygen separation in the conversion of natural gas to syngas.

Fig. 11 Oxygen pump

Introduction

20

Oxygen sensor

In an oxygen sensor (fig. 12), the difference between the oxygen partial

pressure of a standard gas, p0O2, and of an unknown pO2 is obtained by

measuring the open-circuit voltage VOC between the electrodes. The voltage is

related to the difference in the oxygen free energies on opposite sides of the

electrolyte.

Such sensors are extensively used to measure oxygen content in molten

metals particularly liquid steel and also to measure the oxygen content of

gaseous atmospheres like furnace flue gas or car exhaust helping to maintain

air-fuel ratio at an optimum level [9, 22].

Fig. 12 Oxygen sensor

Theory

21

Theory

The electrical conductivity, σ, the proportionality constant between the

current density j and the electric field E, is given by

j/E = σ =

€

i∑ ci Zi q µi

where:

ci is the carrier density (number/cm3);

µi is the mobility (cm2/V s);

Zi q is the charge of the ith charge carrier (q = 1.6 ⋅ 10-19 C).

The differences in σ between metals, semiconductors and insulators generally

result form differences in c rather than µ. On the other hand, the higher

conductivities of electronic versus ionic conductors are usually due to the

higher mobilities of electronic versus ionic species.

Typical ionic solids possess limited numbers of mobile ions, hindered in their

motion because of being trapped in relatively stable potential walls (fig. 13)

so that ionic conduction easily falls below 10-10 S/cm between room

temperature and 200°C.

Theory

22

An activated jump process describes the motion of ions, for which the

diffusion coefficient D is given by

D = D0 e –∆G/kB T = γ (1 – c) Z a2 ν0 e ∆S/kB e –Em/kB T

where:

γ is a constant including geometric and correlation factors;

c is the fractional occupation;

Z is the number of nearest neighbours, so that the (1 – c) Z factor defines the number of neighbouring unoccupied sites;

a is the jump distance;

ν0 is the attempt frequency;

Em is the migration energy.

Since the ion mobility is defined by

µi = Zi q Di / kB T

Fig. 13 Example of potential energy variation along diffusion path

Theory

23

and ci, the density of carriers of Zi q charge, is N c, with N density of ion sites

in the sublattice of interest, the ionic conductivity becomes

σion = N

€

(Zi q)2

kB T γ c (1 – c) Z a2 ν0 e ∆S/kB e –Em/kB T

This expression shows that σion is nonzero only when the product c (1 – c) is

nonzero. Since in a perfect crystal all normal sites are fully occupied (c = 1)

and all interstitial sites are empty (c = 0), this is expected to lead to insulating

characteristics. The classical theory of ionic conduction in solids is thus

described in terms of creation and motion of atomic defects, that is vacancies

and interstitial sites.

The oxygen ion conductivity is therefore given by the sum of oxygen vacancy

and interstitial partial conductivities, though in all oxygen ion electrolytes of

interest the interstitial one does not appear to make significant contribution.

Consequently ionic conductivity can be expressed by

σion ≈ [V

€

O••] 2q µV

with [V

€

O••] concentration of oxygen vacancies doubly positive charged.

Optimized levels of conductivity hence require a combination of high charge

carrier density and mobility without neglecting a continuous pathway of

partially occupied and energetically equivalent ionic sites through the

structure.

Ionic mobilities are greatly enhanced by doping with an aliovalent ion that

introduces mobile vacancies into the occupied sites: this is the method that

has classically been used to induce high charge carrier density in solids.

Another method consists in introducing oxygen atoms in interstitial position

where oxide ions can diffuse.

Theory

24

Furthermore migration energy Em must be minimized. To achieve this task it

should be considered the barrier energy required for a mobile ion to move

from an occupied site to the common interface (fig. 14) between donor and

acceptor sites, respectively occupied site and vacancy.

This energy is minimal where the sum of the ionic radii of the mobile and an

interface ion matches the distance Rb from the centre of the interface to an

interface ion; the effective bottleneck distance Rb is enlarged where the

interface ions are easily polarized and/or the mobile ion can also change its

shape to a prolate ellipsoid.

It must be said that solid oxide electrolytes are purely ionic conductors only

within a specific range of temperatures, oxygen partial pressures and

compositions, beyond which electronic (n- or p- type) contribution becomes

significant. For that reason another important parameter used to describe their

properties is the transport number, a measure of the fraction of the total

Fig. 14 Interface between occupied site and vacancy

Theory

25

current carried by ions: to consider a solid electrolyte useful it must be ≥ 0.99

over a wide range of oxygen partial pressures and temperatures.

It has been observed that at high dopant concentration, which is usually the

case for oxygen ion conductors, the transport number is relatively insensitive

to composition. On the other hand, it is strongly dependant on temperature

and oxygen partial pressure [9, 22, 23].

Experimental

26

Experimental

Sample preparation

All phases were prepared by solid-state reaction in air, starting from oxide

precursor or carbonate ones: Bi2O3 (Riedel - De Haen, 99.5%), Ta2O5

(Aldrich, 99%), SrCO3 (Sigma Aldrich, 98%), La2O3 (Alfa Aesar, 99.9%),

Yb2O3 (Alfa Aesar, 99.9%), CaCO3 (Sigma Aldrich, 99.95%), Y2O3 (Sigma

Aldrich, 99.99%), Pr6O11 (Alfa, 99.9%), Nb2O5 (Aldrich, 99.9%).

Before starting the syntheses, every single reactant was characterized by

X-ray diffraction and, when necessary, thermally treated to obtain the pure

compound.

The appropriate stoichiometric amounts were mixed in an agate mortar, then

put into a crucible and heated up with a heating rate of 200°C/h and

intermediary grinding between the heating steps.

