SYNTHESIS AND CHARACTERIZATION OF POLYCRYSTALLINE FERROELECTRIC BULK CERAMICS

Ph. D Thesis: The Islamia University of Bahawalpur, Pakistan

Submitted to: Department of Chemistry

Submitted by: Nasira Sareecha (49/IU.Ph.D/09)

Ph. D (2017)

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In the Name of Almighty Allah the Merciful, the Compassionate

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i

Student declaration

I Mrs. Nasira Sareecha, Reg. No. 49/IU.PH.D/09, hereby declare that I have produced

the work presented in this thesis during the scheduled period of study. I also declare that I

have not taken any material from any source except referred to wherever due that amount

of plagiarism is within acceptable range. If a violation of HEC rules on research has

occurred in this thesis, I shall be liable to punishable action under the plagiarism rules of

the HEC.

Nasira Sareecha Dated: 20.02. 2017

ii

Supervisor declaration certificate

It is certified that Mrs. Nasira Sareecha , 49/IU.Ph.D/09 has carried out all the work

related to this thesis under my supervision at the Department of Chemistry, Islamia

University of the Bahawalpur, Bahawalpur and the work fulfills the requirement for

award of Ph. D degree.

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Dedication

Dedicated to the sublime personality Mohammad (peace be upon him) that received heavenly

mysterious message in the cave of Hira destined to deliver the morally starved and suffering

humanity from the dreadful dungeon of alluring bestiality.

Ceaseless gratitude’s to the torch barriers of the sublime path that are following the foot prints

of Holy prophet (peace be upon him) towards the edifying path of mystical realization and

spiritualization.

Dedicated to my father; incredible gratefulness to whom that with him my boat of life is

dwelling in the right direction.

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Acknowledgement

All praises for Almighty Allah, the merciful and compassionate, the sustainer and

revealer of Holy Scripture to Muhammad (peace be upon him).

I am pleased to acknowledge Higher Education Commission, Islamabad, Pakistan for

awarding indigenous Ph. D fellowship.

I am grateful to my reverend research supervisor Prof. Dr. Muhammad Latif Mirza for

his benign guidance, caring and promising attitude; my Co-supervisor Prof. Dr. Wajid

Ali Shah for gentle, temperate and considerate long term conduct to replenish the aspects

of research work.

I feel like thanking Prof. Dr. Asghar Hashmi, Dean of faculty sciences and Prof. Dr.

Faiz-ul-Hassan Nasim, Chairman of chemistry department for facilitating research

work. I am very much gratified to the professors of physical chemistry for benevolent

attitude and all others in the department.

I feel like expressing incredible gratitude’s to Prof. Dr. Ashraf Tahir, School of

Chemical and Material Engineering (SCME), National university of Science and

Technology, Islamabad behind the all concern. Prof. Dr. Asghari Maqsood’s (Nano

Scale Physics Laboratory, Department of Physics, Air University, PAF Complex E-9,

Islamabad, Pakistan) efforts headed for the work of analysis and valuable support are

acknowledged. I am grateful to Dr. M. Anis-ur-Rehman for providing reassured access

at Applied Thermal Physics Laboratory, Comsats Institute of Information Technology

(CIIT), Islamabad. Prof. Dr. M. Sher (Department of Chemistry , University of

Sargodha) and Dr. Saif-Ullah Awan (Ibn-e-Sina Institute of Technology (ISIT), H-11/4,

Islamabad, Pakistan are acknowledged for the work of analysis.

I do not forget the remembrance of my dear colleagues Mrs. Asia Akhtar and Ms.

Zubia Iram for their placid cooperation. The names of predominantly Awais Siddique

Saleemi (CIIT) and Shahid Ameer (SCME) shall always be educed for their co-

operation throughout.

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Services of M. Boota, Senior Technician and M. Shah Zaman, Furnace Operator, Glass

and Ceramics Research Centre PCSIR Laboratories Complex, Ferozpur Road, Lahore,

Pakistan are acknowledged as well.

Pleasant and pretty serene persistent efforts of my husband, Sameen Ahmad made me

endured towards the completion of work. He solaced me at the all grounds; online

submission of research articles; from early processing to the final stage of publishing. He

very benignly relived my nonexistence at domestic level. The passion for higher studies

is nothing but a source of continuous intellectual nourishment from my father,

Mohammad Suleiman Salama; he always craved and prayed for my succeeding. At this

stage; I recall my deceased mother, Sabira Salama for her passion towards higher

education. Umar Tahir, my niece is acknowledged for long term association and

cooperation. I feel like expressing my paramount biddings to my brothers and sisters,

family members and beloved ones for their prayers.

I am gratified to almighty Allah for sending Humera Warasat to my home for lovely

caring. I am also thankful to all those that united me and helped me at all levels.

NASIRA SAREECHA

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Publications:

(1) Nasira Sareecha *, W. Ali. Shah, M. Anis-ur-Rehman, M. Latif Mirza, M. Saif-

Ullah Awan, “Electrical investigations of BaTiO3 ceramics with Ba/Ti contents

under influence of temperature”, J. Solid State. Ionics. 303 (2017) 16-23.

(2) Nasira Sareecha *, W. Ali. Shah, A. Maqsood, M. Anis-ur-Rehman, M. Latif

Mirza, “Fabrication and electrical investigations of Pb-doped BaTiO3 ceramics” J.

Mat. Chem. Phys. 193(2017) 42-49.

(3) Nasira Sareecha *, W. Ali. Shah, M. Latif Mirza, Electrical investigations of Bi-

doped BaTiO3 ceramics under influence of temperature, under Review J. Solid

State. Chem. (2017)

(4) Nasira Sareecha *, W. Ali. Shah, M. Latif Mirza, Solid state sintering of PbTiO3

ceramics with Pb/Ti contents to be submitted to J. Solid State. Ionics.

(5) Nasira Sareecha *, W. Ali. Shah, M. Latif Mirza, Electrical investigations of

PbTiO3 ceramics with Pb/Ti contents as a function of temperature, to be submitted

to the Ceram. Int.

Additional Publications

(1) Javeed Iqbal, M. Latif Mirza and Nasira Sareecha, Cation exchange separation of

transition metals and calcium with Zirconium Phosphate, J. Chem. Soc. Pak. Vol.

16, No.1, 1994 ( M. Sc Dissertation).

(2) Abdul Majeed, M. Mansoor Mustafa, R.N. Asma and Nasira Sareecha,

Spectrophotometric determination of Copper with Ascorbic Acid, J. Chem. Soc.

Pak. Vol. 18, No.2, 1996.

(3) Abdul Majeed, Munir Ahmad, M.S. Khan, R.N. Asma and Nasira Sareecha,

Spectrophotometric determination of Vanadium with Ascorbic Acid as a

chromogenic reagent, J. Pure and Appl. Sciences. Vol. 15, No.1, 1996.

Presentation in Conferences

(1) N. Sareecha, W.A. Shah, M.L. Mirza, Paper presentation in1st International

conference on Mathematics and Physics, “Temperature dependent electrical

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properties of BaTiO3 ceramics with Ba/Ti contents” Air University, E-9 PAF

Complex, Islamabad, Pakistan, Feb. 14-16, 2017: Got 1st Prize by Chairman

Higher Education Commission (HEC), Islamabad dated: Feb. 16, 2017

(2) N. Sareecha, M.L. Mirza, Paper presentation in 8th International and 20th National

Chemistry conference, “Biosorption of Crystal Violet from aqueous solution on

Tamarix Aphilla”, Department of Chemistry, Quaid-e-Azam University

Islamabad, Pakistan Feb.15-17, 2010 (M. Phil Dissertation).

(3) N. Sareecha, J. Iqbal, Paper Presentation in the 4 th National Conference of

Chemistry, “Cation exchange separation of transition metals and calcium with

Zirconium Phosphate”, Department of Chemistry, Islamia University of the

Bahawalpur, Bahawalpur, Pakistan Dec.21-24, 1992 ( M. Sc Dissertation).

viii

List of Abbreviations:

BT Barium titanate, BaTiO3

PT Lead titanate PbTiO3

BiT Bismuth titanate, BiTiO3

TC Curie temperature

MLCC Multilayer ceramic capacitors

ABO3 Perovskites

BLSF Bismuth layer structured ferroelectrics

CNO Cadmium niobate ,Cd2Nb2O7

PBT Lead doped Barium titanate, Pb: BaTiO3

BBT Bismuth doped Barium titanate, Bi: BaTiO3

BLT Bismuth lithium titanate, Bi0.5 Li0.5 TiO3

BNT Bismuth sodium titanate Bi0.5 Na0.5 TiO3

BNT Bismuth potassium titanate, Bi0.5 K0.5 TiO3

NTCR Negative temperature coefficient of resistivity

PTCR Positive temperature coefficient of conductivity

MW Microwave

SPS Spark Plasma sintering

V''Pb Lead vacancies

V o⦁⦁ Oxygen vacancies

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Ea Activation energy

Abstract

BaTiO3 and PbTiO3 ceramics were prepared through solid state sintering reaction. The

studies were attempted at the phase pure and crack free preparation and electrical

investigations of BaTiO3 ceramics with Ba/Ti molar ratio (0.98 and 0.94) and PbTiO3

ceramics (1.00, 0.98 and 0.94) Pb/Ti molar ratio in the wide range of temperatures (40–

700°C) at 1kHz frequency of the ceramics perhaps for the first time. Studies were

attempted to find the understanding of the conduction process and useful implementation

of the controlling parameters. Thermogravimetric and Differential scanning calorimetric

analysis (TGA-DSC) revealed melting temperatures, weight losses and variations in the

enthalpy of crystallization of the as ground powders. Ceramics with all precursor

composition were perovskite (ABO3), ferroelectric materials. Cubic structures (Pm-3m)

and tetragonal (P4mmm, P4mm, P4MM) crystal structures were indicated. Curie

temperature (Tc) increased from120-130°C with decreasing Ba/Ti contents. With Pb and

Bi doping, Curie temperature (TC) was shifted from 120 -200°C and 120 -160°C

respectively. Pb/Ti contents did shift the Curie temperature. Broad dielectric constant

peaks and pronounced dielectric anomalies with relaxor like behavior were observed in

the paraelectric regions. Resistivity decreased with increasing temperature, all specimens

showed semiconductor behavior with negative temperature coefficient of resistivity

(NTCR) characteristics. Mobility of electrons increased with thermal activation due to

hopping of charge carriers from one site to another. Ohmic conductivities and associated

activation energies were evaluated by impedance spectroscopy. Conductivity followed

Arrhenius Law with Ea values lying in the range of single ionized and doubly ionized

oxygen vacancies and Pb vacancies; ionic conduction was supposed to be responsible.

Well defined hysteresis P-E loops under electric fields showed ferroelectric

characteristics.

Key Words:

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Ferroelectric materials; Sintering regime; Curie temperature; Perovskites; Paraelectric

regions; Ionic conduction.

Table of contents:

Chapter 1: Introduction 1

1.1 History of ferroelectricity 1

1.2 Ferroelectricity and domains 2

1.3 Ferroelectric Materials

1.3.1 Ferroelectric hysteresis loop 5

1.4 Categories of Ferroelectric materials 6

1.4.1 Bismuth layer group 6

1.4.2 Tungsten bronze ceramics 7

1.4.3 Pyrochlore group 8

1.4.4 Perovskite group 9

1.4.4.1 BaTiO3 9

1.4.4.1.1 Structural phase transitions in BaTiO3 10

1.4.4.1.2 Dielectric and piezoelectric properties of BaTiO3 12

1.4.4.2 PbTiO3 12

1.4.4.2.1 Anisotropy and solid state sintering of PbTiO3 13

1.4.5 Defects in perovskite structure 13

1.5 Preparation methods for polycrystalline BaTiO3 and PbTiO3 ceramics:

Merits and Demerits 14

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1.5.1 Co-precipitation method 14

1.5.2 Sol- gel method 15

1.5.3 Hydrothermal method 15

1.5.3.1 Disadvantages of Hydrothermal method 16

1.5.4 Solid state reaction method 16

1.5.4.1 Sintering process 16

1.5.4.2 Advantages of Solid state sintering method over Hydro-thermal method 18

Chapter 2: Literature review 19

2.1 Objectives of the research work 26

2.2 Scope of the work 28

Chapter 3: Experimental work 29

3.1 Preparation of BaTiO3 and PbTiO3 ceramics 29

3.1.1 Preparation equipment and source of errors 29

3.1.2 Preparation of BaTiO3 ceramics 29

3.1.3 Preparation of Pb- doped BaTiO3 ceramics 29

3.1.4 Preparation of Bi- doped BaTiO3 ceramic 30

3.2 Preparation of PbTiO3 ceramics 30

3.3 Experimental parameters and their influence 31

3.4 Characterization Techniques 31

3.4.1 Thermal Characterization 31

3.4.2 Structural characterization 33

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3.4.2.1 X-ray diffraction method 33

3.4.2.1.1 Crystallite size 34

3.4.2.1.2 Measured Density 34

3.4.2.1.3 X-Ray density 35

3.4.2.1.4 Porosity 35

3.4.3 Scanning electron microscopy 35

3.4.4 Electrical characterization 37

3.4.4.1 Electrical ac measurements 37

3.4.4.2 Electrical dc measurements 41

3.4.4.2.1 Electrical dc resistivity 41

3.4.4.2.2 Drift mobility 42

3.5 Electric polarization 43

Chapter 4: Results and discussions 45

4.1. Thermal analysis of BaTiO3 ceramics 45

4.1.1 Thermal analysis of PbTiO3 ceramics 47

4.2. Structural analysis 49

4.2.1. Preparation analysis and structural properties of BaTiO3ceramics 49

4.2.2. Preparation analysis and structural properties of Pb-doped BaTiO3 ceramics 53

4.2.3 Preparation analysis and structural properties of Bi-doped BaTiO3 ceramics 55

4.2.4 Preparation analysis and structural properties of PbTiO3 ceramics 57

4.3 Microstructural analysis 62

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4.3.1 Microstructural analysis of BaTiO3 ceramics 62

4.3.2 Microstructural analysis of Pb-doped BaTiO3ceramics 63

4.3.3 Microstructural analysis of Bi-doped BaTiO3 ceramics 64

4.3.4 Microstructural analysis of PbTiO3 ceramics 65

4.3.4.1 Microstructural analysis of PbTiO3 (1.00) ceramics 65

4.3.4.2 Microstructural analysis of PbTiO3 (0.98) ceramics 67

4.3.4.3 Microstructural analysis of PbTiO3 (0.94) ceramics 68

4.4 Electrical properties 70

4.4.1. Dielectric studies of BaTiO3 ceramics 70

4.4.1.1 Ac conductivity studies 74

4.4.2 Dielectric studies of Pb-doped BaTiO3 ceramics 76

4.4.2.1 Ac conductivity studies of Pb-doped BaTiO3 ceramics 78

4.4.3 Dielectric studies of Bi-doped BaTiO3 ceramics 79

4.4.3.1 Ac conductivity studies of Bi-doped BaTiO3 ceramics 82

4.4.4 Dielectric studies of PbTiO3ceramics 84

4.5 Electrical dc resistivity and dc conductivity studies of BaTiO3 ceramics 89

4.5.1 Dc mobility studies of BaTiO3 ceramics 93

4.5.2 Electrical dc resistivity and dc conductivity studies of Pb-doped BaTiO3

ceramics 94

4.5.2.1 Dc mobility studies of Pb- BaTiO3ceramics 97

4.5.3 Electrical dc resistivity and dc conductivity studies of Bi-doped BaTiO3 98

xiv

4.5.3.1 Dc mobility studies of Bi-doped BaTiO3 ceramics 100

4.6. Electrical dc resistivity and dc conductivity studies of PbTiO3 ceramics 100

4.7. Electric polarization studies 104

4.7.1 Electric polarization studies for BaTiO3 ceramics 104

4.7.2 Electric polarization studies for Pb-doped BaTiO3ceramics 105

4.7.3 Electric polarization studies for Bi-doped BaTiO3 ceramics 106

4.8 Electric polarization studies for PbTiO3ceramics 107

Conclusions 109

References 111

xv

List of Figures:

Fig. 1.1 Phase transition of BaTiO3 from a cubic paraelectric structure to ferroelectric

tetragonal structure.

Fig. 1.2 The perovskite structure of PbTiO3, having a cubic structure in the paraelectric

phase and tetragonal structure in the ferroelectric phase.

Fig. 1.3 Illustration of P-E hysteresis loop in ferroelectrics

Fig. 1.4 Illustration of Bismuth layered Perovskite structure.

Fig. 1.5 The perovskite structure of Tungsten bronze group.

Fig. 1.6 Illustration of Pyrochlore group structure.

Fig. 1.7 The cubic perovskite type structure ABO3 positions.

Fig. 1.8 Phase transition, spontaneous polarization vectors and unit cell dimensions with

temperature in BaTiO3 crystals.

Fig. 1.9 BaO-TiO2 equilibrium diagram of Rase and Roy

Fig. 3.1 TGA and DSC curve in air for BaCO3, TiO2, and ZrO2 powder mixture.

Fig. 3.2 (a) Process of characteristics X-rays from an atom (b) EDS spectrum.

Fig. 3.3 Relaxation time (τ ) of masses after turning off the applied electric field (b)

Impedance plane of real capacitor (c) Response of real ε 'and imaginaryε ' 'parts of

permittivity.

xvi

Fig. 3.4 ε 'and ε ' ' as a function of frequency with associated different polarization

mechanisms.

Fig. 3.5 Hysteresis loop of typical ferroelectric material at room temperature.

Fig. 3.6 (a) Schematic of Sawyer and Tower method (b) and Block diagram for P-E

measuring system

Fig. 4.1 (a and b), TGA and DSC thermograms of the as ground BaTiO3 specimens, 5a

(0.94) and 5b (0.98)

Fig. 4.1.1 (a, b and c), TGA and DSC thermograms of the as ground PT powders, a

(1.00), b (0.98) and c (0.94).

Fig. 4.2.1.1 XRD Patterns of BaTiO3 ceramics (0.94) sintered at 1300°C/2h, pre fired at

1100 - 1200°C/4h (a- b), 1300°C/1-4h (c- f).

Fig. 4.2.1.2 XRD Patterns of BaTiO3 ceramics (0.98) sintered at 1300°C/2h, pre fired at

1300°C/1-4h (a - d).

Fig. 4.2.1.3. Tetragonality (A) and crystallite size inset (B) of BaTiO3 ceramics sintered

at 1300°C/2h, pre fired at their respective temperatures (a and b) 1300°C/1–4h and (c) at

1100 - 1300°C/4h.

Fig. 4.2.1.4 Density of BaTiO3 ceramics (A) sintered at 1300°C/2h, pre fired at the

temperatures (a and b) 1300°C/1–4h, Inset B shows the powders (a) pre fired at 1100 -

1300°C/4h.

Fig. 4.2.2.1 XRD Patterns of pure and Pb-doped BaTiO3 ceramics sintered at 1200°C/2h.

Fig. 4.2.2.2 Tetragonality and crystallite size of pure and Pb-doped BaTiO3 ceramics

Fig.4.2.3.1 XRD Patterns of the pure and doped BaTiO3 ceramics sintered at 1150°C/2h,

inset illustrates the diffraction peaks of BaBi4Ti4O15 at 2θ= 27.490° and 30.186° in

magnified version at 0.100 mole % doping

Fig.4.2.3.2 Crystallite size and density of the pure and doped BaTiO3 ceramics sintered

at 1150°C/2h.

xvii

Fig. 4.2.4.1 XRD Patterns of PbTiO3 ceramics sintered at 1190°C/1h; (1.0, a- b) pre-

fired1000 - 1100°C /2h and (0.94, c-e pre fired at 1000 - 1190°C± 5C /2h, 1000 - 1000 -

1190°C± 5C /2h respectively

Fig. 4.2.4.2 XRD Patterns of PbTiO3 ceramics (0.98) sintered at 1190°C± 5C /1h, pre

fired at 1190°C/1-4h (a-d) –and 1000- 1190°C/2h respectively

Fig. 4.2.4.3 Tetragonality of PbTiO3 ceramics sintered at 1190°C± 5C /1h with precursor

compositions Vs pre-firing temperature, inset A shows teragonality of powders (0.98) Vs

pre-firing duration, 1190°C± 5C /1-4h.

Fig. 4.2.4.4 Density of PT ceramics sintered at 1190°C± 5C /1h, pre fired at the

temperatures1000-1190°C ± 5C /1h, inset A shows the powders (0.98) pre fired at

1100°C 1-4h.

Fig. 4.3.1 SEM Micrographs of BT powders and pallets. Powders a and b (0.94) were pre

fired at 1100°C-1200°C/4h , BT pallets c (0.94) and d (0.98) at 1300°C/4h. Inset shows

magnified version of grains at 1µm, headed arrows reveal the fractured surfaces.

Fig. 4.3.2 SEM Micrographs of PBT pallets sintered at 1200°C/2h, inset at (a) shows

undoped BT.

Fig. 4.3.3 SEM Micrographs of the pure and doped BT pallets sintered at 1150°C/2h,

inset at micrographa (a) shows undoped BaTiO3

Fig. 4.3.4.1 FE-SEM micrographs for PT (1.00) pallets sintered at 1190± 5C, pre-fired at

1100C/2h.

Fig. 4.3.4.1.1 EDS spectrum of the PT (1.00) pallets sintered at 1190± 5C, pre-fired at

1100C/2h.

