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Synthesis and Characterization of Nano Sized
Pure and Doped Barium Titanate Powders
Prepared by Sol-Gel Emulsion Technique
THESIS
Submitted in partial fulfilment
of the requirements for the degree of
DOCTOR OF PHILOSOPHY
by
AGA ZUBEDA BI HAIDER
Under the Supervision of
Prof. Sutapa Roy Ramanan
BIRLA INSTITUTE OF TECHNOLOGY AND SCIENCE, PILANI
2013
ii
BIRLA INSTITUTE OF TECHNOLOGY AND SCIENCE, PILANI
CERTIFICATE
This is to certify that the thesis entitled Synthesis and characterization of nano sized pure and
doped barium titanate powders prepared by sol-gel emulsion technique and submitted by Ms.
Aga Zubeda Bi Haider ID No P2007PHXF445G for award of Ph.D. of the Institute embodies
original work done by her under my supervision.
Signature of the Supervisor
Name in capital letters Prof. SUTAPA ROY RAMANAN
Designation Professor
Date:
iii
ACKNOWLEDGEMENT
I begin by thanking the Almighty who has always been with me guiding my path towards
new learning, growth and prosperity.
I thank BITS, Pilani – K. K. Birla Goa Campus for allowing me to carry out my work and
for the providing financial assistance.
Among the many who I wish to thank for helping me complete this work, the first in the
list is my supervisor Prof. Sutapa Roy Ramanan for her supervision and guidance through
the course of my Ph.D.
I wish to thank Dr. Srinivas Krishnaswamy, H.O.D., Dept. of Chemical Engineering and
Dr. K. N. Ponnani, Dept. of Chemical Engineering for their valuable suggestions during
the tough times that gave me the required push to move ahead with my work. I am
grateful to the Doctoral Advisory Committee (DAC) members Dr. P. Nandkumar and Dr.
Toby Joseph, Dept. of Physics for reviewing my research work. Departmental Research
Committee (DRC) of the Dept. of Chemical Engineering is also appreciated for keeping a
timely check of the progress of my research work. I am also thankful to Dr. N. N. Ghosh,
Dept. of Chemistry, Dr. Prita Pant and Dr. Bhanudas Naik, his Ph.D. scholars for helping
me to use the characterization facilities available at their end. The faculty and staff of
Dept. of Chemical Engineering are acknowledged for their support and help provided in
the time of need. Faculty members involved in conducting the courses for Ph.D.
qualifying exam are also valued for the efforts taken. Institute staff concerning other
Departments is also recognized for their help in completing the official formalities.
I would like to thank my lecturers Mr. Rajendra and Dr. Efrem D’Sa, Dept. of Physics,
Carmel College of Arts, Science and Commerce for Women for motivating me to pursue
Ph.D. Degree. Always remembered is Dr. K. R. Priolkar, Dept. of Physics, Goa
iv
University who guided me for my M.Sc. Physics thesis. Without his training I would not
have been able to carry out this work. I wish to express my gratitude to Prof. E. Desa and
Dr. R. B. Tangsali, Dept. of Physics, Goa University and Dr. Rahul Mohan, N.C.A.O.R.
Goa for permitting me to avail their characterization facilities.
Lastly, I would like to thank my family, friends and the research fellows from the bottom
of my heart for their support, courage, kindness, love and above all respect that they have
bestowed upon me. Especially to my parents without whose support I would not even
have had applied for this Degree. I will always be indebted to them for giving me the best
upbringing I could have had ever received. I would like to acknowledge my siblings
Ms. Afreen Bi and Ms. Amreen Bi who have stood by my side in every aspect of life. I
also wish to mention my nephew Master Abdullah Sheikh who brought loads of smiles to
me. Finally, I wish to thank my husband Mr. Zikriya Shaikh who supported me during
the last few days of my thesis writing to make sure I complete it with ease.
v
This Thesis is dedicated to my parents
Mrs. Massura Bi Aga and Mr. Haider H. Aga
vi
ABSTRACT
Barium titanate (BaTiO3) due to its ferroelectric and electrical properties is used in a
variety of applications such as multi layer ceramic capacitor (MLCC), sensors, actuators,
current limiters, PTCR thermistors, constant temperature heaters, dynamic random access
memories (DRAMs), etc. Search for lead (Pb) free materials due to environmental
concern, led to further interest in the study of BaTiO3 owing to its simple crystal
structure, ability to accommodate various dopants and ease of being prepared in
polycrystalline ceramic form. Besides these, the other advantages of BaTiO3 are higher
dielectric constant, low dielectric loss and piezoelectric properties in the tetragonal phase.
Current trend towards smaller size electronic components have gained momentum in
developing BaTiO3 ceramics with an average grain size < 100nm. With reduction in size,
the structure and properties are known to deviate from those reported for the bulk. This
phenomenon known as the size effect affects the dielectric properties and the crystal
structure which is also sensitive to factors such as stoichiometry, defects, impurities,
electrical boundary conditions, strain and stress. Miniaturization of components to
achieve higher capacitance and reliability led to the demand for fabricating dense BaTiO3
dielectric layers with uniform size nano grains. Due to this, the traditional sold-state
reaction synthesis route needed to be replaced with wet chemical routes such as sol-gel,
hydrothermal, co-precipitation, etc. since they offer relatively better compositional
control and homogeneity. Various dopants like Sr, La, Mg, Ce, Nb, Pb, Zr, etc. are added
in BaTiO3 to modify its properties by shifting the TC from its reported value, increasing
dielectric constant, broadening the temperature range of maximum dielectric constant,
inducing PTCR effect, etc.
In the present work, nano-size BaTiO3 powders were synthesized via. sol-gel emulsion
technique. Water-in-oil (w/o) type emulsion used comprised of BaTiO3 sol as the water
phase and cyclohexane as the oil phase, which was stabilized using span 80 and span 20
surfactants. Sr, La, Ce, Mg, Li and K were used to dope BaTiO3. Effect of particle size
vii
and the added dopants on the structural and electrical properties of the powders were
studied. Emulsion employed in the synthesis of BaTiO3 produced spherical particles for
span 80 and rod-like particles for span 20. An average size of 57nm was achieved for
pure BaTiO3 powders synthesized using 5% span 80, 69.99nm using 20% span 80 and
66nm using 20% span 20. Increase in the concentration of surfactant from 5% to 20%
distorted the particle shape. XRD patterns confirmed dominance of cubic phase with the
presence of tetragonal phase in small amount that resulted in the pseudocubic symmetry
of the synthesized powders. Lattice parameters calculated for the synthesized pure and
doped BaTiO3 powders matched with those calculated using the empirical formula
proposed by Jiang et. al. for the six-fold co-ordination assumption with an error of <1%.
PTCR effect with the room temperature resistivity of ~ 107 !" - 10
8 !" was obtained
in the synthesized undoped BaTiO3 powders due to adsorbed oxygen at the grain
boundaries with the TC noted at 75°C. With the different dopants used, a shift in TC was
noted from 55°C - 95°C. However, the range of resistivity obtained was similar to that of
undoped BaTiO3 (~ 107 !" – 10
11 !"). Relaxor-type dielectric behavior was achieved
for these small sized particles with disordered surface. Dielectric measurements
confirmed dopants such as as Sr, Mg, Li and K synthesized using span 80 and Sr, La and
Ce synthesized using span 20 to show the presence of core-shell structure for the powders
synthesized using both span 80 and span 20. The two dielectric peaks corresponds to the
two transition temperatures associated with the core and the shell phase separately.
Dopants used in the synthesized powders replaced Ba in the crystal structure. The PTCR
effect with the resistivity values noted and the dielectric behavior was similar to that of
pure BaTiO3. Smaller particle size achieved in the present work plays a dominant role in
controlling the behavior of the material prepared.
viii
TABLE OF CONTENTS
Page
CERTIFICATE ii
ACKNOWLEDGEMENT iii
ABSTRACT vi
TABLE OF CONTENTS viii
LIST OF TABLES xi
LIST OF FIGURES xiii
LIST OF ABBREVIATION xvii
LIST OF SYMBOLS xx
CHAPTER ONE: INTRODUCTION
1.1 Barium Titanate (BaTiO3) 1
1.2 Objective 5
1.3 Scope of research work 6
1.4 Thesis breakup 6
CHAPTER TWO: LITERATURE REVIEW
2.1 History of ferroelectrics 8
2.2.1 Perovskite Oxides (ABO3)
2.2.2 Some Other Ferroelectrics
12
17
2.3 Domain and Polarization 21
2.4 Dielectrics
2.4.1 Classes of Dielectrics
25
29
2.5 BaTiO3: Its Importance and Modification 32
2.6 Techniques Employed in the Synthesis of BaTiO3
2.6.1 Solid-State Reaction Method
37
38
ix
2.6.2 Hydrothermal Method
2.6.3 Coprecipitation Method
2.6.4 Polymeric Precursor Method
2.6.5 Thermal Decomposition Method
2.6.6 Mechanochemical Method
2.6.7 Sol-Gel Method
2.6.8 Refined Sol-Gel Method
39
40
41
41
42
43
43
2.7 Role of Surfactant in the Emulsion 45
2.8 Current Research on BaTiO3 in Last Five Years 50
CHAPTER THREE: MATERIALS AND METHOD
3.1 Precursors 54
3.2 Experimental Procedure 55
3.3 Thermal Analysis (TA)
3.3.1 Thermogravimetry Analysis (TGA) and Differential Thermal
Analysis (DTA)
3.3.2 Differential Scanning Calorimetry (DSC)
58
58
59
3.4 X-Ray Diffraction (XRD) 60
3.5 Fourier Transform Infrared Spectroscopy (FTIR) 62
3.6 Electron Microscopy
3.6.1 Transmission Electron Microscopy (TEM)
3.6.2 Scanning Electron Microscopy (SEM) and Energy Dispersive
Spectroscopy (EDS)
63
63
63
3.7 Two-Probe Resistivity Analyzer 64
3.8 Impedance Spectroscopy (IS)/ Dielectric Analyzer (DEA) 65
CHAPTER FOUR: RESULTS AND DISCUSSIONS
Outline 67
4.1 TGA, DTA and DSC Analysis 67
4.2 X-Ray Analysis 73
x
4.3 FTIR Analysis 95
4.4 TEM, SEM and EDS Analysis 101
4.5 Resistivity Analysis 105
4.6 Impedance Spectroscope Analysis (IS) 130
CHAPTER FIVE: CONCLUSION AND FUTURE SCOPE OF WORK
5.1 Conclusion 141
5.2 Future Scope of Work 150
REFERENCES 151
LIST OF PUBLICATIONS 182
BIOGRAPHY 184
xi
LIST OF TABLES
Page
2.1 Criterion for perovskite structure formation 16
2.2 Class II/ III dielectric capacitor coding for temperature and
capacitance range. EIA codes D – R belong to class II and S – V
belong to class III dielectrics
31
2.3 Highly used synthesis methods of preparing BaTiO3 and the
properties achieved.
46
2.4 Surfactant packing parameters and geometry of self-assemblies in
water
49
4.1 Average crystallite size ±0.001 (nm) calculated for 1:3 sol:support
solvent ratio and various surfactant concentrations
76
4.2 Ionic Radiu#$ %&'($ )#*+$ ,-$ ./*01*.,!23$ !23!)32.,0-$ 04$ 32..,!*$
parameters.
82
4.3 Ba1-xDxTiO3 crystal parameters deduced from XRD patterns
(Figure 4.4) for powders prepared using 1:3 sol: support solvent
ratio with 5% span80 and calcined at 750°C.
89
4.4 Ba1-xDxTiO3 crystal parameters deduced from XRD patterns
(Figure 4.5) for powders prepared using 1:3 sol: support solvent
ratio with 5% span20 and calcined at 750°C.
90
4.5 Experimentally deduced lattice parameter of the pc Ba1-xDxTiO3
powders synthesized using 1:3 sol: support solvent ratio with 5%
span80, calcined at 750°C compared with that derived
theoretically using the empirical formulae with (a) and without
(a5) assuming the six-fold coordination for the ions.
92
4.6 Experimentally deduced lattice parameter of the pc Ba1-xDxTiO3
powders synthesized using 1:3 sol: support solvent ratio with 5%
span20, calcined at 750°C compared with that derived
93
xii
theoretically using the empirical formulae with (a) and without
(a5) assuming the six-fold coordination for the ions.
4.7 Average particle size via SEM and crystallite size via XRD of
750°C calcined pure BaTiO3 powders synthesized as a function of
span 80 concentration.
106
4.8 Resistivity values of Ba1-xDxTiO3 pellets during the heating cycle
with varying D and x in powders synthesized using span 80.
118
4.9 Resistivity values of Ba1-xDxTiO3 pellets during the cooling cycle
with varying D and x in powders synthesized using span 80.
119
4.10 Resistivity values of Ba1-xDxTiO3 pellets during the heating cycle
with varying D and x in powders synthesized using span 20.
126
4.11 Resistivity values of Ba1-xDxTiO3 pellets during the cooling cycle
with varying D and x in powders synthesized using span 20.
127
xiii
LIST OF FIGURES
Page
2.1 Division of point groups on the basis of Symmetry and Properties. 11
2.2 Two different views of the unit cell structure of the ideal cubic perovskite
(ABO3).
14
2.3 (a) Displacement of B-site ion in the perovskite (ABO3) structure resulting
in the formation of cubic, tetragonal and the rhombohedral structure.
(b) Phase transformation in BaTiO3.
16
2.4 Phase diagram of PZT ceramic. 18
2.5 (a) Schematic representation of the Aurivillius layer structure.
(b) SrBi2Ta2O9 structure projected along (110) axis.
22
2.6 Ideal domain configuration in a single crystal of cubic ferroelectric material,
where the
(a) coupling of strain is negligible.
(b) strain effects are important.
(c) schematic diagram of 180° and 90° domains in BaTiO3.
22
2.7 (a) Surface charge associated with spontaneous polarization.
(b) formation of domains to minimize electrostatic energy.
24
2.8 Polarization vs. electric field curves for a single crystal ferroelectric,
polycrystalline ferroelectric and a ferroelectric in the paraelectric state.
24
2.9 (a) Capacitance of the parallel plate capacitor with free space between the
plates.
(b) charge build up when dielectric is being placed.
(c) capacitance when dielectric is placed between the plates.
27
2.10 (a) Dielectric placed in an electric field.
(b) electric field gives rise to bound polarization charges.
(c) representing the whole dielectric in terms of its surface polarization
charges +QP and –QP.
27
2.11 (a) BaTiO3 unit cell structure. 33
xiv
(b) BaTiO3 polymorph distortion of the perovskite structure.
(c) pseudo-binary phase diagram of BaO-TiO2 system.
3.1 Flow diagram for synthesis of nano size BaTiO3 powders via sol-gel
emulsion technique.
57
4.1 BaTiO3 powder synthesized for 5% span 80 (a) TGA (weight %) and DTA
(heat flow) curves, (b) DSC curve for the powder calcined at 750oC.
70
4.2 TGA-DTA curves of Ba0.9D0.1TiO3 powder synthesized using 5% surfactant
and dried at 100°C with D = Sr (a) and La (b), Ce (c), Mg (d), Li (e) & K
(f).
70
4.3 X-ray diffraction pattern of BaTiO3 synthesized (a) using 5% span 80
calcined at (i) 400°C, (ii) 500°C, (iii) 600°C, (iv) 700°C, (v) 750°C, (vi)
800°C, (vii) 900°C, and (viii) 1000°C; (b) using 5% span 80 calcined and
sintered at 750°C; and (c) using 5% span 80 and 20 as the surfactant
calcined at 1000°C (* - BaCO3)
74
4.4 X-ray pattern of Ba1-xDxTiO3 powder calcined at 750oC, for D = Sr, La, Ce,
Mg, Li & K with varying x synthesized using 5% span 80 (* - BaCO3).
77
4.5 X-ray pattern of Ba1-xDxTiO3 powder calcined at 750oC, for D = Sr, La, Ce,
Mg, Li & K with varying x synthesized using 5% span 20 (*-BaCO3)
78
4.6 Maximum intensity (110) diffraction peak of Ba1-xDxTiO3 powders calcined
at 750oC for D = Sr, La, Ce, Mg, Li & K with varying x as 0( ),
0.001 ( ), 0.01 ( ), 0.05 ( ) and 0.1 ( ) synthesized using 5% span 80.
81
4.7 Maximum intensity (110) diffraction peak of Ba1-xDxTiO3 powders calcined
at 750oC for D = Sr, La, Ce, Mg, Li & K with varying x as 0 ( ),
0.001 ( ), 0.01 ( ), 0.05 ( ) and 0.1 ( ) synthesized using 5% span 20.
84
4.8 (002)/(200) diffraction peak splitting of Ba1-xDxTiO3 powders calcined at
750oC for D = Sr, La, Ce, Mg, Li & K with varying x as 0 ( ),
0.001( ), 0.01 ( ), 0.05 ( ) and 0.1 ( ) synthesized using 5% span 80.
85
4.9 (002)/(200) diffraction peak splitting of Ba1-xDxTiO3 powders calcined at
750oC for D = Sr, La, Ce, Mg, Li & K with varying x as 0 ( ),
0.001( ), 0.01 ( ), 0.05 ( ) and 0.1 ( ) synthesized using 5% span 20.
86
xv
4.10 BaTiO3 powder synthesized using 5% span 80 with varying calcinations
temperature at 500oC, 750
oC and 1000
oC
96
4.11 FTIR spectrum of Ba1-xDxTiO3 synthesized using 5% span 80 for the
powder calcined at 750oC with varying x as 0 ( ), 0.001 ( ), 0.01( ) and
0.1 ( ).
98
4.12 FTIR spectrum of Ba1-xDxTiO3 synthesized using 5% span 20 for the
powder calcined at 750oC with varying x as 0 ( ), 0.001 ( ), 0.01( ) and
0.1 ( ).
100
4.13 TEM-BaTiO3 morphology for powder calcined at 750°C using, a) 5% span
80, b) 10% span 80, c) 15% span 80, d) 20% span 80, and e) 20% span 20.
102
4.14 SEM-BaTiO3 morphology for powder calcined at 750°C using, a) 5% span
80, b) 10% span 80, c) 15% span 80, d) 20% span 80, and e) 20% span 20.
103
4.15 Particle morphology of the pellets (5% span 80 powder) sintered at (a)
750°C for 1h soaking, (b) 750°C for 12h soaking and (c) 1200°C for 2h
soaking (d) EDS spectrum (750°C for 1h soaking).
106
4.16 PTCR effect in BaTiO3 heat treated at 750°C synthesized using 5% (a) span
80 and (b) span 20 during the heating and cooling cycle.
107
4.17 Resistivity profile of Ba1-xDxTiO3 synthesized using 5% span80 for the
powder calcined at 750oC with varying x measured during heating (H) and
cooling (C) cycle.
(a) D = Sr
(b) D = La
(c) D = Ce
(d) D = Mg
(e) D = Li
(f) D = K
112
113
114
115
116
117
4.18 Resistivity profile of Ba1-xDxTiO3 synthesized using 5% span20 for the
powder calcined at 750oC with varying x measured during heating (H) and
cooling (C) cycle.
(a) D = Sr 120
xvi
(b) D = La
(c) D = Ce
(d) D = Mg
(e) D = Li
(f) D = K
121
122
123
124
125
4.19 BaTiO3 dielectric response for 5% span80 sample sintered at 750°C (12
hours soaking), measured at (a) fixed frequencies, (b) 1 kHz frequency and
(c) 3 MHz frequency.
131
4.20 BaTiO3 dielectric response for 5% span20 sample sintered at 750°C (1 hour
soaking), measured at (a) fixed frequencies, (b) 1 kHz frequency and (c) 3
MHz frequency.
133
4.21 Dielectric response of Ba1-xDxTiO3 pellets synthesized using 5% span 80 at
3 MHz frequency with v216,-7$8$2-+$9$2#$:;::<$%=(>$:;:<$%?), 0.05 (@) and
0.1( ).
136
4.22 Dielectric response of Ba1-xDxTiO3 pellets synthesized using 5% span 20 at
3 MHz frequency with v216,-7$8$2-+$9$2#$:;::<$%=(>$:;:<$%?), 0.05 (@) and
0.1( ).