Table 2 summarizes the experimental conditions for every phase.

reaction type of crucible temperature (°C)* reaction time (h)

3/2 Bi2O3 + 1/2 Ta2O5 → Bi3TaO7 alumina 900 130

(3–x)/2 Bi2O3 + x SrCO3 + 1/2 Ta2O5 → Bi3–xSrxTaO7–x/2 + x CO2 [x = 0.05 and 0.1] alumina 900 274

(3–x)/2 Bi2O3 + x La2O3 + 1/2 Ta2O5 → Bi3–xLaxTaO7 [x = 0.05 and 0.1] alumina 900 106

3/2 Yb2O3 + 1/2 Ta2O5 → Yb3TaO7 alumina 1600 164

2.9/2 Yb2O3 + 0.1 SrCO3 + 1/2 Ta2O5 → Yb2.9Sr0.1TaO6.95 + 0.1 CO2 alumina 1600 58

2.95/2 Yb2O3 + 0.05 CaCO3 + 1/2 Ta2O5 → Yb2.95Ca0.05TaO6.975 + 0.05 CO2 alumina 1550 104

2.9/2 Yb2O3 + 0.1 CaCO3 + 1/2 Ta2O5 → Yb2.9Ca0.1TaO6.95 + 0.1 CO2 alumina 1600 130

3/2 Y2O3 + 1/2 Ta2O5 → Y3TaO7 alumina 1600 164

2.9/2 Y2O3 + 0.1 CaCO3 + 1/2 Ta2O5 → Y2.9Ca0.1TaO6.95 + 0.1 CO2 alumina 1550 118

2.9/2 Y2O3 + 0.1 SrCO3 + 1/2 Ta2O5 → Y2.9Sr0.1TaO6.95 + 0.1 CO2 alumina 1550 118

1/2 Pr6O11 + 1/2 Ta2O5 → Pr3TaO7 + 1/2 O2 alumina 1250 480

1/2 Pr6O11 + 1/2 Nb2O5 → Pr3NbO7 + 1/2 O2 alumina 1250 480

1/2 La2O3 + 3/2 Ta2O5 → LaTa3O9 alumina 1600 164

1/2 La2O3 + 3/2 Ta2O5 → LaTa3O9 platinum 1700 78

Tab. 2 Sample preparation: experimental conditions * it represents the maximum temperature reached for the syntheses

Experimental

28

XRD characterization

X-ray powder diffraction (XRD) is an analytical technique primarily used for

phase identification of a crystalline material that can also provide information

on unit cell dimensions.

X-ray diffraction is based on the interaction of the incident X-rays with a

crystalline sample that produces constructive interference (and a diffracted

ray) when conditions satisfy Bragg's Law

n λ = 2d sin θ

being λ the wavelength of incident radiation, θ the diffraction angle, d the

lattice spacing in a crystalline sample.

These diffracted X-rays are then detected, processed and counted.

By scanning the sample through a range of 2θ angles (the geometry of an

X-ray diffractometer is such that the sample rotates in the path of the

collimated X-ray beam at an angle θ while the detector is mounted on an arm

to collect the diffracted rays and rotates at an angle of 2θ), all possible

diffraction directions of the lattice should be attained due to the random

orientation of the powdered material. Conversion of the diffraction peaks to d-

spacings allows identification of the crystalline compound because each

compound has a unique set of d-spacings.

On all phases XRD analyses were performed using a D8 Advance

diffractometer (Bruker AXS), using CuKα radiation (λ1 = 1.5406 Å and

λ2 = 1.54439 Å). Data were collected in the 2θ range 10-100°, with a step of

0.02° and a scan time of 0.2 s per step. The diffractometer was equipped with

a θ-θ goniometer and a LynxEye linear detector.

Experimental

29

For Bi3TaO7 and Yb3TaO7 data were even collected at various temperatures in

air, from 50°C to 900°C and 1100°C respectively, every 25°C in order to

follow the evolution of the cubic lattice parameter, in the 2θ range 25-95°,

with a step of 0.02° and a scan time of 0.2 s per step, with a Pt sample holder.

XRD characterizations of reactants were performed on a Guinier G670

HUBER automated diffractometer. It was equipped with an image plate and a

monochromator acting on the source radiation, hence irradiating the samples

with a nearly pure CuKα1 light (λ = 1.5406 Å). That instrument was also

equipped with a multiple sample-holder and worked on transmission.

All diffraction diagrams were evaluated using the software DIFFRACplus

EVA (Bruker AXS).

Data refining

The refining of the collected XRD data was made by the Rietveld method, a

quantitative analysis method that takes into account the contribution of all

diffraction planes for every considered phase.

It was performed the full pattern matching, that is the determination of a

mathematical XRD pattern profile in order to justify the experimental one.

To obtain such profile the parameters taken into account were:

1. the sample displacement;

2. the lattice parameters;

Experimental

30

3. the profile function (a pseudo-Voigt type), defined by the FWHM of the Bragg’s peaks;

4. the mixing factor η of the pseudo-Voigt profile function and the asymmetry.

That is the sequence by which the refining is carried out.

Five refinement cycles were performed with the refined parameters

considered as the starting ones in the next cycle.

By the Rietveld method refining means to minimize the so-called reduced χ2,

the weighted square difference between the observed and the calculated

pattern, respectively yi and yci, divided by the number of degree of freedom:

€

χv2 =

wi yi – yci( )2i∑

N – P +C

with wi =

€

1 σi2 ,

€

σi2 being the variance of the observed intensity yi and

N the number of the points on the diagram;

P the number of parameters taken into account;

C the number of constrained parameters (distance, angle, etc).

The quality of the agreement between observed and calculated profile is

measured by a set of conventional factors:

Profile factor

€

Rp =100yi – yci

i∑

yii∑

Weighted profile factor

€

Rwp =100wi

i∑ yi – yci

2

wi yi2i∑

Expected weighted profile factor

€

Rexp =100 N − P +Cwi yi2

i∑

Experimental

31

Thus, the reduced χ2 can be expressed as:

€

χ2 =RwpRexp

⎡

⎣ ⎢ ⎢

⎤

⎦ ⎥ ⎥

2

Thermal analyses

Heat flow differential scanning calorimetry analyses were performed using a

SETARAM TAG 2400 instrument in air with a heating rate of 5°C/min.

Differential scanning calorimetry (DSC) is a technique by means of which the

difference between the heat flow through a sample and a reference (air,

alumina, SiC, glass etc.) is measured, during a temperature programme. It is

used to study the thermal behaviour of organic and inorganic compounds and

one of the most important applications concerns the study of phase transitions.

Raman scattering spectroscopy

Raman scattering spectra were obtained using a DYLOR XY800 instrument

with a Kr-ion laser light source, equipped with a CCD (charge-coupled

Experimental

32

device) detector. Samples were irradiated with a light of wavelength equals to

647.1 nm and 10 mW of power.

Raman spectra have their origin in the electronic polarization caused by

ultraviolet, visible, and near-IR light and are obtained by measuring the

radiation scattered at a certain angle, usually 90°. This spectroscopic

technique is used to observe the vibrational transitions occurring as a

consequence of the interaction between light and matter, and is a powerful

means for theoretical and structural studies of inorganic systems as well as

organic ones.