Fig. 4.3.4.2 FE-SEM micrographs for PT (1.00) pallets sintered at 1190± 5C, pre-fired at

1100C/2h; inset at micrograph (b) shows PT ceramics at magnified version.

xviii

Fig. 4.3.4.2.1 EDS spectrum of the PT (1.00) pallets sintered at 1190± 5C, pre-fired at

1100C/2h.

Fig. 4.3.4.3 FE-SEM micrographs for PT (0.94) pallets sintered at 1190± 5C, pre-fired at

1100C/2h.

Fig. 4.3.4.3.1 EDS spectrum of the PT (1.00) pallets sintered at 1190± 5C, pre-fired at

1100C/2h.

Fig. 4.4.1. Temperature dependence of dielectric constant (ԑ΄) and loss tangent (δ) of the

BaTiO3 ceramics C (f and d) at 1k Hz frequency. Insets (A and B) show their pre firing

temperature Vs time.

Fig. 4.4.1.1 Variance of ac conductivity for BaTiO3 ceramics (C) with temperature at 1k

Hz frequency. Insets (A and B) describe pre firing temperature of specimens Vs time.

Fig. 4.4.2.1 Temperature dependence of dielectric constant (ԑ΄) and dielectric loss tangent

(δ) for pure and Pb-doped BaTiO3 ceramics at 1k Hz frequency.

Fig. 4.4.2.1.1. Variance of ac conductivity for pure and Pb-doped BaTiO3 ceramics with

temperature at 1k Hz frequency. Inset A shows magnified version of the selected area

confirming the phase transition in accordance with dielectric studies.

Fig.4.4.3 Temperature dependence of dielectric constant (ԑ΄) for the pure and Bi doped

BT ceramics at 1k Hz frequency; inset A, shows the enlarged view of the selected area

corresponding to phase transition.

Fig. 4.4.3.1. Loss tangent, tan (δ) of the BBT ceramics as a function of temperature at 1k

Hz frequency

Fig. 4.4.3.1.1 Conductivity of the pure and doped BaTiO3 ceramics Vs temperature at 1k

Hz frequency. Inset A, the magnified version of the selected area confirms the phase

transition in agreement with dielectric studies.

Fig. 4.5.1.2 ( a) Variation in conductivity (σ dc) with inverse of absolute temperature for

BT ceramics in the temperature range200- 360°C.

Fig. 4.5.1.3 Variation in conductivity (σ dc) with inverse of absolute temperature for BT

ceramics in the temperature range 480- 700°C.

xix

Fig. 4.5.2.1 Variation in conductivity (σ dc) with inverse of absolute temperature for

pure and doped PBT ceramics

4.5.2.1 Dependence of drift mobility with inverse of temperature for PBT ceramics.

Fig. 4.5.3.1 Resistivity of the pure and BBT ceramics Vs temperature in the ferroelectric

and paraelectric regions at 1k Hz frequency.

Fig. 4.5.3.2 Variation in conductivity (σ dc) with inverse of absolute temperature for the

pure and doped BBT ceramics

Fig. 4.6.1 Resistivity of the pure and PT (1.00) ceramics Vs temperature in the

ferroelectric and paraelectric regions at 1k Hz frequency, inset A shows variation in

conductivity (σ dc) with inverse of absolute temperature.

Fig. 4.6.2 Resistivity of the pure and PT (0.98) ceramics Vs temperature in the

ferroelectric and paraelectric regions at 1k Hz frequency, inset A describes variation in

conductivity (σ dc) with inverse of absolute temperature.

Fig. 4.6.3 Resistivity of PT (0.98) ceramics Vs sintering duration in the ferroelectric and

paraelectric regions at 1k Hz frequency, inset A shows variance in conductivity (σ dc) with

inverse of absolute temperature.

Fig. 4.6.4 Resistivity of PT (0.94) ceramics Vs temperature in the ferroelectric and

paraelectric regions at 1k Hz frequency, inset A displays variation in conductivity (σ dc)

with inverse of absolute temperature.

4.7.1 P-E loops of BT ceramics sintered at 1300°C/2h, pre fired at 1300°C/4h.

Fig. 4.7.2 P-E loops of PBT ceramics sintered at 1200°C/2h, inset (B) shows P-E loops

for undoped BaTiO3.

Fig. 4.7.3 P-E loops of BBT ceramics sintered at 1150°C/2h, inset B shows P-E loop for

undoped BaTiO3.

Fig. 4.8 P-E loops of PT ceramics (1.00, 0.98, 0.94) sintered at 1190°C± 5° C /1h.

xx

Chapter 1

Introduction

Ferroelectric perovskite (ABO3) are also very important in material physics and earth

sciences because they exhibit excellent physical properties which make them special

candidates for a wide application range in the electronic industry (Lines and Glass, 1977).

Their phase transition strongly affects their physical and chemical properties.

For the perovskite structure-type systems (ABO3), the partial substitution of A- or B- site

ions (carrier doping) promotes the activation of several conduction mechanisms. In the

last few years, extensive studies have been carried out A- or B- site (Buscaglia et al.,

2000; Kingery et al., 1975). Substitution of different elements having dissimilar

electronic configuration can lead to dramatic effects associated with the electronic

configuration mismatch between the ions located at the same A- or B- site.

1.1 History of ferroelectricityThe discovery of ferroelectricity (Haertling, 1999) extends as far back as the mid-1600s

when Rochelle Salt (sodium potassium tartrate tetra hydrate) was prepared by Seignette

in La Rochelle, France, for medicinal purposes. Approximately 200 years later, this water

soluble crystalline material was investigated for pyro-electric (thermal polar) properties;

after another half century, its piezoelectric (stress polar) properties remained uncovered,

after another 40 years Joseph Valasek (Valasek, 1921) discovered ferroelectricity in this

material. Since then, many other robust ferroelectric oxides, including lead titanate

(PbTiO3) (Shirane et al., 1950a) bismuth titanate (Bi4Ti3O12) (Subbarao, 1961; Uitert and

Egerton, 1961) lithium niobate (LiNbO3) and lithium tantalate (LiTaO3) (Matthias and

Remeika, 1949) and so forth, have been discovered, rendering the theoretical and

applicational studies on ferroelectrics being one of the most active research areas in the

field of solid state materials.

1

1.2 Ferroelectricity and domains

Ferroelectric materials possess a spontaneous electric polarization that can be reversed by

the application of an external electric field (Lines and Glass, 1977). The term is used in

analogy to ferromagnetism, in which a material exhibits a permanent magnetic moment

(Taganstev et al., 2010). A ferroelectric is defined as a crystal, in which some accessible range of environmental

conditions has two or more equilibrium intrinsic lattice electric polarization states in the

absence of an electric field and in which the spontaneous intrinsic lattice electric

polarization can be switched between those orientations by a realizable , appropriately

oriented electric fields (Damjanovic, 1998; Dawber et al., 2005)

According to the crystalline symmetry with respect to a point, there are 21 non-

centrosymmetric classes. However, only 10 classes among these 21 classes have a unique

polar axis. Crystals belonging to these classes are called polar since they possess a

spontaneous polarization, Ps (Nye, 1985). In general, Ps decreases with increasing

temperature to disappear either continuously or discontinuously at a Curie point, TC.

Usually the phase above TC is termed as paraelectric phase. In general, Ps decreases with

increasing temperature to disappear either continuously or discontinuously at a Curie

point, TC. Usually the phase above TC is termed as paraelectric phase.

In ferroelectrics below TC and in the absence of external electric field, a Ps develops at

least along two directions. In order to minimize the energy of depolarizing fields, crystal

splits into polar regions along each of these directions. Each volume of uniform

polarization is called a domain. Crystallographically, domain structure is a type of

twinning (Arlt, 1990). The resulting domain structures usually results in a nearly

complete compensation of macroscopic polarization. The polarization directions of

domains are basically high temperature symmetry axes, such as (001), (110) or (111).

Angles between the dipoles of adjacent domains are those between such symmetry axes

are for example 90, 180

and 71

respectively. The boundaries separating domains are

referred to as domain walls. In ferroelectrics, the domain walls could be shifted or

switched by an external electric field. The polarization- electric field relationship in

2

ferroelectrics is often nonlinear and hysteretic, due to the domain wall motion and

switching.

Barium titanate is such a material that shows polarization in the tetragonal phase.

However, in the cubic phase the central titanium atom serves as an inversion center; the

polarization is not possible. Only with the occurrence of a tetragonal deformation, where

the positively charged barium and titanium ions are displaced with respect to the six

negatively charged oxygen ions, a polar axis is formed in the direction of the tetragonal

deformation, which marks the direction of spontaneous polarization.

PS = 0 PS ≠ 0

Cubic paraelectric phase Tetragonal ferroelectric phase

Fig. 1.1 Phase transition of BaTiO3 from a cubic paraelectric structure to ferroelectric

tetragonal structure (Collaboration: Authors and editors of the volumes, 2000).

The spontaneous electric polarization, PS, in a ferroelectric is an intrinsic lattice

polarization resulting from a spontaneously formed electric dipole moment, and its

orientations are determined by the crystal structure. On cooling from high temperatures,

ferroelectrics usually exhibit a structural phase transition from a paraelectric (cubic)

phase into a ferroelectric (tetragonal) phase. The symmetry of ferroelectric phase is

always less than the paraelectric phase.

PbTiO3 possesses the highest room temperature spontaneous as high as PS = 81 μCcm-2

(Tuttle et al., 1980; Venevtsev et al., 1959) among the perovskite ferroelectrics which is

caused by high tetragonal distortion and ionic displacements.

3

In ABO3 ferroelectric perovskite such as BaTiO3 and KNbO3, the anisotropic distortion of

tetragonal phase from cubic phase arises from the covalent nature of B–O bonds (Ti-O

bonds). Whereas covalence of Pb-O in tetragonal PbTiO3 i-e the 6s2 lone-pair effect from

Pb2+ gives rise to additional distortion (Kuroiwa et al., 2001), which results in high

spontaneous polarization in the tetragonal phase in comparison to BaTiO3.There are two

distinct atomic positions for oxygen in the crystal structure of tetragonal PbTiO3. The

Pb–O2 bonds (parallel to the c-axis) are covalent, whereas the Pb–O1 bonds (normal to the

c-axis) are ionic.

Fig. 1.2 The perovskite structure of PbTiO3, having a cubic structure in the paraelectric

phase and tetragonal structure in the ferroelectric phase (Damjanovic, 1998).

1.3 Ferroelectric materials Ferroelectric materials are a subgroup of spontaneously polarized pyroelectric crystals,

and are characterized by the presence of a spontaneous polarization. This polarization is

reversible under the application of an electric field of magnitude less than the electric

breakdown of the material itself (Jaffe et al., 1971; Rabe et al., 2007).

Most ferroelectric materials undergo as structural phase transition from a high

temperature paraelectric phase to a low temperature ferroelectric phase. The temperature

at which phase transition occurs is known as Curie temperature (TC). In the ferroelectric

phase, the displacement of the central B ion on the application of an electric field to the

unit cell causes the reversal of polarization. The areas with the same polarization

orientation are referred to as domains, with domain walls existing between areas of unlike

4

polarization orientation. The switching of many adjacent unit cells is referred to as

domain reorientation or domain switching. The ionic movement leads to a macroscopic

change in the dimensions of the unit cell and the ceramic as a whole.

In ferroelectric ceramics, domains are randomly oriented and the net polarization is zero

because of their cancellation effect. Therefore the as prepared ceramics are neither

piezoelectric nor pyroelectric. Polycrystalline ferroelectric ceramics must be poled at a

strong DC electric field (1-10 KV/mm). The domains can be made easily switchable at

elevated temperatures.

1.3.1 Ferroelectric hysteresis loopThe peculiar feature of ferroelectrics is the spontaneous polarization that can be

reoriented by an electric field (Lines and Glass, 1977).

Fig. 1.3 Illustration of P-E hysteresis loop in ferroelectrics (Vijatović et al., 2008 ).

Spontaneously polarized regions, with a single direction of polarization, are called

domains. Orientation relationships between domains are directed by the crystal

symmetry. Ferroelectricity is displayed by a ferroelectric hysteresis loop (Fig.1), a plot

polarization versus electric field.

5

On the application of electric field, dipoles which are already oriented in the direction of

the field will remain so aligned, but those which are oriented in the opposite direction

will show a tendency to reverse their orientation.

An applied electric field with amplitude above Ec is needed to reverse the ferroelectric

polarization. In a real ferroelectric material P-E hysteresis loop is formed in the switching

process through the growth of existing domains antiparallel to the applied field by

domain-wall motion, or through the nucleation and growth of new antiparallel domains

(Lines and Glass, 1977) . Consequently, domains are favorably oriented with respect to

the applied electric field. The resulting electric polarization at zero applied electric field

is defined as the remnant polarization Pr, which is always smaller than PS because of the

existence of domains and back-switching. The polarization obtained at the maximum

applied field is the maximum polarization, PS.

1.4 Categories of Ferroelectric materialsFerroelectric materials are divided into four categories: the bismuth layer group, the

tungsten bronze group and the pyrochlore group and the most important perovskite group

(ABO3).

1.4.1 Bismuth layer groupThe bismuth layer structured ferroelectric (BLSF) compounds were first studied by

Aurivillius; belong to the family of bismuth titanate (Bi4Ti3O11: BiT) (Isupov, 2006).They

possess pseudo perovskite layers (An-1BnO3n+1)2- stacked between (Bi2O2) 2+ layers, only

BiT possess monolclinic structure.

6

Fig. 1.4 Illustration of Bismuth layered Perovskite structure (Bhalla et al., 2000).

Aurivillius phases are described by the general formula (Bi2O2)2+ - ((An-1BnO3n+1)2-, An-1, B

is 12 coordinated monovalent, divalent or trivalent cation. B is octahedrally co-ordinated

quadric, penta or hexavalent metal ion, n is an integer representing the number of

perovskite layers that can range from 1 to 8 (Shulman et al., 1996). They have high Curie

temperature (980°C) and good piezoelectric properties. There are some critical issues that

concern the processing like reproducibility of the properties and narrow range of sintering

temperatures. Mixed Aurivillius phases have potentially enhanced properties (Sanson and

Whatmore, 2005).

1.4.2 Tungsten bronze ceramics They have a general formula of (A1)4 (A2)2 (C)4(B1)2(B2)8O30. B-type cations occupy

A1, A2 and C sites. C-type cations occupy the B1 and B2 octahedral sites (Jamieson et

al., 1968) . In the formula, A1, A2, C and B are 15-, 12-, 9-, and 6-fold co-ordinated sites

in the crystal lattice structure. Generally, A1,A2 sites can be filled by Na, Li, K, Ca, Sr,

Ba, Pb, Bi and some rare earth Sm, Nd, Dy and Ce cations. The smallest interstice C is

often empty, and the formula for the filled tungsten bronze structure is A6B10O30.

7

Fig. 1.5 The perovskite structure of Tungsten bronze group (Bhalla et al., 2000)

The metal cations distribution in the different sites plays a crucial role in tailoring

physical and functional properties. Tungsten bronze ferroelectric ceramics show electro-

optic, non-linear optic, piezoelectric and pyroelectric properties. Coupling of the most

important members of this family like barium sodium niobate (BNN), potassium

lanthanum niobate (KLN), strontium sodium niobate (SNN) leads to the modification of

the properties of tungsten bronz ceramics.

1.4.3 Pyrochlore groupThe ceramics belonging to pyrochlore group are represented by the stoichiometry

A2B2O7. A and B sites are tetravalent or pentavalent, trivalent or divalent species

respectively (Weller et al., 2004).

Fig. 1.6 Illustration of Pyrochlore group structure (Bhalla et al., 2000).

The pyrochlore structure is commonly composed of two interpenetrating networks

without common constituents. Cadmium pyroniobate Cd2Nb2O7 (CNO) belongs to this

8

group, it is ferroelectric at low temperatures and exhibits three dielecric anomalies in the

narrow temperature range from 195-205K, above which it is cubic (Ang et al., 2004). The

ferroelectric behavior disappears above 185K; at the same temperature, anomalies in

dielectric constant and specific heat are exhibited (Fischer et al., 2008). The frequency

dependence of the dielectric constant for the pyrochlore group in this temperature regime

is similar to that of typical relaxor materials that indicates the presence of polar clusters

in CNO.

1.4.4 Perovskite group

1.4.4.1 BaTiO3

Among ferroelectric materials, the perovskites (ABO3) are by far the most important and

the widely studied, well known BaTiO3 and PbTiO3 belong to this category. BaTiO3 takes

its name from the mineral perovskite CaTiO3. Perovskite crystals are represented by the

general formula ABO3. The perovskite structure (ABO3) is the most wonderful structure.

A and B are metallic cations and O, an ion usually oxygen.

(a) (b)

Fig. 1.7 The cubic perovskite type structure ABO3 positions (Vijatovic et al., 2008)

This structure forms a network of corner-linked oxygen octahedra, with the smaller

cation filling the octahedral holes and the large cation filling the dodecahedral holes.

Through this model, the paraelectric-ferroelectric (P-F) and ferroelectric-ferroelectric (F-

F) phase transitions could be described as the distortion of the unit cell. All cations and

anions may move with respect to the equilibrium position in the cubic perovskite unit

cell. Substitution at A-sites results in the large family of simple perovskite ferroelectrics.

9

In many ferroelectrics ceramics, the families of BaTiO3 and Pb-based solid, Ti4+ and Zr4+

occupy the B-site while Pb 2+ and Ba2+ ions occupy the A-site (Bhalla et al., 2000). The

perovskite structure can be considered as a three dimensional frame work of BO6

octahedra (Fig.1a) but it can be regarded as a cubic close packed arrangement of A and O

ions, with the B ions filling the interstitial positions (Fig.1.4b).

The unit cell of the cubic perovkite type lattice is shown in Fig. 1.4b. It can be detected

that the co-ordination number of cation A is 12 and that of B is 6.Variations of the corner

linked like tilt or rotation results in new families of the ferroelectrics like Tungsten bronz

and bilayer structures.

1.4.4.1.1 Structural phase transitions in BaTiO3

BaTiO3 important is the first discovered, the extensively studied and the most widely

used simple ferroelectric oxide. It has an ideal cubic perovskite structure above the Curie

temperature (TC = 120C) and undergoes three successive ferroelectric phase transitions

on lowering the temperature.

By decreasing the temperature, the cubic paraelectric phase could transfer into other

phases. Tetragonal, orthorhombic, and rhombohedral phases are the three symmetries

which frequently occur. At TC BaTiO3 has a paraelectric to ferroelectric phase transition

from the cubic to a tetragonal room temperature ferroelectric phase which is stable from -

5 to 120C. In tetragonal phase, the cubic unit cell of the perovskite structure elongates

along the c axis, i.e. the [001] direction, and results in the a = b < c (Fig.1.7). BaTiO3 has

two ferroelectric to ferroelectric phase transition at -5C from tetragonal to orthorhombic

and at -90C from orthorhombic to rhombohedral. The orthorhombic distortion occurs

between -90 to -5C; in orthorhombic phase the unit cell elongates along a face diagonal

(the [110] direction); whereas a = c > b and b, which is the angle between the a axis and c

axis, is < 90. The rhombohedral structure is stable below -90C (Jona and Shirane, 1993;

Koelzynski and Tkacz-Smiech, 2005). In rhombohedral phase, the unit cell distorts along

the [111] direction with a = b = c and b < 90. In each phase, the dipole is generated by

the displacement of the B-site ion along the same direction of the distortion and the Ps in

those phases is parallel to the [001], [110], and [111] direction, respectively.

10

Most ferroelectric materials exhibit a paraelectric-ferroelectric phase transition at TC. But

it is not necessary for all the ferroelectric materials to experience one or several

ferroelectric ferroelectric phase transitions below the TC. For example, PbTiO3 keeps its

tetragonal symmetry below its TC (490C), whereas BaTiO3 transfers from the cubic

paraelectric phase, through tetragonal ferroelectric phase and orthorhombic ferroelectric

phase.

Fig. 1.8 Phase transition, spontaneous polarization vectors and unit cell dimensions with

temperature in BaTiO3 crystals (Collaboration: Authors and editors of the volumes,

2000).

11

1.4.4.1.2 Dielectric and piezoelectric properties of BaTiO3

BaTiO3 is the first ferroelectric ceramic (Guo et al., 2006). which is widely used in the

electronic industry due to its excellent dielectric, ferroelectric and piezoelectric

properties.

As a dielectric material, it is used for manufacturing dielectric ceramics capacitors,

multilayer capacitors due to its high dielectric constant and low dielectric loss. The values

of the dielectric constant depend on the synthesis route; where purity, density, and grain

size play an important role (Guo et al., 2006). Dielectric constant depends on

temperature, frequency and dopants. Fig. 1.7 shows the temperature dependence of the

dielectric constant. Frequency and temperature dependence of dielectric constant are

discussed in details in the characterization section and in the results discussion sections.

BaTiO3 is also broadly used for its strong piezoelectric characteristics. The word

“piezoelectricity” is derived from the Greek “piezein”, which means to squeeze or press;

hence, piezoelectricity is the generation of electricity as a result of mechanical pressure.

A necessary condition for piezoelectricity to exist is noncentrosymmetry in the crystal.

Two effects are operative in piezoelectric crystals, in general, and in ferroelectric

ceramics, in particular. The direct effect (designed as a generator) is identified with the

phenomenon whereby electrical charge (polarization) is generated from a mechanical

stress, whereas the converse effect (designated as a motor) is associated with the

mechanical movement generated by the application of an electric field. The piezoelectric

properties of ferroelectric ceramics are characterized by kP, k33, d33, d31 and g33. The

piezoelectric efficiency is measured in terms of a coupling coefficients (kP, k33),

indicating the friction of applied mechanical force converted into an electric voltage.