139
xvii
LIST OF ABBREVIATIONS
Ag Silver
Ba Barium
BaCO3 Barium Carbonate
BaO Barium Oxide
BaTiO3 Barium Titanate
BTNP BaTiO3 Nanoparticles
BSA Bovine Serum Albumin
BST Barium Strontium Titanate
C Curie-Weiss constant
C.S. Crystallite Size
CaTiO3 Calcium Titanate
CCD Charge-Coupled Device
Ce Cerium
CMC Critical Micelle Concentration
CPP Critical Packing Parameter
DEA Dielectric Analyzer
DRAMs Dynamic Random Access Memories
DSC Differential Scanning Calorimetry
DTA Differential Thermal Analysis
E Electric Field
EC Coercive Field
EDS Energy Dispersive Spectroscopy
EIA Electronic Industrial Alliance
EPD Electrophoretic Deposition
FTIR Fourier Transform Infrared Spectroscopy
FWHM Full Width at Half Maxima
xviii
GB Grain Boundary
HF Heat-Flux
HLB Hydrophile-Lipophile Balance
IBLC Internal Boundary Layer Capacitors
K Potassium
KDP Potassium Dihydrogen Phosphate
La Lanthanum
Li Lithium
LiNbO3 Lithium niobate
LiTaO3 Lithium tantalate
MEMS Microelectromechanical Systems
Mg Magnesium
MgTiO3 Magnesium titanate
MPB Morphotropic Phase Boundary
MLCC Multilayer Ceramic Capacitors
NFeRAM, FRAM Nonvolatile Ferroelectric Random Access Memories
Ni Nickel
NTC Negative Temperature Coefficient
O Oxygen
p Dipole Moment
Pb Lead
PbNb2O6 Lead niobate
PbTiO3, PT Lead Titanate
PbZrO3 Lead Zirconate
PC Power Compensation
Pd Palladium
PNR polar nano-regions
PLZT Lead Lanthanum Zirconate Titanate
PMN Lead Magnesium Niobate
xix
Pt Platinum
PTCR Positive Temperature Coefficient of Resistivity
PPT Polymorphic Phase Transitions
PZT Lead Zirconate Titanate
R Relative Displacement
SEM Scanning electron microscopy
Sr Strontium
SBT Strontium bismuth tantalate
TEM Transmission electron microscopy
TGA Thermogravimetry Analysis
Ti Titanium
TiO2 Titanium dioxide
W Tungsten
w/o water-in-oil
WEEE Waste Electrical and Electronic Equipment
XRD X-ray Diffractometer
xx
LIST OF SYMBOLS
A Area
a Lattice Constant
Optimal Head Group Area
Diffraction Plane Spacing
D Electric Displacement
AB Enthalpy of Transition
Permittivity of Free Space
C5, K Dielectric Constant
C5max Permittivity Maximum
Cr Relative Permittivity
Critical Chain Length
n aggregation number
n Refractive Index
NNumber of Dipoles Per Unit
Volume
o/w oil-in-water
P Polarization
pc Pseudocubic
pav
Average Dipole Moment Per
Unit Molecule
Pr remanant polarization
Ps Spontaneous Polarization
R Resistance
ptotal Net Dipole Moment
rA Ionic Radius of A ion
xxi
rB Ionic Radius of B ion
rO Ionic Radius of O ion
Dp Surface Charge Density
t Tolerance Factor
Thickness
TC Curie Temperature
Td Burns temperature
Tm Maximum Temperature
Q Charge
Qo Charge in Vacuum
Qp Surface Polarization Charges
E>$.2-E Loss Angle
AB
Enthalpy of Transition
Diffraction Angle
F Resistivity
Wavelength
G Volume
H Electric Susceptibility
Grain Boundary Barrier
e Electron Charge
NsDensity of Trapped Electrons at
the Gain Boundaries
Nd Charge Carrier Concentration
Cgb
Relative Permittivity of the
Grain Boundary Region
IJJO Oxygen Vacancies
IKBa Barium Vacancies
1
Chapter 1 Introduction
1.1 Barium Titanate (BaTiO3)
Barium Titanate, was the first discovered piezoelectric transducer ceramic that sufficed
the need of a high dielectric constant material for capacitor application during World War
II in mid-1940s [Gene, 1999]. From the time of its discovery, it has been widely used in
the electronic industry owing to its high dielectric constant and ferroelectric behavior
[Takeuchi et al., 1997]. The work carried out by Wul and Goldman [1945] in USSR and
von Hippel’s group at Massachusetts [Hippel et al., 1946] in 1945-1946 ascertained the
ferroelectric nature of polycrystalline BaTiO3 to be responsible for its dielectric
properties. This was affirmed when the dielectric properties of single crystal BaTiO3
were found to be originating from its ferroelectric nature [Gene, 1999]. BaTiO3
undergoes five different phase transitions known as the polymorphic phase transitions
(PPT), namely; hexagonal (>1460°C), paraelectirc (>120°C up to 1460°C), tetragonal
(>5°C up to 120°C), orthorhombic (>-90°C up to 5°C) and rhombohedral (below -90°C)
[Kasap, 2007; Moulson and Herbert, 2003; Vijatovi et al., 2008; Chandler et al., 1993;
Joshi et al., 2006; Jamal et al., 2008]. Melting point of BaTiO3 is around 1623°C and it
has a density of 6.02g/cc [Aleksander, 2001; Jeon et al., 2005]. BaTiO3 is an insulator at
room temperature having resistivity above 1010!cm and a band gap of ~3eV [Heywang,
1971; Panwar and Semwal, 1991]. Majority of the study conducted on it is based on its
tetragonal phase specifically due to its high permittivity which results in a higher
dielectric constant value and piezoelectric properties. Its tetragonality results from the
outward displacement of Ti4+ ions from the centrosymmetric position in the TiO6
octahedra, categorizing it as a displacive type ferroelectric [Antonio et al., 2006]. BaTiO3
is highly stable chemically and mechanically. Its ferroelectric nature at and above room
temperature has attained great importance as it can be easily prepared and used in the
form of ceramic polycrystalline sample. It belongs to the perovskite (ABO3) family of
compounds, which are significant electronic materials. Owing to its high dielectric
2
constant and low loss characteristics, barium titanate has been used in applications such
as capacitors and multilayer ceramic capacitors (MLCC) [Vijatovi et al., 2008; Vijatovi
M. et al., 2008; Min et al., 2007]. Besides these, the wide variety of electrical phenomena
exhibited by BaTiO3 makes it applicable as sensors, actuators, resonators, filter-
duplexers, voltage-controlled oscillators, antennas, current limiters, constant temperature
heaters, thermistors, dynamic random access memories (DRAMs), etc. [Gene, 1999; Lee
and Su, 2007; Tohma et al., 2002; Frey and Payne, 1996; Vijatovi et al., 2008; Panwar
and Semwal, 1991; Fiorenza et al., 2009].
Particle size is known to strongly influence the final microstructure of ceramics [Yanan et
al., 2011]. With a decrease in the grain diameter to 1"m, an increase in room-
temperature dielectric constant has been reported [Ihlefeld et al., 2007]. Materials having
either one dimension (quantum well), two dimensions (quantum wire) or all three
dimensions (quantum dot) in nanometer range (between 1 to 100nm) are classified as
nanomaterials [Poole, 2003; Ratner and Ratner, 2003]. Distinct and fascinating properties
of nanostructured materials have led to advancement towards nanoscale devices [Shubin
et al., 2008; Bernadette et al., 2002]. Increased relative surface area and quantum effect
are two principle factors that cause the properties of nanomaterials to differ from its bulk
form [Poole, 2003; Ratner and Ratner, 2003]. Among the nanostructured materials,
ferroelectric nanostructures are of particular interest due to their high sensitivities and
coupled responses to external inputs that facilitate their applications as ultrasensitive
sensors, transducers, MLCCs, ferroelectric thin-film memories, electro-optic modulators,
microwave electronics, microelectromechanical systems, etc. [Shubin et al., 2008; Mishra
and Mishra, 2012; Huang et al., 2007; Kholkin et al., 1997]. Lead zirconate titanate
(PZT) is the most widely used ferroelectric that contains ~60 – 70% lead (Pb) by weight
which is highly hazardous and can cause environmental pollution when exposed to air
[Anal et al., 2012; Hirofumi, 2012; Aksel and Jones, 2010; Kholkin et al., 1997]. Search
for environmentally friendly materials to replace Pb-based compounds lead to an increase
in interest in BaTiO3 owing to its simple crystal structure, ability to accommodate
different types of dopants ensuing easy tailoring of its properties for specific
3
technological applications and also due to the ease with which it can be prepared in the
polycrystalline ceramic form [Hirofumi, 2012; Aksel and Jones, 2010; Zanetti et al.,
2004; Wada, 2009; Mishra and Mishra, 2012]. In its ferroelectric tetragonal form, room
temperature dielectric constant value of BaTiO3 (~1000) is higher than that possessed by
other common dielectric materials [West et al., 2004]. The above mentioned phase
transitions and properties are characteristic of BaTiO3 in the bulk form [Huang et al.,
2007]. Downsizing of electronic devices resulted in bulk BaTiO3 to be replaced with
thick films which were then followed by the use of thin films [Park et al., 2010; Huang et
al., 2007; Wang et al., 2002; Stojanovi et al., 2002]. BaTiO3 films, consisting of fine
grains of sub-micron to nano meter size, are used in the advanced electronic devices such
as MLCC, DRAM, nonvolatile ferroelectric random access memories (NFeRAM),
piezoelectric microactuators, etc. [Tohma et al., 2002; Frey and Payne, 1996; Huang et
al., 2007; Park et al.,2010]. Market demand towards small sized electronic components
has resulted in an interest towards developing BaTiO3 ceramics with average grain size <
100nm [Wang et al., 2006]. Thin films are advantageous due to their smaller size, light
weight, easy integration, lower operating voltage, higher speed and unique structure
[Wang et al., 2002]. Thin film structure and properties deviate from those of bulk or
single crystal BaTiO3, a phenomenon referred to as the size effect [Tohma et al., 2002;
Frey and Payne, 1996]. The most prominent result of size effect is on the dielectric
properties and crystal structure [Guo et al., 2012; Tohma et al., 2002; Ihlefeld et al.,
2007; Xiangyun et al., 2006]. This is because the dielectric properties of BaTiO3 are
extremely sensitive to various factors such as stoichiometry, defects, impurities, electrical
boundary conditions, stress and strain [Hirokazu et al., 2010]. For BaTiO3, applications
as capacitors require it to have its Curie Temperature (TC) around room temperature
along with a broad permittivity maximum (#$max) [West, 2004; Yoon et al., 2003]. To
enable miniaturization of components, achieve higher capacitance and reliability, it is
necessary to develop techniques to fabricate dense BaTiO3 dielectric layers that consist of
uniform nano-size grains which is difficult to achieve with the traditional solid-state
reaction method without grain growth [Avinash et al., 2011; Zhenxing et al., 2010;
Hirokazu et al., 2010]. Hence, wet chemical methods such as sol-gel, hydrothermal, co-
4
precipitation, etc., offering relatively better compositional control and homogeneity, are
employed in synthesizing nanoparticles [Gust et al., 1997; Moon et al., 2012; Pithan et
al., 2006].
Various dopants such as Sr, Nb, La, Mg, etc. have been incorporated in BaTiO3 and their
effect on its properties studied [Sutham, 2008; Viviani et al., 2004; Dey and Majhi, 2005;
Da-Yong et al., 2006; Li et al., 2007; Fukuda et al., 2007; Min et al., 2007; Narang and
Kaur, 2009; Beltrán et al., 2004; Zeng et al., 2006]. On being doped with donors, e.g.
trivalent ions at Ba site and pentavalent ions at Ti site, BaTiO3 becomes semiconducting
[Panwar and Semwal, 1991; Senlin et al., 2009; Masó et al., 2008; Gheno et al., 2007;
Kim, 2002; Mancini and Filho, 2006; Darko et al., 2003; Burcu, 2012]. This has been
associated with the presence of Ba2+ vacancies or the convertion of Ti4+ to Ti3+ to create a
charge balance [Panwar and Semwal, 1991; Kim, 2002]. The compensation mechanism
also leads to increase in conductivity thereby rendering insulating pure BaTiO3
semiconducting on doping. Semiconducting BaTiO3 is also characterized with a sudden
increase in the resistivity magnitude by 3 to 6 orders at the TC known as the positive
temperature coefficient of resistivity (PTCR), which has been attributed to the presence
of surface layers of acceptor charges due to segregated acceptor ions or adsorbed oxygen
ions at grain boundaries [Panwar and Semwal, 1991; Ruitao et al., 2004; Kim, 2002;
Kareiva et al., 1999]. These materials find application as PTC thermistor and have been
produced with room temperature resistivity values of ~10 - 100!cm with a 4-6 orders
resistivity rise above TC [Vijatovi et al., 2008; Vijatovi M. et al., 2008; Panwar and
Semwal, 1991; Fiorenza et al., 2010]. PTCR effect noted in bulk and nanometer size
BaTiO3 is reversible upon appropriate heat treatment [Fiorenza et al., 2009]. Apart from
inducing PTCR effect, dopants are also incorporated in BaTiO3 to modify its properties,
such as increasing or decreasing the TC, increasing #$, broadening the temperature range
in which #$max exists, etc. The present work, hence, aims at developing nano-size BaTiO3
powders by a wet chemical method and study the effect of particle size and dopant
addition on the properties of the synthesized powders.
5
1.2 Objective
The objective of this work was
1. To synthesize BaTiO3 nano-powders by sol-gel emulsion technique and study the
size effect on its properties by:
! Employing two different surfactants, namely span 80 and span 20.
! Carrying out a thermal treatment process by which fully crystalline material
may be achieved at a lower temperature.
2. To study the effect of doping in BaTiO3 with six different dopants, namely
Strontium (Sr), Lanthanum (La), Cerium (Ce), Magnesium (Mg), Lithium (Li)
and Potassium (K) for different concentrations. Sr, La, Ce and Mg doping in
BaTiO3 has been reported to decrease the TC towards room temperature. Sr doped
BaTiO3, owing to their high dielectric constant and low losses, are used in the
microelectronic devices. La doping in BaTiO3 increase the dielectric constant and
also induces semiconductor properties giving rise to PTCR effect. Ce doped
BaTiO3 also has high dielectric constant and are used in MLCC. Like K, Li have
been co-doped in BaTiO3, to study the effect of its addition on the structure and is
observed to be similar to that observed with other dopants. However, it is less
studied for their dielectric and resistivity properties.
3. To study the structural and electrical properties of the synthesized compounds
analyzed using TGA/DTA (thermogravimetry and differential thermal analysis)
instrument, Differential Scanning Calorimetry (DSC), X-ray Diffractometer
(XRD), Scanning electron microscopy (SEM), Transmission electron microscopy
(TEM), Fourier Transform Infrared Spectroscopy (FTIR), Two probe resistivity
meter, LCRQ meter.
6
1.3 Scope of research work
Decrease in size of BaTiO3 towards nanometer range is known to produce a change in the
properties from that of bulk BaTiO3.The present work intended to prepare pure and doped
BaTiO3 nano-powders using simplified sol-emulsion-gel method for synthesis.
Previously used sol-emulsion-gel method required to peptize the precursor and needed to
heat the precursor solution to form a translucent colloidal sol [Chatterjee et al., 1999;
Despina et al., 1994; Chatterjee et al., 2003]. Parameters such as temperature at which the
precursor was heated and the amount of peptizing agent added to form a translucent sol
had not been reported with accuracy [Chatterjee et al., 1999; Despina et al., 1994]. In this
work the synthesis process was simplified by avoiding the peptization step during the sol
formation. The effect of surfactants was investigated as the micelles formed in the
emulsion acts as templates for synthesis of the nanopowders. Role of dopants in
modifying the properties of the synthesized nanoparticles were also studied.
1.4 Thesis breakup
The thesis is split into five chapters, including the present one.
Chapter 2 presents the current literature that was reviewed and compiled. It consists of
sections dedicated to the history of BaTiO3, its structure (perovskite) and its properties. It
also includes the other types of ferroelectrics known besides BaTiO3. As BaTiO3 is a
dielectric, a section is dedicated to dielectrics explaining the concept of polarization in
these materials. The various methods used for synthesizing BaTiO3 are also discussed in
this chapter.
Chapter 3 gives a detailed description of the sol-gel-emulsion technique used in the
present work for synthesizing BaTiO3 nano-powder. It gives the particulars of the
materials used along with the flow diagram of the synthesis method explaining every step
7
clearly. The heat treatment schedules involved in this study is also elaborated. The
different characterization methods used in the present work are mentioned and discussed.
Chapter 4 presents the results obtained for the synthesized BaTiO3 nano-powders. This
chapter is divided into sections according to the characterization methods used. It starts
with the discussion over surfactants, its use and its importance with respect to emulsions
formation. Each section involves a comparative study of the results obtained for pure and
doped BaTiO3. The structural parameters obtained from the XRD characterization, the
FTIR profiles detailing the various functional groups detected, the resistivity
characteristics as a function of temperature and the dielectric properties of the powders as
function of temperature and frequency are detailed in this chapter. The effects of
synthesis technique, thermal treatment and dopants on the properties of the synthesized
powders are elaborated.
Chapter 5 includes the conclusions of the present work.. The summary includes the
advantages of the synthesis method employed and the properties of the materials as a
result of it. The effect of particle size and the modifications by doping on the properties
of the synthesized BaTiO3 nanopowders are presented.
8
Chapter 2 Literature Review
2.1 History of ferroelectrics
Ferroelectrics are a class of polar dielectrics having permanent electric dipoles oriented in
a specific direction even in the absence of an external electric field [Vijaya and
Rangarajan 2004; Dragan, 2005]. Ferroelectricity was first discovered by Joseph Valasek
in the year 1921 in single crystal potassium-sodium tartarate tetrahydrate also known as
Rochelle salt, which was then known as Seignette electricity. Later it was found to exist
in polycrystalline barium titanate (BaTiO3) ceramic by the mid-1940s [Nalwa, 1999;
Valasek, 1971; Gene, 1999; Cross and Newnham, 1987]. Following the discovery of
ferroelectricity in BaTiO3, more ferroelectric materials were found, contributing to
various applications commercially [Gene, 1999].
Rochelle salt/ Seignette salt dominated the studies carried out in the first decade of
ferroelectricity research from 1920-1930. It was easily available and facilitated
fabrication of large single crystals having high optical quality [Valasek, 1971; Gene,
1999; Cross and Newnham, 1987]. However, its instability specifically due to its water
solubility led to in the discovery of ferroelectricity in Potassium Dihydrogen Phosphate
(KDP) family of compounds. KDP compounds were symmetric crystals with the
ferroelectricity below -150°C. Nevertheless, this family of materials was worked upon for
another decade from 1930-1940 before the investigation of ferroelectricity in BaTiO3
produced positive result [Nalwa, 1999; Gene, 1999; Cross and Newnham, 1987]
The decade from 1940-1950 referred to as the “early barium titanate era” by Cross and
Newnham [1987] was followed by the period of emergence from 1950-1960, which
marked the discovery of many new ferroelectric materials confirming 25-families of
ferroelectric, more than 20 perovskite compounds, and many solid solutions of the then
known ferroelectrics in the early 1960s. During World War-II in the early 1940s, the
9
demand for higher dielectric-constant capacitors compared to those available from
materials such as mica, Titanium dioxide (TiO2), Magnesium titanate (MgTiO3), etc.,
resulted in the extensive research in ferroelectricity and piezoelectricity of ceramic
materials leading to BaTiO3 as a ceramic capacitor with dielectric constant, % >1100
[Gene, 1999; Nalwa, 1999]. This study was carried out by Thurnauer, Wainer and
Solomon which was then unpublished [Thurnauer, 1977; Coffeen, 1974, 1975]. With the
World War-II nearing its end in the mid-1940s, it was evidenced from the openly
accessible literature publications that BaTiO3 was the material of study for its dielectric
properties in various countries such as United States, United Kingdom, USSR and Japan.
Independent studies carried out by Wul and Goldman [1945] and von Hippel’s group
[Hippel et al., 1946] in 1945 and 1946 demonstrated the ferroelectric properties of
BaTiO3 to originate from its high dielectric constant, which was consequently affirmed
with similar findings in single-crystal BaTiO3 [Gene, 1999]. First poled BaTiO3
piezoelectric ceramic transducer was operated by R. B. Gray in 1945 [Gray, 1949]. This
confirmed that the orientation of domains occurred with application of an external
electric field giving rise to a ceramic showing single-crystal type behavior having
ferroelectric and piezoelectric properties [Nalwa, 1999; Gene, 1999; Cross and
Newnham, 1987]. Successful application of BaTiO3-ceramic as a transducer in the civil
and the military applications ushered the entry of various ferroelectrics raising the
number from 3 to 25 families by early 1960. Of these, the ones discovered earlier were
Lithium niobate (LiNbO3) and Lithium tantalate (LiTaO3) in 1949, Lead zirconate
titanate (PZT) solid-solution systems in 1952, Lead niobate (PbNb2O6) in 1953, Alkali
niobates in 1955, etc [Gene, 1999; Cross and Newnham, 1987].
All ferroelectric materials are pyroelectric and all pyroelectric materials are piezoelectric
in nature [Gene, 1999; Gopalan et al., 2007; Vijaya and Rangarajan, 2004]. The 32 point
groups structured from the 7 crystal systems are divided into centrosymmetric and non-
centrosymmetric classes as seen in Figure 2.1 [Gene, 1999; Gopalan et al., 2007; Vijaya
and Rangarajan, 2004]. The elements used by crystallographers to define symmetry about
a point in space are, (1) a center of symmetry, (2) axes of rotation, (3) mirror planes, and
10
(4) combination of all. Non-centrosymmetry is a necessary condition for piezoelectricity
to exist. However, of the 21 non-centrosymmetric classes known, only 20 are
piezoelectric and one (cubic class 432) is not due to the other combined symmetric
elements [Gene, 1999; Gopalan et al., 2007; Kholkin et al., 2008; Kang et al., 2006].
Lack of a center of symmetry is the key to producing electric dipoles i.e. polarization due
to the net movement of the positive and the negative ions with respect to each other when
homogeneous stress is applied to the crystal. Piezoelectrics are again divided into
pyroelectrics having spontaneous polarization (Ps) and non-pyroelectrics that do not
posses spontaneous polarization [Gene, 1999]. For non-pyroelectric piezoelectric
materials, stress is the only means of generating dipoles. Piezoelectricity is a linear and
reversible effect in which the magnitude of polarization and the sign of charge produced
depend on the magnitude and the type of stress (tensile or compressive) [Gene, 1999].
The phenomena of generating electric charge (polarization) on application of stress and
mechanical movement (strain) on application of electric field are characteristic of
piezoelectric crystals. The generation of electric charge is known as direct piezoelectric
effect (generator) and that of strain is known as indirect piezoelectric effect (motor)
[Gene, 1999; Vijaya and Rangarajan, 2004].
The birth of piezoelectricity in 1880 and that of ferroelectricity in 1920 was due to the
pyroelectric phenomena observed in the mineral tourmaline as early as 4th Century BC
[Lang, 1999]. Pyroelectricity or pyroelectric effect is the change in spontaneous
polarization of certain anisotropic solids with a change in temperature [Vijaya and
Rangarajan, 2004; Kasap, 2007; Sidney, 2005; Gopalan et al., 2007; Lukas, 1999].
Spontaneous polarization is the dipole moment per unit volume of the material. It exists
in a pyroelectric material (non-zero) even in the absence of an applied electric field
[Sidney, 2005]. A subgroup of pyroelectrics constitutes that of ferroelectrics, which are
spontaneously polarized. Their polarization varies with temperature and can be reversed
by application of an electric field [Gene, 1999; Rabe et al., 2007; Vijaya and Rangarajan,
2004]. Ferroelectric ceramics become piezoelectric when electrically poled. Poling is a
process of domain alignment by application of electric field at temperature just below the
11
Fig
2.1
. Div
isio
n of
poi
nt g
roup
s on
the
bas
is o
f S
ymm
etry
and
Pro
pert
ies.
32
Sym
me
tric
Po
int
Gro
up
s
11
Ce
ntr
osym
me
tric
2
1N
on
ce
ntr
osym
me
tric
20
Pie
zo
ele
ctr
ic
•P
ola
rize
d u
nd
er
str
ess
10
Pyro
ele
ctr
ic
•Sp
on
tan
eo
usly
Po
lari
ze
d
Fe
rro
ele
ctr
ic
•Sp
on
tan
eo
us P
ola
rize
d
•Po
lari
za
tio
n R
eve
rsib
le
No
n-F
err
oe
lectr
ic
10
No
n-P
yro
ele
ctr
ic
1 N
on
-Pie
zo
ele
ctr
ic
12
ferroelectric Curie point [Newnham, 1989]. Various known ferroelectric families include
tungsten bronze (AxWO3, where A is a metal most commonly an alkali and x is a
variable < 1, W is tungsten and O is oxygen), oxygen octahedral (ABO3, where A, B are
cations and O is oxygen), pyrochlore, bismuth-layer structure, boracites, etc. [Gene,
1999; Rabe et al., 2007; Dickens and Whittingham, 1968]. Of these, the oxygen
octahedral family of ferroelectrics having perovskite structure is highly important as it
constitute the major portion of the manufactured ferroelectric ceramics in the form of
BaTiO3, PbTiO3 (Lead titanate), PZT (Lead zirconate titanate) and PLZT (Lead
lanthanum zirconate titanate).