All spectra were evaluated by means of the software LABSPEC.

Scanning Electron Microscopy

SEM images were obtained by using a HITACHI S-3800N instrument on

gold-sputtered samples.

By means of the scanning electron microscope (SEM), images of surfaces can

be obtained exploiting the interaction of an electron beam with matter.

Thermoionically emitted electrons are accelerated by means of an applied

voltage (experimental images were obtained with accelerating voltages of 5

kV and 7 kV) and focalised on a sample through a system of electrostatic

lenses. While the electron beam scans the surface of the sample, from the

latter are emitted secondary, backscatterd and Auger electrons along with X

Experimental

33

photons. The secondary electrons are produced at a depth ranging from 50 to

500 Å and they are, thus, sensible to the topography of the sample's surface:

by detecting those ones it can be built an image of the surface.

Electron Probe Microanalyses

The electron probe microanalyzer is used to determine, locally (irradiating

volume ~ 1 µm3), the chemical composition of a sample. The X photons

emitted by means of the interaction of a highly accelerated electron beam (20

kV) with matter are characteristic of the chemical element constituting the

sample that emits them.

Analyses were performed using a CAMECA SX100 instrument equipped

with three wavelength dispersive spectrometers.

Experimental

34

Electrical measurements

Electrical parameters were determined by impedance spectroscopy using a

Solartron 1260 frequency response analyzer in the frequency range 1 Hz to

1 MHz, under a 500 mV voltage. Samples for impedance measurement were

tablets (5 mm diameter, 2-3 mm thickness) sintered at 800°C for 12 hours,

with gold-coated parallel faces: the coating was made in two steps, by

sputtering and by brushing using an organic suspension of gold powder.

Pellets were shaped by uniaxial pressing and compacted by a 5 minutes

isostatic hydraulic pressing at 1700 bars. Impedance spectra were recorded in

air during heating and cooling ramps between 200°C and 800°C at stabilized

temperatures separated by 20°C steps.

The impedance spectroscopy technique is based on the characterization of the

sample’s response to an alternative tension of variable frequency. The electric

impedance of a system, denoted by Z(ω), represents the ratio between the

applied sinusoidal tension, U(ω), and the resulting electrical current, I(ω),

where

ω = 2 π ϕ

and ϕ is the frequency. Z(ω), U(ω) and I(ω) are complex quantities that can

be represented on a Cartesian plot:

Z(ω) = ZR + j ZI = |Z| e j θ

with ZR representing the real part of Z(ω), ZI its imaginary part, |Z| its modulus

and θ its argument. There are many ways impedance spectroscopy data may

be plotted. Conventionally one plots

–ZI = ƒ (ZR)

Experimental

35

to give a complex-plane impedance plot (Cole-Cole plot).

Through Z the resistivity R of a material can be determined by modelling the

physicochemical processes occurring while applying the tension, by means of

an appropriate equivalent electrical circuit.

In the simplest case of an homogeneous tablet of 100% of relative density or a

monocrystalline solid, perfectly in contact with the electrodes, the sample

itself can be described as a resistor in parallel with a capacitor

(fig. 15):

The total impedance of such circuit is:

Multiplying it for its complex conjugate, we obtain:

€

ZR (ω) =R

1+ ( j Cω)2 (‡) and

€

ZI (ω) =−R2ω C

1+ ( jCω)2

Combining these two expressions, we get:

€

ZR − R2( )2

+ ZI2 = R

2( )2

!

Z(") =R

1+ j R C "

Fig. 15 a. Electrical circuit equivalent to the sample; b. related impedance plot

Experimental

36

That is the equation of a circumference centred in (R/2 ; 0) and of diameter R.

From (‡) it can be seen that Z approaches to the intrinsic resistivity of the

material R as the frequency approaches to zero and that Z approaches to zero

as the frequency approaches to infinity. The amplitude maxima for Z is

obtained when

ω = ω0 = 2 π ϕ0 =

€

1RC

where ω0 is the characteristic resonance frequency of the system.

Conductivity is defined by:

€

σ = lR S

where:

l is the thickness of the sample tablet;

S is its surface area;

R is the resistivity.

Given the Arrhenius law

σ =

€

σ0T e–∆E/NA kB T ⇒

⇒ log (σT) = log σ0 –

€

ΔE2.3 ⋅1000 ⋅NA ⋅ kB

⋅1000T

by plotting log (σT) versus 1000/T one can get macroscopic data about the

sample, the activation energy related to the oxygen ions drifting ∆E.

Every impedance plot obtained is evaluated by means of software through

which data are collected and transformed in order to get a table of log σ

versus T.

Results and discussion

37


The first task of the merely laboratory work was, once collected the

bibliographic references about the compounds to synthesize, trying to obtain

the pure phases by means of a solid-state reaction. As always, when facing to

experimental experiences, it is hard to fulfil every task.

It will be shown, first, the results of the experiences that didn't give positive

answers or, at least, not the hoped ones. Then the struggle for bringing order

from chaos started giving better results, upon which continuing the

investigation was therefore possible.

LaTa3O9

The LaTa3O9 phase is orthorhombic with space group Pnma. In the structure

(fig. 16), ribbons of pentagonal TaO7 bipyramids are recognizable parallel to

the a-direction. These ribbons are connected with each other by TaO6

octahedra: the tunnels formed in this way are occupied by lanthanum atoms

[24]. This structure was taken into account because such tunnels might


38

represent a path through which oxygen ions could flow: the substitution of

lanthanum for zirconium might provide some interesting clues about such

possibility by introducing oxygen vacancies within the structure.

Unfortunately it was impossible to obtain that phase with a simple solid-state

reaction in air since other phases are formed.

By using an alumina crucible at 1600°C the phase obtained was La0.33TaO3,

primitive tetragonal with space group P4/mmm. Then it was tried to carry out

the reactionin a platinum crucible: by heating the mixed powders at 1700°C,

different phases coexist, mainly La0.33TaO3 and LaTa5O14, besides

La4.67Ta22O62. Figures 17 and 18 show the XRD diagrams related to the two samples.

Fig. 16 La3TaO9 structure


39

Fig. 17 XRD diagram of the Al2O3/1600°C sample constituted by La0.33TaO3


40

Pr3TaO7 and Pr3NbO7

Both phases are described in the literature [25] as fluorite-related

orthorhombic structures with space group Cmcm: for that reason it was

interesting to investigate the possibility of doping them by substituting Pr3+

ions with lower valence cations, such as Sr2+, therefore creating oxygen

vacancies and allowing so oxide ions conduction. Morevover PrIII can be

Fig. 18 a. XRD digram of the Pt/1700°C sample; b. particular of the diagram showing the phases La0.33TaO3 (x), LaTa5O14 (v) and La4.67Ta22O62 (o)


41

oxidized to PrIV and this might add mixed electronic/ionic conducting

properties.