1.4.4.2 PbTiO3

In the search of new perovskite type ferroelectrics of the general formula ABO3 following

the discovery of barium titanate, Shirane, Hoshino and Suzuki (Shirane et al., 1950b;

Shirane et al., 1950a) studied the lead titanate (PbTiO3) ceramic and reported its

ferroelectricity on the basis of the structural analogy between both compositions.

12

Smolenskii reported (Smolenskii and Fiz., 1950) similar study. Subsequently, Shirane

and Hoshino (Shirane et al., 1950a) reported data on a structural phase transition that

taking place at 490C above which structure was cubic perovskite and below which it was

tetragonally distorted perovskite. PbTiO3 has unique properties like high transition

temperature (TC), low dielectric constant, and low ratio for the planner to thickness

coupling factor and a low aging rate of dielectric constant. In addition, PbTiO3 ceramics

are good candidates as piezoelectric and pyroelectric devices for high temperature or high

frequency applications.

1.4.4.2.1 Anisotropy and solid state sintering of PbTiO3

Anisotropic thermal expansion during cooling from a high sintering temperature creates

large internal stresses in the material, which is destroyed by microcracking. The

expansion is caused by the phase transition from a cubic paraelectric to tetragonal

ferroelectric with a relatively large c/a ratio 1.065. Therefore PbTiO3

ceramics are prepared through solid state sintering only after modification with suitable

dopants. Moreover, PbTiO3 possesses the highest room temperature spontaneous as high

as PS= 81 μCcm-2 (Venevtsev et al., 1959) among the perovskite ferroelectrics which is

caused by high tetragonal distortion and ionic displacements.

1.4.5 Defects in perovskite structureIt has been established that important defects in the perovskite structure are directly

related to vacancies of all three sub-lattices, electrons, holes and substitution impurities

(Lines and Glass, 1977). Such chemical defects strongly depend on the crystal structure

and chemical properties of the constituent’s species. The structure influences the types of

lattice defects and mobility’s of the defects and chemical species. The charges and size of

the ions affect the defects and their ability to be oxidized or reduced determines the

direction and amount of non-stoichiometry and the resulting enhanced electronic carrier

concentrations.

Defects in the crystal lattice generally cause deformation of the surrounding volume and

modification of the local fields (Lines and Glass, 1977; Damjanovic, 1998). which have

considerable influences on the dielectric properties and switching behavior of

ferroelectrics. Defects, including oxygen vacancies, space charges, etc., and their

13

distribution in a ferroelectric have been considered to play important roles in the

ferroelectric domain-wall pinning, polarization fatigue, etc.

It is well known, oxygen vacancies are major structural defects in the barium titanate

(Kang and Choi, 2002; Lemanov et al., 2000), play an important role in conduction. They

are generated due to loss of oxygen during sintering at high temperature in accordance

with following relation, a process defined by Kröger-Vink notation (Ang et al., 2000a)

V O → 12

O2+V ΄΄O+2e¯

1.5 Preparation Methods for Polycrystalline BaTiO3 and PbTiO3

Ceramics: Merits and Demerits

Various methods have been utilized for the fabrication of ferroelectric powders and bulk

ceramics. Co-precipitation, sol-gel processing, hydrothermal technique and solid state

reaction methods are more generally employed.

1.5.1 Co-precipitation method

Co-precipitation process is a widely studied process (Potdar et al., 1999; Prasadarao et

al., 1999; Simon-Seveyrat et al., 2007). Although by using by using this method,

chemical homogeneity can be achieved through mixing of constituent ions on the

molecular level under controlled conditions, yet in the case of oxalate route; it is difficult

to achieve optimal conditions where precipitation of both Ba and Ti ions occurs

simultaneously. This is because titanium is precipitated as titanyl oxalate at PH ≤ 2 in the

presence of alcohol and barium precipitation as BaC2O4 requires PH ≥4. Formation of

anionic species with titanium like TiO (C2O4)2 2- in the PH range 2-4 affect the

stoichiometry (Ba: Ti ratio) during simultaneous precipitation. Co-precipitation of barium

and titanium in the form of individual oxalates has rarely been attempted (Potdar et al.,

1999). However, the simultaneous co-precipitation of Ba and Ti in the form of oxalates

can be achieved by monitoring the chemical conditions such as PH (Prasadarao et al.,

1999), reagent concentration, reaction medium, chelating properties of oxalic acid,

complex formation with metal ions and their stability.

14

1.5.2 Sol- gel methodSol-gel method is used for preparing metal oxide glasses and ceramics by hydrolyzing a

chemical precursor to form a sol followed by gel. Gel gives an amorphous oxide on

drying, evaporation and pyrolysis gives. Crystallization is induced upon further heat

treatment. Three basic steps are involved; in the first step, partial hydrolysis of metal

alkoxide occurs to form reactive monomers. In the second step, polycondensation of the

monomers forms colloid like oligomers (sol). In the third step additional hydrolysis

promotes polymerization that leads to three dimensional matrix (gel)

As polymerization and cross-linking progress, the viscosity of the sol gradually increases

until the sol-gel transition point is attained, where viscosity abruptly increases and gelatin

formation takes place.

In the sol-gel technique, the structural and electrical properties of the final product are

dependent upon the nature of precursor solution, deposition conditions and the substrate

(Lazarevic et al., 2005).

Stearic acid gel and acetic acid gel wet-chemistry synthesis methods involving use of

stearic acid and glacial acetic acid have been employed (Wang et al., 2007). Li et al (Li et

al., 2002) described the oxalic acid precipitation method very similar to the sol-gel

acetate method but only acetic acid was replaced by oxalic acid.

Polymeric precursor method is also another polymeric route for BaTiO3 and PbTiO3

powders, where a solution of ethylene glycol, citric acid and metal ions is polymerized to

form a polyester type resin (Cho and Hamada, 1998).

1.5.3 Hydrothermal methodThe hydrothermal method that belongs to the category of liquid phase reactions

characteristically produces very fine particles with a narrow size distribution maintaining

a spherical morphology (Vivekanandan and Kutty, 1989). BaTiO3 and PbTiO3 powders

have been prepared by this method. This technique utilizes heating of an aqueous solution

of insoluble salts in an autoclave at moderate temperature and pressure for the

crystallization of a desired phase. The advantages of the hyrothermal method are the

15

reduced energy costs due to the modest temperatures for the reaction, less pollution,

simplicity of the equipment and the enhanced rate of the precipitation reaction. Powders

produced by the hydrothermal systhesis and solid state reactions are most commonly used

in multilayer ceramic capacitors (MLCCs).

1.5.3.1 Disadvantages of Hydrothermal method. Although, hydrothermal method is useful with reduced energy costs owning to the

modest temperatures sufficient for the reactions, less pollution and simplicity in the

process equipment and enhanced rate of the precipitation reaction. Extremely find

particles usually less than 200mm, yet some shortcoming are associated with this method.

During processing condition of a high water pressure, large amounts of protons and

hydroxyl ions tend to be incorporated into the BaTiO3 lattice. Considerable enlargement

of the unit cell volume occurs that leads to the superession of the tetragonal distortions of

the perovskite unit cell and bloating occurs in the final stage of powder sintering, which

gives particle density lower than the ideal density. Moreover, additional weight losses

occur due to the release of hydroxyl ions (OH-), protons (H+) and carbonate ions (CO3)-2.

1.5.4 Solid state reaction method

1.5.4.1 Sintering processThe sintering process (Kingery, 1992) has been known for thousands of years. It is a

primary operation in the production of the most traditional and advanced ceramics. Many

sintering techniques such as solid state sintering, liquid phase sintering and pressure

assisted sintering have been used widely and investigated extensively. In the solid state

sintering, the initial sintering stage, the rapid growth of inter-particle necking takes place

usually during heating. In the intermediate stage, grain growth is followed by

simultaneous pore rounding and densification. The final stage is characterized by pore

closure and final densification (German, 1996). Formation of liquid phase during

sintering increases the sintering rate and application of pressure accelerates the

densification process and ensures the elimination of residual pores. For most of the

inorganic powders, mass transport mechanisms including surface diffusion, volume

diffusion, grain boundary diffusion, dislocation climbs, viscous flow, plastic flow,

solution re-precipitation and even vapor transport from surfaces are involved. Traditional

16

solid state sintering has been well summarized by Germans (German, 1996) at theoretical

and practical grounds.

The most commonly used process for the powder synthesis is based on the thorough

mixing of the starting oxides or carbonates followed by solid state reaction at high

temperatures. The solid state reaction is the most traditional method for preparing

BaTiO3 as well as PbTiO3 powders by mixing the starting materials, usually titanium

dioxide (TiO2) and Barium carbonate (BaCO3). Calcination is carried out at elevated

temperatures (1100°C - 1400°C).

Barium titanate is produced from the reaction between TiO2 and BaCO3 (Amin et al.,

1983). After mixing raw materials are treated at high temperatures and then is BaTiO3

produced. Formation of BaTiO3 in the following steps (Beauger et al., 1983b)

Step (I)

Formation of BaTiO3 takes place at the cost of TiO2

BaCO3→ BaO + CO2 (1.1)

BaO + TiO2→ BaTiO3 (1.2)

Reaction takes place rapidly at the contacts surfaces of reactants

Step (II)

When BaTiO3 is formed at the surfaces, the reactants are separated by a product layer; the

course of reaction becomes diffusion controlled. Barium ions must diffuse through

BaTiO3 and penetrate into TiO2 grains. However, on reaching the BaTiO3 interface

barium can react as follows

BaTiO3 + BaO → Ba2TiO4 (1.3)

The formation of Ba2TiO4 takes place between primarily formed BaTiO3and BaO

Step (III)

Finally, Ba2TiO4 react with TiO2 to produce BaTiO3

Ba2TiO4+TiO 2→BaTiO3 (1.4)

In some respects, the sintering behavior of BT represents a rather classic case.

17

Fig. 1.9 The BaO-TiO2 equilibrium diagram of Rase and Roy (Rase and Roy, 1955)

At low temperatures, application of the classical models indicates the anion oxygen

vacancy grain boundary diffusion mechanism. At higher temperatures sintering is heavily

influenced by the presence of liquid phase. BT powders can be Ti-rich, either

deliberately, or due to dissolution of Ba from particle surfaces (Anderson et al., 1988).

The BaO-TiO2 phase equilibrium of Rose and Roy (Rase and Roy, 1955) indicates the

presence of a 1317C eutectic close to BaTi2O5. Hence, during sintering, a liquid phase is

always present to accelerate the densification.

1.5.4.2 Advantages of Solid state sintering method over Hydro-thermal

methodIncorporation of large amounts of protons and hydroxyl ions into the BaTiO3 lattice leads

to tetragonal distortions of the perovskite unit cell and consequently lowers the particle

density than ideal density. While these defects are not incorporated in the solid state

reaction method. Moreover, no additional weight losses are observed in the solid state

sintering reaction method.

18

Chapter 2Literature Review

The development of ferroelectric bulk materials is still under extensive investigation, as

new and challenging issues are growing in relation to their wide spread applications.

Apart from the growing number of new compositions, interest in the first ferroelectrics

like BaTiO3or PZT materials is far from dropping (Galasi, 2011). Among ferroelectric

materials, the perovskites (ABO3) are by far the most important and the widely studied,

well known BaTiO3 and PbTiO3 belong to this category.

Ferroelectric materials (Jaffe et al., 1971; Jona and Shirane, 1993; Lines and Glass, 1977)

make a variety of vital contributions to the digital world, from the capacitor

manufacturing to highly sophisticated piezoelectric systems with growth rates of more

than 15% per annum. Their successful exploitation over the last 60 years has been the

result of multidisciplinary efforts across fundamental and applied physics and chemistry,

through material science and ceramic engineering at an electronic platform.

Various methods including co-precipitation (Mulder, 1970; Stockenhuber et al., 1993)

sol-gel processing (Hwang and Han, 2000) hydrothermal technique (Kumazawa et al.,

1995) and solid state reaction method (Templeton and Pask, 1959) are utilized for the

fabrication of ferroelectric powders and bulk ceramics. Although, nanosized ferroelectric

powders have been synthesized by above mentioned wet chemical methods and

significant progresses have been achieved. Yet sol-gel process uses metal alkoxides

which are very expensive and sensitive to environmental conditions such as light and

heat. Moisture sensitivity necessitates the use of dry boxes or clean rooms for

experimental procedures.

Co- precipitation involve repeated washing in order to eliminate the anions coming from

the employed precursors salts that makes the process complicated and time consuming. It

is also difficult to produce large batches of chemical solutions through various processing

routes. Therefore, alternative methods for the preparation of ferroelectric ceramics are of

still technological and scientific significance. However, powders produced by the

19

hydrothermal synthesis and solid state reactions are most commonly used in multilayer

ceramic capacitors (MLCCs).

Hydrothermal method is also suitable but additional weight losses (Hennings et al., 2001)

due to the incorporation of the large amounts of protons and hydroxyl ions into the

BaTiO3 lattice during hydrothermal processing make solid state sintering reaction

advantageous .

Since the discovery of BaTiO3; it is regarded as corner stone and prototype of the largest

family of useful ferroelectrics: the oxide perovskite. The combination of a large A-site

and smaller B-site ion in this structure allows an infinite variety of simple perovskite

(ABO3) and complex perovskite and their solid solutions (Buscaglia et al., 2000; Lines

and Glass, 1977). This wide compositional and the strong composition dependence of

ferroelectric phase transition have allowed perovskite ferroelectric ceramics to be

exploited in a wide range of applications.

BaTiO3 is the first ferroelectric ceramic (Alrt et al., 1985; Bergstrom et al., 1997;

Haertling, 1999; Maurice et al., 1987) which is widely used in the electronic industry due

to its excellent dielectric, ferroelectric and piezoelectric properties. BT is a definite

chemical compound possessing relatively high stability components that make it easier to

sinter while maintaining good chemical stoichiometry (Haertling, 1999). Ferroelectric to

paraelectric phase transition occurs around 120OC (Curie temperature, Tc), above which

BaTiO3 is cubic, below the Curie temperature, the structure is slightly distorted, taking a

tetragonal symmetry (Zhao et al., 2004).

Significant features of ceramic processing science include phase equilibrium, sintering

mechanisms, microstructure control and defect chemistry. Both the stoichiometry and

composition control (Galassi, 2011) are important parameters for controlling the

ferroelectric properties.

The Ba/Ti ratio has a dramatic influence on the sintering behavior and microstructure

evolution; a small excess of TiO2 and (Ba/Ti < 1) is often added as a sintering aid to

improve the dielectric properties of BT based systems (Erkalfa et al., 2003). Since it has

long been believed that very small amounts of both BaO and TiO2 can be dissolved in

20

BT, < 100 ppm for BaO (Hu et al., 1985; Hwang and Han, 2000) and about 300 ppm for

TiO2 (Hwang and Han, 2000; Sharma et al., 1981) as excess of the precursor composition

can lead to the formation of secondary phases such as Ba2TiO4, Ba2Ti5O12 (Erkalfa et al.,

2003).

W.P. Chen et al (Chen et al., 2008) studied the correlation of crystal structure and Curie

point with Ba/Ti ratio (0.96-1.04). Curie temperature ranging from 98- 120°C was

attained; dielectric studies were made at 150°C. The pallets sintered with

Ba/Ti ratio of 0.98, 1.00 and 1.02 were almost single phased; revealed peaks of

Ba2Ti5O12. With 0.96 precursor’s composition, peaks of Ba2Ti5O12 were detected as

impurity phases. Ba2TiO4 and BaTi4O9 were present in the diffractograms of the obtained

specimens. Perovskite tetragonal structures were indicated. With increasing Ba/Ti ratio,

TC was shifted from 120°C to 98°C. Around TC, the strong temperature dependence of

permittivity restricts the application of BaTiO3 in MLCCs. Pb and V increase the Curie

temperature over 120°C (Haertling, 1999; Vijatovic Petrovic et al., 2011). Although,

interest in the Pb-based materials has decreased due to its toxicity (Ma et al., 2012; Zhou

et al., 2012), yet; Pb-modified BaTiO3 ferroelectric ceramics are still investigated for

applications in many types of electronic devices, such as transducers, actuators, sensors,

hydrophones, electro-optical modulators, infrared sensors and piezoelectric actuators

specially for high frequency and high temperature applications (Arya et al., 2003; Vold et

al., 2001).

BaTiO3 ceramics have too high sintering temperatures (1300°C). It is severely needed to

lessen their sintering temperature for a thermal budget. For this purpose, addition of low

melting glass or oxides, chemical processing and use of ultra-fine raw materials

(Drofenik et al., 1998; Kim et al., 2007; Xu et al., 2007a) are suggested. Bi2O3 is

commonly used to lower the sintering temperature of BaTiO3 ceramics. Pb can also

effectively lower the sintering temperature of BaTiO3 as a thermal budget.

In addition, Pb-doped materials can be attractive and acceptable because of their low

anisotropic field and fast crystallization process. Microstructure and dielectric properties

of lead barium titanate ceramics are reported in the literature (Chaimongkon et al., 2011)

and mostly Pb had been utilized to acquire dense materials; electrical properties of

21

BaTiO3 ceramics were investigated over the restricted range of 30-350°C. The firing

temperature strongly influenced on the phase formation, microstructure, tetragonality and

dielectric properties of the PBT ceramics. Tetragonality of the ceramics was reported to

increase with increasing sintering temperatures. Average particle and average grain size

also increased and dense PBT ceramics were obtained; dielectric constant responded in

accordance.

Over the past decade the use of Bi0.5 Li0.5 TiO3 (BLT), Bi0.5 Na0.5 TiO3 (BNT) and Bi0.5 K0.5

TiO3 (BKT) in BaTiO3 (BT) based systems made significant progress in developing the

ceramics with high Tc. Huo et al (Huo and Qu, 2006) first reported BaTiO3-Bi0.5 Na0.5

TiO3 (BT-BNT) as promising candidates above 130°C. Later, Leng and Yuan (Leng et al.,

2009; Yuan et al., 2013) indicated BaTiO3-Bi0.5 K0.5 TiO3 (BT-BKT) systems for Tc

higher than 130°C. In addition BaTiO3-Bi0.5 Li0.5 TiO3 (BT-BLT) ceramics have been used

to raise the Curie temperature over 30°C (Pu et al., 2010).

Undoped BT is an electrical insulator with its resistivity varying between 109 and 1012 Ω

cm at room temperature (Nowotny and Rekas, 1991). In the past it had been tried to be

converted into a semiconductor after partial substitution of Ba+2 by trivalent cations such

as Y+3, La+3 or Ti+4 by pentavalent Nb+5 and Ta+5 cations (Chen and Yang, 2011; Urek et

al., 2000).

BaTiO3 ceramics have been tailored with variety of dopants in order to modify or

optimize properties such high or low resistivity and temperature of the ferroelectric Curie

transition as well. In many cases doping leads to the adjustment of the Curie temperature

(Tc) in BaTiO3 which is rationalized by the ion–size effects and changes in the tolerance

factor. Doping also results in the variance of conduction. Acceptor doping has the effect

of fixing the charged vacancies through ionic compensation (Yoon et al., 2010). Whereas

donor doped BaTiO3 associated with electron compensation usually increases its

conductivity (Yuan et al., 2012).

Besides, Tc is an important parameter for PTCR applications but BaTiO3 ceramics have

Curie temperature around 120°C (Chen et al., 2008). PTCR effect in polycrystalline

BaTiO3 is characterized by dramatic increase in the electrical resistivity across the

tetragonal (ferroelectric) to cubic (paraelectric) transition across the Tc (Nowotny and

22

Rekas, 1991). Heywang explained PTCR effect in terms of the double Schttky barriers at

the grain boundaries. According to this model, these barriers result from electron trapping

by the acceptor states at the interfaces (Heywang, 1961). Later on Jonker (Jonker, 1964)

extended the model considering the influence of polarization on the resistivity below the

Curie point.

Y. Luo et al (Luo et al., 2006) studied the PTCR effects of BaBiO3 doping at the

molecular ratio of 0.1- 0.5 on BaTiO3 based ceramics. The perovskite BaBiO3 contains

bismuth ions of different oxidation states (+3 and +5) in an ordered arrangement. As the

elements for achieving semiconducting grains, Bi+3 substitutes for barium ions in the

lattice and Bi+5 substitutes for titanium ions (Steigelmann and Goertz, 2004). At small

concentrations, donor incorporation by electronic compensation leaded to high

conductivity. The average dopant concentration had an important effect; the dopant

incorporation at the grain boundaries was shifted from electronic to vacancy

compensation, resulting in the formation of highly resistive layers

Li and B. Bergman (Li and Bergman, 2005) studied the electrical properties of doped

BaTiO3 ceramics by single dopants including bismuth, hafnium, samarium, lanthanum

and ceriumat 0.1-0.6 mol% doping; relationship between the resistivity change and aging

time was investigated at the room temperature. The resistivity of Bi- and Ta- doped

samples decreased at the initial ageing stage, and then increased with ageing time. The

electrical properties and ageing characteristics were related to interior oxidation and

reduction reactions. The electro negativities of the atoms were supposed as a primary

factor for the reduction reaction.

Y. Leyet et al (Leyet et al., 2010) studied the relaxation dynamics of the conductive

process of BaTiO3 ceramics at high temperatures from room temperature to 630°C at

1kHz, 10 kHz and 100 kHz frequencies. Dielectric constant was both temperature and

frequency dependent. Maximum dielectric constant was obtained at 1kHz frequency.

