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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
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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).
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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
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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.
64
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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