During the synthesis processes it was observed the persistent presence of

grains of unreacted praseodymium (III, IV) oxide which required a 20 hours

ball milling treatment and a following heating step to disappear: a preliminary

reduction of the particles size is therefore desirable before starting the

syntheses. The XRD diagrams of the niobate phase and of the tantalate one

are very similar as it can be expected by the resemblance of the ionic radius of

the two ions (Nb+5, 69 pm; Ta5+, 68 pm).

Figures 19 and 20 show the XRD diagram and the refined profile for the

phase Pr3TaO7 while figures 21 and 22 show the diagrams related to Pr3NbO7.

Unfortunately the lack of time made the carrying on of the investigation

impossible to fulfil so the report on these phases stops with the definition of a

procedure for the syntheses.

Fig. 19 XRD diagram of Pr3TaO7 at room temperature


42

Fig. 20 Fitted profile for Pr3TaO7, orthorhombic, base-centered, space group Cmcm Lattice parameters a = 10.97392(12) Å, b = 7.51560(4) Å, c = 7.67241(4) Å Rp = 16.7%; Rwp = 23.4%; Rexp = 8.00%; χ2 = 8.56

Fig. 21 XRD diagram of Pr3NbO7 at room temperature


43

Fig. 22 Fitted profile for Pr3NbO7, orthorhombic, base-centered, space group Cmcm Lattice parameters a = 10.96504(10) Å, b = 7.51192(3) Å, c = 7.66809(3) Å Rp = 15.3%; Rwp = 20.6%; Rexp = 8.41%; χ2 = 5.99


44

Y3TaO7 and related phases

Y3TaO7

Y3TaO7 phase is said to be cubic above 1600°C and orthorhombic at 1400°C

[26], but many essays showed that the cubic defect-fluorite phase at 1600°C is

in equilibrium with a small quantity of orthorhombic phase.

In fig. 23 the experimental XRD diagram and in fig. 24 an enlargement

showing the pattern of the orthorhombic phase.

Fig. 23 X-ray powder diffraction pattern for Y3TaO7 at room temperature


45

This first positive result allowed to continue for the second step, the

substitution of Y3+ ion for a lower valence one in order to create oxygen

vacancies in the fluorite-related phase, therefore enhancing the conditions for

conducting oxide ions. Calcium and strontium were tried, attempting at

obtaining a Y3–xMexTaO7–δ solid solution with x = 0.1.

A very first sign of a successful substitution is a change in the lattice

parameters, compared to those of the non-doped compound: in accordance

with Vegard's law, it exists a linear relation between the crystal lattice

parameters of a solid solution and the concentration of the doping agent. It

must be said that this is an empirical rule so it is not sufficient to confirm the

substitution but it is anyway a good start.

Fig. 24 XRD pattern for Y3TaO7; enlargement showing the orthorhombic phase pattern


46

The experimental pattern was therefore refined in order to find a value for the

cubic lattice parameter that was influenced to a lesser extent by the

experimental conditions. For Y3TaO7 both phases were considered by

imposing the Fm

€

3 m space group for the cubic one and the C2221 space group

for the orthorhombic one; fig. 25 shows the refined pattern while results are

displayed in legend.

Fig. 25 Fitted profile for Y3TaO7: cubic, space group Fm m and orthorhombic, space group C2221. Cubic phase - Lattice parameter a = 5.25739(4) Å Orthorhombic phase - Lattice parameters a = 10.42639(36) Å, b = 7.42912(12) Å, c = 7.42020(12) Å Rp = 14.8%; Rwp = 23.5%; Rexp = 6.05%; χ2 = 15.1 !

3


47

Y2.9Ca0.1TaO6.95

For the Y2.9Ca0.1TaO6.95 phase some results were positive but further

investigations are needed because although the Ca2+ ion seems to be entered

within the structure, since the cubic lattice parameter has changed (from

5.25739 Å to 5.25814 Å), an unexplained reflection occurred in the XRD

pattern.

The XRD pattern of the Ca2+-added Y3TaO7 phase was refined considering

only the cubic phase since it is the one to consider in discussing the cubic

lattice parameter variation; the fitted profile is shown in fig. 26. Figure 27

shows an enlargement of the experimental XRD diagram with the

unexplained peak.

Fig. 26 Fitted profile for Y2.9Ca0.1TaO6.95, cubic, space group Fm m. Lattice parameter a = 5.25814(9) Å; Rp = 15.0%; Rwp = 22.1%; Rexp = 6.06%; χ2 = 13.3

!

3


48

Y2.9Sr0.1TaO6.95

For the Y2.9Sr0.1TaO6.95 phase something can instead be stated: the solid

solution was not obtained because during the synthesis a phase separation

occurred and the Y(Sr0.5Ta0.5)O3 phase was formed (fig. 28 and 29), polluting

the main fluorite-type compound. For the experimental conditions in which

the reaction was carried out the substitution was impossible so the solid

solution doesn't exist or, at least, a change in the experimental conditions is

needed.

Fig. 27 XRD pattern for Y2.9Ca0.1TaO6.95. Enlargement showing the unexplained peak


49

Fig. 28 XDR patterns for Y3TaO7 (black) and the Sr2+-added phase (red, up-shifted); a. entire dagrams and b. enlargement showing some of the reflections due to the Y(Sr0.5Ta0.5)O3 phase

Fig. 29 Reflections due to the Y(Sr0.5Ta0.5)O3 phase


50

Yb3TaO7 and related phases

Yb3TaO7

The Yb3TaO7 phase has a cubic fluorite-type structure, an fcc lattice with

fully occupied anion sites. Fig. 30 and 31 show the experimental XRD pattern

and the fitted profile, respectively. In fig 32 the XRD diagrams with raising

temperature, in fig. 33 the DSC performed on the compound and in fig. 34 the

thermal variation of the cubic lattice parameter. Data collected show that the

phase is stable at least until 1300°C with no phase transition occurring and

that it has a lattice parameter that changes linearly with the raising

temperature. Proofs of its stability until 1700°C can be found in the literature

[27]. Having a pure compound opens the door on the possibility of doping the

phase with lower valence cations, such as Sr2+ and Ca2+.