Phase transition was followed at 130°C; above 220°C the dielectric response was

changed to a typical relaxor behavior. The values of the activation energy (Ea

0.72-0.8) calculated in the investigated temperature interval suggested the existence of

23

two conductive relaxation mechanisms corresponding to the movement of single ionized

oxygen vacancies and electrons of the oxygen vacancies in second ionization state.

In perovskites, the partial substitution of A- or B- site ions (carrier doping) promotes the

activation of several conduction mechanisms. In the last few years, extensive studies have

been carried out A- or B- site. Substitution of different elements with dissimilar

electronic configuration can result in dramatic effects in association with the electronic

configuration.

Dielectric dispersions related to a conductivity phenomenon which obeys the Arrhenius´s

dependence have been reported for Ba1-xPbxTiO3 ceramics (Bidault et al., 1994). At high

temperature regions, dielectric anomalies have been reported for BaTiO3(Pb,La)TiO3 and

(Pb,La) (Zr,Ti)O3 systems (Kang et al., 2003b). The dielectric anomalies were related to

the competition phenomenon of the dielectric relaxation and the electrical phenomenon

by the oxygen vacancies which are extrinsically formed in the lattice (Ang et al., 2000a).

Among the perovskite materials, PbTiO3 (PT) is another popular member with a relatively

high phase transition temperature 490°C and a low dielectric constant of about 200, which

make it attractive for high temperature and high frequency applications (Ikegami et al.,

1971; Takahashi, 1990) and large polarization and high pyroelectric coefficients find

pyroelectric, electro-optical and non-linear optical materials applications. Above phase

transition, the structure is cubic and below, it is tetragonally distorted perovskite.

However, PT has undesirable mechanical properties due to its large tetragonal strain,

anisotropy ( c/a = 1.064) (Shirane et al., 1956). On cooling through the Curie

temperature (TC), the large anisotropy of the material becomes fragile that renders

difficult the sintering of pure lead titanate. Moreover, it has large coercive field that make

poling difficult. In this perovskite, substitution of Pb by isovalent cations such as as

Ca2+ Ba2+ Cd2+ into the Pb2+ sites or off-valent ions including Sm3+ Gd3+ Y3+ Nd3+ La3+

produces a decrease in the coercive field and lattice anisotropy of the material (Shenglin

et al., 1995; Takeuchi et al., 1982). Risk of micro-cracking is reduced and samples

become dense and considerable values of remnant polarization are obtained. Addition of

Mn2+ (Takeuchi et al., 1982) also reduces the lattice anisotropy of the ceramics. Above

mentioned elements have been used as dopants in PT bulk ceramics.

24

G.S. Forrester et al (Forrester et al., 2004) synthesized PbTiO3 ceramics by long term

sintering 24h in air through conventional mechanical alloying. Monolithic ceramics were

obtained at 700°C whereas sintering above 800°C leaded to the development of the

cracks whereas spontaneous fracturing was observed at 1050°C. Development of micro-

cracking was not influenced by the c/a ratio; as the value did not differ among the

samples sintered at different temperatures.

Low melting additives such as glass and Si have been used by the researchers to sinter the

lead based ceramics. V.L Palkar et al (Palkar et al., 2000) employed 2 mol % Si as

sintering aid for the densification. Si acted as a binder in the matrix and prevented the

crumbling of the samples by sustaining the large stain developed during phase

transformation on cooling. In addition, Si has a tendency to form a silicate, generally in

the form of viscous glass, at the firing temperature. This viscous glass has the capability

of flowing into the pore regions and lead to major densification. Crystal structure and

teragonality of PbTiO3 ceramics were not affected, physical properties of the materials

were retained as well. However, dielectric constant of 1100 was achieved at

phase transition.

In the last decade, many new sintering techniques such as microwave sintering (MW),

laser sintering and spark plasma sintering (SPS) have been developed to improve

sintering (German, 1996; Kingery, 1992; Thomas et al., 1996).

Spark plasma sintering (SPS), a pressure assisted sintering method involving the use of

microscopic electrical discharges among particles (Gao et al., 2008; Omori, 2000;

Takeuchi et al., 1999); produces dense ceramics in typically in few minutes at low

temperatures, several hundred degrees Centigrade less than that of conventional solid

state sintering reaction method (Wu et al., 2002; Yoshida et al., 2008) . It has been

successfully used to sinter nanostructured materials, organic composites etc. (Li et al.,

2006; Nygren and Shen, 2003; Ran and Gao, 2008; Shen et al., 2002; Wu et al., 2005; Xu

et al., 2007b; Zhan et al., 2003).

In the last five years more than one thousand research papers have been published under

practical implementation of this technique. However, a fundamental understanding of

SPS mechanism is yet to be realized due to the complexity of the thermal, electrical and

mechanical processes being involved during SPS (Bernard-Granger and Guizard, 2009;

25

Chaim, 2007). Although, several different mechanisms such as vaporization-

condensation, plastic deformation, surface, grain boundary and volume diffusions have

been assumed for the sintering and densification for SPS, yet fundamental understanding

of the SPS mechanism is quiet under debate (Chaim and Margulis, 2005; Groza and

Zavaliangos, 2000; Tokita, 1997).

There has also been growing interest in the microwave sintering (MS) of ceramics to get

the dense and fine grains with uniform grained structures; it also includes the potential

advantages such as rapid heating and penetrating radiations involving the heating of

sample from inside to outside (Gopalan and Virkir, 1999; Thostenson and Chen, 1999).

S. Singh et al (Singh et al., 2005) reported the properties of microwave sintered Sm

modified Pb CaTiO3 (PCT) ceramics. Microstructures and dielectric properties of

conventionally and microwave sintered PCT ceramics. The grain size distribution for MS

specimens was more uniform as compared to conventionally sintered specimens; average

grain size was also smaller. However, the impedance spectroscopy studies revealed that

the values of dielectric constant were slightly less for MS sintered PCT as compared to

conventionally sintered specimens.

2.1 Objectives of the research workSince the discovery of ferroelectricity in single crystalline materials (Rochelle Salt) in

1921 and its subsequent extension during early to mid - 1940s into the realm of

polycrystalline ceramics (barium titanate) led to investigation of new materials and

technological development with number of industrial and commercial applications, being

continued to the present day (Haertling, 1999). Among ferroelectric materials, the

perovskites (ABO3) are the most important and the extensively studied, well

acknowleged BaTiO3 and PbTiO3 belong to this category (Damjanovic, 1998)

At present researchers are mainly focusing on the processing structure and properties of

these materials. An affective structure-property relationship is persistently built up for

finding the several processing parameters to attain the desired properties (Guo et al.,

2006).

As stoichiometry and composition control (Kosec et al., 2001a) are important parameters

for controlling the ferroelectric properties. The Ba/Ti ratio has a dramatic influence on

the sintering behavior and microstructure evolution; a small excess of TiO2 and (Ba/Ti <

26

1) is often added as a sintering aid to improve the dielectric properties of BT based

systems (Erkalfa et al., 2003). Since it has long been believed that very small amounts of

both BaO and TiO2 can be dissolved in BT, < 100 ppm for BaO (Hu et al., 1985; Hwang

and Han, 2000) and about 300 ppm for TiO2 (Hwang and Han, 2000; Sharma et al., 1981)

as excess of the precursor composition can lead to the formation of secondary phases

such as Ba2TiO4, Ba2Ti5O12 (Erkalfa et al., 2003).

As some manufacturers might use Ti- excess for reducing the sintering temperature of

BT; hence, studies were attempted with 0.94 and 0.98 Ba/Ti molar ratio to find the

controlling parameters such as Curie temperature (TC) and novel resistivity of the

obtained ceramics for their manipulation to PTCR effects. Electrical properties of

undoped BT in the wide range of temperatures (40 – 700°C) at 1kHz frequency were

accomplished to find the understanding of conduction process at higher temperatures; as

BT based systems are currently investigated for energy harvesting applications (W.-B. Li

et al., 2016; Xu et al., 2016). Besides; crystal structure, morphology, thermal and

dielectric behavior coupled with ferroelectric measurements were explored for finding

structure-property relationship.

Another center of this research was the fabrication of PT ceramics with Pb/Ti contents

through solid state sintering reaction method. Objectives hold the sintering of PT

ceramics; as their large c/a ratio makes fragile and renders difficult their sintering. They

were electrically investigated as well in the wide spectrum of temperatures (40–700°C) at

1kHz frequency by impedance spectroscopy for assortment of applications.

Studies can be summarized as follows

Control of the impurity phases under various processing parameters under

composition control, sintering temperature and sintering duration.

Study the dielectric constant, loss tangent, ac conductivity, dc resistivity and

conductivity in the wide range of temperatures (40 -700°C) at 1kHz frequency.

Finding of the controlling parameters i-e Curie temperature with Ba/Ti and Pb/Ti

contents for useful implementation of the work.

Investigation of the dielectric and electrical response in the ferroelectric and

paraelectric regions at elevated temperatures.

Following of conduction process.

27

Development of the conduction process by the use of dopant at A-site of BT

ceramics for practical application.

Finding of the Arrhenius dependence of Conduction process of un-doped and

doped BT and PT ceramics

Study of the ferroelectric characteristics of the studied materials.

We have performed this work in three systematic sections.

In the section one, two molar precursors of Ba/Ti contents were subjected to varying

processing parameters to obtain phase pure BT ceramics. We had optimized various

conditions under the influence of sintering temperature, sintering duration; almost phase

pure BT ceramics were prepared and exposed to number of characterizations.

In the section two, almost phase pure BT ceramics were doped; TC, dielectric properties,

conductivity and ferroelectricity were increased was increased.

In section three, preparation of PT ceramics was attempted with three molar precursors

of Pb/Ti contents under varying processing parameters for many times. Pre-fired PT

ceramics 1000 C 1100C, 1180± 10C were subjected to sintering. In the long run crack

free PT ceramics were sintered successfully with all precursor compositions after a

period of four months.

But this is a very minute contribution toward next exploitation of the ferroelectric PT

materials.

2.2 Scope of the WorkElectrical properties of BT has potential applications for electronic devices such as

multilayer ceramic capacitors (MLCC), thermistors and varistors as well as sensors and

piezoelectric actuators, embedded capacitance in circuit boards, under water transducers,

electroluminescent panels and energy storage systems.

PT may be used for high temperature and high frequency applications and electro-optical

materials applications. Pyroelectric PT ceramics have been found useful for thermal

imaging and human detection.

Chapter 3Experimental Work

28

3.1 Preparation of BaTiO3 and PbTiO3 ceramics

Polycrystalline ceramics of BaTiO3 and PbTiO3 ceramics were prepared through solid

state reaction method.

3.1.1 Preparation equipment and source of errors

For the preparation of barium titanate, lead titanate and their derivatives, the required

equipment, controlling parameters and their influence on the properties are given. The

apparatus used were digital balance with high precision, agate mortar, die, hydraulic

press and open air furnace. Since open air furnace was employed, oxygen vacancies

might be encountered during sintering process.

3.1.2 Preparation of BaTiO3 (BT) ceramics

BT ceramics were prepared from commercial grade precursors barium carbonate (BaCO3,

Fluka 98.5 % pure) and titanium oxide (TiO2, Sigma Aldrich 99% pure) by solid state

sintering method. Two batches in compositions with Ba/Ti ratio 0.94(a) and 0.98(b) were

weighed accordingly and hand agate mortared intensively for eight hours. In the first

step, mixed powders (a) were fired at 1100°C and 1300°C for four hours, in the second ,

powders of both batches (a) and (b) were fired at 1300°C for 1 - 4 hours separately. The

fired powders were milled in agate mortar for two hours to obtain chemical homogeneity

(Haertling, 1999) and sieved in laboratory test sieve - stainless steel with mesh no. 150.

Sieved powders were granulated and binded with 5wt % polyvinyl alcohol (PVA),

pressed into circular pallets of 2mm thickness and 12mm diameters at 20 MPA pressure.

They were fired in an open air furnace at 1300°C for 2 hours .

3.1.3 Preparation of Pb-doped BaTiO3 (PBT) ceramics

BT ceramics in composition with Ba/Ti (0.98) was prepared from commercial grade

precursors barium carbonate (BaCO3, Fluka 98.5 % pure), titanium oxide (TiO2, Sigma

Aldrich 99% pure) and lead carbonate (PbCO3 Fluka 98.5 % pure).The powders were

weighed accordingly and agate mortared intensively continuously for eight hours. BaTiO3

was prepared through solid state reaction at sintering regime of 1300°C/4h. Fired BaTiO3

29

powder was agate mortared for two hours and five compositions were prepared at x=

0.00, 0.025, 0.050, 0.075 and 0.100% PbCO3 doping (0.00– 0.1 mole %). Samples were

fired at 1200°C for one hour in an open air furnace. All fired powders were milled

intensively in agate-mortar for two hours to obtain chemical homogeneity (Haertling,

1999) and sieved in laboratory test sieve - stainless steel with mesh no. 150 (BS 410,

Endecott's Ltd, London, England). Sieved powders were granulated and binded with 5 wt

% polyvinyl alcohol (PVA).They were pressed into circular pallets of 2mm thickness and

12 mm diameters at 20 MPA pressure and fired in an open air furnace at 1200°C for 2h.

3.1.4 Preparation of Bi- doped BaTiO3 (BBT) ceramics

BaTiO3 in composition with Ba/Ti (0.98) was prepared from commercial grade precursors

barium carbonate (BaCO3, Fluka 98.5 % pure) and titanium dioxide (TiO2, sigma Aldrich

99% pure) and bismuth nitrate penta-hydrate (Bi (NO3)3 .5H2O sigma Aldrich 98%

pure).The powders were weighed accordingly and agate mortared intensively

continuously for eight hours. BaTiO3 was prepared through solid state reaction at

sintering regime of 1300°C/4h. Fired BaTiO3 powders were agate mortared for two hours

and five compositions were prepared at x= 0.00, 0.025, 0.050, 0.075 and 0.100% Bi

(NO3)3 .5H2O doping (0.00– 0.1 mole %). Samples were fired at 1200°C for one hour in

an open air furnace. All fired powders were milled intensively in agate-mortar for two

hours to obtain chemical homogeneity (Haertling, 1999) and sieved in laboratory test

sieve - stainless steel with mesh no. 150 (BS 410, Endecott's Ltd, London, England).

Sieved powders were granulated and binded with 5wt % polyvinyl alcohol (PVA).They

were pressed into circular pallets of 2mm thickness and 12 mm diameters at 20 MPA

pressure and fired in an open air furnace at 1200°C for 2h.

3.2 Preparation of PbTiO3 (PT) ceramics

PT ceramics were prepared from commercial grade precursors barium carbonate (PbCO3,

Fluka 98.5 % pure) and titanium oxide (TiO2, sigma Aldrich 99% pure) by solid state

sintering method. Three batches in compositions with Pb/Ti ratio 0.94(a) and 0.98(b) and

1.00 were weighed accordingly and hand agate mortared intensively for eight hours. In

the first step, mixed powders (a), (b) and (c) were fired at 1000-1190°C for two hours.

30

In the second, powders of batches (b) were fired at 1190°C for 1 - 4 hours separately. All

fired powders were milled in agate mortar for two hours to obtain chemical homogeneity

(Haertling, 1999) and sieved in laboratory test sieve - stainless steel with mesh no. 150.

Sieved powders were granulated and binded with 5wt % polyvinyl alcohol (PVA),

pressed into circular pallets of 2mm thickness and 12mm diameters at 20 MPA pressure.

They were fired in an open air furnace at 1190°C for 1hour.

3.3 Experimental parameters and their influence

As ground BT, PT precursors were exposed to numeral sintering temperatures, sintering

durations to get specimens at better performances. Further, the role of cations at A-site of

the obtained specimens was explored under various characterizations for practical

implantation of the work.

3.4 Characterization Techniques

Characterization of the sintered pallets was required to observe various parameters like

phase identification, crystallite size, lattice parameters, dielectric constant, transition

temperatures, electrical resistivity and conductivity and electrical polarization. The as

ground and sintered specimens were subjected to thermal, structural and electrical

characterization as follows.

3.4.1 Thermal characterization

Under this branch, sample properties are investigated under the influence of temperature.

Properties like, phase transition temperatures, variation in mass, dimension, optical

properties mechanical stiffness and dielectric permittivity are measured employing

techniques like Differential Scanning Calorimetry (DSC), Thermo Gravimetric Analysis

(TGA), Thermo Mechanical Analysis (TMA), Thermo Optical Analysis (TOA), Dynamic

Mechanical Analysis (DMA) and Dielectric Thermal Analysis (DEA) respectively.

We had employed TGA/ DSC on the as ground precursor’s composition powders before

subjecting them to heat processing treatment. In thermogravimetric analysis, the changes

in the physical and chemical properties of materials are measured as a function of

31

increasing temperature at constant heat rates (10C/min). As the temperature increases,

various components of the sample decompose and the weight percentage of the each

mass change can be observed. The change in the weight loss can immediately be seen in

the TGA curve as a trough, or as a shoulder or tail to the peak.

Fig. 3.1 TGA and DSC curve in air for BaCO3, TiO2, and ZrO2 powder mixture

While, DSC provides information about different phase transition temperatures like

antiferromagnetic to paramagnetic temperature (TN), crystallization temperature,

rhombohedral to orthorhombic (α to β ) orthorhombic to cubic (β to γ) or ferroelectric to

paraelectric(TC) and melting temperatures (Tm ) (Karthic et al., 2012). DSC monitors the

differential heat flow between the sample and reference material during phase transition

or chemical reaction as a function of temperature. The temperature of both the sample

and reference material is maintained at the same value and is increased at the same rate

during the whole scan. Since DSC analysis is performed at a constant pressure; heat flow

(Q) equivalent to enthalpy (H). The differential rate of heat flow∆ (dH/dt) between

sample and reference can be either positive or negative according to exothermic and

endothermic reactions. The area under the peak in thermogram is directly proportional to

the heat evolved or absorbed by the reaction. Height of the curve is directly proportional

to the rate of reaction.

32

We have performed thermal analysis on TGA/DSC (SDT Q600 V8.0 build 95 model,

Germany) in a dynamic environment of inert gas (Ar) from room temperature to 1300°C.

3.4.2 Structural characterization

3.4.2.1 X-ray diffraction method

X- rays like other electromagnetic rays interact with electron cloud of atoms. Because of

their shorter wave length are X-rays scattered by adjacent atoms in the crystal which can

interfere and give rise to diffraction effects. X-ray is the most suitable nondestructive and

general method for studying the crystal structure. It requires no elaborate sample

preparation(Schroder, 1998). It provides information on structures, phases, preferred

crystal orientation (texture) and other structural parameters such as average grain size,

crystallinity and crystal defects. A beam of the monochromatic light when falls upon a

crystal, is scattered in all directions within it. But owning to the arrangement of the atoms

in certain directions, the scattered waves interfere constructively with one another while

in the others they interfere destructively. Constructive interference takes place only

between those scattered rays that obeys Bragg’s law λ=2 d sinθB, where λ is the

wavelength of X-ray beam of incident incident upon a crystal at angle θB with the family

of planes whose spacing is ‘d. In case of rotating crystal method, a crystal is mounted

with one of its axis, or with or more of its crystallographic direction, normal to the

monochromatic X-ray beam. As the crystal rotates, a particular set of lattice planes will

make the correct Bragg angle for reflection of the monochromatic incident beam. The

peak intensities are determined by the atomic decoration within the lattice planes.

Consequently, the X-ray diffraction pattern is the finger print of periodic atomic

arrangements in a given material. The radiations in our case were Cu K α (λ = 1.5406 Å).

Powder X-ray diffraction measurements were made on XRD PANalytical, X’pert pro,

Netherlands.

3.4.2.1.1 Crystallite size

33

The experimental diffraction patterns can be compared to those in JCPDS ( joint

committee on powder diffraction standards) for phase identification. The broadening of

the experimental diffraction peak can give the additional information about crystallite

size, grain size that have pronounced effect on the many of the properties of

polycrystalline materials. The crystallite sizes were calculated using Scherrer formula

(George et al., 2006) by taking into account the full width at half maximum (FWHM)

values of the indexed peaks in the X-ray patterns

t= 0.9 λβ cosθB

(3.1)

Where, t is the crystallite size, λ is the wavelength of the incident X-rays (1.5406 Å), θB

is the Bragg’s diffraction angle andβ is the full width at half maximum (FWHM) at θB in

radians.

3.4.2.1.2 Measured density

Measured density is an intrinsic property of the material; it denotes the relationship

between its mass and unit volume. This parameter is difficult to be characterized because

it can be affected by temperature, pressure and amounts of the substitution of the

different elements. It is used as an indexed property or an independent variable to predict

the other properties of the materials. The measured density (ρm ¿ was determined by

measuring the dimensions and mass of the samples and calculations were made by

employing the equation (Vijatovic et al., 2011)

ρ= 4 md2

hπ (3.2)

Where, m is the mass, d is the average diameter and h is the height of the sintered

samples.

3.4.2.1.3 X-Ray density

34

The x-ray density of the materials was calculated by employing the following relationship

ρ ¿ nMa2 c

(3.3)

Where, M is the molecular weight, ‘a’ and ‘c’ are the lattice constants and n denotes the

number of formula units in the unit cell of BaTiO3 and PbTiO3 ceramics.

3.4.2.1.4 Porosity

The storage capacity of any material is referred to as porosity; it depends upon the shape

size of the grains and the degree of their storing and packing. Porosity (P) of the samples

was calculated by the underlying relationship

P = 1−ρm

ρ x (3.4)

Where, ρm and ρx are the measured and calculated densities.