2.2.1 Perovskite Oxides (ABO3)
From the various families of ferroelectrics known, perovskite oxide has been highly
explored. This was due to the structural simplicity and practical use of the first developed
perovskite ferroelectric i.e. BaTiO3 owing to its chemical and mechanical stability at
room temperature and easy synthesis in the ceramic or single crystal form.
Ferroelectricity in this compound was known to be highly structure sensitive and
ultimately became the cause of interest in materials with ABO3 structure for practical
applicability [Nalwa, 1999; Cross and Newnham, 1987]. BaTiO3 was also the first
ferroelectric to display a paraelectric high temperature phase, having the highest
symmetry i.e. centrosymmetric. This paraelectric phase is also called as the prototype
phase. BaTiO3 is also polymorphic as it is found to have more than one ferroelectric
phase [Nalwa, 1999; Gopalan et al., 2007].
Perovskite compounds have their name inferred from the naturally occuring mineral
calcium titanate (CaTiO3) and ABX3 as their general formula [Schwartz, 1997; Jana et
al., 2008; Rabe et al., 2007; Chonghe et al., 2004]. Ideal perovskite is cubic in structure
with Pm3m space group having 5 atoms in the basis. However, the actual perovskite
structure is not cubic but pseudocubic (distorted cubic structure) in nature with a majority
13
being formed as oxides or fluorides [Jana et al., 2008; Schwartz, 1997; Rabe et al., 2007;
Lufaso, 2002; Kumar et al., 2009]. As our interest lies in the oxides of these perovskite,
the formula considered is ABO3, where the anion is oxygen occupying the face-centers of
the unit cell. It consists of two cations wherein the larger A cation occupies the vertices
of the unit cell and the smaller B cation occupies the body-center position in the unit cell
as seen in Figure 2.2. Hence, the perovskite structure is formed with a network of corner-
linked oxygen octahedra with the A-ion occupying the dodecahedral holes in the 12-fold
cuboctahedral coordination and the B-ion occupying the octahedral holes surrounded by
6-fold oxygen in octahedral coordination [Femina and Sanjay, 2012; Wang et al., 2002;
Anjana et al., 2008; Peña and Fierro, 2001]. Perovskite structure is known to
accommodate most of the metallic ions of the periodic table [Lufaso, 2002; Chonghe et
al., 2004]. Distortions in the symmetry of the perovskite structure results in the diverse
physical properties making them useful in various industrial applications [Jana et al.,
2008; Lufaso, 2002; Schwartz, 1997]. As they are capable of holding a large number of
oxygen vacancies, certain perovskite make oxygen ionic conductors [Jana et al., 2008].
Electrical conductivity can also be achieved when small B-site transition element exhibit
multi-valance on exposing the lattice to different conditions. Various perovskite oxides
are known which exhibit such ionic and conducting behavior [Jana et al., 2008]. An
example of such perovskite is BaTiO3 which is known to have applications like
capacitors, PTC thermistors, piezoelectric devices, optoelectronic elements, etc. [Jana et
al., 2008].
Formation and stability of perovskite-type compounds have been related to its tolerance
factor (t) by Goldschmidt in early 1920s as given in eqn. 2.1 [Chonghe et al., 2004;
Kumar et al., 2009; Sullivan, 2009].
! "#$%$"&'($)"*$%$"&+ ---(2.1)
14
Fig
2.2
. Tw
o di
ffer
ent
view
s of
the
uni
t ce
ll s
truc
ture
of
the
idea
l cu
bic
pero
vski
te (
AB
O3)
[R
abe
et a
l., 2
007]
.
15
where, rA, rB and rO are the ionic radii of A, B and O ion respectively. For an ideal
perovskite t=1, having the ratio of bond length A-O to that of B-O as ', : 1, where the
bond length is considered to be the sum of two ionic radii in consideration. However, it
was observed by Goldschmidt that experimental values of t = 0.8 – 0.9 corresponds to the
cubic perovskite structure. The perovskite compounds known till date have‘t’ ranging
from 0.75 – 1.0 [Chonghe et al., 2004; Sullivan, 2009; Kumar et al., 2009]. There have
also been reports of systems with t = 0.8 – 0.9 that do not take the perovskite form.
Hence, an additional parameter known as the octahedral factor (rB/rO) is considered
along with the tolerance factor to predict the perovskite formation, the criterion for which
is as mentioned in the Table 2.1.
Tolerance factor, t = 1 would imply that the ideal cubic perovskite structure constitute of
spherical ions that are closely packed such that the nearest neighbors are in contact
[Sullivan, 2009; Rick, 2007]. Deviation from this value distorts the cubic structure to the
tetragonal, orthorhombic or rhombohedral types [Sullivan, 2009]. For perovskites with t
< 1, the B-O bond gets compressed due to relatively larger B-cation and smaller A-cation
as compared to the ideal values which creates tension in the A-O bond. To relieve the A-
O bond stress, tilting of BO6 octahedra takes place optimizing the cation co-ordination
environment. GdFeO3 is an example of such perovskite with t = 0.81 having an
orthorhombic structure [Johnsson and Peter, 2005]. Similarly, for perovskites with t > 1,
the ionic radius of B-cation is smaller and that of A-cation is larger than the ideal values
[Johnsson and Peter, 2005; Rick, 2007; Sullivan, 2009]. BaNiO3 is an example of such
perovskite with t = 1.13 having an hexagonal structure [Johnsson and Peter, 2005].To
improve the B-O bonding in this case, the B-cation displaces from the center of BO6
octahedra either to the apex, the equatorial edge or towards the face of the octahedra
giving rise to a tetragonal, an orthorhombic or a rhombohedral displacement respectively
(Figure 2.3) [Sullivan, 2009; Dragan, 2005; Newnham, 1989]. Besides the cationic radii,
anionic non-stoichiometry and Jahn-Teller distortions could also result in deviations from
the cubic symmetry in perovskites [Sullivan, 2009].
16
Table 2.1. Criterion for perovskite structure formation [Chonghe et al., 2004; Kumar et
al., 2009; Wang and Kang, 1998].
Parameters Values
-.-/ 0.414 – 0.732
! -0$1$-/',$)-. $1 $-/+ 0.857 – 1.032
-.-/ 2$3 4 -0 1 -/', )-. 1 -/+5 1 6
For A = 4.317, B = 3.912 and for A = 1.327, B = 1.781
Fig 2.3. (a) Displacement of B-site ion in the perovskite (ABO3) structure resulting in the
formation of cubic (left), tetragonal (center) and the rhombohedral (right) structure, (b)
Phase transformation in BaTiO3 [Dragan, 2005; Newnham, 1989].
17
2.2.2 Some Other Ferroelectrics
Lead Based Materials: After successful application of BaTiO3 transducers, studies were
carried out to investigate other ferroelectric compounds having perovskite structure for its
applications. PbTiO3 with a transition temperature (TC) at 490°C was among the first few
to undergo detailed examination. It was reported to have the highest room temperature
spontaneous polarization and a large structural distortion with c/a = 1.064 as compared to
that of 1.01 for BaTiO3 [Nalwa, 1999; Dragan, 2005]. Its electromechanical properties
made it applicable at high temperature and frequency [Nalwa, 1999]. PbTiO3 has varied
applications owing to its ferroelectric, pyroelectric and piezoelectric properties [Lifeng et
al., 2008]. High strain in PbTiO3 developed cracks during single crystal growth.
Polarization degradation, increase in electrical conductivity, crack formation was also
observed in pure PbTiO3 due to polarization reversal resulting in studies replacing Pb2+ to
be carried out. Though this produced crack-free ceramics, the tetragonality was however
reduced [Nalwa, 1999]. PZT (Lead zirconate titanate) (PbTiO3:PbZrO3) solid solution
system (Figure 2.4) with a composition dependent phase diagram, formed the basis for
PZT as the most desirable piezoelectric. The highly advantageous temperature
independent morphotropic phase boundary (MPB) at 52:48 mole fraction composition of
PbTiO3:PbZrO3 helped scientist to tailor phase change over the entire temperature range
of the poling process at which their dielectric, ferroelectric and piezoelectric properties
gets enhanced [Cross and Newnham, 1987; Moulson and Herbert, 2003; Kornev et al.,
2006; Mu et al., 2007; Chamola et al., 2011; Neumeister and Blake, 2011; Devemy et al.,
2009; Gene, 1999]. Easy formation of solid-solution systems with varied constituents by
substituting higher (donors) and lower (acceptors) valence ions expanded the range of
properties achieved with PZT material which has its TC at 390°C [Chamola et al., 2011;
Gene, 1999; Chu et al., 2004]. Of the modified PZTs, PLZT (lead lanthanum zirconate
titanate) produced when Pb2+ was substituted with La3+ enhanced densification rate and
produced pore-free homogeneous microstructure [Santos et al., 2001; Cerqueira et al.,
2000]. PLZT ceramics incorporated all the dielectric, piezoelectric, pyroelectric,
18
Fig
2.4
. Pha
se d
iagr
am o
f P
ZT
cer
amic
[N
ewnh
am, 1
989]
.
19
ferroelectric and electrooptic properties associated with the PZT ceramics [Gene, 1999;
Sabat, 2012]. Depending on the concentration of La, PLZT properties demonstrate a shift
from ferroelectric transducer applications to electrooptic applications because of their
optical transparency from the visible to the near-infrared region and high refractive index
(n & 2.5) [Cerqueira et al., 2000; Sabat, 2012]. PLZT materials are categorized as relaxor
ferroelectrics due to their frequency-dependent diffused ferroelectric-paraelectric phase
transition [Sabat, 2012]. Properties of these materials changed gradually with the
ferroelectric- paraelectric transition occurring over a wide range of temperatures referred
to as permittivity maximum temperature (Tm) [Park et al., 1998]. PMN (Lead magnesium
niobate) relaxor ferroelectrics are known to have better dielectric and electrostrictive
properties than other relaxor materials [Park et al., 1998]. Diffused phase transition in
these ferroelectrics results from the existence of polar nano-regions (PNR) at Burns
temperature (Td) at which PNR starts forming, which lies few hundred degrees above TC
[Burns and Dacol, 1983; Desheng et al., 2009]. Td lies at 327°C (600K) for PMN [Burns
and Dacol, 1983; Desheng et al., 2009; Eudes, 2011; Guangyong, 2011]. Relaxor
ferroelectric systems with Pb(Mg1/3 Nb2/3)O3 – PbTiO3 (PMN-PT) composition possess
high sensitivity, high-power piezoelectric, high dielectric and electromechanical behavior
near the MPB with 35% PT [Hana and Marija, 2011; Urisic et al.,2012 Edwards et al.,
2006; Yaoyao et al., 2011]. Synthesizing PMN-PT as high-density ceramics is difficult
due to undesired lead niobate-based pyrochlore phase formation in the initial processing
stage that degrades its dielectric properties [Eudes, 2011]. Besides, with high
concentration of PT in PMN-PT, the relaxor properties gradually disappear and are
overpowered by the ferroelectric properties [Guangyong, 2011; Eudes, 2011].
With all the advantages associated with these materials, the presence of Pb in PZT which
was most widely used ferroelectric contains ~60% – 70% lead (Pb) by weight which is
hazardous and can cause environmental pollution when exposed to air [Anal et al., 2012;
Hirofumi, 2012; Aksel and Jones, 2010; Kholkin et al., 1997]. According to the Waste
Electrical and Electronic Equipment (WEEE) passed in 2003 by the European parliament
and revised in June 2006, products introduced in the open market should not include
20
more than 0.1 wt.% of the environmentally harmful element lead to comply with the
environment safeguard [Klaus, 2010; Hirofumi, 2012]. Besides being highly toxic, these
materials when prepared in the form of thin films show polarization fatigue i.e.
degradation of its ferroelectric properties, especially when coated on metallic electrodes
[Anal et al., 2012; Zanetti et al., 2004; Yan et al., 1999; Pintilie and Alexe, 1999; Kholkin
et al., 1997; Panda et al., 2004]. To overcome this fatigue problem in PZT, oxide
electrodes are to be used which becomes cost ineffective for IC technology [Zanetti et al.,
2004]. Hence, the current trend is towards preparing Pb-free materials so as to replace
Pb-based compounds in the future [Hirofumi, 2012; Aksel and Jones, 2010; Zanetti et al.,
2004; Wada, 2009]. Currently this forms the basis for the ongoing revived interest in
BaTiO3 compounds.
SrBi2Ta2O9 (Strontium bismuth tantalate, SBT): SBT belongs to the Aurivillius
family of bilayered pseudocubic perovskite oxide material that exhibit ferroelectricity at
room temperature [Anal et al., 2012; Ke et al., 2011; Pintilie and Alexe, 1999]. SBT has
an orthorhombic structure at room temperature with perovskite-type [SrTa2O7]-2 units and
[Bi2O2]+2 layers alternately stacked along the c-axis as seen in Figure 2.5. Spontaneous
polarization in SBT material arises due to the distortion of [Bi2O2]+2, TaO6 octahedral
layers and the atomic displacement in the a-axis which results in highly anisotropic
electrical properties [Anal et al., 2012; Moret et al., 1998]. Polarization fatigue-free
behavior, low leakage current, low coercive field and large polarization makes SBT
highly useful in memory devices in replacing PZT materials [Anal et al., 2012; Alexei
and Yuji, 1998; Kholkin et al., 1997; Yan et al., 1999]. However, high processing
temperature and low remnant polarization makes it non compatible with the current IC
technology and has led to the search for its replacement [Shyu and Lee, 2003; Moert et
al., 2002; Zanetti S. et al., 2004; Shimakawa and Kubo, 1999].
21
2.3 Domain and Polarization
Ferroelectric materials possess regions with uniformly oriented spontaneous polarization
referred to as ferroelectric domains/domains. When two domains with different
polarization directions exist alongside each other, they are separated by a region called as
the domain wall [Vijaya and Rangarajan, 2004; Solymar and Walsh, 2004; Kasap, 2007;
Jiles, 2001; Gopalan et al., 2007; Dragan, 2005; Gene, 1999; Seymen, 2009]. A
polarization change across the domain wall is continuous and steep [Dragan, 2005;
Solymar and Walsh, 2004]. Oppositely oriented polarizations are separated by 180°
domain walls and those with perpendicular polarization orientation are separated by 90°
domain walls [Dragan, 2005; Seymen, 2009]. Figure 2.6 shows the domain orientation
for the 90° and the 180° domains. Mobility of 180° walls is higher than the 90° walls as
mechanical strain is associated only with the 90° walls [Jiles, 2001; Newnham, 1989].
However, the types of domain walls separating the domains are dependent on the crystal
symmetry. For example, the tetragonal phase only consists of the 180° and 90° domain
walls, the orthorhombic phase consists of the 60°, 90°, 120° and 180° domain walls and
the rhombohedral phase consists of the 71°, 109° and 180° domain walls [Seymen, 2009;
Gopalan et al., 2007; Newnham, 1989; Gene, 1999]. Application of an electric field
transforms the multi-domain state into a single domain state by orienting the domain
polarization in the direction of the applied field [Vijaya and Rangarajan, 2004; Moulson
and Herbert, 2003; Seymen, 2009]. During this process of poling the domains with the
polarization oriented in the field direction grows at the expense of those that are
oppositely directed until a single domain orientation exists [Vijaya and Rangarajan, 2004;
Moulson and Herbert, 2003; Newnham, 1989]. Hence, domain walls become an
important source of dielectric loss due to energy dissipation during the process of domain
orientation [Newnham, 1989].
Application of mechanical stress along the polar axis induces switching only through the
90° domains and not through the 180° domains [Moulson and Herbert, 2003]. 90°
domains are the strain-producing domains and the 180° domains do not produce strain.
22
Fig
2.5
. (a)
Sch
emat
ic r
epre
sent
atio
n of
the
Aur
ivil
lius
lay
er s
truc
ture
and
(b)
SrB
i 2T
a 2O
9 st
ruct
ure
proj
ecte
d al
ong
(110
) ax
is.
[Dra
gan,
200
5; Y
an e
t al
., 19
99]
Fig
2.6
. Id
eal
dom
ain
conf
igur
atio
n in
a s
ingl
e cr
ysta
l of
cub
ic f
erro
elec
tric
mat
eria
l, w
here
the
(a)
cou
plin
g of
str
ain
is
negl
igib
le a
nd (
b) s
trai
n ef
fect
s ar
e im
port
ant,
and
(c)
sch
emat
ic d
iagr
am o
f 18
0° a
nd 9
0° d
omai
ns i
n B
aTiO
3 [R
abe
et a
l.,
2007
; M
ouls
on a
nd H
erbe
rt, 2
003]
.
23
Hence, macroscopic changes in the dimensions of the ferroelectric material occur only
when the strain-producing domains are switched [Jiles, 2001; Gene, 1999]. Onset of
spontaneous polarization in a crystal results in the appearance of surface charge density
and the depolarizing field as seen in Figure 2.7a [Vijaya and Rangarajan, 2004; Moulson
and Herbert, 2003; Gopalan et al., 2007]. The energy of this depolarizing field is
minimized by twinning the crystal into many oppositely polarized domains, Figure 2.7b
(The crystal is divided into many 180° domains) [Moulson and Herbert, 2003; Gopalan et
al., 2007].
While aligning the polarization of a ferroelectric material with application of an electric
field, the ferroelectric domains closer to the applied field are favored, which grow due to
nucleation or due to the movement of existing domain walls giving rise to hysteresis
[Gopalan et al., 2007; Moulson and Herbert, 2003]. A ferroelectric hysteresis loop
displays the average polarization reversal with application of the electric field in the
material as seen in Figure 2.8 [Gopalan et al., 2007]. Initially when very small electric
field is applied to the ferroelectric, the polarization increases linearly similar to the linear
dielectric behavior. With increase in the electric field (i.e. at higher fields), nucleation of
new domains in the direction of the field and the irreversible domain wall motion results
in faster polarization response than in the linear stage. The field at which the orientations
of all dipoles align with the field is the saturation point and the ferroelectric behaves like
a linear dielectric again. On achieving a state of saturation polarization, with further
application of the field if an increase in polarization is induced, then this increase
corresponds to the electronic and ionic processes and is not due to domain reversal.
Reversal in the field after saturation point does not return the ferroelectric polarization to
its original value and a remanant polarization (Pr) exist in the material at zero electric
field. The field E, at which equal number of dipoles having opposite polarity corresponds
to zero net polarization, is referred to as the coercive field (EC). On changing the polarity
of the field and applying a field E > EC the material undergoes a similar behavior with the
direction of polarization reversed [Gopalan et al., 2007; Dragan, 2005; Seymen, 2009].
24
Fig
2.7
. (a
) S
urfa
ce c
harg
e as
soci
ated
wit
h sp
onta
neou
s po
lari
zati
on,
(b)
form
atio
n of
dom
ains
to
min
imiz
e el
ectr
osta
tic
ener
gy [
Mou
lson
and
Her
bert
, 200
3].
Fig
2.8
. Pol
ariz
atio
n vs
. ele
ctri
c fi
eld
curv
es f
or a
sin
gle
crys
tal
ferr
oele
ctri
c, p
olyc
ryst
alli
ne f
erro
elec
tric
and
a f
erro
elec
tric
in
the
para
elec
tric
sta
te [
Sey
men
, 200
9].
25
2.4 Dielectrics
Materials with high electrical resistivity are specified as dielectrics and insulators
[Moulson and Herbert, 2003]. Owing to its large energy gap between the valence and
conduction bands, the absence of free electrons in an insulator does not generate flow of
conduction current through it when subjected to an electric field [Vijaya and Rangarajan,
2004; Stuerga, 2006]. On the application of an electric field, an insulator can accumulate
electric charge and hence electrostatic energy [Stuerga, 2006]. The principle use of an
insulator is to hold the conducting elements in position and avoid contact between them
[Moulson and Herbert, 2003]. Dielectric cover a wide range of materials including
electrolytes and are used in circuit functions in which their permittivity and dissipation
factor plays an important role [Stuerga, 2006; Moulson and Herbert, 2003]. A good
dielectric is necessarily a good insulator, but the converse is not true [Moulson and
Herbert, 2003].
Application of high-frequency electromagnetic energy heats up an insulating material as
polarization induced in the material cannot follow the rapid reversal of electric field
[Stuerga, 2006; Elissalde and Ravez, 2001]. Polarization in dielectrics originates from
local reorganization of linked and free charges [Stuerga, 2006]. Electrical component of
the electromagnetic wave is known to induce current of free charges (electronic or ionic
origin) and also to induce local reorganization of linked charges (dipolar moments).
Molecules that do not possess a center of symmetry called as polar molecules have
permanent electric dipole moment and are electrically neutral due to equal amount of
positive and negative charges on them. Thermal motion of the polar molecule reduces the
alignment of permanent dipoles by the electric field resulting in a rapid declination of the
polarization orientation. Being associated with the chemical bonds, any motion in the
dipoles induces a correlative motion of molecular bonds. As dielectric properties are
group properties, they cannot be associated by an interaction between a single dipole and
electric fields [Stuerga, 2006].
26
When a neutral atom is placed in external dc electric field (E), the positive and the
negative charges gets displaced. The motion of the nucleus (positively charged) is in the
direction of the field and the surrounding electrons move in the opposite direction
maintaining a distance between the center of positive and negative charges that undergoes
a net relative displacement (R). This neutral atom subsequently acquires an electric
dipole moment (p) proportional to E as given in eqn. 2.2 [Vijaya and Rangarajan, 2004].
p = ' E ---(2.2)
where, ' is known as the polarization of the atom (or molecule).
Hence, when a dielectric material is placed between two current carrying conducting
plates, charged stored on the plates (Q) is higher than that when the dielectric is replaced
with air/vaccum (Qo). This increases the capacitance i.e. the charge storage ability per
unit voltage of the system from Co to C as depicted in Figure 2.9. The dielectric constant
also known as the relative permittivity (#r) is the ratio of the capacitances with and
without the dielectric medium written as (eqn. 2.3)
7 ! $ 889 !$ :
:9 ---(2.3)
The increase in stored charge is a result of polarization build in the dielectric by the
electric field [Kasap, 2007]. Polarization is the separation of positive and negative
charges that induces dipole moment. Polarization process called electronic polarization
result from a dipolar moment induced by distortion of electron shells. Atomic
polarization is a process in which the polarization results from the dipolar moment
induced by distortion of nuclei positions when an electromagnetic field in the infrared
region induces atomic vibrations in molecules and crystals. These two polarization
processes can be connected together in distortion polarization. In the microwave range,
however, the electromagnetic fields lead to rotation of polar molecules or charge
redistribution that corresponds to the process of orientation polarization [Stuerga, 2006].