Fig. 30 X-ray powder diffraction pattern for Yb3TaO7 at room temperature


51

Fig. 31 Fitted profile for Yb3TaO7, cubic, space group Fm m. Lattice parameter a = 5.19448(4) Å; Rp 10.2%; Rwp 13.6%; Rexp 6.57%; χ2 = 4.1

!

3

Fig. 32 a. XRD diagrams for Yb3TaO7 with raising temperature in air; b. enlargement of the red highlighted Bragg’s peaks (the firsts on the left)


52

-1

4

9

14

19

24

29

34

0 200 400 600 800 1000 1200 1400

Hea

t flo

w (µ

V)

Temperature (°C)

→

5.19

5.20

5.21

5.22

5.23

5.24

5.25

0 200 400 600 800 1000 1200

Latti

ce p

aram

eter

(Å)

Temperature (°C)

Fig. 33 DSC performed on Yb3TaO7 in air, heating rate 5°C/min

Fig. 34 Thermal variation of the cubic lattice parameter, a, in Yb3TaO7


53

Yb2.9Sr0.1TaO6.95

In the same attempt at creating oxygen vacancies, it was tried to make the

strontium ion enter the structure of Yb3TaO7 by replacing the Yb3+ ion. What

happened is that even for Yb2.9Sr0.1TaO6.95 a phase separation occurred during

the synthesis and the SrYb0.5Ta0.5O3 phase was formed (fig. 35 and 36),

polluting the main fluorite-type compound: the solid solution Yb3–xSrxTaO7–x/2

cannot exist if one adopts the experimental conditions that were successful for

the synthesis of the non-doped Yb3TaO7.

Fig. 35 XDR patterns for Yb3TaO7 (black) and the Sr2+-added phase (red, up-shifted);

a. entire diagrams and b. enlargement showing some of the reflections due to the SrYb0.5Ta0.5O3 phase


54

Yb3–xCaxTaO7–x/2

To replace Yb3+ ions for the Ca2+ ones a change in the experimental

conditions was tried. Two degree of substutition were tried: for x = 0.05 the

temperature raises between the heating steps were made smaller and the phase

was formed at 1550°C; for the x = 0.1 compound the same conditions used for

the Sr-doped one gave a positive answer. Fig. 37 shows the XRD diagrams of

the three phases: x = 0, x = 0.05 and x = 0.1. Data were, then, refined and the

results are shown in fig. 38 and 39.

Fig. 36 Reflections due to the SrYb0.5Ta0.5O3 phase


55

Fig. 37 XRD patterns for the Yb3–xCaxTaO7–x/2 solid solution; x = 0 in black, x = 0.05 in red, x = 0.1 in blue

Fig. 38 Fitted profile for Yb2.95Ca0.05TaO6.975, cubic, space group Fm m. Lattice parameter a = 5.19672(9) Å; Rp = 12.2%; Rwp = 15.9%; Rexp = 8.52%; χ2 = 3.5

!

3


56

From the XRD diagrams of the Ca2+ doped phases it cannot be seen any phase

separation but the variation of the lattice parameter as the presence of calcium

ions becomes higher doesn't follow the Vegard's law (fig. 40). Anyway the

Vegard's law is only an empirical rule so further investigations are needed.

What the experiments with the yttrium and the ytterbium phases show is that

the Sr-doped compounds are harder to obtain than the Ca-doped ones. This

can be explained by the larger ionic radius of the Sr2+ ion compared to that of

the Ca2+ ion, 112 pm versus 99 pm. While calcium ion seems to replace the

yttrium ion (ionic radius 89.3 pm), with which the diagonal relationship also

might be invoked, and the ytterbium one (85.8 pm) in the fluorite-related

lattice, the strontium ion seems too big to fit their place.

Fig. 39 Fitted profile for Yb2.9Ca0.1TaO6.95, cubic, space group Fm m. Lattice parameter a = 5.19662(4) Å; Rp = 10.8%; Rwp = 14.2%; Rexp = 6.68%; χ2 = 4.5

!

3


57

Bi3TaO7 and related phases

δ-Bi2O3 has a disordered cubic fluorite structure and exhibits one of the

highest known oxide ion conductivities, 2.3 S cm–1 at 800°C. The reasons for

that relatively high conductivity are [28]:

a. 1/4 of the oxygen sites are vacant in the fluorite-type lattice;

b. the electronic structure of Bi3+ is characterized by the presence of 6s2 lone pair electrons, leading to high polarizability of the cation

network, which in turn leads to oxide ion mobility;

c. the ability of the Bi3+ to accommodate highly disordered surroundings.

Fig. 40 Evolution of the cubic lattice parameter a with increasing Ca2+ presence. Error is reported as 3σ


58

This phase is stable only between 730°C and 825°C but the substitution of

bismuth cations by others can lead to its room temperature stabilization.

Ta5+ is such a cation, giving, by a 25% molar substitution, the phase Bi3TaO7,

a yellow compound with a fcc fluorite-type structure belonging to the Fm

€

3 m

space group. Abrahams et al. [6] proposed a structure in which there are

chains of apical sharing TaO6 octahedra, explaining that as a consequence of

the presence of tantalum: such chains would make the structure similar to a

weberite (fig. 41), at which we can look as a defect-fluorite by considering

only the cationic array.

Bi3TaO7

Figure 42 shows the X-ray diffraction pattern for Bi3TaO7 at room

temperature while figure 43 shows the refined diffraction diagram: red dots

for the experimental diagram, black profile for the calculated one and blue

profile for the difference between them.

Fig. 41 The structure of the weberite Na2MgAlF7


59

Fig. 42 X-ray powder diffraction pattern for Bi3TaO7 at room temperature

Fig. 43 Fitted profile for Bi3TaO7, cubic, space group Fm m. Lattice parameter a = 5.45761(9) Å; Rp 16.0%; Rwp 21.0%; Rexp 7.53%; χ2 = 7.8

!

3


60

The difference between the experimental diagram and the calculated one,

represented by the blue profile on the refined pattern, might seem not

negligible but such difference arises from the presence of preferential

orientations due to the preparation of the sample for the XRD analysis.

The teflon sample holder gives a parasite reflection at 2θ ≅ 18°: given the

absence of Bragg’s reflections related to the main phase in those proximities,

the refinement was made by excluding the region around that peak; the blank

space in the diagram is due to that reason.