3.4.3 Scanning electron microscopy

Scanning electron microscopy (SEM) is considered to be a nondestructive technique and

gives sample images at high resolution. It includes applications in medical science,

forensics and material sciences. Electrons interact with the sample and reveal information

about surface morphology, particle size, their distribution and shape, orientation (EBSD)

and chemical composition (EDS). In SEM, electrons from gun are accelerated and

focused with different lens condensers for faster scanning of the sample. These electrons

have sufficient amount of kinetic energy to interact with the material. The energy of

electrons is is dissipated to the material and variety of signals are produced like

secondary electrons, backscattered electrons, characteristic X-rays, cathode luminescence

and heat. Secondary electrons are used to produce SEM images for morphology and

topography. Contrast in multiphase composition can be illustrated by backscattered

electrons (BSE). In-elastic collision of incident electrons causes characteristic X-rays for

elemental analysis (EDS). Cathode- luminescence is used for compositional maps based

on trace elements. For SEM analysis, materials require minimal sample preparation. But

35

there are limitations like size constraint, outgassing and bon conducting specimens.

Charge accumulation on the surface of insulating materials causes blurring of the images.

Hence, a layer of conductive coating is required for such kind of samples. We have

observed morphological features using scanning electron microscope (JSM 6490)

equipped with energy dispersive X-ray spectroscopy (EDX).

Energy dispersive X-ray spectroscopy (EDX/EDS) is normally integrated with the

SEM system. High energy incident electrons from SEM source generate X-rays on

inelastic collision with the specimen. These electrons kick out the inner electrons of the

atom within the sample. Characteristic X-rays are evolved on occupying the higher

energy level electrons. EDS has a cooled liquid Nitrogen equipped with highly sensitive

X-ray detector to convert these X- rays into an energy spectrum as shown in Fig. 3.

Different peaks correspond to theKα, K β, Lα, Lβ etc. the X-rays of different elements.

(a)

36

(b)

Fig. 3.2 (a) Process of characteristics X-rays from an atom (b) EDS spectrum

3.4.4 Electrical characterization

3.4.4.1 Electrical ac measurements

Electrical ac measurements as a function of temperature were followed at 1k Hz

frequency.

Dipolar dynamics of materials can be elaborated by dielectric spectroscopy technique;

impedance spectroscopy is a powerful tool to investigate the electrical properties of the

ceramics at different frequencies and temperatures. It helps to understand the conduction

process and dielectric relaxation mechanism. These studies have been focused intensively

due to industrial and technological applications (Jonscher, 1999). Charge separation in

solids due to external applied electric field is the origin of dielectric response. There are

four basic mechanism of polarization

Interfacial polarization or space charge polarization: Orientation of charges in

response of external field at grain boundaries, inter phase boundaries and at the

surfaces, contribute towards the dielectric response of the material.

Dipolar or orientation polarization: Materials which have natural or

spontaneous dipoles like H2O are randomly oriented due to thermal agitation and

37

have no net polarization. In response to the applied field, these dipoles align to

some extent and contribute in the polarization.

Ionic polarization: Ionic crystals like NaCl, has built in dipoles which exactly

cancel each other and unable to spin. But external field can displace them slightly

from their mean position and can induce net polarization.

Electronic or atomic polarization: Almost all materials have this type of

polarization. Overall the atom with its electrons has a spherical symmetry. An

external field can displace the electrons with respect to the nucleus and a dipole

moment is formed.

The total volume polarization P is the vector sum of all above mentioned dipole moments

per unit volume. Current will lead the voltage by 90o when varying electric field is

applied on two parallel plates in vacuum. But if a dielectric material is placed between

these plates then the polarization can be related as

P =ε ° χ E (3.5)

Where ε° and χ are the vacuum permittivity and susceptibility respectively. All the above

mentioned polarization mechanisms will respond the applied field by moving masses

(charges). This mechanical response of masses will take some time to follow the applied

frequencyvas shown in fig. 3a. This time is called the relaxation time, τ and due to

which current leading agle of 90o for an ideal capacitor reduces by a factor of δ due to the

lossy part of the material response and is mentioned as equivalent series resistance (ESR)

in Fig.3b. This factor loss is defined as the dielectric dissipation factor (DF) or tan(δ )

DF = tan(δ ) = 1Q

=¿ ESRXC

(3.6)

Where, Q is the quality factor, XC is capacitive reactance,ε is the permittivity of the

material when a dc field is applied to the material

ε = ε r ε° (3.7)

Where, εris the relative permittivity and is also known as dielectric constant

38

Fig. 3.3 Relaxation time (τ ) of masses after turning off the applied electric field (b)

Impedance plane of real capacitor (c) Response of real ε 'and imaginaryε ' 'parts of

permittivity.

But the permittivity becomes a complex function ἐ when electric field E is frequency

dependent (E α e−iwt ) , it can be written as

ε ¿ = ε '(ω¿ + ε ' '(ω¿ (3.8)

Where ωis the angular frequency, ε ' is real part of permittivity and is known as dielectric

constant (related as the energy storage) of the material.ε ' ' is the imaginary part of

permittivity and relates to the dissipation or loss of energy. ε 'and ε ' 'are out of phase as

shown in Fig. 3c.

All the four polarization mechanisms contribute towards the dielectric constant at lower

frequencies of applied electric field. The movement of charges in polarization

mechanisms cannot follow the higher frequencies and their contribution vanishes one by

one. The dielectric constant decreases with increase of frequency and at each relaxation

frequency; there is a peak of dissipation factor or relative loss factor as shown in Fig. 3

39

Fig. 3.4 ε 'and ε ' ' as a function of frequency with associated different polarization

mechanisms (Murphy and Morgan, 1938).

Dielectric constant was determined from the following equation (Yasin Shami et al.,

2011)

ε ' = CdƐ o A (3.9)

Where d is the thickness of pallet, C is the capacitance, Ɛo is the permittivity of free space

and A is the cross-sectional area of the pallet. We have used WYNE KERR precision

component analyzer (6440B) connected with homemade temperature programmable

furnace was used for recording Capacitance and resistance at 1kHz frequency in the

temperature range of 40 - 700°C; heating rate was 5°C/ min.

In general, dielectric constant (ԑ΄), at different frequencies increases with increase in

temperature and is influenced by dipolar, interfacial, ionic and electronic polarizations

(Singh et al., 2002). Dielectric constant increases at lower frequencies while decreases at

higher frequencies which is a normal behavior of the ferroelectric materials. Since at

40

higher frequencies, the dipoles cannot follow the applied AC electric field; decrease in

the dielectric constant is observed. At lower frequencies, the dipolar and interfacial

polarization contribute significantly to the dielectric constant (Lines and Glass). Both of

these are temperature sensitive; hence, the dielectric constant increases at higher rates at

1k Hz as compared to 10 kHz, 100 kHz and at M Hz frequencies (Silveira et al., 2013).

Therefore, dielectric measurements were carried out at 1kHz frequency in the above

mentioned range of temperatures.

The ac conductivity was evaluated from dielectric constant (ε ') by employing the

following relation (Sharma et al., 2014)

σac= ω Ɛ0 ε ' tan (δ) (3.10)

Where, σ ac is the ac conductivity and ω = 2 πf , is the angular frequency. Near the Curie

temperature, the domain structure break up, carriers become free and take part in

conduction mechanism.

3.4.4.2 Electrical dc measurements

Depending upon the resistivity, generally two methods; two probe and four probe, are

used to measure the resistivity. If the resistivity of the specimen is higher (kΩ cm and M

Ω cm) than the contact resistance, then two probe method is used. On the other hand, if

the contact resistance and probe resistance is comparable to the resistivity of device under

test (DUT), then four probe method is employed. Since, the resistivity of our samples was

in the range of (MΩ cm); therefore two probe method was employed.

3.4.4.2.1 Electrical dc resistivity

Electrical dc resistivity (ρ) measurements were made using two probe method (Soitah and

Yang, 2010) at room temperature and from 40 -700°C by employing circular pallets of

2mm thickness and 12mm diameters. Pressure contacts equal to the pallet size were

applied after polishing the surfaces. Conductivity of the specimens was evaluated from

the impedance spectrum by using the relation (Xu et al., 2016)

41

σ dc= dA× R

(3.11)

Where, d is the sample thickness, A is the electrode area; R is the resistance of bulk

ceramics. Specimens are generally resistive at room temperature; ionic conduction as a

function of temperature is mostly followed by oxygen vacancies. Activation energy (Ea)

required for the process was calculated by the Arrhenius equation

σ = σ₀ exp (−EaKB T ) (3.12)

Where σ₀ is the pre-exponential factor, Ea is the activation energy needed to release an

electron from an ion for a jump to neighboring ion to give rise to electrical conductivity;

KB is the Boltzmann’s constant and ‘T’ is the absolute temperature. The temperature

dependence of conductivity arises only due to mobility. Arrhenius equation gives the

temperature dependence of reaction rates. This equation has vast and important

application in determining rate of chemical reactions and for the calculation of energy of

activation. It can be used to model the temperature variation of diffusion coefficients,

population of crystal vacancies and many other thermally induced processes and

reactions. Currently, it has been principally employed in various ferroelectric systems for

energy of activation calculations (W.-B. Li et al., 2016; Xu et al., 2016).

3.4.4.2.2 Drift mobility

Drift mobility, a thermal transport property of the material; is in fact a proportionality

factor between the drift velocity of the charge carriers in semiconductor and electric field.

Drift velocity of charge carriers in semiconductor under the electric field, as opposed to

the carriers in free space carriers in semiconductor, are not infinitely accelerated by the

electric field and thus, they reach a finite velocity regardless of the period of time at

which the field is acting at a given electric field. Drift velocity is determined by the

mobility of charge carriers. We calculated drift mobility, μd at the temperatures (480 -

700°C) using the following relationship

μd=1

neρ (3.13)

42

Where is ρ resistivity, e is the charge on electron and n is the carrier concentration.

Value of n was calculated by following famous equation (Faraz and Maqsood, 2012)

n=N A Dm PBa

M (3.14)

NA , Dm, M denote the Avogadro’s no., mass density and molecular weight of BT and PBa

denotes the number of barium atoms in the formula unit of BT or PT.

3.5 Electric polarization

BaTiO3 and PbTiO3 are ferroelectric materials. The polarization versus electric field (P-

E) is one of the most important electrical characteristic of ferroelectric ceramics. The

term ferroelectric is derived from ferromagnetic; the iron (Ferro) based compounds in

which magnetization was first noticed based compounds. Ferroelectric usually do not

have iron (Fe) but show analogous electric behavior; they have permanent electric

polarization that can be reversed by the application of the electric field. Hysteresis loop

of ferroelectric materials is similar to the magnetic loop (magnetization versus magnetic

field) of ferromagnetic material.

Fig. 3.5 Hysteresis loop of typical ferroelectric material at room temperature

43

P-E loop measurement method is generally based on the Sawyer Tower method as shown

in Fig. 3. A sinusoidal high voltage is applied across the sample and a large capacitor

which integrates the current into charge and the signal is fed to X-axis channel of

oscilloscope. A resistor work also attenuate high voltage and both signals are plotted as

P-E loop on oscilloscope through X and Y channels. Integrating capacitors have been

replaced by the current to voltage converter; operational amplifiers and a PC can control

and record the data. For a ferroelectric material, we have saturation polarization Ps,

remanent polarization Pr, and coercive field Ec as shown in labelling (Fig. 3). High

remanent polarization values (Pr) relates to the high internal polarizability, strain,

electromechanical coupling and electrooptic activity whereas switching field (Ec) is an

indication of grain size for agiven material; lower Ec accounts for largeer grain size and

higher Ec for smaller grain size. We had made Ferroelectric measurements on

ferroelectric tester and keithly (8512 A) electrometer.

(a) (b)

Fig. 3.6 (a) Schematic of Sawyer and Tower method (b) and Block diagram for P-E

measuring system

44

Chapter 4

Results and discussion

Polycrystalline ceramics, BT with Ba/Ti contents, Pb-doped BT, Bi-doped BT and PT with

Pb/Ti contents were prepared under varying processing parameters by the conventional

solid state sintering reaction method as discussed in chapter 2. In this chapter, stepwise

preparation all the specimens have been given. Optimization parameters at the sintering

of the specimens were studied. The sintered pallets were exposed to a number of

characterizations.

4.1. Thermal analysis of BT ceramics

Thermal behavior, crystallization temperatures of the as ground raw materials were

studied using high temperature differential scanning calorimeter (DSC) well equipped

with thermogravimetric analysis (TGA). The samples were scanned from room

temperature to 1300°C at the heating rate of 10°C/min under the dynamic atmosphere of

insert (Ar) Gas. Fig.5 a, b describes the DCS/TGA thermograms of the as ground raw

powders for both precursor compositions. The powders exhibited broad endothermic

curves in the entire temperature range of about 1170-1190°C and then responded to

exothermic behavior.

Two conspicuous regions were observed in thermograms around 449°C and 810°C that

might correspond to decomposition and melting temperatures of BaCO3.

45

Fig. 4.1 a

Fig. 4.1 b

Fig. 4.1 (a and b), TGA and DSC thermograms of the as ground BaTiO3 specimens, 5a (0.94) and 5b (0.98)

46

First crystallization was initiated at 965°C, the other at 1100-1170°C (Fig 5.b) and then

responded to exothermic region at 1300°C. DCS/TGA thermo gram of both samples were

almost similar with the exception of small endothermic peak at 1100-1170°C that reveals

better crystallization of the powders in accordance with XRD patterns Fig 2. (d)

4.1.1 Thermal analysis of PT ceramics

Thermal behavior, crystallization temperatures of the as ground PT raw materials were

studied using high temperature differential scanning calorimeter (DSC) well equipped

with thermogravimetric analysis (TGA).

Fig. 4.1.1 a

The samples were scanned from room temperature to 1200°C at the heating rate of

10°C/min under the dynamic atmosphere of insert (Ar) Gas. Fig.4.1.1a,b and c describe

the DCS/TGA thermograms of the as ground raw powders for 1.00, 0.98 and 0.94

47

precursor compositions. The powders showed three weight losses in the entire

temperature range.

Fig. 4.1.1 b

Three distinct weight losses can be observed around 269°C, 315°C and 360 °C with

± 5° C for all compositions. For all specimens, DSC curves clearly indicated endothermic

peaks around 269-320°C± 5° C that exhibit the associated decomposition and melting of

PbCO3. In all thermograms above 800°C, well defined upward trends indicate the solid

state reaction between PbO and TiO2 for PbTiO3 formation. For 0.94 and 0.98

compositions, endothermic peaks were followed around 875°C and 1055°C respectively.

Inset A at (inset A, Fig. 4.1.1 b DTGA) reveals endothermic peak at 900°C followed by

two other peaks points around 1000°C and 1100°C. Formation of endothermic peaks

might indicate the gradual increase in the glassy phase for the studied specimens.

48

Fig. 4.1.1 c

Fig. 4.1.1 (a, b and c), TGA and DSC thermograms of the as ground PT powders, a

(1.00), b (0.98) and c (0.94).

4.2. Structural analysis

The microstructure corresponds to grain size, shape, size distribution, porosity size and

porosity distribution and anisotropy.

4.2.1. Preparation analysis and structural properties of BT ceramics

W.P. Chen et al (Chen et al., 2008) investigated correlation of crystal structure and Curie

point with Ba/Ti ratio (0.96-1.04). Phase formation, tetragonality and crystallite size and

porosity of the obtained ceramics were determined from the indexed XRD diffraction

patterns. Ba2TiO4 was detected as secondary phase in both compositions 0.94 and 0.98

(Fig.4.2.1.1 (a – f), Fig. 4.2.1.2 (a - b). An excess of BaO or TiO2 in the precursor

49

compositions usually result in the formation of extra phases like Ba2TiO4, Ba2Ti5O12

(Erkalfa et al., 2003).

Beauger et al (Beauger et al., 1983a) suggested that BaTiO3 is easily formed at the

surface of TiO2 particles that act as catalyst for BaCO3 decomposition, meanwhile the

kinetics are governed by the barium and oxygen ions that diffuse through surface layer of

BaTiO3 into virgin TiO2 phase. The excess of barium and oxygen in the surface layer

generally contribute to the initial formation of Ba2TiO4. Finally the reaction between

Ba2TiO4 and TiO2 result in BaTiO3 formation. Increasing sintering time contributed to

increase in crystalinity, however minor peaks of Ba2TiO4 were detected analogous to the

other reports (Buscaglia et al., 2005) With composition 0.94, besides Ba2TiO4, minor

peaks of TiO2 were observed around 2θ= 43° (fig.4., a - f).

Fig. 4.2.1.1 XRD Patterns of BaTiO3 ceramics (0.94) sintered at 1300°C/2h, pre fired at

1100 - 1200°C/4h (a- b), 1300°C/1-4h (c- f).

50

All the peaks were labeled by comparing the XRD data with JCPD card No. 01-074-

1964, 01-075-2120, 01-083-1880, the perovskite barium titanate synthesis. Ceramics with

both the precursor composition showed tetragonality > 1 and were indexed to the

perovskite (ABO3) ferroelectric materials (Kong et al., 2008b). Cubic (Pm-3m) and

tetragonal (P4mmm, P4MM) crystal structures were indicated. Splitting of 002/200 peaks

around 2θ= 45° that supported the tetragonal symmetry (Fig. 4.2.1.1, e and Fig. 4.2.1.2, c

- d). The widely used indirect method based on XRD data was employed for calculations,

lattice parameters ‘c’ and ‘a’ were taken into account. Lattice parameter c increased with

increasing temperature and time that increased tetragonality of both compositions (Fig.

4a).

Fig. 4.2.1.2 XRD Patterns of BaTiO3 ceramics (0.98) sintered at 1300°C/2h, pre fired at

1300°C/1-4h (a - d).

51

However, cubic perovskite BT was obtained with composition 0.94 (Fig. 4.2.1.1, f)

owning to increase in lattice parameters ‘a, with increasing sintering time. Chekcell

software was used to calculate the lattice parameters. The crystallite size of the BT

ceramics was calculated using Scherrer formula. All the peaks were employed in the

calculations. The density was obtained by measuring the dimensions and mass of the

samples and calculations were made by employing the equation no. (2). Density

increased owning to decrease in porosity with increasing sintering temperature and

sintering time for both compositions (Fig. 4.2.1.5 A and B). Maximum densities were

4.67 and 5.02 g/cm3 with considerable corresponding porosities 22.87 and 16.82 %.

Values of density are relatively less due to the ununiformed coarse large grain structure []

and to the presence of porosity in the obtained specimens (SEM Fig. 4.3.1 c-d).

Crystallite size increased as well owing to coalescence of crystallites (Ramping et al.,

2009) (Fig. 4.2.1.4 B)

Fig. 4.2.1.3. Tetragonality (A) and crystallite size inset (B) of BaTiO3 ceramics sintered

at 1300°C/2h, pre fired at their respective temperatures (a and b) 1300°C/1–4h and (c) at

1100 - 1300°C/4h.

52

Fig. 4.2.1.4 Density of BaTiO3 ceramics (A) sintered at 1300°C/2h, pre fired at the

temperatures (a and b) 1300°C/1–4h, Inset B shows the powders (a) pre fired at 1100 -

1300°C/4h.

4.2.2. Preparation analysis and structural properties of Pb-doped BaTiO3 (PBT)

ceramics

In our previous studies, BT ceramics with precursor composition at 0.98 Ba/Ti molar

ratio remained almost phase pure; depicted better crystallization and were more dense.

They were further employed for doping in order to avoid the development of undesirable

impurity phases. The effect of lead contents on the XRD patterns of BaTiO3 ceramics can

be seen in Fig. 4.2.2.1

53

Fig. 4.2.2.1 XRD Patterns of pure and Pb-doped BaTiO3 ceramics sintered at 1200°C/2h.

XRD patterns indexed at room temperature were employed for determining phase

development, tetragonality and crystallite size. All the peaks were labeled by comparing

the XRD data with JPCD Card No. 01-074-2491 and 01-074-2492. Lattice parameter a

decreased while c increased with increasing lead contents that resulted in the increase of

tetragonality, c/a ratio (Fig. 4.2.2.2) analogous to other results (Chaimongkon et al.,

2011). Tetragonal P4mm crystal structures were indicated. The splitting of 002/200 peaks

around 2θ= 45° supported the tetragonal symmetry.

54

Fig. 4.2.2.2 Tetragonality and crystallite size of pure and Pb-doped BaTiO3 ceramics

PBT ceramics showed tetragonality > 1 and indexed to the perovskite ABO3 ferroelectric

material (Kong et al., 2008a). The lattice parameters were calculated using Chekcell

software. Indirect method based on XRD data was used for calculations; lattice

parameters ‘a’ and ‘c’ were used in calculations. Crystallite size of PBT ceramics were

calculated by employing the Scherrer formula. Crystallite size increased as well owing to

coalescence of crystallites (Ramping et al., 2009). Density increased with increasing lead

contents maximum obtained density was 5.17 g/cm3 (inset A, Fig. 4.2.2.2).

4.2.3 Preparation analysis and structural properties of Bi-doped BaTiO3 (BBT)

ceramics

Since the BT ceramics with precursor composition 0.98 at the present studies showed

better results; they were doped with bismuth nitrate penta-hydrate (Bi (NO3)3 .5H2O. All

55

the indexed peaks were labeled by comparing the XRD data with JPCD No. 01-079-

2265, 01-074-1962 and 01-074-1963.