In a dielectric, under the application of an electric field when all dipoles align head to tail,
the charge within the bulk gets nullified. However, the charges on the surface do not get
27
Fig
2.9
. (a)
Cap
acit
ance
of
the
para
llel
pla
te c
apac
itor
wit
h fr
ee s
pace
bet
wee
n th
e pl
ates
, (b)
cha
rge
buil
d up
whe
n di
elec
tric
is
bein
g pl
aced
, and
(c)
cap
acit
ance
whe
n di
elec
tric
is
plac
ed b
etw
een
the
plat
es [
Kas
ap, 2
007]
.
Fig
2.1
0.
(a)
Die
lect
ric
plac
ed i
n an
ele
ctri
c fi
eld,
(b)
ele
ctri
c fi
eld
give
s ri
se t
o bo
und
pola
riza
tion
cha
rges
, an
d (c
)
repr
esen
ting
the
who
le d
iele
ctri
c in
ter
ms
of i
ts s
urfa
ce p
olar
izat
ion
char
ges
+Q
P a
nd –
QP [
Kas
ap, 2
007]
.
28
cancelled and are bound due to polarization of the molecules as seen in Figure 2.10.
These charges are called as surface polarization charges (Qp) related to the dipole
moment of the material as
p = Qp d
where, d is the total distance between the surface charges on either plates.
Polarization (P) of the dielectric medium is defined as net dipole moment (ptotal) per unit
volume.
; ! $ <=>=?@ABCDEF ! G$<?H !$IJ3 ! $KJ
where pav is the average dipole moment per unit molecule, N is the number of dipoles per
unit volume, and (p is the surface charge density [Kasap, 2007].
For a linear isotropic dielectric, the polarization is proportional and parallel to the applied
field E. The electric displacement D is related to P and E as
D = E + 4) P = # E
where, # is the scalar dielectric constant of the medium.
# = 1 + 4) *
where, * = P/E is the electric susceptibility, and #, * are scalar quantities dependent on the
molecular properties of isotropic dielectric.
For anisotropic dielectrics, #, * and ' are tensors.
LM !$N7OM $PMQ
MRS$$$$$$$$$$$$$$$$$$$$$$$$$$$$$$$$$)T ! UV ,V W+
Therefore, when the dielectric is subjected to an alternating field, both D and P vary
periodically with time. However, due to the inertial and energy dissipation effects
(losses), these parameters will not respond to the field instantaneously and will lag in
phase with the field E.
29
i.e. if E = E0 cos (+t)
D = D0 cos (+t – ,)
where, , is the loss angle independent of E0 but dependent on frequency.
In the presence of dielectric losses and relaxation effects, # becomes a complex quantity
composed of charging (real) and loss (imaginary) component.
i.e. #* = #- – I #ý
and the loss angle related to these values is given in eqn. 2.4 as
XYZ$[$ ! $ \]\^ ---(2.4)
Loss angle is related to the Q-factor of the dielectric as Q = 1/ tan, [Rout, 2006; Kasap,
2007; Vijaya and Rangarajan, 2004; Seymen, 2009]. On application of alternating
current, the dipole orientation gets out-of-phase with the electric field reversal at certain
frequencies (relaxation frequency). This effect, referred as the dielectric relaxation, is
characterized by a decrease in the real and an increase in the imaginary component of the
dielectric constant [Elissalde and Ravez, 2001].
2.4.1 Classes of Dielectrics
Current trend towards increased inline voltages due to transmission network, high
frequency application due to advances in telecommunication systems, higher computing
power due to improved computer technology, development of satellite and mobile
communications, makes it essential for the behavior of the insulating ceramics used for
these applications to meet their respective requirements. One of the major applications of
a dielectric is in capacitors. Ceramic dielectrics and insulators cover wide range of
properties with relative permittivity ranging from 6 to > 20,000 [Moulson and Herbert,
2003]. Dielectrics are divided into four classes. Class I dielectrics include low-
permittivity ceramics with #r <15 and medium- permittivity ceramics with loss angle
(dissipation factor) < 0.003, and #r ranging from 15 – 500. The medium-permittivity
30
ceramics have stable temperature coefficient of resistivity. Class II/ III dielectrics consists
of high-permittivity ferroelectric ceramics. The value of #r of these classes ranges
between 2000 – 20,000 and the dissipation factor usually < 0.03, which may increase in
certain temperature ranges and on application of high electric field. Table 2.2 gives the
temperature coding and capacitance variation for this class of dielectrics. These classes of
dielectrics render higher variation of material properties with temperature, field strength
and frequency than class I dielectrics. These classes of dielectrics possess high
volumetric efficiency. BaTiO3-based dielectric and relaxor ferroelectric (e.g. PMN) are
an important example of class II/ III dielectrics. Class IV dielectrics contains a
conductive phase which reduces the thickness of dielectric by at least one order of
magnitude in the capacitor. However, they suffer high losses and their application is
limited to low working voltages typically between 2V - 25 V. Barrier layer capacitors are
examples of the class IV dielectrics which are rarely used [Moulson and Herbert, 2003].
To be of commercial use, ceramic capacitors must either be low in cost so as to compete
with polymer film units or must possess special and superior qualities. These
requirements govern the structure and method of manufacture of a capacitor. Being the
lowest in cost, disc and tubular shape capacitors are used for all classes of dielectrics. The
working capacitance range for a class I dielectric is from 0.1pF – 1000pF, for class II/ III
dielectrics is 1000pF – 100,000pF and for class IV is from 0.1"F - 2"F. Class I, II and III
dielectrics can be employed safely up to voltages of 100V, although the maximum
applied voltage in electronic circuits is ~ 10V. The thickness of dielectric ranges from
50"m - 2mm with the thicker units being able to withstand mains supply directly. Disc
diameter measured without encapsulant ranges from 2mm - 30mm, while tubes may be
5mm – 60mm long and 1mm – 10mm in diameter [Moulson and Herbert, 2003].
Class I ceramics with #r < 15 are low-permittivity dielectric which are used for insulation,
wherein their mechanical properties are considered over their dielectric properties. Large
scale usage of these requires them to be low-cost material. The dielectric properties of
these are of importance when used as substrate for components. It is used as capacitors
31
Table 2.2. Class II/ III dielectric capacitor coding for temperature and capacitance range.
EIA codes D – R belong to class II and S – V belong to class III dielectrics [Moulson and
Herbert, 2003].
Example: For an EIA code X5D, capacitance in the temperature range -55°C to +85°C
will either increase or decrease no more than 3.3%.
EIA Code Temperature range
(°C)EIA Code
Capacitance
change (%)
X5 -55 to +85 D ±3.3
X7 -55 to +125 E ± 4.7
X8 -55 to +150 F ± 7.5
Y5 -30 to +85 P ± 10
Z5 +10 to +85 R ± 15
S ± 22
T +22 to -33
U +22 to -56
V +22 to -82
32
when very small capacitances at high frequencies or high current passing capacitors are
needed. Since medium-permittivity class I dielectric have dissipation factor i.e.
tan, < 0.003, most ferroelectric compounds are ruled out as their losses are high (i.e.
tan, > 0.003) especially when subjected to high a.c. fields. These dielectrics find
application in high-power transmitter capacitors for 0.5MHz – 50MHz frequency
range, stable capacitors for general electronic use in 1kHz – 100MHz frequency range
and microwave resonant devices in 0.5GHz – 50GHz frequency range. Materials that
belong to class I dielectrics include TiO2, SrTiO3, CaTiO3, MgTiO3, etc.
2.5 BaTiO3: Its Importance and Modification
BaTiO3, the classic ferroelectric has tetragonal structure as its room temperature stable
phase. The critical temperature above which BaTiO3 loses its ferroelectric property is
known as Curie temperature (TC) [Kasap, 2007]. Below TC BaTiO3 is spontaneously
polarized, and this spontaneous polarization is accompanied with the crystal structure
distortion due to the displacement of Ti ion [Kasap, 2007]. Also, the ease with which
BaTiO3 can be switched back and forth between different ferroelectric phases led to
scientific and technological interest in it. These ferroelectric instabilities arise from the
covalent hybridization between Ti and O ions [Sanna et al., 2011]. Figure 2.11 shows the
unit cell structure of BaTiO3, its ferroelectric transitions and the phase diagram of BaO-
TiO2 system [Moulson and Herbert, 2003; Rase and Roy, 1955; Avrahami, 2003; Lee et
al., 2007]. BaTiO3 was the first piezoelectric transducer developed in 1945 [Gene, 1999].
Piezoelectric materials are one of the highly used materials in sensor and actuator
technologies due to its unique ability of coupling electrical and mechanical
displacements. These materials offer high pressure per density ratio for actuator devices,
high environmental and chemical stability compared to other electromechanical
transduction technologies for example microelectromechanical systems (MEMS) [Aksel
and Jones, 2010]. The wide range of applications as multilayer ceramics capacitors
(MLCC), piezoelectric sensors, transducers, actuators, non-volatile ferroelectric random
33
Fig
2.1
1. (
a) B
aTiO
3 un
it c
ell
stru
ctur
e, (
b) B
aTiO
3 po
lym
orph
dis
tort
ion
of t
he p
erov
skit
e st
ruct
ure
and
(c)
pseu
do-b
inar
y
phas
e di
agra
m o
f B
aO-T
iO2
syst
em.
34
access memories (FRAM), dynamic random access memories (DRAMs) electro-optic
devices and positive coefficient of resistance (PTCR) thermistors makes BaTiO3 one of
the highly desired electroceramic of the various known ferroelectrics [Kumar et al., 2009;
Parveen et al., 2009; Mahajan et al., 2009; Gene, 1999; Wang et al., 2002; Xiaodong et
al., 2007; Gust et al., 1997; Park et al., 2010]. Also, the simple crystal structure, high
stability, extremely high dielectric constant at the transition temperature, low leakage
current and anisotropic optical behavior adds up to the importance of BaTiO3 [Sabina,
2001]. In its tetragonal phase, BaTiO3 can be used in many electronic devices as
underwater transducers, sensors and heaters owing to its ferroelectric properties.
However, in the cubic form i.e. the paraelectric phase, BaTiO3 has high dielectric
constant which makes it suitable as capacitors such as MLCC [Mohammad et al., 2009].
Owing to its important applications, modifying BaTiO3 to improve its electrical
properties, by controlling the microstructure and composition, is of great interest [Kumar
et al., 2009; Gene, 1999; Mahajan et al., 2009]. The microstructure of BaTiO3 can be
modified either by using additives to prohibit grain growth for obtaining highly dense
ceramics, or by employing novel processing techniques [Kumar et al., 2009; Mahajan et
al., 2009]. Substituting additive ions in BaTiO3 structure at the A- or B- site cations can
be achieved using isovalent or aliovalent cations [Mahajan et al., 2009]. Aliovalent
cations can be grouped as donors (substituting ion having valence higher than the host
ion) and acceptors (substituting ion having valence lower than the host ion). Highly
studied dopants in BaTiO3 are, Sr2+, La3+ [Mahajan et al., 2009; Dey and Majhi, 2005;
Zhang et al., 2007; Panwar and Semwal, 1991; Cheng and Chang, 1995]. Based on the
transducer application, Sr2+ is used to reduce TC below 120°C, whereas Pb2+ is used to
vary TC above 120°C, Ca2+ is doped to increase temperature stability range of the
tetragonal phase and Co2+ to reduce the high-electric field losses without affecting the
piezoelectric constants [Gene, 1999]. For BaTiO3application as a capacitor, TC shifters
such as SrTiO3, CaZrO3, PbTiO3 and TC depressors such as Bi(SnO2)3, MgZrO3, CaTiO3
are used as additives [Gene, 1999]. With TC shifters, lower TC is desired so as to achieve
high permittivity values near room temperature or the temperature of operation, and with
35
TC depressors, the flatter permittivity- temperature profile is desired by depressing the
sharpness of the permittivity at TC so as to have a wider temperature range in which
maximum permittivity is obtained rather than having it at one particular temperature itself
[Gene, 1999]. Well known is the fact that a pure BaTiO3 is an insulator at room
temperature, and can be doped to bring its resistivity in the semiconducting range. Many
dopants such as La3+, Y3+, Sb3+ in place of Ba2+ and Nb5+, Ta5+, Sb5+ in place of Ti4+
referred to as donors are known to increase the conductivity of BaTiO3 inducing n-type
semiconductivity. This results in a sudden resistivity rise at TC by 3 to 6 orders of
magnitude showing positive temperature coefficient of resistivity (PTCR) [Panwar and
Semwal, 1991; Ruitao et al., 2004; Kim, 2002; Kareiva et al., 1999]. These materials
have many applications such as current limiters, constant temperature heaters, thermal
sensors, also used in investigation of heat exchange processes in plants and living
organisms, due to which they are widely studied [Kareiva et al., 1999].
BaTiO3 is rarely used in its pure form and is doped to modify the properties according to
the specific application [Liang et al., 2004; Arya et al., 2003; Sutham, 2008; Vargas-Ortíz
et al., 2012; Da-Yong et al., 2006; Mahajan et al., 2009; Sun et al., 2007]. Of the various
dopants incorporated in BaTiO3, Sr is vastly searched. Shift of TC towards room
temperature in Sr doped BaTiO3 can be tuned by adjusting Ba-to-Sr ratio [Viviani et al.,
2004; Sutham, 2008; .miga et al., 2009; Czekaj et al., 2010; Valant and Suvorov, 2004].
Ba1-xSrxTiO3 (BST) ceramics are used in the fabrication of electronic components having
wide range of applications such as electromechanical sensors, transducers, actuators, etc.
[Sutham, 2008; .miga et al., 2009; Dudley et al., 2006]. High dielectric constant, low
dielectric loss and non-linear behavior of dielectric properties with respect to the applied
dc voltage makes BST materials suitable for various microwave devices such as phase
shifters, tunable filters and high-Q resonators operating at room temperature [Czekaj et
al., 2010; .miga et al., 2009; Chen et al., 2009; Qiwei et al., 2010; Liang et al., 2004;
Wodecka-Du/ et al., 2007; Xue et al., 2007]. Very few studies have been carried out on
the PTCR behavior of BST. These work employed high temperature sintering (>1300°C)
of BST [Sutham, 2008; Cheng and Chang, 1995]. Donor doping in BaTiO3 is known to
36
produce semiconducting properties in the otherwise known insulating BaTiO3 [Panwar
and Semwal, 1991]. La3+ and Ce3+ doped at the Ba2+ site acts as donors as their valence is
higher than that of Ba [Panwar and Semwal, 1991; Beltrán et al., 2004; Da-Yong et al.,
2006; Masó et al., 2008; Sabina, 2001; Park et al., 1998]. Both these dopants have been
reported to shift the paraelectric to ferroelectric transition temperature (TC) towards lower
temperature value [Da-Yong et al., 2006; Beltrán et al., 2004; Panwar and Semwal, 1991;
Masó et al., 2008; Park et al., 1998; Sabina, 2001]. Doped BaTiO3 when sintered in air is
known to exhibit high dielectric constant and show PTCR effect, which was verified with
La incorporation in BaTiO3 [Burcu, 2012; Da-Yong et al., 2006; Beltrán et al., 2004;
Panwar and Semwal, 1991; Masó et al., 2008]. However, Beltrán et al. [2004] showed
that there exists a substitution limit of La in BaTiO3, above which the semiconducting
properties disappear and the material again becomes insulating in nature. Owing to its
dielectric properties, La doped BaTiO3 have been used as internal boundary layer
capacitors (IBLC) and as PTCR thermistors [Yanxia et al., 2007; Beltrán et al., 2004]. Ce
doped BaTiO3 ceramics have been used in MLCCs, meeting high capacitance
requirements [Park et al., 1998]. Increasing use of nickel electrodes in MLCCs required
annealing in reduced atmosphere, resulting in semi-conductivity of the material due to
oxygen vacancy. To reduce the effect of these defects on the dielectric properties of
BaTiO3-based devices, Mg2+ as an acceptor was added, to replace Ti4+ for charge
compensation [Min et al., 2007; Li et al., 2007; Fukuda et al., 2007]. Acceptors are often
present in the form of impurities but are sometimes added as sintering aids and also to
modify the PTCR properties of BaTiO3 [Fisher et al., 2006]. Mg doping has been
reported to achieve small grains with high density [Min et al., 2007]. Mg Doping
decreases TC towards lower temperature and has been shown to possess resistivity values
in the insulating range [Min et al., 2007; Li et al., 2007]. Hetero-valent doping of BaTiO3
with alkali ions like Li+ and K+ have been analysed using FTIR to study the effect of
doping on the structure [Sun et al., 2007; Jin et al., 2009]. K doping in modified BaTiO3
for 20µm particles was reported to have properties that make them a potential PTCR
material and also to increase the dielectric constant of nanometer size BST material
[Senlin et al., 2009; Xiaohua et al., 2006]. However, very few studies have been carried
37
out on the effect of incorporating these ions on its dielectric properties and PTCR effect
[Senlin et al., 2009; Dai et al., 2010; Xiaohua et al., 2006].
2.6 Techniques Employed in the Synthesis of BaTiO3
BaTiO3 properties are well known to be dependent on the synthesis and the sintering
routes employed, which set the basis for the vast studies carried out in these areas,
besides those carried out to learn the effect of barrier layers, boundary layers, external
fields, and space charge layer. The recent shift in interest from bulk BaTiO3 to
synthesizing nanoscale BaTiO3 is due to the fact that the electrical properties are strongly
dependent on the grain size and the crystalline structure (microstructure and composition)
[Un-Yeon et al., 2004; Xinhua et al., 2005; Dazhi et al., 2007; Gang et al., 2010; Kumar
et al., 2009]. With a decrease in the grain diameter towards 1"m, an increase in room-
temperature dielectric constant has been reported [Ihlefeld et al., 2007]. BaTiO3 films
consist of fine grains of sub-micron to nano meter size and are used in the advanced
electronic devices such as MLCC [Tohma et al., 2002; Frey and Payne, 1996; Huang et
al., 2007]. Thin films are advantageous due to their smaller size, light weight, easy
integration, lower operating voltage higher speed and unique structure [Wang et al.,
2002]. However, superior dielectric properties associated with the tetragonal structure of
BaTiO3 disappears with decreasing crystallite size transforming it to the cubic structure at
the critical grain size [Mohammad et al., 2009; Huarui and Lian, 2003; Blake, 2004].
Powder processing method depends on two factors, firstly on the cost and second, more
importantly on the end application [Gene, 1999; Vijatovi et al., 2008]. Since the
ferroelectric properties are highly dependent on the grain size, domain structure and
composition; controlled processing conditions of these materials, both at the powder
synthesis stage and subsequent densification to a solid component needs to be maintained
[Pavlovi et al., 2002; Valdivieso, 1996]. Particle agglomeration influences the
development of microstress, microstructural defects and agglomerated grains affecting
the microstructural development of a sintered material. Also, agglomeration of particles
38
depends on the size distribution, degree of agglomeration and relative density [Pavlovi
et al., 2002].
2.6.1 Solid-State Reaction Method
Synthesis of BaTiO3 has been carried out with solid-state reaction method since the time
of its discovery in 1940s, which involved ball milling of BaCO3 or BaO with TiO2 at
temperature >1000°C commercially [Darko et al., 2003; Li et al., 2007; Chie et al., 2006].
This method produces large particles (2 - 13"m) with irregular uncontrolled morphology,
having high impurity contents (e.g. BaCO3 and other intermediate phases formed) due to
high reaction temperature and heterogeneous solid phase reaction, limiting the electrical
properties of the sintered ceramics [Un-Yeon et al., 2004; Xinhua et al., 2005; Huarui and
Lian, 2003; Darko et al., 2003; Aparna et al., 2001; Ruitao et al., 2004; Mahajan et al.,
2009; Dazhi et al., 2007]. However, Chie et al. [2006] in their effort towards improving
solid-state process that require high processing temperature showing unavoidable
consequences of grain growth, synthesized pure BaTiO3 with an average particle size of
120 nm, by the solid state route with BaCO3 and TiO2 as the starting materials, to which
bovine serum albumin (BSA) was added to aid the reaction process. However, calcination
of the precursor mixture was carried out at 1000°C and sintering of the sample pellets
was carried out at 1300°C. The particles produced were tetragonal in nature, and the
tetragonality decreased with increasing BSA content. To outdo the limitations of the
conventional method, various wet-chemical routes, such as sol-gel process [Radonji et
al., 2008; Barik et al., 2009; Nahum et al., 2008; Beltrán et al.,2004; Hiromitsu and
Atushi, 2003], microemulsion-mediated emulsion route [Pithan et al., 2006],
coprecipitation process [Blake, 2004], hydrothermal route [Xinhua et al., 2005; Huarui
and Lian, 2003; Qi et al., 2005], and polymeric precursor route [Durán et al., 2001] have
been developed to generate highly pure, homogeneous, reactive ultrafine barium titanate
powders at low temperatures [Vijatovi M. et al., 2008; Un-Yeon et al., 2004; Xinhua et
al., 2005; Srimala et al., 2008].
39
2.6.2 Hydrothermal Method
Hydrothermal synthesis of BaTiO3, due to the combined effects of solvent, temperature,
and pressure on the ionic reaction, stabilizes desirable products, inhibiting the formation
of undesirable compounds such as BaCO3. Using this method, synthesis of BaTiO3 in a
single step without using sophisticated apparatus or expensive reagents has been
achieved. Barium chloride/hydroxide, titanium chloride and a mineralizer for colloid
forming is used as the starting materials [Qi et al., 2005; Huarui and Lian, 2003; Xinhua
et al., 2005; Min et al., 2007; Miao et al., 2007]. Tetragonal BaTiO3 have been
synthesized at temperatures about 240°C to temperatures greater than 400°C, using this
method but the particles obtained were of submicrometer or even millimeter size [Wu et
al., 1996; Kajiyoshi et al., 1991; Christensen, 1970]. Nanometer-sized BaTiO3 have also
been prepared with cubic structure as reported from XRD and DSC analysis [Zhu et al.,
1997; Dutta et al., 1994; Pinceloup et al., 1999; Hu et al., 2000; Lu et al., 2000].