The X-ray powder diffraction profile reveals a superlattice ordering (fig. 44),

better shown by the neutron diffraction pattern (fig. 45), arising from a three-

dimensionally modulated incommensurate structure:

Fig. 44 X-ray powder diffraction for Bi3TaO7. Enlargement showing a superlattice reflection (green mark)


61

The second order Bragg peaks were indexed by using a single modulation

parameter, ε = 0.3780(3) (courtesy of Pascal Roussel). Considering that the

reflections due to that incommensurate structure are observed with greater

intensity on the neutron diffraction profile, we can say that most likely this

kind of structure involves oxygen atoms.

Such incommensurately modulated structure, with an underlying average

fluorite unit cell that has the space group symmetry Fm m, is known in the

literature as Type II [29].

A careful analysis of the diffraction profile reveals, besides, the presence of

another phase (fig. 46a and 46b), Bi7Ta3O18, with peaks of slight intensity that

grow as temperature rises (fig. 47).

!

3

Fig. 45 Neutron diffraction pattern for Bi3TaO7 at room temperature in air. Black marks indicate the first order Bragg reflections, green ones the second order reflections due to the superlattice structure


62

Fig. 46a X-ray powder diffraction for Bi3TaO7 and enlargement showing the Bragg’s peaks related to the Bi7Ta3O18 phase

Fig. 46b Enlargement showing a portion of the diffraction pattern of Bi7Ta3O18


63

The presence of the phase Bi7Ta3O18 can be attributed to the sublimation of

Bi2O3 (melting point 825°C) as suggested by the lowering of the Bi/Ta molar

ratio, equals to 3 in Bi3TaO7 and to 2.3 in Bi7Ta3O18.

The Bi7Ta3O18 phase is labelled in the literature as a Type II* structure, a

commensurately modulated δ-Bi2O3-related phase. This is not, anyway, the

right way to refer at it strictu sensu. The space group symmetry is triclinic C1

but is very close to monoclinic C2/m. The metal atom array is approximately

face-centred cubic, punctuated by regularly spaced displacement faults.

The coordination environments of the Bi atoms appear to be related to the

cubic (eightfold) coordination found in fluorite, distorted by the presence of

TaO6 octahedra: this leads to pyramidal and trigonal prismatic coordination

Fig. 47 Growing of the peaks related to the phase Bi7Ta3O18 as temperature rises: syntesized compound (black), 3 hours sintering at 950°C (red) and 3 hours sintering at 1000°C (blue)


64

environments, typically observed around cations with stereochemically active

electron lone pairs, such as Bi3+ [30].

Bi3TaO7 reveals so its instability with rising temperature; conversely it can be

said that both phases and their relative abundances are stable from room

temperature up to 900°C (the synthesis temperature) being the diffraction

diagrams the same after every further thermal treatment: they do not make

Bi2O3 to sublime any more nor any phase transition occurs as it is shown by

the thermal analysis (fig. 48). The XRD diagrams (fig. 49a and 49b) show that

the Bi7Ta3O18 phase is still there, though its main reflection is progressively

hidden by the left shift of the Bi3TaO7 phase's peak at 2θ ≅ 28.2°, produced

by the raise of temperature.

Fig. 50 shows the thermal variation of the cubic lattice parameter, that reveals

a small degree of non-linearity above ca. 500°C.

-4 -2 0 2 4 6 8

10 12 14 16

0 200 400 600 800 1000

Hea

t flo

w (µ

V)

Temperature (°C)

→

Fig. 48 DSC performed on Bi3TaO7 in air, heating rate 5°C/min


65

Fig. 49a a. XRD diagrams for Bi3TaO7 with raising temperature in air; b. enlargement of the red highlighted Bragg’s peaks (the first on the left)

Fig. 49b The phase Bi7Ta3O18 is still visible from room temperature up to 900°C


66

Bi3–xLaxTaO7

Despite the presence of the Bi7Ta3O18 phase, whose presence is almost

negligible but still identifiable, the substitution of the bismuth ion with the

other cations was tried.

The results for the solid solutions with La2O3 were conflicting: although the

XRD diagrams don't show any phase separation and furthermore the

increasing presence of lanthanum seems to inhibit the formation of Bi7Ta3O18

(fig. 51), the evolution of the lattice parameter with the increasing lanthanum

presence doesn't follow the Vegard's law (fig. 52). In figures 53 and 54 the

refined diffraction profiles for Bi2.95La0.05TaO7 and Bi2.9La0.1TaO7

respectively.

5.45

5.46

5.47

5.48

5.49

5.50

5.51

5.52

0 100 200 300 400 500 600 700 800 900 1000

Latti

ce p

aram

eter

(Å)

Temperature (°C)

Fig. 50 Thermal variation of the cubic lattice parameter, a, in Bi3TaO7


67

Fig. 51 a. XRD patterns for the Bi3–xLaxTaO7 solid solution; x = 0 in black, x = 0.05 in red, x = 0.1 in blue b. enlargement showing the disappearing of the main reflection due to Bi7Ta3O18 while lanthanum presence increases

Fig. 52 Evolution of the cubic lattice parameter a with increasing La3+ presence. Error is reported as 3σ


68

Fig. 53 Fitted profile for Bi2.95La0.05TaO7, cubic, space group Fm m. Lattice parameter a = 5.45901(6) Å; Rp 13.7%; Rwp 17.2%; Rexp 7.38%; χ2 = 5.4

!

3

Fig. 54 Fitted profile for Bi2.9La0.1TaO7, cubic, space group Fm m. Lattice parameter a = 5.45821(9) Å; Rp 18.1%; Rwp 23.3%; Rexp 7.28%; χ2 = 10.2

!

3


69

Other evidences that would confirm or refute that the substitution had actually

taken place can be provided by the examination of the Raman scattering

spectra and, in particular, of the band at ca. 748 cm–1 that can be related to the

Ta—O bond stretching.

Fig. 55 compares the portion of interest of the spectra of the La-doped

compounds with that of Bi3TaO7 (courtesy of Annick Rubbens).

The change of the tantalum surroundings, as a result of the introduction of

La3+ cations within the structure, would cause a shift of the emitting

frequencies of the band but none seems to have occurred, thus most likely the

substitution has not taken place.

Fig. 55 Raman scattering spectra of the Bi3–xLaxTaO7 solid solutions


70

Bi3–xSrxTaO7–x/2

The data collected in the case of the Sr-doped compounds suggest that a solid

solution exists for 0 < x ≤ 0.1: no phase separation has occurred (fig. 56);

there is a linear relation between the strontium concentration and the

evolution of the lattice parameter (fig. 57); a slight lower frequencies shift of

the band at ca. 748 cm–1, ascribable to the Ta—O bond stretching, is observed

while increasing molar substitution (fig. 58).