Fig.4.2.3.1 XRD Patterns of the pure and doped BaTiO3 ceramics sintered at 1150°C/2h,

inset illustrates the diffraction peaks of BaBi4Ti4O15 at 2θ= 27.490° and 30.186° in

magnified version at 0.100 mole % doping

Bi doping obviously affected the crystal structure and phase development (Fig.4.2.3.1).

Lattice parameters, ‘a’ and ‘c’ both increased increasing bismuth, BBT materials were

perovskite ABO3 ferroelectric material with c/a ratio = 1. The splitting of 200 (Bragg)

reflection around 2θ= 45° indicated the tetragonal symmetry for undoped BaTiO3.

However, with increasing Bi contents the peak intensity of 002 was subsided; Cubic Pm-

3m crystal structures were noticed. Barium bismuth titanium oxide existed as BaBi4Ti4O15

with tetragonal crystal structures (14/mmm) around 2θ= 27.490° and 30.186° at 0.100

mole % doping. The peaks were labeled with JPCD Card No. 01-073-2184. attice

56

parameters ‘a’ and ‘c’ were used in calculations. Crystallite size decreased with

increasing bismuth contents due to smaller ionic radii of Bi (1.17 Å) as compared to Ba

(1.61 Å). With increasing bismuth contents (0.075 mole %), density increased to 5.25

g/cm3. Bi2O3 being a low melting additive forms liquid-phase and promotes the sintering

reaction (Yu et al., 2008). However further increase in the Bi contents, the structures

appeared to be porous (SEM studies) with lowering of density (5.20 g/cm3).

Fig.4.2.3.2 Crystallite size and density of the pure and doped BaTiO3 ceramics sintered at

1150°C/2h.

4.2.4 Preparation analysis and structural properties of PbTiO3 (PT) ceramics

Phase formation, tetragonality and crystallite size of the obtained ceramics were

determined from the indexed XRD diffraction patterns. All peaks of the precursor’s

composition 1.00, 0.98 and 0.94 were labeled by comparing the XRD data with JCPD

57

card No. 01-075-1605, 01-072-1135, 01-075-0438, 01-077-2002 and 01-078-0298.

Increasing sintering temperature contributed to increase in crystalinity.

Fig. 4.2.4.1 XRD Patterns of PbTiO3 ceramics sintered at 1190°C/1h; (1.0, a- b) pre-

fired1000 - 1100°C /2h and (0.94, c-e pre fired at 1000 - 1190°C± 5C /2h, 1000 - 1000 -

1190°C± 5C /2h respectively

Evidently, sintering temperature influenced the crystallization and phase purity of the

materials. Minor peaks of TiO2 were detected in the all specimens of 0.94 molar ratio

around 2θ 27.531°, 27.534° and 27.479° (Fig. 4.2.4.1, c-e) at all pre-sintering

temperatures. Existence of minor TiO2 might be attributed to the slight excess of titanium

contents and lead loss during sintering process.

With 0.98 molar ratio, very minor TiO2 peaks were also detected at the pre sintering

regime of 2h /1100°C around 2θ 27.477° (Fig. 4.2.4.2, f). However, on extending the pre

sintering regime to 1190°C± 5° C /2h, no noticeable impurity peaks were detected in

PbTiO3 diffractograms that indicates the completion of sintering (Fig. 4.2.4.2, g).

Sintering duration affected the phase formation and crystalinity of the specimens as well.

58

Occurrence of TiO2 at extended duration, 3-4 h /1100°C is considered due to the lead loss

and (Fig. 4.2.4.2, b and d)

Fig. 4.2.4.2 XRD Patterns of PbTiO3 ceramics (0.98) sintered at 1190°C± 5C /1h, pre

fired at 1190°C/1-4h (a-d) –and 1000- 1190°C/2h respectively

Tetragonality, a relative ratio of the lattice parameter of the c- to the a-axis (c/a)

characterizes the lattice structure of the materials. Ceramics with the all precursor

composition showed c/a ratio > 1, they were perovskite (ABO3) ferroelectric materials.

Tetragonal (P4mm and P4/mmm) crystal structures were indicated. Splitting of 002/200

peaks around 2θ= 45° sustained the tetragonal symmetry. Lattice parameter c increased

with increasing temperature and time and contributed to the increase in tetragonality.

Higher c/a values 1.06427 were observed for 1.00 composition at the pre-firing regime of

59

1000°C/2h. For this composition c/a is slightly higher than the reported value i-e

1.064. Composition (1.00) being rich in lead contents displayed high c/a

values.

Fig. 4.2.4.3 Tetragonality of PbTiO3 ceramics sintered at 1190°C± 5C /1h with precursor

compositions Vs pre-firing temperature, inset A shows teragonality of powders (0.98) Vs

pre-firing duration, 1190°C± 5C /1-4h.

However, c/a values were slightly lowered with extension in the pre-firing temperature to

1100°C (Fig. 4.2.4.3). In addition, for this composition, further extension of the prefiring

temperature to 1190°C± 5° C resulted in the melting of precursors due to the low melting

temperature of PbCO3. Hence, the studies were restricted to the prefiring regime of

1100°C/2h for 1.00 precursor composition. Stoichiometry also influenced tetagonality;

specimens with composition 1.00 displayed greater c/a ratio. Inset A, Fig. 4.2.4.3 shows

the effect of sintering duration on tetragonality. Volatility of lead during extended

sintering duration lowered the described ratio.

60

Fig. 4.2.4.4 Density of PT ceramics sintered at 1190°C± 5C /1h, pre fired at the

temperatures, 1000-1190°C ± 5C /1h, inset A shows the powders (0.98) pre fired at

1100°C 1-4h

Density increased with increasing lead contents and increasing sintering temperature for

all compositions. Values were 6.88, 6.86 and 6.55 g/cm3 for composition 1.00, 0.98 and

0.94 respectively. Considerable density was attained with 0.98 composition at the pre-

sintering regime of 1190°C± 5C /2h. For this composition, no impurity phases were

noticed in the XRD patterns at sintering regime of 1190°C± 5C /2h. However, for

composition 1.00 (rich in lead contents); decrease in density with increasing pre sintering

temperature might be taken due to low decomposing temperatures for PbCO3. Inset A at

Fig. 4.2.4.4 describes continuous decrease in density with increasing pre-firing duration

holding to lead loss.

61

4.3 Microstructural analysis

Morphological features of the ceramics were studied using scanning electron microscopy

(SEM) to observe the surface morphology, grain size and their distribution for the studied

specimens. Sintered samples were coated with gold.

4.3.1 Microstructural analysis of BT ceramics

Sintering temperatures and Ba/Ti molar ratio obviously influenced grain formation and

densification. Large irregular agglomerated grains coupled with needle like formation

were observed in the micrographs of the BT powders, prefired at 1100-1200°C (Fig. 4.3.1

a, b).

Fig. 4.3.1 SEM Micrographs of BT powders and pallets. Powders a and b (0.94) were pre

fired at 1100°C-1200°C/4h , BT pallets c (0.94) and d (0.98) at 1300°C/4h. Inset shows

magnified version of grains at 1µm, headed arrows reveal the fractured surfaces.

62

Needle like grains revealed the existence of Ba2TiO4 phase in accordance with XRD

results (Fig .1 a, b). Significant densification and growth formation progressed on

increasing the prefiring temperature and duration to 1300°C /4h. Grain size varied from

0.5µm- 1 µm and 5µm- 10 µm (Fig.4 c-d). With 0.98 Ba/Ti contents, microstructure of

typical polycrystalline materials was observed with almost interconnected angular and

somewhat rounded grains (Fig.4d). However, the grains were coarse and appeared to be

denser and surrounded by the randomly oriented neighboring grains. They were closely

adhered at the grain boundaries; fractured surfaces were also revealed. On decreasing the

Ba/Ti contents to 0.94, less dense specimens (Fig. 3b) with polygonal, rounded and plate

like grains were obtained (Fig.4 c). Difference in the grain sizes and shapes indicates

varying Ba/Ti contents and the lack of uniformity among the specimens.

4.3.2 Microstructural analysis of PBT ceramics

Fig. 4.3.2 illustrates the morphology of Pb-doped BaTiO3 ceramics. Inset at micrograph

(a) shows the undoped BaTiO3, grains of about 2.5- 10 µm can be observed.

63

Fig. 4.3.2 SEM Micrographs of PBT pallets sintered at 1200°C/2h, inset at (a) shows

undoped BT.

Varying concentration of Pb to barium titanate ceramics showed differences in the

morphologies of PBT. Amount of the grains; their formation, loosening and densification

was observed with Pb doping. There was an obvious grain loosening at 0.025 mole%.

Since Pb has Low decomposing temperature; grain growth was followed by the

increasing concentration of Pb due to the formation of liquid phase in the sintering

process. Almost spherical grains developed with well-defined edges. Since, ionic radius

of Pb (1.20) is smaller than that of Ba (1.34); it effectively reduced the grain size (Fig.

4.3.2 c-d). Grains were 0.1- 4.0 µm in size and became more

connected

4.3.3 Microstructural analysis of BBT ceramics

Fig. 4.3.3 describes the morphology of Bi-doped BaTiO3 ceramics. At micrograph (a),

inset shows undoped BaTiO3, grains of varying size can be observed (2.5- 10 µm).

Obviously the grains size, grain formation and their distribution were changed with Bi

doping. Large differences were noticed in the morphologies of BaTiO3 ceramics. Amount

of the grains, loosening and densification was followed with Bi doping. At 0.025 mole%,

loosening of the grains can be noticed (micrograph b). Since Bi assists in the

densification process due to the formation of liquid phase in the sintering process (Yu et

al., 2008); grain growth was followed with increasing Bi contents.

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Fig. 4.3.3 SEM Micrographs of the pure and doped BT pallets sintered at 1150°C/2h, inset at micrograph (a) shows undoped BaTiO3

The density of the grain boundaries changed with increasing Bi contents; denser

specimens were obtained at 0.075 mole % (micrograph c). Grains were more adhered at

the grain boundaries with some random orientations. Since, ionic radius of Bi is smaller

(1.17 Å) than Ba (1.61 Å), doping effectively reduced the grains size from about 10µm-

0.1 µm. Grains of varying sizes (7µm- 5 µm- 0.1 µm) can be noticed. However, they

looked to be porous and were less dense at 0.1 mole% doping (Fig. 4.3.3 d). The porous

specimen formation might be taken due to the formation of Aurivillius BaBi4Ti4O15

ceramics (Fig. 4.2.3.1). Besides, the reduction in crystallite size (Fig. 4.2.3.2) is in

confirmation of the smaller grain morphology. The differences in the sizes might be

attributed to the different grain growth rates during diffusion process.

4.3.4 Microstructural analysis of PT ceramics

4.3.4.1 Microstructural analysis of PT (1.00) ceramics

Field emission scanning electron microscopy (FE-SEM) was employed to observe the

surface morphology, grain size and its distribution in the PT pallets pallets sintered at

1190± 5C. Precursor composition affected the grain growth and morphology. Obviously

65

grains developed with grain boundaries at the mentioned sintering temperature. But the

precursor composition (1.00: Pb/Ti) containing Pb excess revealed the abnormal grain

growth accompanied by melting regions; that might be attributed to low melting

decomposing temperature of lead in accordance with DSC results (Fig. 4.1.1 a) and large

c/a ratio 1.0643 (Fig. 4.2.4.3 Tetragonality, A )

Fig. 4.3.4.1 FE-SEM micrographs for PT (1.00) pallets sintered at 1190± 5C/1h, pre-

fired at 1100C/2h; white headed arrows points out ferroelectric domain formation while

black headed arrows indicates melting regions.

Fig. 4.3.4.1.1 EDS spectrum of the PT (1.00) pallets sintered at 1190± 5C, pre-fired at

1100C/2h.

66

Sintered PT pallets were chemically analyzed by the scanned EDS spectrum (Fig.

4.3.4.1.1); Pb/Ti molar ratio was 1.00; this ratio was found to be less from the

instrumental analysis. In this regard, loss might be related due to the melting of PbCO3

revealed in the FE-SEM micrographs for this composition.

4.3.4.2 Microstructural analysis of PT (0.98) ceramics

Fig. 4.3.4.2 shows FE-SEM micrographs for PT (0.98) pallets sintered at 1190± 5C/1h.

Grain formation and morphology was strongly affected by stoichiometry. With optimum

0.98 composition, the grains developed with well-defined grain boundaries. Grains were

plate like, angular and hemispherical; closely adhered and surrounded by the neighboring

grains. Grains of varying sizes ranging from 0.5µm- 12 µm were observed. Besides,

formation of ferroelectric domains was clearly observed (Fig. 4.3.4.2 b and inset at a).

Fig. 4.3.4.2 FE-SEM micrographs for PT (0.98) pallets sintered at 1190± 5C, pre-fired at

1100C/2h; inset at micrograph (b) shows PT ceramics at magnified version.

67

Fig. 4.3.4.2.1 EDS spectrum of the PT (1.00) pallets sintered at 1190± 5C, pre-fired at

1100C/2h.

(Fig. 4.3.4.2.1) indicates the EDS spectrum of sintered PT pallets; Pb/Ti molar ratio was

0.98; this ratio was found to be within the almost accuracy of the chemical analysis.

However, error encountered in stoichiometry might be considered due to Pb loss during

sintering.

4.3.4.3 Microstructural analysis of PT (0.94) ceramics

Grains of varying sizes were observed with 0.94 precursor composition. Polyheral,

hemispherical; spherical grains were interconnected with more or less-defined

boundaries. Grain boundaries are basically defects and atoms/molecules at the surface are

at high energy states.

68

Fig. 4.3.4.3 FE-SEM micrographs for PT (0.94) pallets sintered at 1190± 5C, pre-fired at

1100C/2h.

With increase of sintering temperature, they gain more activation energy and agglomerate

to reduce the grain boundaries. The smaller grains join together and form bigger grains.

In addition, for specimens of this composition lower c/a values; 1.06157 (Fig. 4.2.4.3

Tetragonality, A) were observed among other specimens of composition 0.98 and 1.00

Fig. 4.3.4.3.1 EDS spectrum of the PT (1.00) pallets sintered at 1190± 5C, pre-fired at

1100C/2h.

69

Fig. 4.3.4.3.1 depicts EDS spectrum of sintered PT pallets; Pb/Ti molar ratio was0.94;

this ratio was found to be less from the chemical analysis. However, error encountered in

stoichiometry might be considered due to Pb loss during sintering.

4.4 Electrical properties

Most of the electrical studies are restricted at the phase transition (TC= Curie temp.) In the

present studies; we have extended our studies onward to TC. W.P. Chen et al studied

correlation of crystal structure and Curie point with Ba/Ti ratio (0.96-1.04).

In our studies, Curie temperature ranging from 98- 120°C was attained; dielectric studies

were restricted to only 150°C; while conduction process remained unexplored. Perhaps

for the first time; we have investigated electrical properties of the solid state sintered

BaTiO3 ceramics at 0.98 and 0.94 Ba/Ti molar ratio in the wide range of temperatures

(40–700°C) at 1kHz frequency. The studies were focused to find the controlling

parameters, resistance and understanding of the conduction process. In this regard our

paper entitled “Electrical investigations of BaTiO3 ceramics with Ba/Ti contents

under influence of temperature” has been accepted in the Journal of Solid State

Ionics on Feb. 3, 2017.

In addition; this Article got 1st prize by the chairman, Higher education commission

(HEC) presented in the 1st international conference Air University, E-9 PAF Complex,

Islamabad, Pakistan, dated Feb. 16, 2017.

For a given composition, TC and the electrical, mechanical and optical properties strongly

depend on the microstructure.

4.4.1. Dielectric studies of BT ceramics

The dielectric parameters such as dielectric constant (ԑ΄), dielectric loss tangent (δ ), of

the sintered ceramics (1300°C /2h) were measured by parallel plate technique in the wide

range of temperatures (40 -700°C ) at 1kHz frequency. Samples were heated in a home-

made programmable furnace at the rate of 2C/min; the temperature was controlled by a

70

thermocouple contacting the sample holder near where the sample was situated.

Calculations were made by using the relation.

ԑ΄ = Cd/ Ɛ0 A

Where d is the thickness of pallet, C is the capacitance, Ɛo is the permittivity of free space

and A is the cross-sectional area of the pallet. Characterization of ԑ΄ as a function of

temperature is the most important tool to observe phase transition in ferroelectric system.

In general at different frequencies, ԑ΄ increases with increase in temperature and is

influenced by dipolar, interfacial, ionic and electronic polarizations (Singh et al., 2002).

Dielectric constant increases at lower frequencies while decreases at higher frequencies

which is a normal behavior of the ferroelectric material. Since at higher frequencies, the

dipoles cannot follow the applied AC electric field; decrease in the dielectric constant is

observed. At lower frequencies, the dipolar and interfacial polarization contribute

significantly to the dielectric constant (Lines and Glass). Both of these are temperature

sensitive; hence, the dielectric constant increases at advance rates at 1kHz as compared to

other higher frequencies like 10 kHz, 100 kHz (Silveira et al., 2013)etc. Hence, dielectric

measurements were carried out at 1kHz frequency in the above mentioned range of

temperature.

Specimens were sintered at 1300°C /2h. In Fig. 4.4.1, insets A (a-f) and B (c-d) describe

the effect of varying sintering temperature and sintering time on the dielectric constant of

BT ceramics. Increasing sintering temperature and prolongation of sintering time resulted

in the crystallization (Fig. 4.2.1.1 a-f and Fig. 4.2.1.2 a-d) and densification of specimens

with conspicuous grain growth (Fig. 4.3.1 a-d) (Kim et al., 2004; Lee et al., 2015; Wang

et al., 2015); dielectric constant enhanced accordingly.

71

Fig. 4.4.1. Temperature dependence of dielectric constant (ԑ΄) and loss tangent (δ) of the

BaTiO3 ceramics C (f and d) at 1k Hz frequency. Insets (A and B) show their pre firing

temperature Vs time.

Temperature dependence of dielectric constant (ԑ΄) can be divided into two temperature

regions, region (I) from 40-200°C corresponding to structural phase transition from

ferroelectric to paraelectric. Region (II) from 200 -700°C correspond to high temperature

region where dielectric anomalies were observed. In region (I), dielectric constant

increased with increasing temperature, maximum dielectric constant was observed at

phase transition (TC = Curie temperature). In the ferroelectric phase, below TC, BT has

tetragonal symmetry with permanent electric dipole. These dipoles are ordered with a

domain structure. At Curie temperature, ionic and electronic polarizability are at the peak

levels and dielectric constant attains the maximum value (Forsbergh, 1949; Merz, 1949).

72

Among all specimens, f (0.94) and d (0.98) showed maximum dielectric constant at TC,

1110 and 1520 that are considerably low as compared to other studies (Chen et al., 2008;

Yasmin et al., 2011) due to stresses generated owning to large grain structures(Wng et

al., 2003). Moreover, as fully dense materials were not obtained; porosity presence also

lowered the permittivity values and resulted in the increase of loss tangent (Fang et al.,

1993; Herbert, 1985).TC was shifted from 120-130°C with decreasing Ba/Ti contents in

accordance with other results (Chen et al., 2008). Specimen f and d were employed for all

further investigations in the wide temperature range of 40 -700°C.

In the paraelectric phase, BT has cubic symmetry with large dipole

moments. With increasing temperature, the randomness of dipoles enhances that

contribute to decrease in dipolar polarization. Above TC, in the cubic structure, the Ti4+

ions oscillate about the centers of the TiO6 octahedra without effective mutual coupling,

but with large dipole moments, because the Ti4+ ion has a tendency to change from ionic

to covalent bonding as its distance to an atom decreases (Loge and Suo, 1996). Dielectric

anomalies were observed for both materials (f and d) in the vicinity of 500°C (Fig. 4.4.1,

C) analogous to other studies for BT system (Leyet et al., 2010).

At 700°C, anomalous increase in ԑ΄and δ with temperature reveals the presence of

thermally activated transport properties in the materials. It is well known, oxygen

vacancies are major structural defects in the barium titanate (Kang and Choi, 2002;

Lemanov et al., 2000), generated due to loss of oxygen during sintering at high

temperature in accordance with following relation, a process defined by Kröger-Vink

notation (Ang et al., 2000a)

V O → 12

O2+V o⦁⦁+2 e¯

Kang et al (Kang et al., 2003b) related dielectric anomalies to the competition

phenomenon of the dielectric relaxation and the electrical conduction by oxygen

vacancies.

73

4.4.1.1 Ac conductivity studies

The ac conductivity was evaluated from dielectric constant (ԑ΄) and dielectric loss tangent

tan (δ) in the same frequency range (1kHz) by employing the following relation

σ ac = ω Ɛ0 ԑ΄ tan(δ)

Where, σ ac is the ac conductivity and ω = 2 πf , the angular frequency, ԑ΄ is the dielectric

constant and tan (δ) is the loss tangent. Insets A and B (Fig. 4.4.1.1) show ac conductivity

from 40-120°C-130°C at varying processing parameters. Ac conductivity increased with

increasing sintering temperature and prolongation of sintering time in accordance with

dielectric studies. Maximum value ac conductivity around the vicinity of phase transition

were 2.41×10−6 Sm−1 and 2.57 ×10−6 Sm−1, for f and d specimens respectively. As fully

dense materials were not obtained; morphology of the specimens indicated the presence

of porosity. The presence of pores might provide conduction path to electricity; could be

a main contribution of tan (δ) and thus to ac conductivity (Fang et al., 1993; Herbert,

1985). The values are almost in the pact of other results (Silveira et al., 2013).