However, on being analyzed via Raman spectroscopy, these cubic powders revealed
tetragonal structure, in which small particles aggregated to form large particles having
dielectric constant lower than the unaggregated submicrometer-sized particles. These size
and structure variations hinder their application in MLCCs [Huarui and Lian, 2003]. The
formation of cubic structure of nano-powders was attributed to the incomplete transition
from the cubic-to-tetragonal phase (distortion of TiO6 octahedra) when cooling the
sample through the Curie temperature. It was suggested that due to the small size of the
nanocrystals the structural defects in the particles prevented them from completing the
structural transition leading to high strains within the crystals [Xinhua et al., 2005;
Takeuchi et al., 1997]. The hydrothermally produced particle morphology depends on the
alkaline concentration, stirring conditions and reaction temperature. Various authors
achieved cubic phase BaTiO3 on completion of synthesis without calcining at high
temperature [Dutta et al., 1994; Lu et al., 2000; Hu et al., 2000; Woo et al., 2000;
Sasirekha et al., 2008; Maison et al., 2003; Pulugurtha et al., 2010]. However, when
sintered, these powders produce abnormally densified tetragonal grains. Whereas the as-
prepared tetragonal-phase powders obtained by Xu and Gao well densified maintaining a
40
smaller and more uniform grain size on drying at 60°C for 24 hrs in vacuum oven
[Huarui and Lian, 2003]. Hence, it is desired to generate nanocrystalline tetragonal
BaTiO3 with good dielectric properties at temperature as low as 250°C [Huarui and Lian,
2003; Joshi et al., 2010]. This is also supported by Zhu et al. [1997] who claimed that
hydrothermally synthesized tetragonal-BaTiO3 powders have better sintering behavior
than the cubic-phase powders. It was also suggested that the strains developed in the
particle prevent it from assuming the tetragonal form and that the presence of OH-
impurities in the lattice leads to defects and result in micro-strains that keep the particle
cubic [Huarui and Lian, 2003; Zhu et al., 1997; Dutta et al., 1994; Sasirekha et al., 2008;
Joshi et al., 2010; Huarui et al., 2002; Detlev et al., 2001].
2.6.3 Coprecipitation Method
Coprecipitation is a simple convenient method by which homogeneous compounds, due
to mixing of constituent ions on the molecular level under controlled condition, can be
achieved [Vijatovi et al., 2008; Gaikwad et al., 2005]. Among the various wet chemical
synthesis routes, this method is of interest as it avoids the preliminary preparation of solid
or gel precursor [Buscaglia et al., 2004; Her et al., 1996; Kumar, 1999; Leoni et al., 1996;
Viviani et al., 2003]. Stevens Blake [2004] synthesized BaTiO3 using barium hydroxide
and titanium n-propoxide as the precursors. Oxalate process is a coprecipitation process
in which a mixed solution containing metallic cations are added in an oxalate acid
solution. Low solubility of the metal oxalate causes them to precipitate [Park et al., 1997;
Cheung et al., 2001; Vijatovi et al., 2008]. Coprecipitation via oxalate route makes it
difficult to optimize the synthesis conditions at which both Ba and Ti precipitate
simultaneously. This is because only at pH 0 2 does Ti precipitate as titanyl oxalate and
at pH 1 4 Ba precipitates as BaC2O4 which affects the Ba:Ti stoichimetric ratio [Vijatovi
et al., 2008]. Nitrogen gas was used to hinder the formation of BaCO3 and the solution
pH was controlled with NaOH to increase the effectiveness of the surfactant that was
added. This process produced high yield ratio but did not allow the control of particle size
41
and interparticle agglomeration [Park et al., 1997; Gallagher et al., 1963; Northover,
1965; Schrey, 1965; Gallagher and Thomson, 1965].
2.6.4 Polymeric Precursor Method
Polymeric precursor method consists of rigid polyester network with immobilized metal
ions without segregation of the cations during thermal decomposition of organic material
[Vijatovi et al., 2008; Guang et al., 1998; Stojanovi B. et al., 2002; Ashok et al., 2005].
This method employs complexing of cations in an aqueous-organic medium and use of
low cost precursors result in a homogeneous ion distribution at the molecular level [Ries
et al., 2003; Lee et al., 2000; Saravanan et al., 2007]. Limitations in the form of large
process duration and higher requirement of organic materials are associated with this
method [Vijatovi et al., 2008; Ries et al., 2003]. However, it is still used because of its
control over stoichiometry, lower synthesis temperature, and easy introduction of dopants
in comparison to the mixed metal oxide synthesis route [Mirjana et al., 2008; Durán et
al., 2001; Lee et al., 2000; Ashok et al., 2005]. Pseudocubic BaTiO3 particles were
successfully synthesized by this method with barium nitrate, titanium tetrabutoxide, citric
acid, ethylene glycol and nitric acid as the starting reagents and a calcination temperature
of 435°C [Durán et al., 2001].
2.6.5 Thermal Decomposition Method
Thermal decomposition method is the one in which the synthesized powder parameters
are controlled during heat treatment process [Polotai et al., 2004; Peng et al., 2009; Wada
et al., 2008]. A desire to synthesize nanosized powders with calcination temperature
lower than 700°C led to the study of flexible non-isothermal heating regime, as the
heating rate is recognized to allow flexible control of transformation, particle size and
morphology [Polotai et al., 2004]. At low heating rates, the reaction and nucleation rate is
low, with long processing time leading to a normal nuclei growth giving coarser particles,
and at rapid heating rates, faster nucleation results in large nuclei density and coarsening
42
due to coalescence [Polotai et al., 2004]. The starting compounds used are either
commercially available or as synthesized barium titanyl-oxalate/glycolate/peroxide
precursor [Polotai et al., 2004; Panwar and Semwal, 1991; Huang et al., 2006; Zhang et
al., 2006; Ischenko et al., 2007]. The particle size attained via this route is reported to
exist between 10- 40nm, with the soft agglomerate of size 200-700nm as achieved by
Polotai et al [2004]. The powders reported by Zhang et al. [2006] with a particle size of
20-30nm were cubic in nature. Wada et. al [2008] synthesized defect free BaTiO3 with
particle size between 10-300nm using two-step thermal decomposition method and
reported a dielectric constant of ~30,000 for BaTiO3 particle size between 60 and 85nm.
2.6.6 Mechanochemical Method
Mechanical treatment of ceramic powders reduces particle size making it possible to
obtain nano-structured powders [Vijatovi et al., 2008]. Mechanical activation influences
different structural changes such as phase transitions, dislocation generation and crystal
lattice microstrains, the driving force of which is enthalpy as a result of activation in
high-energy mills [Stojanovi et al., 1999; Pavlovi et al., 2002; Pavlovi et al., 2008;
Marinkovi et al., 1999]. During this process, heat is released initiating solid-state
reaction due to the formation of new surfaces and increasing concentration of crystal
lattice defects. This is accompanied by an increase in surface energy of the particles,
increasing the reactivity of the starting mixture and lowering the sintering temperature
[Vijatovi et al., 2008; Pavlovi et al., 2002; Pavlovi et al., 1999; Rojac et al., 2005].
Pavlovic et al. observed a decrease of crystallite size, agglomeration and tetragonal c/a
ratio of BaTiO3 with an increase of mechanically mixing time. Agglomeration is known
to influence the development of microstress, microstructural defects and exaggerated
grains which affects the microstructural development of a sintered material, the effect of
which depends on the size distribution, degree of agglomeration and relative density
[Pavlovi et al., 2002].
43
2.6.7 Sol-Gel Method
Barium titanate synthesized by this method is obtained by hydrolyzing the chemical
precursor, usually barium acetate and titanium iso-propoxide to form a sol and then a gel,
which on drying and pyrolysis gives its amorphous oxide that is further heat treated to
induce crystallization [Kareiva et al., 1999; Nahum et al., 2008; Beltrán et al., 2004].
Various combinations of calcination and sintering temperatures of 1000°C and1400°C
are used in synthesizing BaTiO3 with this method. Partial hydrolysis of metal alkoxide
form reactive monomers and polycondensation of these monomers form colloidal-like
oligomers. Additional hydrolysis promotes polymerization and cross-linking leading to a
3-D matrix (gel). To crystallize barium titanate, the hydrolysis product is usually calcined
at temperatures above 500°C [Srimala et al., 2008]. Beltrán et al. [2004] have reported a
sol-gel route to produce La doped BaTiO3 ceramics, using barium acetate, lanthanum
acetate and tetra-isopropyl orthotitanate as the starting materials to obtain a clear solution
which was dried and gelled at room temperature. These gels were calcined at 1100°C,
powdered, pelletized and sintered again at 1100°C. They reported these powers to be
more resistive than the ones prepared by the solid-state reaction route. Barik et al.
employed the sol-gel hydroxide route in the synthesis of BaTiO3 wherein barium
hydroxide and tetra-isopropyl orthotitanate were used as the starting materials producing
tetragonal particles with an average size of 44 nm at 750°C as the calcination temperature
[Barik et al., 2009].
2.6.8 Refined Sol-Gel Method
Sol-precipitation, sol-crystal, microemulsion, sol-emulsion-gel, gel-sol process are
recently developed refined sol-gel processes [Vijatovi et al., 2008; Takeuchi et al., 1997;
Pithan et al., 2006; Chatterjee et al., 1999; Un-Yeon et al., 2004]. Barium
metal/acetate/chloride/nitrate and titanium iso-propoxide are used as the starting
materials. Pithan et al. [2006] in order to characterize the structural and morphological
evolution of the as-prepared powders upon calcinations, annealed the prepared raw
44
powders in air at temperature between 100°C and 1300°C using different cycles for a
soaking time of one hour at a heating rate of 5°C/min. In their study, the as-synthesized
powders revealed strong broadening Bragg reflection corresponding to phase-pure and
crystalline BaTiO3 with minor impurities in the form of BaCO3. Impurity level increased
with calcination temperature, as seen from the diffraction pattern of the calcined powders
at 400°C to that of 600°C wherein high amount of BaCO3 was noted. Calcination at
700°C marked the onset of particle growth with a pseudo-cubic structure due to reduced
breadth of the diffraction peak observed. Distinct presence of (002) and (200) Bragg
reflections above this temperature associated with tetragonal BaTiO3 was observed
[Pithan et al., 2006]. This implies that BaTiO3 formed in the as-synthesized powders were
cubic in structure producing broad Bragg reflection peaks with small amount of BaCO3.
Takeuchi et al. [1997] proposed that the presence of organic compounds, occluded water
and BaCO3 may interrupt the ordering of the tetragonal domains leading to a pseudocubic
phase in the as prepared samples or the small domain size may force the unit cell to have
pseudocubic symmetry.
Pure BaTiO3 particles, with a size of 15nm to 230nm have been achieved by these
methods. BaTiO3 is an example of a multiple oxide synthesized by sol-emulsion-gel
method [Dibyendu, 2005]. Particles below a critical size transformed from tetragonal to
pseudo-cubic structure [Pithan et al., 2006; Huang et al., 2006; Zhu et al., 1997; Dutta et
al., 1994; Clark et al., 1999; Wada et al., 2008]. Reduction in particle size to nanometer
can be achieved by enhancing heterogeneous nucleation and by restricting crystal growth
during particle formation [Pithan et al., 2005, 2006]. These elementary processes can be
controlled during the hydrolysis reaction by utilization of water-in-oil (w/o)
microemulsion as reaction media. Microemulsion is a thermodynamically stable
dispersion of a polar and a non-polar solvent in an optically isotropic transparent liquid.
Typically water or an aqueous solution and a liquid hydrocarbon is used as a polar and
non-polar solvent respectively [Pithan et al., 2006; Ahmad et al., 2005; Tokeer and
Ashok, 2004; Ashok et al., 2005]. Addition of a surface-active agent (surfactant) having
amphiphilic character consisting of a hydrophilic head group and one or two
45
lipophilic/hydrophorbic (C-H-chain) tail group adsorb at the interface of the two liquids
that are immiscible in normal condition stabilized as 1-100nm size droplets. In the w/o
microemulsions with spherical nano-sized aqueous micelles, dispersed in an oil matrix,
the aqueous droplet can be used as nano-reactors and templates for the preparation of
solid nano-particles [Pithan et al., 2005, 2006; Tokeer and Ashok, 2004].
Of the various synthesis processes described, Table 2.3 summarises the parameters
studied using the two most used synthesis methods, i.e. solid state method and the
hydrothermal method.
2.7 Role of Surfactant in the Emulsion
Surfactants are amphiphilic compounds having hydrophilic head and a hydrophobic tail,
widely used for their surface activity in lowering the surface/interface tension [Manisha
et al., 2009; Milan et al., 2006; Schramm, 2000; Salager, 2002]. They are also known as
wetting agents and foam formers. The term surfactant which is a contraction of surface-
active-agents was devised in the year 1950 [Manisha et al., 2009; Schramm, 2000].
Based on the nature of the polar head group, surfactants are classified as anionic
(negative charge), nonionic (no-charge), cationic (positive charge), and
zwitterionic/amphoteric (both positive and negative charge) [Schramm, 2000; Manisha et
al., 2009]. Nonionic surfactants are effective over the pH range of 3-10. They do not
produce ions in the aqueous solution, are less sensitive to electrolytes than the ionic
surfactants, excellent grease/oil removers and emulsifiers, and can be used with high
salinity and hard water [Manisha et al., 2009; Salager, 2002]. Surfactants as emulsifiers
play a key role in the emulsion preparation. However, selection of an appropriate
emulsifying agent from the hundreds available proves to be a time consuming task. Close
examination of parameters such as HLB (Hydrophile-Lipophile Balance), CMC (Critical
Micelle Concentration), and CPP (Critical Packing Parameter) makes this selection
easier. HLB value assigned to an emulsifier is the molecular balance of hydrophilic and
46
Table 2.3. Highly used synthesis methods of preparing BaTiO3 and the properties
achieved.
Synthesis
Methods Product Parameters Reported
Solid State
Reaction Method
BaTiO3
[Chie et al., 2006]
Tetragonality (c/a) = 1.0085
Particle size = 120nm
Particle shape - spherical particles
Mg doped BaTiO3
[Li et al., 2007]
Resistivity range 1010 – 1012 !cm
Dielectric Constant range 1200 - 2800
F doped BaTiO3
[Darko et al., 2003] Resistivity range 101 – >109 !cm
BaZr0.1Ti0.9O3
[Mahajan et al., 2009]
Particle size = 13"m
Dielectric Constant range 5000 - 25000
Ba1-xLaxTiO3
[Aparna et al., 2001]
Tetragonality = 0.999 - 1.002
Resistivity range 1.2x1011 – 2.66x1012 !cm
Dielectric Constant at TC range 1176-3325
Hydrothermal
Method
BaTiO3
[Sasirekha et al.,2008]
Particle size = 80 - 90nm
Particle shape - spherical particles
BaxSr1-xTiO3
[Miao et al., 2007]
Particle size = 20 -40nm
Dielectric Constant at 5.7°C range 1920-6000
BaTiO3
[Xinhua et al., 2005]
Particle size = 65nm
Particle shape - spherical particles
Mg doped BaTiO3
[Dong et al., 2007]
Particle size = 60 - 125nm
Particle shape - spherical particles
Dielectric Constant at TC range 7500-10000
BaTiO3
[Huarui et al., 2003]
Particle size = 70nm
Dielectric Constant at TC > 10000
47
the lipophilic/hydrophobic groups of the emulsifier [ICI, 1980; Inderjit, 2004]. An
emulsifier with HLB value below 9 is lipophilic (oil soluble), above 11 is hydrophilic
(water-soluble), and in-between 9-11 is intermediate in nature [ICI, 1980]. HLB value is
also indicative of the type of emulsion, i.e. water-in-oil (w/o) in-between 4-6, or oil-in-
water (o/w) between 8-18 emulsion and act as a solubilizer in-between 10-18 for certain
oils [ICI, 1980].
The ability of the surfactants to self-aggregate into micelles in solvents has made these
compounds attractive for study in the surface and colloidal chemistry [Sujit et al., 2003].
These micelles are characterized by CMC (concentration at which micelles are formed),
their aggregation number (n), and their size [Sujit et al., 2003]. Near CMC, the micelle
takes the spherical shape [Lok et al., 2011; Chatterjee et al., 2000;]. However, for
surfactant concentration far above CMC the micelle shape changes from spherical to rod-
like to uneven stumpy forms [Bourel and Schecter, 1998; Dickinson, 1994; Zhang et al.,
2003; Shrestha et al., 2011]. Non-spherical micelle structures have relatively small
surface contact of the hydrophobic tail with the water phase and are more stable [Jian et
al., 2005]. Those micelles formed above CMC in the w/o type emulsion are known as
“reverse micelles” [Lok et al., 2011; Chatterjee et al., 2000; Minati and Amitava, 2001].
Reverse micelles having smaller aggregation number exhibit spherical geometry near
CMC, wherein, the polar head groups form the micellar core and the non-polar tail
groups are dispersed in the oil phase [Lok et al., 2011]. Hence, CMC plays an important
role in determining the water (aqueous sol) droplet size and consequently the final
particle size. Detailed studies have shown that CMC value depends on the nature of the
oil phase (organic solvent) [Chatterjee et al., 2000;]. Particle size and distribution control,
and enhanced reaction rates are the important advantages of micellar routes for synthesis
over the conventional methods [Upendra et al., 1996].
CPP determines the structure of these aggregates in the aqueous media which in turn is
derived from the simple geometrical considerations [Sujit et al., 2003]. The relation
between CPP (N) calculated from equation (2.5) and the geometry of self-assembly is
48
shown in Table 2.4 [Hamley, 2000].
G !$ _?9@` --- (2.5)
where,$2 is the volume of the hydrocarbon chain (assumed as fluid), ab the optimal head
group area, and Cc the critical chain length (maximum effective length that the chain can
assume) [Sujit et al., 2003]. Length of the hydrophobic tail of a surfactant plays an
important role in the micelle formation. Higher length of the hydrophobic tail lowers the
CMC value of the surfactant [Inderjit, 2004].
Surfactant concentrations much below CMC results in relatively large droplets in an
emulsion which significantly reduces in size when the surfactant concentration just
reaches or exceeds CMC [Bourel and Schecter, 1998; Dickinson, 1994]. In sol-gel
emulsion system (w/o) this phenomenon offers a scope to tailor the shape and size of the
gel particles [Zhang et al., 2003]. Size of the water pools varies with the aggregation
number of the surfactant molecules in the micelle and the relative quantity of the water
phase [Bourel and Schecter, 1998; Dickinson, 1994]. The micelles in w/o emulsions are
known to have much smaller aggregation numbers than those in o/w types [Dickinson,
1994]. For surfactant concentrations exceeding CMC, the shape changes have been
reported to vary from spherical to rod-like to uneven stumpy forms [Bourel and Schecter,
1998; Dickinson, 1994; Shrestha et al., 2011; Paul and Mitra, 2005]. Span 80 or sorbitan
monooleate (C24H44O6) and span 20 or sorbitan monolaurate (C18H34O6) were used as the
non-ionic surfactants in the emulsion during synthesis to enhance the emulsion formation.
The Hydrophilic-Lipophilic-Balance (HLB) that gives the total hydrophilic content in the
surfactant molecule of these fatty acids are 4.3 and 8.6 for span 80 and span 20
respectively [Schramm, 2000] . This indicates that span 80 is more soluble in organic
solvents than span 20 [Smitha et al., 2008; Inderjit, 2004]. The reverse micelles act as
templates to the sol droplets which are trapped within resulting into gel droplets after
addition of a gelling agent triethylamine [Paul and Mitra, 2005; Natarajan et al., 1996;
Lok et al., 2011]. The advantage of using reverse micelles over normal micelle is that the
49
Tab
le 2
.4. S
urfa
ctan
t pa
ckin
g pa
ram
eter
s an
d ge
omet
ry o
f se
lf-a
ssem
blie
s in
wat
er [
Ham
ley,
200
0; Z
hang
et
al.,
2003
].
50
size of its aqueous micellar core is comparatively smaller with smaller aggregation
number and it enhances the reaction rates [Paul and Mitra, 2005; Beck et al., 2001].
2.8 Current Research on BaTiO3 in Last Five Years
Research in BaTiO3 materials is focussed on its properties which were explored via
various means to be compatible with its applications concerning current technology.
Among those, synthesis methods employed to produce BaTiO3 in nanodimension, low
temperature heat treatment to reduce the cost of synthesis, doping with different periodic
elements and also with known compounds to modify its properties inorder to achieve
desired result are most exploited [Gorelov et al., 2011; Wei et al., 2010; Meng-Fang et
al., 2011; Andrea and Giuseppe, 2011; Yanan et al., 2011; Katsuki and Komarneni, 2011;
Radonji et al., 2008; Wei et al., 2008; Yanxia et al., 2007; Niimi et al., 2007].
Improvement in the properties of BaTiO3 for its use as a dielectric in MLCC has been
reported [Andrea and Giuseppe, 2011; Kessel et al., 2010; Yanxia et al., 2007]. Desired
decrease in the dielectric layer thickness resulted in its dimensions approaching a value as
small as 0.5µm [Andrea and Giuseppe, 2011; Kessel et al., 2010]. Increase in the
dielectric permittivity to a maximum of 160,000 ~ 280,000 with low dissipation factor of
0.1 measured at 1kHz has been achieved [Kessel et al., 2010; Yanxia et al., 2007].
However, longer life of the dielectric components is essential which depends on the
intrinsic materials property [Andrea and Giuseppe, 2011]. Inorder to control the failure
and increase the life cycle of these components, the cause of its failure which depends on
both dielectric and mechanical breakdown needs to be controlled. The root cause of this
breakdown is migration of point defects like oxygen vacancies, disruptive charges and
increase in stress due to domain switching and chemical fluctuations, cracks at the edges
or corners of internal electrodes, etc. [Andrea and Giuseppe, 2011; Kessel et al., 2010;].
Also, the decrease in ferroelectricity with decreasing particle size needs to be taken care
of. Controlling these parameters turns out to be the main processing problem.
51
PTCR-effect was confirmed to be a resultant of grain boundaries and outer grain-shells of
individual grain with a width of ~450 – 550nm and ~10% of the volume fraction of the
grains [Fiorenza et al., 2009]. Ca-doping of semiconducting BaTiO3 between 1.0050
(Ba+Ca+La)/Ti 0 1.010 was noted to show better PTCR characteristics even on being
oxidized at low temperature of 800°C in air than the conventional semiconducting
BaTiO3 [Niimi et al., 2007]. A smaller resistivity value in a PTC thermistor is desired in
BaTiO3, which can be achieved with multilayer PTCR structure. This imposes a
constraint over the fabrication process as co-firing BaTiO3 with PTCR characteristics and
the internal electrods is difficult [Niimi et al., 2007].
Besides the MLCC and PTCR thermistor, other applications of BaTiO3 have also been
exploited. Ciofani et al. [2010] studied bio-application of BaTiO3 for in vitro
investigations into cytocompatibility and cell interactions of BaTiO3 nanoparticles
(BTNP) stabilized using two non-covalent functionalizations with glycol-chitosan and
doxorubicin in aqueous environment. They observed optimal cytocompatibility of glycol-
chitosan bound BTNP even at concentrations as high as 100µg/ml and noted the
efficiency of the widely used chemotherapy drug doxorubicin to have enhanced due to
complexation with BTNPs. Their study suggested that BTNPs could be applied in various
biomedical applications. However, they also observed the cytotoxic activity of Dox-
BTNP to have enhanced as compared to that of doxorubicin alone [Ciofani et al., 2010].