It was observed also that the Bi7Ta3O18 phase does not appear in the Sr-doped

phases (fig.59), despite the preparation was almost the same than that of

Bi3TaO7 (up to 900°C).

In figures 60 and 61 the refined diffraction profiles for Bi2.95Sr0.05TaO6.975 and

Bi2.9Sr0.1TaO6.95 respectively.

Fig. 56 XRD patterns for the Bi3–xSrxTaO7–x/2 solid solution; x = 0 in black, x = 0.05 in red, x = 0.1 in blue


71

Fig. 57 Evolution of the cubic lattice parameter a with increasing Sr2+ presence. Error is reported as 3σ

Fig. 58 Raman scattering spectra for the Bi3–xSrxTaO7–x/2 solid solution


72

Fig. 59 XRD diagrams for the Bi3–xSrxTaO7–x/2 solid solution: x = 0 in blue, x = 0.05 in red, x = 0.1 in green. The Bi7Ta3O18 phase (dark blue marks) is not present in the Sr2+-doped compounds

Fig 60 Fitted profile for Bi2.95Sr0.05TaO6.975, cubic, space group Fm m. Lattice parameter a = 5.45631(4) Å; Rp 10.5%; Rwp 13.4%; Rexp 7.87%; χ2 = 2.9

!

3


73

How the strontium ion influences the conducting properties of Bi3TaO7

Once verified the existence of the Bi3–xSrxTaO7–x/2 solid solution, the second

step was the preparation of the samples for the electrical measurements, that

is making the tablets.

Tablets must be compact and their relative density has to approach as much as

possible the theoretical one, in the case of Bi3TaO7 equals to 9.478 g/cm3 [6].

Fig. 61 Fitted profile for Bi2.9Sr0.1TaO6.95, cubic, space group Fm m. Lattice parameter a = 5.45474(8) Å; Rp 14.7%; Rwp 18.4%; Rexp 8.13%; χ2 = 5.12

!

3


74

Attaining the maximum relative density is important because conductivity is

inversely proportional to resistance and the unfilled space within the tablet

gives a contribution to the latter.

Anyway since the objective was to investigate how the strontium ion

influences the conducting properties of Bi3TaO7, powders were grounded only

manually in the agate mortar, then pressed and shaped as tablets.

In table 3 the relative densities of the tablets are summarized.

sample relative density

Bi3TaO7 78.8 %

Bi2.95Sr0.05TaO6.975 78.8 %

Bi2.9Sr0.1TaO6.95 78.5 %

The size of particles ranged from few µm up to 20 µm. Fig. 62 shows some

SEM images of the particles of the Bi3TaO7 phase while in fig. 63 is shown

the surface of the tablet .

Tab. 3 Relative densities of the tablets used for the electrical measurements

Fig. 62 SEM images of Bi3TaO7 particles


75

By looking at the tablet surface, it can be noticed that there is plenty of

scratches and holes.

Another contribution to the resistance is given by the imperfect contact

between the tablet and the electrode, and it seemed unlikely that the

imperfections were formed as a consequence of the size of particles.

Even if the contribution to the resistance related to those imperfections could

have been negligible, since within the first heating cycle the contact becomes

better as the gold-coated surface of the tablet welds to the electrode, by trying

to look for the cause of the phenomenon it was discovered an heterogeneity in

the appearance of the tablet: the scratches were actually produced during the

polishing of the surface by little grains (diameter up to 100 µm) dispersed in

the tablet matrix, formed during the sintering step.

In figure 64 those grains are showed as they appear at the microscope.

An XRD analysis was then performed on a sample of a grounded tablet but

showed no differences.

To investigate the nature of those little grains an analysis with the electron

probe microanalyzer was carried out (courtesy of Séverine Bellayer).

Fig. 63 Bi3TaO7; SEM images of the tablet's surface. Scratches and holes keep the tablet from being in perfect contact with the electrodes


76

Despite the different colour, the analysis confirmed the same composition for

the grains and the matrix (fig. 65 a and b).

Fig. 64 Images of the tablet's surface showing the grains dispersed in the tablet matrix

Fig. 65a EPMA analysis of grains


77

Figure 66 a and b and figure 67 a and b show the particles size and the tablet

surface of Bi2.95Sr0.05TaO6.975 and Bi2.9Sr0.1TaO6.95, respectively. The size of

the particles ranged for both phases from few µm to 20 µm and the surfaces

are scratched in the same manner as the Bi3TaO7 tablet.

Even in those tablets the same grains were present and, by induction, it was

concluded that they had the same composition as the matrix.

All the tablets, thus, presented the same characteristics: their relative densities

were very similar, their particles had similar size, their surfaces displayed the

same imperfections. Given that, the electrical measurements were performed.

Results are shown in the form of an Arrhenius plot in fig. 68 along with

δ-Bi2O3 and YSZ as general terms of comparison.

Fig. 65b EPMA analysis of the tablet matrix


78

Fig. 66 SEM images of the Bi2.95Sr0.05TaO6.975 phase; a. particles and b. tablet surface

Fig. 67 SEM images of the Bi2.9Sr0.1TaO6.95 phase; a. particles and b. tablet surface


79

The slope change in the x = 0 curve, at around 525°C, is reflected in the

temperature variation of the cubic lattice parameter of Bi3TaO7 (fig. 50).

The conductivity exhibited by Bi3TaO7 is lower than that of δ-Bi2O3, as

expected, for the concentration of vacancies decreases when doping Bi2O3

with a pentavalent oxide like Ta2O5.

In fact by expressing the unit cell of Bi2O3 as Bi4O6□2 (where □ represents

oxygen vacancy) the resulting composition is:

2 [(Bi2O3)1–x (Ta2O5)x] → Bi4(1–x)Ta4xO6+4x□2–4x

It may be considered that while the vacancy concentration of pure Bi2O3 is

too high for the δ phase to be stable at low temperature, the filling up of

vacancies with the added oxide stabilizes the phase at room temperature [9].

Doping Bi3TaO7 by adding Sr2+ ions within the structure should have been,

therefore, a way to enhance its conductivity by introducing oxygen vacancies,

Fig. 68 Arrhenius plot of total conductivity for the Bi3–xSrxTaO7–x/2 solid solution and some terms of comparison


80

whose concentration is directly proportional to the conductivity itself but the

results showed the contrary.