However, conductivity continued to increase with increasing temperature, owning to

increase in dielectric constant and dielectric losses. Fig. 4.4.1.1, C shows the variance in

conductivity with temperature. At higher temperatures, like dielectric response, electrical

anomalies were observed. Conductivity approached to about order jump of

1.9×10−3 Sm−1−2.85× 10−3 Sm−1 for specimen f and d at 700°C, high values indicate the

long range movement of the charge on thermal excitation.

Ang et al (Ang et al., 2000a) described that doubly ionized oxygen vacancies can move

due to thermal excitation. Conduction electrons created by the ionization of oxygen

vacancies can cause hopping of electrons between Ti 4+ and Ti 3+. Enhanced hopping of

electrons at higher temperatures might have decreased the resistance of grains thereby

increasing the probability of electrons reaching the grain boundaries.

74

Fig. 4.4.1.1 Variance of ac conductivity for BaTiO3 ceramics (C) with temperature at 1k

Hz frequency. Insets (A and B) describe pre firing temperature of specimens Vs time.

Increased polarization had possibly elevated the ac conductivity of the studied materials.

Hence, conduction could occur through an electron-hopping mechanism (La Course and

Amarakoon, 1995) at low temperatures and at higher temperatures, σ ac tends to the value

of dc conductivity on thermal activation. Moreover, interpretation of different theoretical

models (Funke, 1993; Jonscher, 1996; Ngai, 1993) concludes that ac conductivity

originates from migration of ions by hopping between neighboring potential wells at

lower temperatures which eventually give rise to dc conductivity at high temperatures at

lower frequencies.

75

4.4.2 Dielectric studies of PBT ceramics

In our previous studies, BT ceramics with composition 0.98 showed more dielectric

behavior. This composition was sintered at 1300°C/4h, was employed for lead doping.

Pre-sintering was adopted to avoid the development of impurity phases. The studies were

investigated in the wide range of temperatures (40–700°C) at 1kHz frequency for PBT

ceramics as well. PBT tablets were sintered at 1200°C /2h. These studies entitled

“Fabrication and electrical investigations of Pb-doped BaTiO3 ceramics” has been

accepted in the Journal of Materials Chemistry and Physics on Jan. 31, 2017.

Fig. 4.3.2 displays the dielectric constant of undoped BaTiO3 and PBT ceramics with

temperature. The temperature dependence of ԑ΄ can be divided into two regions. Region

one from 40-200°C correspond to structural phase transition from ferroelectric to

paraelectric phase. In region one, dielectric constant increased with lead contents. With

lead contents dielectric constant increased owning to grain formation (Kim et al., 2004;

Lee et al., 2015; Wang et al., 2015). Curie temperature was shifted from 120-200°C (TC =

Curie temp). The specimens showed dielectric constant in the vicinity of TC, 1500- 1730,

the maximum value was obtained with 0.1 mole %. However, the values are considerably

low as compared to other studies (Mudinepalli et al., 2014) possibly due to the stresses

generated owning to large grain structures (Wng et al., 2003). Ferroelectric to paraelectric

phase transition peaks were more diffused and broadened. The diffused phase transition

occurs (Ye, 2002) mainly due to compositional fluctuation and/or substitutional disorder

in the arrangement of cations in one or more crystallographic sites of the perovskite

structure, which leads to microscopic heterogeneity in the compound with different local

Curie point. However, in region two (200- 700°C), above 200°C, the dielectric response

changed to a relaxor behavior with increasing temperature and dielectric anomalies were

observed. Occurrence of the dielectric relaxation at low frequency can be related to the

space charges in association with the oxygen vacancies (Mudinepalli et al., 2014) that can

be trapped at the grain boundaries of the electrode-sample interface.

76

Fig. 4.4.2.1 Temperature dependence of dielectric constant (ԑ΄) and dielectric loss tangent

(δ) for pure and Pb-doped BaTiO3 ceramics at 1k Hz frequency.

The specimens, d and e showed dielectric anomalies, this characteristic was more

pronounced at 0.1 mole % Pb doping. Increase in ԑ΄ and δ with temperature reveal the

presence of thermally activated transport properties in the materials.

It is known, oxygen vacancies are major structural defects in the barium titanate (Kang

and Choi, 2002; Lemanov et al., 2000) generated due to loss of oxygen during sintering at

high temperature. Moreover, volatilization of PbO occurs during stages of powders

calcination and sintering of Pb-based compounds at high temperatures. Such

volatilization provides both fully- ionized cationic lead vacancies (V''Pb) and anionic

oxygen vacancies (V o⦁⦁) (Eyraud et al., 1984; Eyraud et al., 2002). At low temperatures,

the lead vacancies are quenched defects, difficult to be activated. They could become

mobile with activation energy, Ea values around and above 2 eV (Guiffard et al., 2005).

77

Dielectric anomalies at the high-temperature region have been reported for BaTiO3,

(Pb,La)TiO3 and (Pb,La)(Zr,Ti)O3 systems system (Kang et al., 2003b); anomalies were

related to the competition phenomenon of the dielectric relaxation and the electrical

conduction by oxygen vacancies. The role of oxygen and lead vacancies would be

clarified in the dc conduction studies.

At the high temperatures, increase in tangent (δ) in the paraelectric region for PBT

ceramics are taken due to thermally activated conduction losses.

4.4.2.1 Ac conductivity studies of PBT ceramics

Insets A (Fig. 5) shows the enlarged view of the selected area of graph ‘A’ from 40-

200°C that corresponds to the phase transition. The ac conductivity increased in the

vicinity of phase transition from 2.57 ×10−6- 1.7 ×10−4 Sm−1with lead doping for all the

specimens; the specimen with 0.1 mole % doping showed maximum conductivity. Near

the Curie temperature, the domain structure break up, carriers become free and take part

in conduction by trapping mechanism. After down fall at phase transition Conductivity

decreased (Fig. 4.4.2.1). Above 200°C, ac conductivity began to increase gradually;

electrical anomalies were observed with the rising temperatures; where loss tangent, tan

(δ) effectively contributed to the conduction process. Interpretation of different

theoretical models (Funke, 1993; Jonscher, 1996; Ngai, 1993) concludes that ac

conductivity originates from migration of ions by hopping between neighboring potential

wells at lower temperatures which eventually give rise to dc conductivity at high

temperatures. Conductivity approached to about two order jump of 2.85 ×10−3 Sm−1 -

1.83 ×10−2 Sm−1with lead doping (0.00- 0.100 mole %) at 700°C. High values indicate

the long range movement of the charge carriers on thermal excitation. Oxygen vacancies

are the most mobile charge carriers in oxide ferroelectrics and play an important role in

the conduction process in most dielectric ceramics (Liu et al., 2011; Rani et al., 2013).

Around the oxygen vacancy, long range potential wells may be formed, there can be large

number of titanium centers within each potential well surrounding the oxygen vacancy.

Conduction electrons created by the ionization of oxygen vacancies can cause hopping of

electrons between Ti 4+ and Ti 3+. Thus dc conductivity may be associated with the

hopping between the long range potential wells created by the oxygen vacancies, while

78

the ac conduction at low temperatures may occur through the charge carrier motion over

a short range distance between sites in the potential well.

Fig. 4.4.2.1.1. Variance of ac conductivity for pure and Pb-doped BaTiO3 ceramics with

temperature at 1k Hz frequency. Inset A shows magnified version of the selected area

confirming the phase transition in accordance with dielectric studies.

4.4.3 Dielectric studies of BBT ceramics

BT ceramics with composition 0.98 was employed for fabricating doping. This

composition was sintered at 1300°C/4h, was employed for lead doping. The studies were

investigated in the wide range of temperatures (40–700°C) at 1kHz frequency for PBT

ceramics as well. BBT tablets were sintered at 1150°C /1h Fig. 4.3.2 shows the variance

of dielectric constant of BaTiO3 and BBT ceramics with temperature.

79

Fig.4.4.3 Temperature dependence of dielectric constant (ԑ΄) for the pure and Bi doped

BT ceramics at 1k Hz frequency; inset A, shows the enlarged view of the selected area

corresponding to phase transition.

The dielectric constant increased with increasing Bi contents and reached the maximum

value at 0.075 mole % doping owning to the increase in densification. Bismuth ions

doped in the BaTiO3 normally occupies the Ba sites due to nearly similar ionic radii and

electro negativities. Therefore, substitution of Bi+3 for Ba+2 takes place according to

Kröger-Vink notation (Ang et al., 2000b)

Bi2O3 → 2Bi⦁Ba + V ˝ Ba + 3Oox (1)

Where,2 Bi⦁Bais an ionized Bi donor, V ˝ Ba a doubly ionized barium vacancy andOox

stands for neutral oxygen atom on oxygen site. The bismuth ions located at A-sites of

BaTiO3 lattice will take positive charge which in turn, be compensated by electron and/ or

80

barium vacancies to keep the charge balance. Curie temperature increased from 120-

160°C (TC= Curie temp.) with Bi doping. Dielectric constant increased from 1520-2205

upto 0.075 mole % doping.

When Bi3+ ion substitutes for Ba2+ ion in BaTiO3, ion volume of A-site decreases due to

the barium vacancy, which makes a bigger active space for Ti4+ addition. For increase of

electrovalence from +2 to +3 of A-site, a residual positive charge appears and mutual

effect between A and B sites increases strongly. Consequently, polarization of Ti+4 is

enhanced which results in the increase of dielectric constant. The values of dielectric

constants were lower as fully dense materials were not obtained. Sharp dielectric constant

peak was observed for undoped BaTiO3 (inset B) and transformed to broad diffused

peaks at increasing doping levels. With increasing Bi doping, formation ofBi⦁Ba and V ˝ Ba

increases that resultantly enhances the inhomogeneity among the specimens. Therefore,

diffused nature of phase transition on Bi doping can be attributed(Vugmeister and

Glinchuk, 1990; Ye, 2002) mainly due to compositional fluctuation and/or substitutional

disorder in the arrangement of cations in one or more crystallographic sites of the

perovskite structure, which leads to microscopic heterogeneity in the compound with

different local Curie points.

D. Gulwade et al (Gulwade and Gopalan, 2009) described that the partial substitution of

Ba or Ti ions with Bi causes the appearance of ferroelectric relaxor like behaviour. With

further increase in temperature, above 200°C, a relaxor like dielectric behavior continued

to be displayed; dielectric anomalies were observed particularly at doping levels of 0.075

and 0.100 mole %. At 700°C, this effect was predominantly pronounced at 0.075 – 0.1

mole % doping. However, the dielectric constant decreased owning to the existence of

BaBi4Ti4O15 at 0.1 mole % doping. Although, trivalent Bi ion could be compensated by

the barium vacancies, yet oxygen vacancies can easily be created due to the loss of

oxygen from the crystal lattice during sintering at high temperatures (Ang et al., 2000b).

Oxygen vacancies can be neutralV o , singly ionized V o⦁, and doubly ionizedV o

⦁ ⦁.

Ionization of oxygen vacancies create conduction electrons process may bond to Ti 4+ and

cause reductionof valence Ti 4+ → Ti 3+. In the paraelectric regions, hopping of electrons

between Ti 4+ and Ti 4+, might lead to the conduction process; the combination of single

81

ionized oxygen vacancies and electrons of the oxygen vacancies in the second ionization

state (Ang et al., 2000a); doubly ionized oxygen vacancies, V o⦁⦁ are considered the most

mobile charge carriers in most perovskite mostly in titanates; play important role in

conduction.

Oox → ½ O2 +¿ V o

⦁⦁ +¿ 2 e−¿¿

Fig. 4.4.3.1 Loss tangent, tan (δ) of the BBT ceramics as a function of temperature at 1k

Hz frequency

Fig. 4.3.3.1 gives the effect of temperature on loss tangent (δ), tan (δ). At high

temperature, in the paraelectric regions; thermally activated conduction losses might be

attributed to the increase in tan (δ) for BBT ceramics.

4.4.3.1 Ac conductivity studies of BBT ceramics

The ac conductivity almost increased with bismuth doping (0.025- 0.075 mole %),

maximum conductivity in the vicinity of phase transition was 1.92 ×10−6- 4.8 × 10−6 Sm−1

. However, at 0.1 mole % doping, conductivity decreased due to the existence of

82

BaBi4Ti4O15 in agreement with dielectric studies. Afterwards with phase transition,

conductivity continued to decrease with a classical behavior most probably due lower

loss tangent, tan (δ) values. Insets, A (Fig. 4.3.3.1) shows magnified view of the selected

area of the graph corresponding to the phase transition (40-200°C).

Fig. 4.4.3.1.1 Conductivity of the pure and doped BaTiO3 ceramics Vs temperature at 1k

Hz frequency. Inset A, the magnified version of the selected area confirms the phase

transition in agreement with dielectric studies.

No appreciable increase in tan (δ) was observed up to 300°C. Above 320°C; conductivity

began to increase with further rising temperatures, electrical anomalies with a relaxor like

behavior were observed. Increase in ԑ΄ and δ with temperature revealed the presence of

thermally activated transport properties in the materials. Ac conductivity may originate

from migration of ions by hopping between neighboring potential wells at lower

temperatures which eventually give rise to dc conductivity at high temperatures. At

700°C, conductivity increased about three times with Bi doping (0.00- 0.100 mole %)

83

2.85 ×10−3 Sm−1 - 9.03×10−3 Sm−1 in consistence to our previous studies. Details for this

anomalous increase in conductivity had already been explained in the previous sections.

4.4.4 Dielectric studies of PT ceramics

Most of the studies are concerted at the crack free sintering with special additives to get

the denser PT ceramics. Perhaps for the first time; we have tried the solid state sintering

coupled with electrical investigation of the PT ceramics at 1.00 0.98 and 0.94 Pb/Ti molar

ratio in the wide spectrum of temperatures (40–700°C) at 1kHz frequency. The studies

were focused to find the controlling parameters, resistance and understanding of the

conduction process and Arrhenius dependence of the specimens. Specimens were

sintered at 1190°C ± 5C /1h after several attempts.

Fig.4.4.4.1-4.3 displays temperature dependence of dielectric constant (ԑ΄), dielectric loss

tangent (δ) of the sintered PT ceramics for all stoichiometric compositions (1.00, 0.98

and 0.94). Stoichiometry exerted a significant influence on the dielectric constant and

phase transition for all PT ceramics. Among all specimens, ceramics with 0.98 Pb/Ti

molar ratios showed greater values of dielectric constant with conspicuous phase

transition at the pre sintering temperatures, 1190°C± 5° C /2h. Composition 1.00 equally

exposed phase transition predominantly at 490°C. Dielectric constant values at the phase

transition were 1930, 2307 and 884, 1870, 4600 for 1.00 and 0.98 compositions. For

composition 0.94, dielectric constant continued to increase around phase transition; the

values were 1543, 1760 and 549 at their respective pre-sintering temperatures (Fig.

4.4.4.3). Composition (0.94) being lesser in Pb contents showed lower values of ԑ΄ at

higher pre-sintering temperature, 1190°C± 5° C which might be taken due to additional

Pb losses.

Among all compositions, PT ceramics with 0.98 Pb/Ti contents showed higher values at

the pre-sintering temperatures of 1190°C± 5° C. For this composition, X-ray

crystallographic studies revealed no impurity phases as TiO2; denser specimens were

obtained as well (Fig. 4.2.4.4).

84

For all compositions, with rising temperature, anomalous increase in ԑ΄ and ԑ΄΄ values

assures the presence of thermal transport process. Pronounced dielectric anomalies were

noticed all specimens at elevated temperatures.

Fig.4.4.4.1 Temperature dependence of dielectric constant (ԑ΄) for the PT ceramics (1.00)

at 1kHz frequency; inset A, shows the enlarged view of the selected area corresponding

to phase transition.

85

Fig.4.4.4.2 Temperature dependence of dielectric constant (ԑ΄) for the PT ceramics (0.98)

at 1kHz frequency; inset A, shows the enlarged view of the selected area corresponding

to phase transition.

It has already been explained; oxygen vacancies are major structural defects in the most

ferroelectrics (Kang and Choi, 2002; Lemanov et al., 2000); play an important role in

conduction. Dielectric anomalies at the high-temperature region have been reported for

(Pb,La)TiO3 and (Pb,La)(Zr,Ti)O3 systems (Kang and Choi, 2002); anomalies were

related to the competition phenomenon of the dielectric relaxation and the electrical

conduction by oxygen vacancies.

Electrons created by ionization of thermally activated oxygen vacancies captured by Ti 4+

might cause hopping of electrons between Ti 4+ and Ti 3+ in the form Ti 4+ +¿ e¯ → Ti 3+.

The short range hooping of oxygen vacancies might contribute to the dielectric relaxation

for the studied specimens.

86

For (Pb1-xLax) (Zr0.90Ti 0.10)1-x/4O3 (PLZT) compositions, oxygen vacancies were

considered as the most mobile defects (Pelaiz-Barranco et al., 2008b) ; their influence on

the dielectric relaxation processes had been reported.

In addition, volatilization of PbO (Eyraud et al., 1984; Eyraud et al., 2002) during stages

of powders calcination and sintering of Pb-based compounds at high temperatures

provides both fully- ionized cationic lead vacancies (V''Pb) and anionic oxygen vacancies

(V o⦁⦁).The volatility of PbO can also effect the equilibrium and defect choice (non-

stoichiometry).

The role of V o⦁⦁ and V''Pb would be elucidated in the dc conduction studies

Fig.4.4.4.3 Temperature dependence of dielectric constant (ԑ΄) for the PT ceramics (0.94)

at 1kHz frequency; inset A, shows the enlarged view of the selected area of phase

transition.

87

Fig. 4.4.4.1.1. Loss tangent, tan (δ) of the PT (1.00) ceramics as a function of temperature

at 1k Hz frequency

Fig. 4.4.4.2.1 Loss tangent, tan (δ) of the PT (0.98) ceramics as a function of temperature

at 1k Hz frequency.

88

Fig. 4.4.4.3.1 Loss tangent, tan (δ) of the PT (0.94) ceramics as a function of temperature

at 1k Hz frequency

4.5 Electrical dc resistivity and dc conductivity studies of BT ceramics

(0.94 and 0.98)

Two probe method was manipulated to measure the electrical dc resistivity (ρ) at room

temperature and from 40 -700°C by employing circular pallets of 2mm thickness and

12mm diameters sintered at1300°C /2h . Pressure contacts equal to the pallet size were

applied after polishing the surfaces

4.5.1 Electrical dc resistivity and dc conductivity studies of BT ceramics

In Fig. 4.4.1 insets A (a-f) and B (c-d) describe the effect of varying sintering temperature

and sintering time on the resistivity of BT ceramics. Increasing sintering temperature and

prolongation of sintering time resulted in the increase of resistivity (Chen, 2007;

89

Nowotny and Rekas, 1991). Specimen f and d showed the greater resistivity. Resistivity

measured at the room temperature for f (0.94) and d (0.98) were 4.5 × 109Ω cm,

3.5×109 Ωcmrespectively in agreement with the reported literature (Nowotny and Rekas,

1991)

Fig. 4.5.1 Resistivity of BaTiO3 ceramics (C) Vs temperature in the ferroelectric and

paraelectric regions at 1k Hz frequency. Insets (A and B) depict resistivity at 40-200°C.

Resistivity plots can be divided into regions. The first region from 40-120-130°C, up to

phase transition decrease in resistivity was slow that may be due to the ordered state of

ferroelectric phase. The second region from 140-700°C (Fig. 8C) where decrease in

resistivity was more rapid, semiconductor behavior was observed. Oxygen vacancies are

the most mobile charge carriers in oxide ferroelectrics and play an important role in the

90

conduction process in most dielectric ceramics (Liu et al., 2011; Rani et al., 2013).

Oxygen vacancies can be neutral, single and doubly ionized respectively V o, V o⦁, V o

⦁⦁.

Activation energies, Ea of 0.3-0.5 and 0.6-1.2 eV are typically assigned to single ionized

and doubly ionized oxygen vacancies (Ang and Yu, 2000; Ciomaga et al., 2011).

Conductivity of the specimens (f and d) was evaluated from the impedance spectrum by

using the relation (Xu et al., 2016)

σdc= d

A × R

Where d is the sample thickness, A is the electrode area; R is the resistance of bulk

ceramics. Fig. 9a and 9b show the variation in dc conductivity of both samples with

inverse of absolute temperature, which followed the Arrhenius Law (Kang et al., 2015).

Both specimen f and d showed semiconductor behavior with negative temperature

coefficient of resistivity (NTCR) characteristics.

σ = σ₀ exp (−EakB T )

Where σ₀ is the pre-exponential factor, Ea is the activation energy, kB is the Boltzmann’s

constant and T is the absolute temperature in kelvin.

91

Fig. 4.5.1.2 ( a) Variation in conductivity (σ dc) with inverse of absolute temperature for

BT ceramics in the temperature range200- 360°C.

A close agreement between the experimental conductivity data and Arrhenius curves was

obtained for f and d specimens in the range of 200- 360°C and 480- 700°C ( Fig. 4.5.1.2

(and Fig. 4.5.1.3); respective estimated Ea values were 0.2948- 0.3284 and 0.967- 1.189

eV. Obtained Ea values were associated to singly ionized and doubly ionized oxygen

vacanciesV o⦁, V o

⦁⦁.

Hence, with increasing temperature, conduction seemed to be governed by the

combination of single ionized oxygen followed by the electrons of the oxygen vacancies

in the second ionization state for BaTiO3 ceramics (Leyet et al., 2010).