Application in multiple stage memory elements, magnetic valves, filtering devices,
demand the use of materials having coexistence of ferromagnetism and ferroelectricity
[Wei et al., 2010]. Single phase magnetoelectrics like BiFeO3 has its transition
temperatures above room temperature. An attempt to produce room temperature
magnetoelectric by substituting B-site cation in ferroelectric perovskite oxide BaTiO3 by
magnetic cation like Fe was successfully attempted [Wei et al., 2010]. This led to the
tetragonal polymorph to show the presence of ferroelectricity and ferromagnetism,
wherein, the same was achieved in hexagonal polymorph on doping BaTiO3 with a
transition metal dopant [Xu et al., 2009; Ray et al., 2008; Lin and Shi, 2009; Du et al.,
52
2010]. Perovskite ceramics though have high dielectric constant, are brittle and possess
low dielectric strength. Polymers on the other hand are flexible, easy to process with low
processing temperature and possess high dielectric breakdown field. Hence, combining
the advantages of the two was of interest to synthesize a composite dielectric which was
flexible, easy to process with high dielectric constant and high breakdown strength
[Meng-Fang et al., 2011]. In order to achieve this, Meng-Fang et. al [2011] synthesized
Nd doped BaTiO3 via hydrothermal route without use of surfactant or high temperature
sintering. They carried out synthesis at 200°C and dried the sample at 100°C for 12hrs to
produce hollow BaTiO3 nanoparticles in single crystalline form with particle size around
75nm. These hollow spheres were then used to fabricate nanocomposites with
poly(vinyldene fluoride) (PVDF) polymer and have shown to demonstrate high
performance capacitors. These nanocomposites have been proposed to be used in fields
such as energy storage, ceramic capacitors and catalysis. The particles size was reported
to be dependent on the Nd concentration and synthesis time, with longer duration
producing larger particles. Dielectric constant of 480.3 with a dielectric loss of 0.6 was
obtained at 100Hz frequency.
ABE et. al [2010] synthesized Ba0.9Sr0.1TiO3 thin-films using electrophoretic deposition
(EPD) owing to its high dielectric constant . This is an inexpensive route as no vacuum is
employed during film making and it can be made in to a continuous process in
comparison to the other known methods such as sputtering, chemical vapor deposition
and chemical solution deposition. Using the EPD technique, 20nm size BaTiO3 powders
synthesized via alkoxy method were used to develop a 0.54µm thick film with
capacitance density of 73.9nF/cm2 and a loss tangent of 0.021 having leakage current
density of 0.3µm/cm2 at 10V satisfying the requirements for an embedded capacitor for
printed circuit board (PCB). Shen et. al [2008] carried out a comparative study on the
resistive switching behavior of BST thin films using tungsten (W) and platinum (Pt) top
electrodes. They observed improvement in yield, endurance and reliability with W top
electrode. However, in Pt top electrode fast drop in resistivity was noted. With W top
electrode the device could be switched 104 times without degradation due to the
53
reversible oxidation and reduction in WOx layer at the W-BST interface [Shen et al.,
2008]. The effect of water used during sol-gel synthesis of BaTiO3 thin films was studied
and found to influence the gel structure evolution and thin film crystallization [Radonji
et al., 2008].
54
Chapter 3 Materials and Method
Polycrystalline pure and doped barium titanate with general formula Ba1-xDxTiO3 were
synthesized through sol-gel emulsion technique, where D = Sr, La, Ce, Mg, Li and K and
x = 0.001, 0.01, 0.05 and 0.1. The materials synthesis route and thermal sample
characterization were as follows:
3.1 Precursors
Ba(CH3COO)2.H2O (Sisco Research Laboratory Pvt. Ltd., India/ Loba Chemie Pvt. Ltd.,
India), Ti-isopropoxide (Merck, Germany), Sr(NO3)2 (Merck, India), La(NO3)3·6(H2O)
(Merck, Germany), CeN3O9·6(H2O) (S.D. Fine-Chem Limited, India), KNO3 (Merck,
India), LiNO3 (S.D. Fine-Chem Limited, India) and Mg(NO3)2·6(H2O) (Merck, India)
were used as the initial raw materials for synthesizing pure and doped barium titanate
powders. Acetic acid (Merck, India) and acetylacetone (Merck, India) were used as
chelating agents in synthesizing the sol. Emulsion was prepared by dispersing the sol into
support solvent. The support solvent comprised of cyclohexane and a surfactant. Two
surfactants namely span-80 (Loba Chemie Pvt. Ltd., India) and span-20 (S.D. Fine-Chem
Limited, India) were used to stabilize the emulsion.
Sol-gel process is a wet chemical route to produce inorganic and hybrid materials at low
temperature [Chen et al., 2012; Pfaff, 1992; Moreno et al., 1995; Kuo and Ling, 1994;
Kareiva et al., 1999; Vijatovi et al., 2008]. A sol comprise of a colloidal suspension of
solid particles in a liquid [Brinker and Scherer, 1990]. Sols are generally prepared using
metal alkoxide which reacts vigorously with water and precipitate the metal oxide
[Brinker and Scherer, 1990; Arnout and David, 1998; Imhof and Pine, 1997]. To avoid
this, the metal alkoxide is introduced in an appropriate solvent mixture to achieve a clear
homogeneous sol [Despina et al., 1994]. Emulsion however is a heterogeneous system
55
that consists of two immiscible phases prepared by dispersing one liquid in the form of
droplets into another [Brinker and Scherer, 1990; Rosaria et al.; Chen et al., 2012; Arnout
and David, 1998]. Sol-emulsion-gel is a method in which gelling of emulsified sol
droplets takes place. It comprise of two techniques; (1) emulsion-water extraction and (2)
emulsion-proton extraction [Ganguli 2003]. Both these techniques only differ in the
mechanism of forming gel microspheres from the emulsified sol droplets. The present
work employs emulsion-proton extraction technique also known as neutralization
[Ganguli 2003]. Here the H+ ions (protons) in the aqueous sol droplets are extracted with
an organic amine (catalyst) to increase the basicity of the system leading to the formation
of gel microspheres [Ganguli 2003, Vinothan et al., 2001]. Emulsion droplets behave as a
microreactor for the hydrolysis and condensation reactions of the precursors or materials
[Rosaria et al.; Chen et al., 2012]. These droplets are stabilized by surfactants that adsorb
onto the interfaces between the immiscible phases [Chen et al., 2012; Rosaria et al.;
Arnout and David, 1998; Vinothan et al., 2001]. Stirring of the w/o emulsion leads to the
formation of sol droplets inside the emulsion drop protected by the adsorbed surfactant at
the interface of the liquids [Rosaria et al.; Arnout and David, 1998]. Gelled droplets are
separated by evaporation or dissolution in a suitable liquid. This liquid removes oil and
some surfactant and when heat treated removes the solvent and residual organic materials
[Rosaria et al.; Arnout and David, 1998].
3.2 Experimental Procedure
Firstly, titania sol was prepared by dropwise addion of titanium alkoxide to a solution of
acetic acid and acetylacetone under continuous stirring. The molar ratio of acetic acid to
titanium alkoxide was 2.4 and acetylacetone to titanium alkoxide was 0.3. A 1.5M
barium acetate solution in distilled water was added to the titania sol drop wise under
continuous stirring at ambient conditions, maintaining a Ba/Ti atomic ratio of 1:1 to
obtain a clear translucent sol with no traces of precipitate formation.
56
The prepared sol (water medium) was dispersed under continuous stirring in a mixture
called support solvent (oil medium) to form w/o type emulsion. Support solvent consisted
of an organic liquid; cyclohexane and a surfactant; span 80 or span 20. Two different
sol:support solvent ratio, namely, 1:2 and 1:3 was studied for pure barium titanate with
two different types of surfactant (Span 80 and 20). The surfactant concentration during
the synthesis of pure barium titanate was changed as 5 vol%, 10 vol%, 15 vol%, and 20
vol% of the cyclohexane. Triethylamine (Rankem, India) was added to the emulsion as a
gelling agent to transform the sol droplets to gel droplets, until a pH value of 9.0 was
reached. The prepared gel droplets were washed with methanol (Rankem, India), filtered,
then dried at 100oC, and then calcined to remove the volatile components and obtain
crystalline barium titanate. The dried powders were heat treated at 400°C, 500°C, 600°C,
700°C, 750°C, 800°C, 900°C and 1000°C. Calcination cycle at 400°C, 500°C, 600°C and
700°C temperatures was carried out at a heating rate of 4°C/min with two hours soaking
time at the specified temperature. For 750°C, 800°C and 900°C the calcination cycle was
carried out at a heating rate of 4°C/min with one hour soaking time each at 500°C and at
the specified temperature. For 1000°C calcination cycle was carried out at a heating rate
of 4°C/min with one hour soaking time at 500°C, half-an-hour soaking time at 900°C and
again one hour soaking time at 1000°C. Pellets prepared from these powders were
sintered using the same cycle at 750°C. The results obtained were almost same for the
two ratios. Hence, the emulsion composition used in the synthesis of doped barium
titanate was 1:3 as the sol: support solvent ratio with 5 vol% surfactant concentration.
Also from the detailed study of the effect of calcinations temperature carried out for the
pure barium titanate, heat treatment for the calcinations and sintering was fixed at 750oC.
When synthesizing doped BaTiO3, the dopant (D) in the form of nitrate salt was mixed
with Ba-acetate solution, the concentration of which was calculated using the formula
Ba1-xDxTiO3. The dopants Sr, La, Ce, Mg, Li and K were used in the present work.
The calcined powders were analyzed using an X-ray diffractometer (Rigaku, Miniflex II),
Thermogravimetric Analyzer (DTG-60, Shimadzu), Differential Scanning Calorimetry
57
Fig 3.1. Flow diagram for synthesis of nano size BaTiO3 powders via sol-gel emulsion
technique.
58
(DSC-60, Shimadzu), Transmission Electron Microscopy (TEM, Phylips CM200),
Scanning Electron Microscopy (SEM, JEOL JSM-6500F), SEM (JEOL 6360LVSEM),
Energy Dispersive Spectroscopy (EDS, INCA 6360LVSEM), Two-probe resistivity setup
(Scientific Equipment and Services, Roorkee), and Impedance Spectroscopy (Wayne
Kerr Precision Component Analyzer 6440 B). Test parameters employed for the
characterizations have been discussed in the respective sections.
3.3 Thermal Analysis (TA)
Thermal analysis consists of techniques used to monitor the changes in physical or
chemical properties of the material under consideration, with temperature variation
[Earnest, 1998; Grega et al., 2010]. Various methods are used in determining the thermo-
physical properties. Of these the ones employed in the current study include;
thermogravimetry analysis (TGA), differential thermal analysis (DTA) and differential
scanning calorimetry (DSC) [Grega et al., 2010].
3.3.1 Thermogravimetry Analysis (TGA) and Differential Thermal Analysis (DTA)
Thermogravimetry analysis is a technique used to study the weight change of the material
due to evaporation, decomposition and phase change with changing temperature and/ or
time under controlled atmosphere [Lei, 2006; Dmitry, 2008]. Hence the requirement of a
stable and precise weighing balance becomes very important. The sample is placed on a
refractory pan that is positioned in the furnace. The refractory pan is suspended from a
high precision balance such that any change in the sample weight disturbs the balance
equilibrium producing a proportional response to restore the equilibrium by electronic
compensation so as to avoid the motion of the pan when the sample weight changes.
Roberval type balance operates on the unique mechanism that allows high precision
measurements and prevents sensitivity changes from factors like thermal expansion.
59
Differential thermal analysis is a technique in which both physical and chemical
processes occur in relation to the thermal changes. Here the endothermic and the
exothermic changes in the sample can be detected. Thermal analysis is conducted
between the sample and the reference as a function of reference temperature or time,
when both are heated in one furnace (heat-flux type) [Grega et al., 2010; Dmitry, 2008].
In the current study, DTG-60 SHIMADZU, which is a simultaneous differential
thermogravimetric analyzer, was used to measure the weight change alluring the
formation of barium titanate. The DTG-60 combines a heat-flux type heating chamber
DTA with a Roberval type TGA and hence measures the weight change and the changes
in the physical state of a sample as a function of temperature over time. The
measurements were conducted during the heating cycle with a heating rate of 10°C/min
in air with platinum pan to avoid reactions between the pan and the sample.
3.3.2 Differential Scanning Calorimetry (DSC)
Differential scanning calorimetry is a technique that measures the change in the heat flow
rate difference to the sample and the reference when subjected to a controlled
temperature change. Thermal changes that do not comprise of the mass changes are
represented in this technique. Similar to a DTA, DSC helps in determining the
temperature of the phase transitions like melting point, onset of oxidation and heat
capacities, evaporation temperature, crystallization, etc. DSC profile gives the behavior
of heat flux with respect to time or temperature and hence, can be used to calculate the
enthalpy, specific heat, etc and is known to be more sensitive than a DTA. Also, the
necessity of maintaining the sample and the reference under identical temperature regime
is overruled in DSC. The two basic type of DSC are; (1) heat-flux (HF) DSC in which the
sample and the reference are placed in the same furnace, also known as a type of
Boresma DTA and (2) power compensation (PC) DSC where the sample and the
reference are placed in two different furnaces each having an individual heater. In DSC
analysis a constant mass during the measurement of enthalpy change is desired. The DSC
60
analysis were carried out using DSC-60 SHIMADZU at a heating rate of 5°C/min in air
with the sample placed in the aluminum pan.
3.4 X-Ray Diffraction (XRD)
X-Ray diffraction is a powerful characterization technique used for detailed structural
analysis. The atomic arrangement of the structure can be deduced from the XRD patterns
as the diffraction plane spacing (d) is of the order of X-Ray wavelength (e), where the
various orders (f) of reflection occur at precise values of diffraction angle (g) satisfying
the Bragg equation given by [Scott; Willard et al., 1986; Goodhew et al., 2001; Cullity,
1956];
f$e ! ,$dhi@$jTfg$."?kk ---(3.1)
Advantages associated with this technique are, (1) the XRD pattern acts as a blueprint of
the substance under consideration, (2) each component in a mixture produces its
characteristic pattern irrespective of the other component, and (3) both qualitative (e.g.
compound, structure, etc.) and quantitative (e.g. crystal size, d-spacing, lattice parameter,
etc.) analysis of the substance can be achieved.
XRD characterization in the current work was carried out using Rigaku Miniflex II to
confirm the formation of barium titanate, identify the phase formed, determine the
crystallite size and lattice parameters. The measurement was carried out using a Cu-K' (3
= 1.5405Å) X-Ray beam with a 4:24 motion of the sample and the detector maintaining a
step size of 0.02° at a scan rate of 2°/min. On the basis of XRD line broadening at half
maxima of the maximum intensity diffraction peak, crystallite size (C.S.) was estimated
using the Debye-Scherer formula,
lm nm ! $ bmo$pq$c>rs !!!(3.2)
61
where, 3 is the wavelength of the target used, 5 = 6(Bm2 – Bs
2), Bm is the full width at half
maxima (FWHM) in radians of the sample analyzed, Bs is the FWHM of the standard
sample analyzed, 4 is the Bragg angle of the 100% intensity peak.
Intensities of the X-Ray diffraction peaks were used to calculate % tetragonality from the
X-Ray patterns for the synthesized material as described by Takeuchi et al. [1997]. In the
process of calculating the same, the tetragonal peaks were first identified. For example,
(200) and (002) correspond to the tetragonal peaks. The sum of intensities of such
tetragonal peaks when divided by the sum of intensities of all peaks present in BaTiO3
XRD pattern excluding those corresponding to BaCO3 on being multiplied by 100 gives
the % tetragonality in that sample.
Lattice parameters were calculated as specified by B. D. Cullity [1956] for using the
following equations. For n=1, eqn. 3.1 becomes:
e ! ,$d$jTfg !!! (3.3)
Stu !$ v$rOwus
pu !!!(3.4)
Relation between d and the lattice parameters is given as
Stu !$ hu$$
?u 1 iuxu 1 @u
cu !!!(3.5)
where, (hkl) are the Miller indices corresponding to lattice planes and a, b, c are the axial
lengths also known as the lattice parameters.
For a cubic structure, a = b = c and eqn. 3.5 becomes:
! " #$%&'(&')& !!!(3.6)
Substituting eqn. 3.6 in eqn. 3.4 gives:
62
* +,-& ."/& !" %&'(&')&#&
The lattice parameter (a) can be calculated from the diffraction pattern obtained in the
present work using the above formula. As the systematic error in a was observed to
decrease with increasing value, hence, the value of a corresponding to the highest
value was chosen. For tetragonal structure, a = b ! c. To calculate c, those (hkl) values
for which h = k = 0 and l ! 0 have been used.
3.5 Fourier Transform Infrared (FTIR) Spectroscopy
Fourier transform infrared spectroscopy is a technique by which the chemical analysis of
a material can be made as it studies the interaction of infrared light with matter. Above
absolute zero, all universal objects radiate infrared radiation which can be absorbed by
any matter on illumination. This absorption of the infrared radiation causes the chemical
bond in the absorbing material to vibrate. Chemical structural fragments known as the
functional groups (e.g. C-O group, OH group, C-H group, etc.) within molecules
regardless of the molecular structure absorb infrared radiation in the same wave number
range which helps in identifying the molecules [Brian, 2011; Jag, 2004]. The energy
associated with infrared radiation is small and only induces transition between the
vibrational and the rotational energy levels of a molecule. Absorption of infrared
radiation causes bond deformation by either stretching or bending of the functional group
at quantized frequencies.
In the current study FTIR spectrophotometer FTIR- 8900, SHIMADZU was employed in
the transmittance mode from 300cm-1 – 4000cm
-1. The final scan was recorded by
averaging 100 scans at a resolution of 8.0 to obtain the FTIR profiles.
63
3.6 Electron Microscopy
Electron microscopy gives us the general information about the shape and size of the
particles as well as the surface microstructure, defects etc of a dense body. Two most
widely used electron microscopy techniques are transmission electron microscopy (TEM)
and scanning electron microscopy (SEM).
3.6.1 Transmission Electron Microscopy (TEM)
TEM is a technique in which electrons generated by an electron source penetrate a thin
specimen (0.5"m or less) that are then imaged by specific lenses on a fluorescent screen
or in the modern instruments by charge-coupled device (CCD) camera [Egerton, 2005;
FEI, 2006, 2010]. In early TEMs gas discharge was used as the source of electrons which
was then replaced by a V-shaped filament made from tungsten wire which when heated
in vacuum emits electrons [Egerton, 2005]. The emitted electrons are accelerated by
applying high voltage generated by electromagnetic lenses [Egerton, 2005; FEI, 2010].
Higher energy electrons resulting from higher accelerating voltage in the gun can
penetrate comparatively thicher specimens which could otherwise not be imaged as the
electrons would be stopped [FEI, 2006, 2010]. The entire electron path from gun to
camera is maintained under vacuum so as to avoid scattering and absorption of electrons
by the air molecules [FEI, 2010]. TEM has been invaluable as it was capable of giving
information about the structure of materials such as dislocations that validated some
proposed theories. However, one limitation of producing thin specimen for TEM analysis
acts as a constraint [Egerton, 2005]. This resulted in the study of developing electron
microscopes capable of examining relatively thick specimens.
3.6.2 Scanning Electron Microscopy (SEM), Energy Dispersive Spectroscopy (EDS)
SEM is a technique in which the sample surface is scanned and analyzed when the
electrons generated with an electron gun are irradiated over the sample [Egerton, 2005;
64
Rudawska, 2012; Rahul, 2012]. The SEM operates in two modes, namely; the secondary
electron mode and the backscattered electron mode [Reimer, 1985; Egerton, 2005]. In the
secondary mode, operation takes place when the high energy primary (i.e. incident)
electron undergoes small angle scattering (< 90°) which after irradiation re-emerges from
the sample surface. However, in the backscattered mode, the high energy primary
electron undergoes multiple elastic large angle scattering (> 90°) within the specimen
before re-emerging from the sample surface after irradiation. The energy form the
secondary electron is less than that of the backscattered electron which has energy close
to that of the primary electron. Backscattered electrons are rich in surface sensitive
information due to which is highly important to the SEM users [Reimer, 1985].
Energy dispersive spectroscopy is a semi-qualitative technique used in identifying and
quantifying the elemental composition of a sample. On irradiating the sample with
electron beam in techniques like SEM, the atom in the sample radiates characteristic X-
rays [Joy et al., 1986; Reimer, 1985; Egerton, 2005]. These X-rays can be detected by the
detector where electron-hole pair is created using an energy dispersive spectrometer. This
change having energy proportional to the incoming X-ray is amplified and correlated to
the atoms element according to its energy. However, this technique has a relatively poor
energy resolution and is unable to detect lighter elements below sodium (Z = 11).
SEM analysis was carried out using two different instruments; JEOL JSM-6500F and
JEOL 6360L VSEM. The EDS was carried out using INCA 6360L VSEM. The sample
pellets were platinum coated under vacuum before the measurements were carried out.
3.7 Two-Probe Resistivity Analyzer
Two-probe resistivity measurement technique is used to study the electrical resistivity of
highly resistive materials (near insulators) and its behavior as a function of temperature.
Analysis about the materials structure and its application in various fields of technology
65
can be achieved with the information obtained. Two-probe resistance is also known as
spreading resistance and is determined through the depth variation of the resistivity on
application of a voltage between the two probe tips that provide the electrical contact to
the sample pellet [Kopanski et al., 1990]. The measured material resistance is
proportional to its resistivity as
0 !" 1"23 ---(3.7)
where, R is the resistance of the material, # is the resistivity of the material,""4 is the
thickness of the pellet, and A the area of the material.
The measurement of two-probe resistivity was carried out using a set up provided by
Scientific Equipment and Services. The pellets to be analyzed were silver coated after
sintering at 750°C. The silver coated pellet was placed in a sample holder enclosed in an
oven designed for the measurement using two probes. Pellet current was recorded with
the change in temperature at a fixed voltage. Resistance of the pellet was deduced using
Ohms law, and the resistivity thus calculated.
3.8 Impedance Spectroscopy (IS) / Dielectric Analyzer (DEA)
The electrical parameter used in characterizing electronic circuits, components and
materials is its impedance [Barsoukov and Macdonald, 2005; Lei, 2006]. The basic
elements comprise of inductance (L), capacitance (C) and resistance (R). By employing
LCR meters and measuring the current flow through the material under study, all
impedance parameters can be calculated. Thus, impedance spectroscopy is a technique in
which the dielectric properties as a function of temperature, frequency and time are
measured. The measured dielectric constant (property) of a material is directly
proportional to the capacitance of the material. The dielectric properties give an insight to
the capacitive and conductive properties of the insulating material. The capacitive and
conductive properties determine the materials capability to store electric charge and the
66
motion of these charges within the material respectively. Phase transitions can also be
identified using this technique.