In fig. 69 is shown the high temperature portion of the Arrhenius plot of total

conductivity for the Bi3–xSrxTaO7–x/2 solid solution, that is the portion beyond

the point where the slope change is observed.

It is taken for granted that the observed conductivities are due mainly to

mobile O2– ions whose drifting, in Bi3TaO7, has an activation energy of 1.18

eV, in good accordance with the value reported in the literature, 1.16 eV [6].

The lower conductivity showed by the doped samples can partly be explained

with the shrinking of the lattice parameter, which influences the migration

energy of the oxide ions. However more important seems the fundamental

difference between the electronic structure of the Sr2+ ion and of the Bi3+ one:

-1.8

-1.3

-0.8

-0.3

0.2

0.7 0.9 0.95 1 1.05 1.1 1.15 1.2 1.25 1.3

log

(σ T

) (S

cm-1

K)

1000/T (K-1)

x = 0

x = 0.05

x = 0.1

∆E0 = 1.18 eV

∆E0.05 = 1.211 eV

∆E0.1 = 1.198 eV

Fig. 69 Arrhenius plot of total conductivity for the Bi3–xSrxTaO7–x/2 solid solution (high temperatures)


81

the Bi3+ has a 6s2 lone pair that confers to the ion a high polarizability, crucial

to the conducting properties [31].

The role of the latter is even more highlighted at lower temperatures (fig. 70)

in the phase with higher substitution degree. In fact the higher concentration

of vacancies, created by means of the strontium ion, means also a higher

hindrance to the mobility of the oxide ions since that privileged path provided

by the Bi3+ ions becomes more limited. Here, at low temperatures, the loss of

polarizability play a more important role whereas as temperature raises the

contribution of the increasing concentration of vacancies (the intrinsic ones

that arise as temperature is increased) tends to erase the initial disadvantage.

-7

-6

-5

-4

-3

-2

-1

0 1.2 1.3 1.4 1.5 1.6 1.7 1.8 1.9 2 2.1

log

(σ T

) (S

cm–1

K)

1000/T (K–1)

x = 0

x = 0.05

x = 0.1

ΔE0.1 = 1.08 eV

ΔE0 = 0.95 eV

ΔE0.05 = 1.04 eV

Fig. 70 Arrhenius plot of total conductivity for the Bi3–xSrxTaO7–x/2 solid solution (low temperatures)


82

The measurements performed on Bi3TaO7 also showed the possible aptitude

of the phase at conducing protons: in the chart of fig. 71 it can be observed the

difference between the curve of total conductivity recorded during the first

heating cycle and that recorded during the last one, the fourth. The two curves

rejoin at ca. 400°C, probably when all the water leaves the tablet. Given that

the phenomenon is observed only during the first heating cycle while the

other three give the same curves, protons might give some contribution to the

total conductivity, as they do in La3TaO7 [4].

Fig. 71 Arrhenius plot of total conductivity for Bi3TaO7. Curves related to the first heating cycle and to the fourth one

Conclusions

83

Conclusions

This scientific research explored the possibility to develop new electrolyte

materials for solid oxide fuel cells.

The basic research deals with the synthesis and the characterization of new

oxide ions conductors starting from the former experiences in this domain and

trying to answer to the need for more performing materials mainly for what

concerns the lowering of operating temperatures.

This work looked at several phases starting from the synthesis and coming to

the substitution of cations for lower valence ones to obtain new doped phases.

Some syntheses did not succeed, as in the case of LaTa3O9, for which a

different method must be tried since a classic solid-state reaction cannot lead

to the desired compound.

For some others, as in the case of the niobate Pr3NbO7 and the tantalate

Pr3TaO7, the lack of time did not allow to continue through the substitution

step and the investigation stopped with the definition of a procedure for the

syntheses which must involve an accurate reduction of the size of the

reagents' particles in order to obtain the desired phases.

Some other syntheses were successful: for Y3TaO7 and Yb3TaO7 the second

step followed, the attempt at substituting the Y3+ ion and the Yb3+ one,

respectively, for lower valence cations such as Sr2+ and Ca2+. Unfortunately

Conclusions

84

all the substitutions involving the Sr2+ ion did not succeed since phase

separations occurred whereas on the XRD diagram of the phase

Y2.9Ca0.1TaO6.95 a peak has remained unexplained . The only solid solution

that seems to exist is Yb3–xCaxTaO7–x/2 although further investigations are

needed to confirm it. What seems to be a preference for the calcium ion can

be explained by considering its size that is smaller than that of the strontium

one, therefore it can enter within the structure.

A deeper investigation was carried out on the Bi3TaO7 phase, a yellow

compound with a fluorite-type structure.

Abrahams et al. [6] proposed a structure in which there are chains of TaO6

octahedra that would make the structure similar to a weberite, at which we

can look as a defect-fluorite by considering only the cationic array.

Preux [4] investigated the weberite-related phase La3TaO7 and its solid

solution La3–xSrxTaO7–x/2, showing that for x = 0.1 we get a material with

substantially enhanced conducting properties.

Two substitution, thus, were tried: Bi3+ for La3+, to explore the possibility to

obtain a solid solution in order to find an answer to the compatibility issues of

Bi-based compounds in the reducing conditions of SOFC devices; Bi3+ for

Sr2+, to see if the latter could enhance the conducting properties of Bi3TaO7 as

it does for the La3TaO7 phase.

While the existence of the Bi3–xLaxTaO7 solid solution is questionable, the

data collected by means of the X-ray diffraction and the Raman scattering

spectroscopy confirmed the existence of a Bi3–xSrxTaO7–x/2 solid solution for

0 < x ≤ 0.10.

Conclusions

85

Such result allowed to go into the next step: how the presence of the strontium

ion influences the conducting properties of Bi3TaO7.

By means of the impedance spectroscopy technique it was discovered that

doping the Bi3TaO7 phase by adding Sr2+ ions within the structure doesn't

lead, as hoped, to a material with enhanced oxide ions conducting properties.

In fact, despite the increase of oxygen vacancies concentration, the doped

phases are worst conductors. This is due mainly to the differences in the

electronic structure between the strontium ion and the bismuth one: the latter

has a 6s2 lone pair, high polarizable, which is of central importance for the

conducting properties of the non-doped phase and, in general, of all δ-Bi2O3

related phases.

The electrical measurements performed on Bi3TaO7 also showed the possible

property of the phase at conducing protons, therefore further investigations in

this direction might be carried out.

References

86

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Documents

Possible electrolytes for SOFC: synthesis and characterization of Bi3–xMexTaO7–δ and others fluorite-related phases