92

Fig. 4.5.1.3 Variation in conductivity (σ dc) with inverse of absolute temperature for BT

ceramics in the temperature range480- 700°C.

We propose that ionic conduction may be responsible for the conduction process for the

studied ceramics. Although both specimens depicted semiconductor behavior, yet

specimen, f was conspicuously more resistive than d in the entire temperature range.

Ceramics with more Ti contents, specimen, f (0.94) showed greater resistivity as

compared to that with more Ba contents, specimen, d (0.98) in agreement with other

studies (T-Falin et al., 1990).

4.5.1 dc mobility studies of BT ceramics

Drift mobility is in fact, a proportionality factor between the drift velocity of the charge

carriers in semiconductor and electric field. Drift velocity is determined by the mobility

of charge carriers. Drift mobility, μd was calculated at the temperatures (480 - 700°C)

using the following relationship

μd=1

neρ

93

Where is ρ resistivity, e is the charge on electron and n is the carrier concentration. Value

of n was calculated by following famous equation Where is ρ resistivity, e is the charge

on electron and n is the carrier concentration. Value of n is calculated by following

equation (Faraz and Maqsood, 2012)

n=N A Dm PBa

M

NA , Dm, M denote the Avogadro’s no., mass density and molecular weight of BT and

PBadenotes the number of barium atoms in the formula unit of BT.

4.5.1.4 Dependence of drift mobility with inverse of absolute temperature for BT

ceramics.

Fig. 4.5.1.4 displays temperature dependence of drift mobility for materials f and d. Our

Ea values calculated at the temperatures 200-360°C and 480-700°C, were in the range of

singly and doubly ionized oxygen vacancies V o⦁⦁. Since doubly ionized oxygen

vacancies can move due to thermal activation (Ang and Yu, 2000), mobility of electrons

increased with rise in temperature due to hopping of charge carriers from one site to

94

another owning to decrease in resistivity for both materials. Ceramics with more Ba

contents (specimen, 0.98 ) showed greater drift mobility as compared to that with more Ti

contents (specimen, 0.94) in accordance with conductivity studies (Fig. 4.3.1 and 4.4.1).

4.5.2 Electrical dc resistivity and dc conductivity studies of PBT

ceramics

Resistivity decreased with lead doping (Nowotny and Rekas, 1991), PBT materials with

0.1 mole% doping were observed to be less resistive.

Fig. 4.5.2 Resistivity of pure and Pb-doped BaTiO3 ceramics Vs temperature in the

ferroelectric and paraelectric regions at 1k Hz frequency.

Resistivity plots can be divided into two regions. In the first region from 40-200°C

decrease in resistivity was slow that may be due to the ordered state of ferroelectric

phase. The second region from 200-700°C, the paraelectric region where decrease in

resistivity was more rapid, all specimens responded to semiconductor behavior (Fig.

95

4.5.2). Since oxygen vacancies are formed in the process of sintering (Kang et al., 2003a)

and are the most mobile charge carriers in oxide ferroelectrics and play an important role

in the conduction process in most dielectric ceramics (Liu et al., 2011; Rani et al., 2013).

Oxygen vacancies can be neutral, single and doubly ionized respectivelyV o, V o⦁, V o

⦁ ⦁. At

room temperature, the oxygen vacancies exhibit low mobility, whereby the ceramics

indicate enhanced resistance (Pelaiz-Barranco et al., 2008a). On activation with

increasing temperature, resistance decreased with the observed electrical behavior. Fig.

4.5.2 show the variation in dc conductivity of both samples with inverse of absolute

temperature, which followed the Arrhenius Law (Xu et al., 2016).

Fig. 4.5.2.1 Variation in conductivity (σ dc) with inverse of absolute temperature for pure

and doped PBT ceramics

All specimens showed semiconductor behavior with negative temperature coefficient of

resistivity (NTCR) characteristics. A close agreement between the experimental

conductivity data and Arrhenius curves was obtained for all materials in the range of 480-

700°C (Fig.6); respective estimated Ea values were 1.187, I.361, 1.184, 1.172, 1.169 eV

96

(0.00- 0.100 mole % doping). Steinsvik et al (Steinsvik et al., 1997) reported, the

activation energy for ABO3 perovskites decreases with the increase of oxygen vacancies

contents, our studies followed the reported pattern. Ea values of 0.6-1.2 eV are commonly

associated to doubly ionized oxygen vacancies (Ang and Yu, 2000). Our estimated Ea

values are within the said range. Thus, values of Ea for the obtained PBT ceramics

suggest that electrical conduction in the high temperature range was ionic due to doubly

ionized oxygen vacancies; not to the lead vacancies (V''Pb).

4.5.2.1 dc mobility studies of PBT ceramics

Drift mobility, μd was calculated at the temperatures (480 - 700°C).

4.5.2.1 Dependence of drift mobility with inverse of temperature for PBT ceramics.

Temperature dependence of drift mobility for PBT ceramics (0.00- 0.100 mole %) can be

seen in Fig. 4.5.2.2. Marked dc mobility was observed at o.1 mol% doping (Fig. 8) in

agreement with the dc resistivity studies (Fig. 4.4.2). Our Ea values calculated at the

temperatures 480-700°C lied in the range of doubly ionized oxygen vacancies V''O. Since

97

doubly ionized oxygen vacancies can move due to thermal activation (Ang and Yu, 2000)

most probably mobility of electrons increased with rising temperature due to long range

hopping of charge carriers from one site to another owning to decrease in the resistivity.

4.5.3 Electrical dc resistivity and dc conductivity studies of BBT

ceramics

Dependence of resistivity of BT ceramics with dopant contents and temperature (ρt) is

shown in the Fig. 4.4.3 Bismuth doping obviously lowered the resistivity of specimens.

Up to phase transition, decrease in resistivity was slow duo to the ordered state of

ferroelectric phase.

Fig. 4.5.3.1 Resistivity of the pure and BBT ceramics Vs temperature in the ferroelectric

and paraelectric regions at 1k Hz frequency.

98

However, in the paraelectric regions; at the present doping level, resistivity decreased

rapidly with the increasing temperature. Near the Curie temperature, the domain structure

break up, carriers become free and take part in conduction mechanism. All specimens

responded to semiconductor behavior with negative temperature coefficient of resistivity

characteristics (NTCR). Conductivity of the specimens was evaluated from the

impedance spectrum by using the equation no. Variation in dc conductivity with inverse

of absolute temperature followed the Arrhenius Law (Fig. 4.4.3.2)

Fig. 4.5.3.2 Variation in conductivity (σ dc) with inverse of absolute temperature for the

pure and doped BBT ceramics

Experimental conductivity data for all materials fitted with close agreement to Arrhenius

curves in the range of 480- 700°C (Fig. 4.5.3.2); respective estimated Ea values were

1.189, I.105, 1.063, 1.2595, 1.157 eV (0.00- 0.100 mole % doping). Steinsvik et al

(Steinsvik et al., 1997) described, the activation energy for ABO3 perovskites decreases

with the increase of oxygen vacancies. It appears that with increasing Bi contents, more

oxygen vacancies are likely to be created. Ea values of 0.6-1.2 eV are commonly

99

connected to doubly ionized oxygen vacancies (Ang and Yu, 2000).Our estimated Ea

values followed the described range. Thus, we propose that ionic conduction might be

responsible for BBT ceramics followed by the contribution of the doubly ionized oxygen

vacancies.

4.5.3.1 dc mobility studies of BBT ceramics

Drift mobility, μd at the temperatures (480 - 700°C).

Fig. 4.5.3.3 Dependence of drift mobility with inverse of temperature for BBT ceramics.

Fig. 4.5.3.3 shows temperature dependence of drift mobility for BBT ceramics (0.00-

0.100 mole %); pronounced dc mobility was observed at o.1 mol % doping (Fig. 6). Our

Ea values calculated reclined the range of doubly ionized oxygen vacanciesV o⦁⦁. Since

doubly ionized oxygen vacancies can move on thermal activation (Ang and Yu, 2000);

100

mobility of electrons amplified with rising temperature due to long range hopping of

charge carriers from one site to another with decreasing resistivity.

4.6. Electrical dc resistivity and dc conductivity studies of PT ceramics Electrical dc resistivity (ρ) of PT ceramics was measured by two probe method at room

temperature and from 40 -700°C by employing circular pallets of 2mm thickness and

12mm diameters sintered at 1190°C± 5° C /1h . Pressure contacts equal to the pallet size

were applied after polishing the surfaces

Fig. 4.6.1-.6.4 displays the resistivity of PT ceramics for all molar compositions (1.00,

0.98 and 0.94). Stoichiometry, sintering temperature and sintering duration essentially

influenced the resistivity of PT ceramics. The measured electrical resistivity was found to

vary with increasing titanium contents. Room temperature resistivity’s (ρ25) were

2.33 ×108 Ωcm, 7.11×108Ωcmand 5.86 ×109Ωcm for 1.00, 0.98 and 0.94 compositions;

resistivities are in pact with the described values for PT (Chen et al., 2007). However,

increasing pre-sintering temperature and prolongation of sintering duration resulted in the

decrease of resistivity; which might be attributed to the lead loss.

Large differences in the dc resistivity in the ferroelectric and paraelectric regions can be

seen in the Fig. 4.6.1-.6.4. In the ordered state ferroelectric phase (40-490°C), there was

gradual decrease in resistivity for all specimens. While in the paraelectric regions,

semiconductor like behavior was observed with increasing temperature.

101

Fig. 4.6.1 Resistivity of the pure and PT (1.00) ceramics Vs temperature in the

ferroelectric and paraelectric regions at 1k Hz frequency, inset A shows variation in

conductivity (σ dc) with inverse of absolute temperature.

102

Fig.4.6.2. Resistivity of the pure and PT (0.98) ceramics Vs temperature in the

ferroelectric and paraelectric regions at 1k Hz frequency, inset A describes variation in

conductivity (σ dc) with inverse of absolute temperature.

Fig. 4.6.3 Resistivity of PT (0.98) ceramics Vs sintering duration in the ferroelectric and

paraelectric regions at 1k Hz frequency, inset A shows variance in conductivity (σ dc) with

inverse of absolute temperature.

103

Fig. 4.6.4 Resistivity of PT (0.94) ceramics Vs temperature in the ferroelectric and

paraelectric regions at 1k Hz frequency, inset A displays variation in conductivity (σ dc)

with inverse of absolute temperature.

Activation energies, Ea of 0.3-0.5 and 0.6-1.2 eV are typically assigned to single ionized

and doubly ionized oxygen vacancies (Ang and Yu, 2000; Ciomaga et al., 2011).

Conductivity of the specimens was evaluated from the impedance spectrum.

Insets A’s at Fig. 4.6.4.1- Fig. 4.6.4.1 show the variation in dc conductivity of PT

ceramics (1.00, 0.98 and 0.94) with inverse of absolute temperature, which followed the

Arrhenius Law. Calculated Ea values were in close agreement with the experimental

conductivity.

Arrhenius curves were obtained in the range of 480- 700°C; respective estimated Ea

values were 2.3265- 2.6269, 0.8302- 0.7246 and 1.7665-0.3889 eV for compositions

1.00, 0.98 and 0.94 respectively. Estimated Ea values were in association of cationic lead

vacancies (V''Pb), and doubly ionized oxygen vacanciesV o⦁⦁. For 1.00 composition,

calculated Ea values followed the range of 2.3265- 2.6269; At low temperatures; V''Pb

vacancies are quenched defects which are difficult to be activated. They could become

104

mobile with Ea values of around and above 2 eV(Guiffard et al., 2005); composition 1.00

being rich in Pb contents followed the said range (2.3265- 2.6269).

Hence, with increasing temperature, conduction seemed to be governed by doubly

ionized oxygen vacancies for 0.98 and 0.94 PT ceramics. We propose that ionic

conduction may be responsible for the conduction process for the studied ceramics.

Although, all specimens depicted semiconductor behavior, yet specimens with 0.94

composition were conspicuously more resistive in comparison to 0.98 and 1.00

compositions.

4.7 Electric polarization studies

4.7.1 Electric polarization studies for BT ceramics

P-E loops were measured at room temperature under different applied voltage on disc

shaped sintered pallets of the specimens. Composition f (0.94) and d (0.98) sintered at

1300°C/2h of BT were employed for measurements.

Fig. 4.7.1 P-E loops of BT ceramics sintered at 1300°C/2h, pre fired at 1300°C/4h.

105

Well defined (P-E) hysteresis loops under electric field were obtained at room

temperature (Fig.11) that indicate spontaneous polarization of BT samples, a

characteristic of typical ferroelectric materials (Xu, 1991). Slight gape existed in the P-E

loop that may be taken as an incompletion of the electrical cycle. Maximum polarization,

remnant polarization and coercive field were estimated to be Pm = 1.615 and 2.872 C/cm2,

2Pr = 0.794 and 1.408, EC = 8.394 and 8.449 respectively for f and d specimens. Low

values of remnant polarization can be related to the low internal polarizability and lower

EC values may account to the large grain structures of the specimens

4.7.2 Electric polarization studies for PBT ceramics

Fig. 4.7.2 displays polarization vs. electric field (P-E) hysteresis loops of the PBT

ceramics at room temperature. Well defined characteristic hysteresis P-E loops indicates

spontaneous polarization of PBT specimens, a characteristic of typical ferroelectric

materials (Xu, 1991).

Fig. 4.7.2 P-E loops of PBT ceramics sintered at 1200°C/2h, inset (B) shows P-E loops

for undoped BaTiO3

106

The values of polarization (Pm) increased from 2.872 to 3.3715-3.891 C/cm2, remnant

polarization (2Pr) increased from 1.408 to 1.844-2.077 at 0.00, 0.075 and 0.1 mole % Pb

doping. Increase in polarization values account for ferroelectricity while increase in the

remnant polarization values indicates the internal polarizability of the materials

(Haertling, 1999). Lower values of coercive field (EC), 4.110-4.283 and 4.77 indicate

still large grain structures (Fig. 4.7.2) that may be due to lack of homogeneity and

uniformity among the grains of the studied materials.

4.7.3 Electric polarization studies for BBT ceramics

In our studies, Bi2O3 as a low melting additive promoted the densification process up to

0.075 mole% doping. However, at 0.100 mole% doping, the structures seemed to be

porous and lowering of density was observed (Fig. 4.3.3 c and d). XRD revealed the

Aurivillius BaBi4Ti4O15 ceramics at 0.1mole% doping (Fig.4.2.3.1, e). Hence, the

ferroelectric studies were made at 0.075-0.100 mole% doping.

Fig. 4.7.3 P-E loops of BBT ceramics sintered at 1150°C/2h, inset B shows P-E loop for

undoped BaTiO3.

107

Fig. 4.7.3 shows polarization Vs electric field (P-E) hysteresis loops of the BBT ceramics

at room temperature. Inset B shows the P-E loop for the undoped BaTiO3. The values of

polarization (Pm), remnant polarization (2Pr) and coercive field were 2.872 C/cm2, 1.5477

C/cm2 and 8.449 kV/cm respectively. Pm and 2Pr increased to 4.416 C/cm2, 2.2029 C/cm2

at 0.075mole % Bi doping. Increase in the Pm and 2Pr values accounts for increasing

ferroelectricity and internal polarizability of the materials. However, with further increase

of Bi contents (0.100 mole %) P-E loop were slightly shifted to x-axis that may be taken

due to the presence of internal bias originating from the polar defects. Pm, 2Pr values were

reduced to 3.338 µC/cm2 1.546 µC/cm2. Lower values of coercive field (EC) 7.918, 6.165

kV/cm indicate large grain structures due to lack of homogeneity and uniformity among

the grains.

4.8 Electric polarization studies for PT ceramics

All PT ceramics the precursor composition’s (1.00, 0.98 and 0.94) showed to be the

ferroelectric characteristics (Xu, 1991); Well saturated hysteresis loops were obtained

that illustrate the typical ferroelectric characteristic for all fabricated PT materials

The values of polarization (Pm) were 5.4076 C/cm2, 8.3049 C/cm2 and 3.542 kV/cm

for1.00, 0.98 and 0.94 composition’s respectively; remnant polarization (2Pr) remained

5.0791C/cm2 3.325 C/cm2, 2.123 C/cm2. The values of remnant polarization (2Pr) are

generally lower than spontaneous polarization (Ps) values. Since tetragonal distortion

(c/a) corresponds to the spontaneous polarization (Ps); PbTiO3 has highest value among

all the ferroelectric perovskite. In our studies, spontaneous polarization (Ps) and remnant

polarization (2Pr) almost increased with increasing Pb contents and tetragonal distortion

(Fig. 4.2.4.3); however for the composition 1.00, the values were comparatively lower

that might be attributed to the abnormal grain growth and melting regions revealed in the

morphological studies (Fig. 5a and 5b).

108

Fig. 4.7.4 P-E loops of PT ceramics (1.00, 0.98, 0.94) sintered at 1190°C± 5° C /1h.

PT ceramics with optimal precursor composition 0.98 showed greater values of

spontaneous polarization. In fact, Ps is often higher in polycrystalline materials due to the

formation of opposite domains. Fig. 4.3.4.1 clearly indicates the formation of

ferroelectric domains.

Subsequently, the electric polarization values for the studied PT specimens are lower than

the reported results (Venevtsev et al., 1959) that may be taken into account for

heterogeneity among the specimens.

109

Conclusion

BaTiO3 ceramics at 0.98 and 0.94 Ba/Ti molar ratio, crack free PbTiO3 ceramics at 1.00,

0.98 and 0.94 Pb/Ti molar ratio ceramics were successfully were prepared through solid

state sintering reaction method. Pb-doped BaTiO3 and Bi-doped BaTiO3 ceramics

prepared with pre-sintered BT ceramics at 1300°C/4h were solid state sintered as well.

Pre-sintering technique was adopted to avoid the development of the undesirable phase

impurities. BT ceramics prepared at 0.98 molar ratio almost remained X-ray phase pure

at the pre-sintering regime of 1300°C/3-4h while PT ceramics prepared at 0.98 molar

ratio were also nearly X-ray phase pure at the pre-sintering regime of 1190°C± 5C /2h

and 1100°C/4h. With 0.94 precursor composition, peaks of TiO2 were detected in the

XRD diffractograms of BT ceramics and PT ceramics.

Ba/Ti and Pb/Ti ratios influenced both the crystal structure and Curie temperature (TC).

All obtained materials were perovskite ferroelectric; showed cubic and tetragonal

structures (Pm-3m, P4MM, P4mm and P4mmm). Curie temperature decreased with

increasing Ba/Ti contents (TC 130 -120°C). For PbTiO3 ceramics, Pb/Ti contents did not

shift Curie temperature. With 1.00 and 0.98 precursor compositions sharp phase

transition points were noticed; yet wit 0.94 composition containing less Pb contents no

characteristic phase transition was observed; which might be glassy phase transition.

With lead doping and its variation, Curie temperature (TC) was shifted from 120 -200°C;

Bi doping shifted it from 120 -160°C.

Enhanced electrical response was observed in the paraelectric regions; at elevated

temperatures BT ceramics depicted dielectric anomalies. For PBT and BBT ceramics;

dielectric and electrical anomalies were observed in the paraelectric regions, the effect

was pronounced at 0.075 and 0.1 mole% Pb and Bi doping. PbTiO3 ceramics also showed

dielectric and electrical anomalies were observed in the paraelectric regions with rising

temperatures. Ac conductivity increased with increasing sintering temperature and

prolongation of sintering time in accordance with dielectric studies.

At higher temperatures, electrical anomalies were observed in accordance with the

postulation of theoretical models; ac conductivity originates from migration of ions by

110

hopping between neighboring potential wells at lower temperatures which eventually

give rise to dc conductivity at high temperatures at lower frequencies.

All specimens showed increased resistivities in the ordered ferroelectric regions after

wards eventually decreased with increasing temperatures in the paraelectric regions.

Semiconductor behavior was depicted with negative temperature coefficient of resistivity

(NTCR) characteristics. Dc mobility of electrons increased with rising temperature due to

long range hopping of charge carriers from one site to another owning to decrease in the

resistivity for all specimens; with increasing Pb and Bi doping pronounced effect of drift

mobility was observed. With increasing temperature estimated values of activation

energies decreased which may be accounted for the more contents of oxygen vacancies (

V o⦁⦁ ).

Conductivities followed Arrhenius Law; with associated activation energies, Ea reclined

the range of single ionized and double ionized oxygen vacancies oxygen vacancies (V o⦁⦁)

for BaTiO3, Pb-doped BaTiO3, Bi-doped BaTiO3ceramics ceramics PbTiO3 ceramics.

However, for the PbTiO3 ceramics with precursor composition 0.98 followed the range of

lead vacancies (V''Pb). Hence ionic conduction was supposed to be responsible for the

conduction process for all obtained ceramics.

All materials showed ferroelectric characteristics with well-defined P-E loops.

Considerable values of polarization (Pm) and remnant polarization (2Pr) were obtained

that account for ferroelectricity and internal polarizability. The values were enhanced

values with Pb and Bi doping except 0.100 mole% Bi doping due to the existence of

Aurivillius BaBi4Ti4O15.

All obtained ceramics showed positive temperature coefficient of conductivity (PTC)

characteristics at high temperatures, with Pb and Bi doping pronounced PTC effects were

obtained; which might be used for BaTiO3 based energy storage systems. Studied

materials with 0.94 composition had Curie temperature 130°C and were found to be more

resistive; they might be employed for high temperature positive temperature coefficient

of resistance (PTCR) application after manipulation with suitable dopants. High

performance materials can be obtained by extending milling time.

111

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