The dielectric measurement in the current study was conducted using a Wayne Kerr
Precision Component Analyzer 6440 B. The capacitance of the equivalent parallel circuit
was recorded in two modes of operation, a constant frequency mode with varying
temperature and a constant temperature mode with varying frequency. The sintered
pellets were coated using silver paste in a manner similar to that done for resistivity
measurements before inserting into the sample holder for the measurements. The
dielectric constant of the material is related to its capacitance as
56 !" 7"289"3 !!!(3.8)
where, :6 is the dielectric constant, C the capacitance of the material, "4 the thickness of
the pellet, :; the permittivity of free space and A the area of the material.
67
Chapter 4 Results and Discussions
Outline
This chapter describes the results obtained from the study on pure and doped BaTiO3
with varying dopant concentration synthesized using span 80 and span 20 as the
surfactants. Six dopants namely strontium (Sr), lanthanum (La), cerium (Ce), magnesium
(Mg), lithium (Li) and Potassium (K) were used. They have been discussed for their
effect on the structural and electrical properties of the powders synthesized with respect
to pure BaTiO3. Doped barium titanate were synthesized using the formula Ba1-xDxTiO3,
where D is the dopant ion and the concentration x is varied as 0.001 (0.1 ± 0.03 at.%),
0.01 (1 ± 0.01 at.%), 0.05 (5 ± 0.02 at.%) and 0.1(10 ± 0.04 at.%).
4.1 TGA, DTA and DSC Analysis
Figure 4.1a shows the TGA and the DTA curve of pure BaTiO3 powder synthesized
using 5% span 80. TGA curves for synthesized powders showed weight loss in three
phases. The first phase was noted, between 35oC and 200
oC, the second phase from
200oC to 450
oC and the third phase from 500
oC to 900
oC. Weight loss in the first phase is
caused by the evolution of water and other volatiles, in the second phase due to the
combined effect of removal of organic components, hydroxyl group along with
decomposition of acetate and the final weight loss is attributed to the decomposition of
residual organics, intermediate phase formation like BaCO3 and their reaction to form
BaTiO3 releasing CO2 [Durán et al., 2001; Zhang et al., 2006; Kareiva et al., 1999;
Newalkar et al., 2001; Asiaie et al., 1996]. As seen from Figure 4.1a, first phase had a
60% weight loss followed by a 22% weight loss in the second phase and 3% weight loss
in the third phase. This weight loss process is endothermic in nature in the initial phase
when the sample starts absorbing heat to release water and other volatiles, and in the third
68
phase when it starts crystallizing into BaTiO3 from the intermediate phases like BaCO3
[Mao et al., 2007; Miao et al., 2007; Zhou et al., 2009]. In the range of 50°C to 390°C,
the process is exothermic involving combustion of organic components leading to the
formation of BaTiO3 [Mao et al., 2007; Miao et al., 2007]. Similar trend was observed
for powders using span 20 as the surfactant, however, the weight loss in the
corresponding temperature regions as that noted in Figure 4.1a was 40%, 21% and 4%
respectively.
Figure 4.1b shows the DSC curve of the powder synthesized using 5% span 80, calcined
at 750 o
C. No endothermic peak characteristic of ferroelectric BaTiO3 was observed near
the transition temperature for the synthesized powders. DSC technique being very
sensitive to the phase change shows an endothermic peak at the tetragonal to cubic
transition temperature at 130°C of BaTiO3 [Takeuchi et al., 1997; Huarui and Lian,
2003]. The relation between the enthalpy of transition ($H) and polarization (P) of the
material is [Asiaie et al., 1996];
<= ! " >?@&AB7 --- (4.1)
where, C is the Curie-Weiss constant.
This shows that an increase in polarization achieved with higher tetragonality content will
increase $H, which will be exhibited as an endothermic peak in the DSC profile [Huarui
and Lian, 2003; Yasukawa et al., 2007]. As reported by Takeuchi et al. [1997] higher
sintering temperature resulted in large particles, increasing the tetragonality in the
crystals. For powders showing tetragonality below 60% no endothermic peak was
observed in the DSC results and only at 64% a weak diffused peak was achieved.
However, above 80% tetragonality a very distinct endothermic peak was obtained, with a
transition enthalpy ($H) value of 200J/mol for 95% tetragonality [Takeuchi et al., 1997].
Xu et al. also reported the DSC peak at ~130°C for particles with a minimum of 80%
tetragonality having $H value of 0.7-0.75J/g [Huarui and Lian, 2003]. A shift of Tc
towards a higher value with increasing tetragonality in the samples was also noted
69
[Huarui and Lian, 2003]. Since no endothermic peak was observed near the transition
temperature for the synthesized powders, it is clear that the percentage tetragonality in the
synthesized powders is below 64%.
Figure 4.2 shows the TGA-DTA curve for doped BaTiO3 powders. Major features
obtained for doped BaTiO3 were similar to that observed in pure BaTiO3 seen in Figure
4.1a. The weight loss is observed in three phases for all dopants and is attributed to the
same phase changes as noted in pure BaTiO3. For Ba0.9Sr0.1TiO3, the first phase of weight
loss was noted between 35oC and 240
oC, the second phase from 240
oC to 550
oC and the
third phase from 550oC to 900
oC. The small shift in the temperature range of the phases
could be attributed to the presence of nitrate ions due to the use of the dopants in its
nitrate form. First phase reported a weight loss of 40%, the second phase of 27% and the
third phase a 4% weight loss. For Ba0.9La0.1TiO3, the three phases were noted between
35oC to 240
oC, from 240
oC to 575
oC and from 575
oC to 900
oC. First phase had a 30%
weight loss followed by a 45% weight loss in the second phase and 10% weight loss in
the third phase. For Ba0.9Ce0.1TiO3, the first phase with a 45% weight loss was noted
between 35oC and 150
oC, the second phase from 150
oC to 350
oC had a 25% weight loss
and the third phase from 350oC to 900
oC showed a 12% weight loss. For Ba0.9Mg0.1TiO3,
the first phase with a 30% weight loss was noted from 35oC to 240
oC, the second phase
from 240oC to 600
oC had a 35% weight loss and the third phase weight loss was 8% from
600oC to 900
oC. For Ba0.9Li0.1TiO3, the first phase with a 35% weight loss was noted
between 35oC and 250
oC, the second phase from 250
oC to 520
oC had a 25% weight loss
and the third phase from 520oC to 900
oC showed a 7% weight loss. For Ba0.9K0.1TiO3 the
first phase noted between 35oC and 250
oC showed a 27% weight loss, the second phase
from 250oC to 630
oC a 30% weight loss and the third phase from 630
oC to 900
oC showed
a 10% weight loss.
From Figure 4.2 a clear difference in the DTA graph with different dopants is visible. For
Sr doped BaTiO3 two exothermic peaks in the temperature range of 250°C - 500°C is
70
Fig
4.1
. B
aTiO
3 p
ow
der
sy
nth
esiz
ed f
or
5%
sp
an 8
0 (
a)
TG
A (
wei
gh
t %
) an
d D
TA
(h
eat
flo
w)
curv
es,
(b)
DS
C c
urv
e fo
r th
e
po
wd
er c
alci
ned
at
75
0 oC
.
F
ig 4
.2.
TG
A-D
TA
curv
es o
f B
a 0.9D
0.1T
iO3 p
ow
der
sy
nth
esiz
ed u
sin
g 5
% s
urf
acta
nt
and
dri
ed a
t 1
00
°C w
ith D
= S
r (a
) an
d L
a
(b).
050
100
150
200
250
b
Endothermic Exothermic
Tem
pera
ture
(0C
)
0100
200
300
400
500
600
700
800
900
0
20
40
60
80
100
-150
-100
-50
050
Heat Flow (mV)
Weight (%)
Tem
pera
ture
(0C
)
a
0100
200
300
400
500
600
700
800
900
20
30
40
50
60
70
80
90
100
-150
-100
-50
050
Heat Flow ( V)
Endothermic Exothermic
Weight (%)
Tem
pera
ture
(0C
)
a
0100
200
300
400
500
600
700
800
900
0
20
40
60
80
100
-150
-100
-50
050
100
b
Tem
pera
ture
(0C
)
Weight (%)
Heat Flow ( V)
Endothermic Exothermic
71
F
ig 4
.2.
con
t. T
GA
-DT
A c
urv
es o
f B
a 0.9D
0.1T
iO3 p
ow
der
synth
esiz
ed u
sing 5
% s
urf
acta
nt
and
dri
ed a
t 100°C
wit
h D
= C
e (c
),
Mg
(d
), L
i (e
) &
K (
f).
0100
200
300
400
500
600
700
800
900
0
20
40
60
80
100
-150
-100
-50
050
100
150
c
Heat Flow ( V)Endothermic Exothermic
Tem
pera
ture
(0C
)
Weight (%)
0100
200
300
400
500
600
700
800
900
0
10
20
30
40
50
60
70
80
90
100
-150
-100
-50
050
d
Heat Flow ( V)Endothermic Exothermic
Tem
pera
ture
(0C
)
Weight (%)0
100
200
300
400
500
600
700
800
900
0
10
20
30
40
50
60
70
80
90
100
-150
-100
-50
050
e
Heat Flow ( V)
Endothermic Exothermic
Tem
pera
ture
(0C
)
Weight (%)
0100
200
300
400
500
600
700
800
900
0
10
20
30
40
50
60
70
80
90
100
-150
-100
-50
050
f
Heat Flow ( V)
Endothermic Exothermic
Tem
pera
ture
(0C
)Weight (%)
72
observed. The presence of exothermic peaks is attributed to the combustion of organic
components due to decomposition of the nitrate and the alkoxide precursors. A similar
behavior of separate exothermic peaks in Sr doped BaTiO3 is noted by other authors and
have been attributed to the decomposition of alkoxide radical and intermediate carbonate
phases separately [Tian et al., 2000; Xiaohua et al., 2006; Kareiva et al., 1999]. La doped
BaTiO3 also shows the presence of two exothermic peaks which again correspond to the
decomposition of alkoxide radicals and the intermediate phase. However, the peaks
observed in La doped BaTiO3 was closer than Sr doped BaTiO3. This indicates that the
decomposition of the precursor and the intermediate compounds takes place in quick
succession in La doped BaTiO3. The small endothermic peak at around 550°C noted in
La doped BaTiO3 suggests that La doping reduces the onset of crystallization of BaTiO3.
Such behavior was observed for K doped Ba0.6Sr0.4TiO3 by Xiaohua et al. [2006]. Ce
doping on the other hand showed one single exothermic peak suggesting the process of
decomposing the precursor and the intermediate phases to occur simultaneously at the
same temperature. Also, a very small endothermic kink observed at around 450°C at
which the weight loss of the sample begins to stabilize implies that the crystallization
starts at a lower temperature with Ce doping in BaTiO3. Mg doping in BaTiO3 showed a
broad exothermic peak in the temperature range of 250°C-450°C merging three separate
exothermic peaks placed very close to each other. Again a very small endothermic peak
observed at 650°C marks the onset of crystallization of the final sample. Li doped
BaTiO3 shows the presence of two separate exothermic peaks which implies that the
decomposition of precursor and the intermediate phases occur at different temperatures.
The endothermic peak due to crystallization of the final compound is also noted at 610°C.
The behavior noted in K doped BaTiO3 is similar to that of Li doped BaTiO3. It can be
noted that the dopants with same valence states, Sr and Mg (+2), La and Ce (+3) and Li
and K (+1) showed similar trend of phase changes.
73
4.2 X-Ray Analysis
Pure BaTiO3 powder synthesized using 5% span80 was heat treated at 400°C, 500°C,
600°C, 700°C, 750°C, 800°C, 900°C, and 1000°C and characterized by XRD to study
their crystalline behavior. Figure 4.3a shows the XRD patterns of these powders. Single
phase BaTiO3 with presence of barium carbonate as an intermediate phase formed during
synthesis was observed up to a calcination temperature of 700oC. Only trace amounts of
BaCO3 was present after calcination at 750°C. Crystalline peaks appeared after heat
treatment at 400oC, and well developed peaks were noted from 600
oC onwards. There
was no distinct splitting of the (200) peak up to a calcination temperature of 900oC
indicating the predominance of the cubic phase [Gorelov et al., 2011; Devaraju et al.,
2006; Xie et al., 2009; Zhu et al., 1997]. However, a shoulder around (002) was noted in
the powders calcined at and above 750oC indicating the onset of transformation from
cubic to tetragonal phase. The (200), (002) peak distinction was clearly visible for the
1000oC calcined powder as seen in the inset of Figure 4.3a. Splitting of this peak
confirmed the presence of tetragonal BaTiO3 in these powders. Figure 4.3b compares
XRD pattern of the powder calcined at 750°C to that of the pellet prepared from this
powder and sintered at 750°C. Trace amount of BaCO3 was present in 750°C calcined
powder and was absent in the pellets sintered at 750°C (Figure 4.3b.). As the resistivity
and dielectric studies were carried out with sintered pellets, 750°C was chosen as the
calcination temperature in all future studies to have a material with average particle size
<100nm (section 4.4) and pure BaTiO3 phase. Fig. 4.3c shows the XRD pattern of
synthesized BaTiO3 powders calcined at 1000°C with different surfactant. Trace amount
of BaCO3 phase was present in the powders synthesized using span 20 but was absent in
the powders synthesized using span 80. The percentage of tetragonality present in the
powders was calculated as the ratio of integrated intensities of the tetragonal peaks to the
integrated intensities of all peaks for a single X-ray pattern as described by Takeuchi et
al. [1997]. For pure BaTiO3, the % tetragonality for powders synthesized using span 80
and span20 was found to be < 50%. The percentage tetragonality calculated for the
1000°C calcined powders was below the amount that can be detected via DSC analysis
74
Fig
4.3
. X
-ray
dif
frac
tion p
atte
rn o
f B
aTiO
3 s
ynth
esiz
ed (
a)
usi
ng 5
% s
pan
80 c
alci
ned
at
(i)
400°C
, (i
i) 5
00°C
, (i
ii)
600°C
,
(iv
) 7
00
°C,
(v)
75
0°C
, (v
i) 8
00
°C,
(vii
) 9
00
°C,
and
(v
iii)
10
00
°C;
(b)
usi
ng
5%
sp
an 8
0 c
alci
ned
an
d s
inte
red
at
75
0°C
; an
d (
c)
usi
ng
5%
sp
an 8
0 a
nd
20
as
the
surf
acta
nt
calc
ined
at
1000°C
(* -
BaC
O3)
cba
20
30
40
50
60
70
80
**
Sp
an
80
Sp
an
20
Intensity (arb. units)
2-T
heta
(d
eg
ree)
20
30
40
50
60
70
80
**
**
*
Sin
tere
d
Un
sin
tere
d
Intensity (arb. units)
2-T
heta
(d
eg
ree)
20
30
40
50
60
70
80
viii vii
vi v iv iii ii i
44.7
45.0
45.3
45.6
**
**
*
(200)
(002)
(301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
Intensity (arb. units)
2-T
heta
(d
eg
ree)
75
(section 4.1). Heat treatment at temperatures higher than 800°C were reported to generate
tetragonality above 60% for particles in the size range of 70nm to several micrometer
[Takeuchi et al., 1997; Huarui and Lian, 2003; Asiaie et al., 1996; Aparna et al., 2001;
Srimala et al., 2008; Kareiva et al., 1999; Krishna et al., 1993; Gorelov et al., 2011;
Chatterjee S. et al., 2003; Frey and Payne, 1996]. Tetragonality in BaTiO3 is dependent
on the particle size, synthesis route, calcination temperature and cycle. With the
nanometer range particle (as observed from SEM discussed in section 4.4) achieved for
the synthesized powders, the smaller domain sizes may have forced the unit cells to have
pseudocubic (pc) symmetry resulting in lower percentage of tetragonal phase and higher
percentage of cubic phase [Takeuchi et al., 1997; Asiaie et al., 1996]. Crystal structure of
an ideal perovskite when distorted, due to coexistence of cubic phase with any other
phase is known to possess pc symmetry [Borah and Mohanta, 2012; Serrate et al., 2007;
Jana et al., 2008].
Average crystallite size calculated for varying surfactant concentration in the 1:3
sol:support solvent ratio is shown in Table 4.1. For 750°C calcined powders the
crystallite size varied from 27-33 nm for span 80 and 24-35 nm for span 20. Crystallite
size variation for 1000°C calcined powders was between 44-45 nm for span 80 and 43-52
nm for span 20. Since the powders prepared with 5% surfactant on an average produced
smaller crystallite size, the synthesis of all samples for the present work was carried out
using 1:3 sol: support solvent with 5% surfactant.
Figure 4.4 and 4.5 shows the X-ray pattern of Ba1-xDxTiO3 powder calcined at 750 o
C, with
varying dopant concentration synthesized using 5% span 80 and span 20 respectively. All
doped samples showed similar XRD pattern to that of undoped BaTiO3 and no secondary
phases corresponding to the dopant were noted. Trace amounts of BaCO3 was also
present in the doped samples.
Lattice parameters a, and c were calculated from the XRD pattern using Bragg’s equation
(section 3.4) taking into consideration the tetragonal splitting of the peaks at ~45°. As the
76
Ta
ble
4
.1.
Av
erag
e cr
yst
alli
te
size
±
0.0
01
(n
m)
calc
ula
ted
fo
r 1
:3
sol:
sup
po
rt
solv
ent
rati
o
and
var
ious
surf
acta
nt
con
cen
trat
ion
s
Em
uls
ion
Con
cen
trati
on
C
alc
ina
tio
n T
emp
era
ture
(75
0°C
)
Ca
lcin
ati
on
Tem
per
atu
re
(10
00
°C)
Sol:
Support
Solv
ent
Rat
io
Su
rfac
tan
t
Vo
lum
e S
pan
80
Span
20
Span
80
Span
20
1:0
3
5%
2
7.9
34
2
4.0
17
4
5.0
00
4
3.0
51
10
%
29
.84
7
30
.92
3
44
.00
5
46
.81
9
15
%
30
.37
2
31
.64
2
45
.70
4
47
.99
3
20
%
33
.03
3
35
.51
3
44
.01
4
52
.00
9
77
Fig
4.4
. X
-ray
pat
tern
of
Ba
1-xD
xT
iO3 p
ow
der
cal
cin
ed a
t 7
50
oC
, fo
r D
= S
r, L
a, C
e &
Mg w
ith v
aryin
g x
synth
esiz
ed u
sing 5
%
span
80
(*
- B
aCO
3).
20
30
40
50
60
70
80x
= 0
.1
x =
0.0
5x =
0.0
1x =
0.0
01
x =
0
Ce
(301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
**
**
*
*
Intensity (arb. units)
2T
heta
(0)
20
30
40
50
60
70
80
x =
0.1
x =
0.0
5x =
0.0
1x =
0.0
01
x =
0
Mg
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
(301)
**
**
*
*
Intensity (arb. units)
2T
heta
(0)
20
30
40
50
60
70
80x
= 0
.1x =
0.0
5x =
0.0
1x =
0.0
01
x =
0
La
(301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
**
**
*
*
2 T
heta
(0)
Intensity (arb. units)
20
30
40
50
60
70
80x
= 0
.01
x =
0.1
x =
0.0
5
x =
0
Sr
x =
0.0
01
(301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
**
**
*
*
Intensity (arb. units)
2T
heta
(0)
78
Fig
4.4
. co
nt.
X-r
ay p
atte
rn o
f B
a1-xD
xT
iO3 p
ow
der
cal
cin
ed a
t 7
50
oC
, fo
r D
= L
i &
K w
ith
var
yin
g x
sy
nth
esiz
ed u
sin
g 5
%
span
80
(*
- B
aCO
3).
Fig
4.5
. X
-ray
pat
tern
of
Ba
1-xD
xT
iO3 p
ow
der
cal
cin
ed a
t 7
50
oC
, fo
r D
= S
r &
La
wit
h v
ary
ing
x s
yn
thes
ized
usi
ng
5%
sp
an 2
0
(*-B
aCO
3).
20
30
40
50
60
70
80
x =
0.1
x =
0.0
5x =
0.0
1x =
0.0
01
x =
0
K
(301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
**
**
*
*
Intensity (arb. units)
2T
heta
(0)
20
30
40
50
60
70
80
x =
0.1
x =
0.0
5x =
0.0
1x =
0.0
01
x =
0
Li
**
*
(301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)*
*
**
**
Intensity (arb. units)
2T
heta
(0)
20
30
40
50
60
70
80
x =
0.1
x =
0.0
5
x =
0.0
1
x =
0.0
01
x =
0
Sr (301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
**
**
*
*
2T
heta (
0)
Intensity (arb. units)
20
30
40
50
60
70
80
x =
0.1
x =
0.0
5x =
0.0
1
x =
0.0
01
x =
0
La (301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
**
**
*
*
Intensity (arb. units)
2 T
heta (
0)
79
Fig
4.5
. co
nt.
X-r
ay p
atte
rn o
f B
a1-xD
xT
iO3 p
ow
der
cal
cin
ed a
t 7
50
oC
, fo
r D
= C
e, M
g,
Li
& K
wit
h v
aryin
g x
synth
esiz
ed u
sing
5%
sp
an 2
0 (
* -
BaC
O3).
20
30
40
50
60
70
80x
= 0
.1
x =
0.0
5x =
0.0
1
x =
0.0
01
x =
0
K
(301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
**
**
**
Intensity (arb. units)
2T
heta
(0)
20
30
40
50
60
70
80
x =
0.1
x =
0.0
5x =
0.0
1x =
0.0
01
x =
0
Li
(301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
**
**
*
*
Intensity (arb. units)
2T
heta
(0)
20
30
40
50
60
70
80
x =
0.1
x =
0.0
5
x =
0.0
1
x =
0.0
01
x =
0
Mg
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
(301)
**
**
*
*
Intensity (arb. units)
2T
heta
(0)
20
30
40
50
60
70
80x
= 0
.1
x =
0.0
5
x =
0.0
1
x =
0.0
01
x =
0
Ce (301)
(300)
(202)
(211)
(201)
(200)
(111)
(110)
(001)
**
**
*
*
Intensity (arb. units)
2T
heta
(0)
80
crystal structure was not 100% tetragonal assuming a pc structure, the cell volume that
was deduced from these lattice parameters was used to calculate the experimental pc
lattice parameter using the formula [Dasgupta et al., 2002]
CDE ! FGH
I ---(4.2)
where, V = abc is the volume of the crystal lattice (for cubic structure, a=b=c and for a
tetragonal structure, a=b!c), Z is the number of ABO3 units in one unit cell of crystal. Z
=1 for the tetragonal (P4mm) structure [Buttner and Maslen, 1992; Sologub et al., 2002].
Figure 4.6 shows BaTiO3 (110) diffraction peak with various dopants and varying dopant
concentration for powders prepared using span 80. It was noted that the maximum
intensity (110) diffraction peak shifted towards higher angle irrespective of the dopant
used. The (110) peak noted at 31.47° for undoped BaTiO3 was observed at 31.59° for Sr,
31.78° for La and Mg, 31.92° for Ce, 31.95° for Li and 31.84° for K, for x=0.1 dopant
concentration. The high angle shift of diffraction peaks has been attributed either to the
decreasing tetragonality and increasing cubic structure of the crystals or to lower values
of lattice spacing [Da-Yong et al., 2006; Mao et al., 2010; Kribalis et al., 2006]. For Sr-
doped BaTiO3, Ouyang et al. and Vargas and co-workers attributed it to the smaller ionic
radius of Sr as compared to Ba (Table 4.2) [Ouyang et al., 2010; Vargas-Ortíz et al.,
2012; Sutham, 2008; Kribalis et al., 2006]. For every 0.1 increase in x of Ba1-xSrxTiO3,
Vargas and co-workers observed a shift of ~0.1° in (110) diffraction peak [Vargas-Ortíz
et al., 2012]. Similar shift of diffraction peak towards higher angle with La3+
doping and
a decrease in lattice constant with Ce3+
doping, when the dopant ion replaced Ba2+
ion has
been reported [Park et al., 1998; Sabina, 2001; Da-Yong et al., 2006]. However, a low
angle shift of the diffraction peaks have been reported when Ce4+
(rCe = 0.87 Aû, CN =
6) or Mg2+
(rMg = 0.72 Aû, CN = 6) replaced Ti4+
in BaTiO3 as their ionic radii are closer
to that of Ti (rTi = 0.605 Aû, CN = 6) than that of Ba (rBa = 1.35 Aû, CN = 6). This
81
Fig
4.6
. M
axim
um
inte
nsi
ty (
110)
dif
frac
tion p
eak o
f B
a 1-x
DxT
iO3 p
ow
der
s ca
lcin
ed a
t 7
50
oC
for
D =
Sr,
La,
Ce,
Mg,
Li
& K
wit
h v
aryin
g x
as
0 (
),
0.0
01 (
), 0
.01 (
),
0.0
5 (
) an
d 0
.1 (
)
synth
esiz
ed u
sing 5
% s
pan
80.
30.8
31.0
31.2
31.4
31.6
31.8
32.0
32.2
32.4
K
Intensity (arb. units)
2T
heta
(0)
30.8
31.0
31.2
31.4
31.6
31.8
32.0
32.2
32.4
32.6
Li
Intensity (arb. units)
2T
heta
(0)
30.6
30.8
31.0
31.2
31.4
31.6
31.8
32.0
32.2
32.4
32.6
Mg
Intensity (arb. units)
2T
heta
(0)
30.8
31.0
31.2
31.4
31.6
31.8
32.0
32.2
32.4
Ce
Intensity (arb. units)
2T
heta
(0)
30.8
31.2
31.6
32.0
32.4
32.8
La
Intensity (arb. units)
2 T
heta
(0)
30.8
31.0
31.2
31.4
31.6
31.8
32.0
32.2
Sr
Intensity (arb. units)
2T
heta
(0)
82
Ta
ble
4.2
. Io
nic
Rad
ius
(Aû)
use
d i
n t
heo
reti
cal
calc
ula
tio
n o
f la
ttic
e p
aram
eter
s [R
ick
, 2
00
7;
Dav
id,
20
05
]
CN
r B
a (
2+
) r T
i (4
+)
r O (
2-)
r Sr
(2+
) r L
a (
3+
) r C
e(3
+)
r Mg
(2
+)
r Li
(1+
) r K
(1
+)
XII
(1
2)
1.6
1
- -
1.4
4
1.3
6
1.3
4
- -
1.6
4
VI
(6)
1.3
5
0.6
05
1.4
1.1
8
1.0
32
1.0
1
0.7
2
0.7
6
1.3
8
II (
2)
- -
1.3
5
- -
- -
- -
83
causes crystal cell expansion due to their larger ionic size as compared to Ti in the TiO6
octahedra [Da-Yong et al., 2006; Mao et al., 2010; Fukuda et al., 2007; Min et al., 2007].
Besides this, the lower angle shift of the diffraction peak may also be due to some dopant
ions being incorporated in Ba-[TiO6] sub lattices. For Sr doping, Ouyang et al. have
reported that Sr-[TiO6] sub-lattice starts forming with increasing Sr concentration in
BaxSr1-xTiO3. Sr-[TiO6] sub-lattice is unable to sustain Ba cation and expel them due to
their larger ionic size, resulting in Sr-rich phases. However, they reported that the smaller
size Sr gets incorporated in the Ba-[TiO6] sub-lattices in the Ba-rich phases due to its
comparatively larger cell volumes. Figure 4.7 shows BaTiO3 (110) diffraction peak with
various dopant and varying dopant concentration for powders prepared using span 20.
This peak was noted either at higher 2! value or the same 2! value within the range of
0.05° as that noted for pure BaTiO3. Also, a small increase in tetragonality was noted in
the doped powders for certain dopant concentrations. Hence, in the present study for
doped BaTiO3 powders, the (110) diffraction peak shift towards higher angle with respect
to that of pure BaTiO3 most probably is due to the smaller lattice spacing as well as the
increase in cubicity with increasing dopant concentration.
Figures 4.8 and 4.9 show the (002)/(200) peak splitting for Ba1-xDxTiO3 powders
synthesized using span 80 and span 20 respectively calcined at 750°C. A shoulder at
(002) was noted for pure BaTiO3 and the % tetragonality calculated was 16.97% for the
powder synthesized using span 80 and 20.04% for that using span 20. Sr, La, Li and K
doped BaTiO3 powders for x=0.1 dopant concentration possessed a single symmetric
(200) peak (Figure 4.8) indicating it to be predominantly cubic. This result is in
corroboration with the peak shift also noted from Figure 4.6 for these powders
synthesized using span 80. However, it was noted that the %tetragonality slightly
increased for low doping concentration, for example, Sr doping in BaTiO3 led to an
increase in % tetragonality up to 21.23% for x=0.01. This may possibly be due to Sr
getting incorporated in the Ba-[TiO6] sub lattice instead of replacing Ba ion owing to
lower concemtartion of Sr, resulting in the c-axis elongation. Further increase in Sr
concentration to x=0.05 and 0.1 showed a single symmetric (200) peak (Figure 4.8)
84
Fig
4.7
. M
axim
um
inte
nsi
ty (
110)
dif
frac
tion p
eak o
f B
a 1-x
DxT
iO3 p
ow
der
s ca
lcin
ed a
t 7
50
oC
for
D =
Sr,
La,
Ce,
Mg,
Li
& K
wit
h v
aryin
g x
as
0 (
),
0.0
01 (
), 0
.01 (
),
0.0
5 (
) an
d 0
.1 (
)
synth
esiz
ed u
sing 5
% s
pan
20.
30.6
30.8
31.0
31.2
31.4
31.6
31.8
32.0
32.2
32.4
Ce
Intensity (arb. units)
2T
heta
(0)
31.0
31.2
31.4
31.6
31.8
32.0
K
Intensity (arb. units)
2T
heta
(0)
31.0
31.2
31.4
31.6
31.8
32.0
Li
Intensity (arb. units)
2T
heta
(0)
31.0
31.2
31.4
31.6
31.8
32.0
Mg
Intensity (arb. units)
2T
heta
(0)
30.8
31.0
31.2
31.4
31.6
31.8
32.0
32.2
32.4
La
Intensity (arb. units)
2 T
heta
(0)
31.0
31.2
31.4
31.6
31.8
32.0
32.2
Sr
Intensity (arb. units)
2T
heta
(0)
85
Fig
4.8
. (0
02)/
(20
0)
dif
frac
tion p
eak s
pli
ttin
g o
f B
a1-xD
xT
iO3 p
ow
der
s ca
lcin
ed a
t 750
oC
for
D =
Sr,
La,
Ce,
Mg,
Li
& K
wit
h
var
yin
g x
as
0 (
),
0.0
01
(
), 0
.01
(
), 0
.05
( )
and
0.1
( )
syn
thes
ized
usi
ng
5%
sp
an 8
0.
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
46.0
46.2
Ce
Intensity (arb. units)
2Th
eta
(0 )
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
46.0
46.2
46.4
Li
Intensity (arb. units)
2T
heta
(0)
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
46.0
46.2
K
Intensity (arb. units)
2T
heta
(0)
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
46.0
Mg
Intensity (arb. units)
2T
heta
(0)
44.4
44.6
44.
845.
045.2
45.4
45.6
45.8
46.0
La
2 T
heta
(0 )
Intensity (arb. units)
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
46.0
46.2
Sr
Intensity (arb. units)
2T
heta
(0)
86
Fig
4.9
. (0
02)/
(20
0)
dif
frac
tion p
eak s
pli
ttin
g o
f B
a1-xD
xT
iO3 p
ow
der
s ca
lcin
ed a
t 750
oC
for
D =
Sr,
La,
Ce,
Mg,
Li
& K
wit
h
var
yin
g x
as
0 (
),
0.0
01
(
), 0
.01
(
), 0
.05
( )
and
0.1
( )
syn
thes
ized
usi
ng
5%
sp
an 2
0.
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
K
Intensity (arb. units)
2T
heta
(0)
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
46.0
Li
Intensity (arb. units)
2T
heta
(0)
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
Mg
Intensity (arb. units)
2T
heta
(0)
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
46.0
Ce
2T
heta
(0)
Intensity (arb. units)
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
46.0
46.2
La
2 T
heta
(0)
Intensity (arb. units)
44.4
44.6
44.8
45.0
45.2
45.4
45.6
45.8
46.0
46.2
Sr
Intensity (arb. units)
2T
heta
(0)
87
confirming the predominance of cubic phase. Similarly for La, Li and K doped BaTiO3
powders, x=0.1 dopant concentration is sufficient for majority of the dopant ion to
replace Ba ion resulting in a predominantly cubic structure. On the other hand, in Ce and
Mg doped BaTiO3 powders synthesized using span 80, presence of (002) peaks or
shoulders produceing asymmetric (200) peak with tetragonal structures was present for
all dopant concentrations. Since Ce can exist as Ce3+
and Ce4+
, the tetragonality may
either be due to Ce3+
getting incorporated in the Ba-[TiO6] sub lattice or some amount of
Ce4+
present in the powders may replace Ti due to its closeness of the ionic radii therby
causing the c-axis elongation. Mg has its ionic radius much closer to Ti as compared to
Ba and hence, is most likely to replace Ti in the synthesized powders. However, as the
(110) peak did not show any significant shift towards lower value, it is possible that Mg
might have also replaced Ba ion. Of the small percentage of tetragonality present in the
synthesized powders, which renders it the final pc structure, tetragonal expansion or
compression was identified by monitoring the intensities of (002) and (200) peaks. As has
been reported, higher intensity of (002) peak over that of (200) peak implies tetragonal
compression [Mao et al. 2010]. The normally observed case wherein the intensity of
(200) peak is higher than that of (002) peak, implies tetragonal expansion along the c-axis
of the crystal structure [Mao et al. 2010]. Figure 4.9 shows the presence of (002) and
(200) diffraction peaks for all dopant concentrations including x=0.1 for powders
synthesized using Span 20, showing higher probability of presence of tetragonality in
these powders as compared to those synthesized using Span 80. This may be due to the
elongated structures (Fig. 14c) formed in powders synthesized using span 20 which might
facilitate elongation along c-axis. Large polarization in response to the applied electric
field of asymmetric tetragonal BaTiO3 unit cell, affects its microstructure and crystalline
behaviour [Shackelford, 2005]. Atomic dipole moments vary as the crystal expand or
contract [Anderson et al., 2003]. Below TC in the absence of applied electric field, there
are at least two directions along which spontaneous polarization can develop. When the
crystal cools through the cubic-tetragonal transition temperature, the cubic cell slightly
expands along one edge due to the movement of the octahedrally coordinated Ti4+
ions to
produce the tetragonal c-axis and is compressed along the other two edges to form
88
tetragonal a and b-axes [Tilley, 2004; Kasap, 2007]. This results in the dipole to point
along c-axis. Since all the six directions parallel to ±x, ±y and ±z are equivalent and no
preference is given as to which of the original cubic axis becomes the polar direction, any
of these six can become the polar direction [Tilley, 2004]. Covalent bond between the
Ti4+
cation is stronger with the four neighbouring oxygen anions in the ab plane, than the
two oxygen anions along the c-axis of the perovskite structure [Jun et al., 2011]. Due to
this, the tetragonal expansion and compression is favorable in the c-axis of the perovskite
structure. For all powders synthesized using span 80 and span20 showing presence of
(002) and (200) peaks, the intensity of (200) peak was higher than (002) peak indicating
tetragonal expansion.
Tables 4.3 and 4.4 reports the parameters of Ba1-xDxTiO3 crystals prepared using 1:3 sol:
support solvent with 5% span 80 and span 20 respectively, calcined at 750°C. These
tables summarize the crystallite size, lattice parameters, c/a ratio, cell volume and
%tetragonality. Crystallite size obtained for the powders synthesized using span 80 varied
from 18.121- 36.318nm, and that using span 20 varied from 15.981- 40.718nm. For
powders synthesized using span 80, crystallite size noted for x=0.1 dopant concentration
were smaller than that noted for the other dopant concentrations. Corresponding to this,
small reduction in the cell volume was also noted. Cell volume was observed to shrink,
between 1.24 and 2.27% for x=0.1 dopant concentration in powders synthesized using
span 80 with reference to that of pure BaTiO3. For the other concentrations the cell
volume was almost the same, varying by 0.83–1.17%. In powders synthesized using span
20, the cell volume varied between 0.02 and 1.76% for all the dopant concentrations with
reference to that of pure BaTiO3. The range of change in % tetragonality for the powders
synthesized using span 80 was 16.97-24.35% and that using span 20 was 15.57-26.90%.
A comparative study of the experimentally determined lattice parameters with those
predicted by mathematical models developed by Jiang et al. [2006] and Rick Ubic [2007]
was done for all the synthesized powders. The ionic radii of all elements corresponding to
the coordination number associated with that of the ions forming the BaTiO3 perovskite
89
Table 4.3. Ba1-xDxTiO3 crystal parameters deduced from XRD patterns (Figure 4.4) for
powders prepared using 1:3 sol: support solvent ratio with 5% span80 and calcined at
750°C.
Ba1-xDxTiO3 Structure
Crystallite
Size ± 0.001
(nm)
a ±0.0009
(A°)
c
±0.0009
(A°)
apc
±0.0006
(A°)
c/a
Cell
Volume
±0.01
(A°3)
%
Tetragonality
BaTiO3 pc 27.934 4.0111 4.0295 4.0171 1.0046 64.83 16.97
Ba0.999Sr0.001TiO3
pc
30.883 4.0005 4.0249 4.0086 1.0061 64.42 17.42
Ba0.99Sr0.01TiO3 31.429 3.9980 4.0241 4.0067 1.0065 64.32 21.23
Ba0.95Sr0.05TiO3 34.334 4.0014 - 4.0014 - 64.07 -
Ba0.9Sr0.1TiO3 22.605 4.0007 - 4.0007 - 64.03 -
Ba0.999La0.001TiO3
pc
28.454 4.0089 4.0275 4.0150 1.0046 64.73 19.38
Ba0.99La0.01TiO3 30.882 4.0055 4.0377 4.0162 1.0080 64.78 19.01
Ba0.95La0.05TiO3 33.512 4.0038 4.0377 4.0151 1.0084 64.73 18.82
Ba0.9La0.1TiO3 29.273 3.9961 - 3.9961 - 63.85 -
Ba0.999Ce0.001TiO3
pc
29.611 4.0241 4.0368 4.0283 1.0032 65.37 24.35
Ba0.99Ce0.01TiO3 30.106 4.0064 4.0343 4.0156 1.0070 64.75 18.63
Ba0.95Ce0.05TiO3 34.848 4.0080 4.0402 4.0187 1.0080 64.90 21.57
Ba0.9Ce0.1TiO3 18.121 3.9739 4.0122 3.9867 1.0076 63.36 20.62
Ba0.999Mg0.001TiO3
pc
36.128 4.0055 4.0317 4.0142 1.0065 64.69 17.89
Ba0.99Mg0.01TiO3 31.567 4.0080 4.0249 4.0137 1.0042 64.66 17.17
Ba0.95Mg0.05TiO3 29.610 4.0114 4.0283 4.0170 1.0042 64.82 17.58
Ba0.9Mg0.1TiO3 26.149 3.9905 4.0072 3.9961 1.0042 63.81 17.49
Ba0.999Li0.001TiO3
pc
27.806 4.0114036 4.028305 4.0093 1.0042 64.82 16.43
Ba0.99Li0.01TiO3 34.675 4.0022 4.0224 4.0089 1.0050 64.43 21.90
Ba0.95Li0.05TiO3 36.318 4.0055 4.0368 4.0159 1.0078 64.77 22.17
Ba0.9Li0.1TiO3 23.833 3.9957 - 3.9957 - 63.79 -
Ba0.999K0.001TiO3
pc
31.151 4.0080 4.0309 4.0156 1.0057 64.75 18.18
Ba0.99K0.01TiO3 33.834 4.0064 4.0292 4.0139 1.0057 64.67 21.44
Ba0.95K0.05TiO3 35.749 4.0106 4.0360 4.0190 1.0063 64.92 15.27
Ba0.9K0.1TiO3 21.180 3.9954 - 3.9954 - 63.78 -
90
Table 4.4. Ba1-xDxTiO3 crystal parameters deduced from XRD patterns (Figure 4.5) for
powders prepared using 1:3 sol: support solvent ratio with 5% span20 and calcined at
750°C.
Ba1-xDxTiO3 Structure
Crystallite
Size ± 0.001
(nm)
a
±0.0009
(A°)
c
±0.0009
(A°)
apc
±0.0006
(A°)
c/a
Cell
Volume
±0.01
(A°3)
%
Tetragonality
BaTiO3 pc 24.017 4.0126 4.0329 4.0194 1.0051 64.94 20.04
Ba0.999Sr0.001TiO3
pc
29.253 4.0072 4.0471 4.0204 1.0100 64.99 16.53
Ba0.99Sr0.01TiO3 30.618 4.0055 4.0343 4.0151 1.0072 64.73 18.87
Ba0.95Sr0.05TiO3 25.243 4.0114 4.0326 4.0184 1.0053 64.89 19.57
Ba0.9Sr0.1TiO3 26.048 3.9947 4.0334 4.0075 1.0097 64.36 18.29
Ba0.999La0.001TiO3
pc
33.530 3.9872 4.0131 3.9958 1.0065 63.80 17.80
Ba0.99La0.01TiO3 40.718 4.0072 4.0360 4.0168 1.0072 64.81 16.62
Ba0.95La0.05TiO3 15.981 3.9930 4.0215 4.0025 1.0071 64.12 19.15
Ba0.9La0.1TiO3 18.579 4.0038 4.0343 4.0140 1.0076 64.67 19.64
Ba0.999Ce0.001TiO3
pc
35.566 4.0114 4.0360 4.0196 1.0061 64.94 17.91
Ba0.99Ce0.01TiO3 29.372 4.0089 4.0394 4.0190 1.0076 64.92 16.84
Ba0.95Ce0.05TiO3 30.750 4.0080 4.0351 4.0170 1.0068 64.82 18.04
Ba0.9Ce0.1TiO3 22.335 4.0122 4.0402 4.0216 1.0070 65.04 19.58
Ba0.999Mg0.001TiO3
pc
35.937 4.0097 4.0351 4.0182 1.0063 64.88 16.45
Ba0.99Mg0.01TiO3 31.424 4.0114 4.0360 4.0196 1.0061 64.94 17.55
Ba0.95Mg0.05TiO3 31.562 4.0122 4.0377 4.0207 1.0063 65.00 18.04
Ba0.9Mg0.1TiO3 31.989 4.0122 4.0334 4.0193 1.0053 64.93 16.94
Ba0.999Li0.001TiO3
pc
27.596 4.0122 4.0266 4.0170 1.0036 64.82 25.36
Ba0.99Li0.01TiO3 28.232 4.0097 4.0351 4.0182 1.0063 64.88 19.97
Ba0.95Li0.05TiO3 29.736 4.0055 4.0317 4.0142 1.0065 64.69 16.27
Ba0.9Li0.1TiO3 37.311 4.0106 4.0292 4.0168 1.0046 64.81 15.57
Ba0.999K0.001TiO3
pc
30.749 4.0064 4.0343 4.0156 1.0070 64.75 26.90
Ba0.99K0.01TiO3 30.748 4.0097 4.0343 4.0179 1.0061 64.86 23.27
Ba0.95K0.05TiO3 31.016 4.0038 4.0377 4.0151 1.0084 64.73 21.44
Ba0.9K0.1TiO3 30.619 4.0097 4.0360 4.0184 1.0065 64.89 16.42
91
structure, necessary for the study is listed in Table 4.2. Table 4.5 and 4.6 compares the
experimental values of the lattice parameters obtained in the pure and doped BaTiO3
powders with the theoretically determined values using Jiang and Rick’s formulation and
the deviation between the values for powders synthesized using Span 80 and Span 20
respectively. The properties of perovskite oxides are strongly linked to their crystal
structure [Zhang et al., 2012; Gorelov et al., 2011; Frey and Payne, 1996; Takeuchi et al.,
1997; Xue et al., 2007; Fong et al., 2004]. Intensive research is carried out to design and
synthesize new cubic perovskites to be used as thin films, substrate or buffer materials
[Zhang et al., 2012; Moreira and Dias, 2007; Kumar et al., 2009]. Hence, predicting the
lattice parameters of these compounds is important as it may assist in developing new
materials [Zhang et al., 2012; Moreira and Dias, 2007; Kumar et al., 2009]. Jiang et al.
[2006] proposed an empirical relationship (eqn. 4.3) relating the ionic radii of the A, B
and O ions of the ABO3 cubic perovskite structure to the lattice constants assuming a six-
fold coordination for all the ions involved. Jiang et al. [2006] assumed the lattice constant
to be ideally equal to 2(JK L"JM) and accounted for the variation in the lattice constant
due to larger A ion by using tolerance factor (t). In an ideal cubic perovskite structure, t
% 1 and lattice constant a = $N"OJ3 "L "JMP = 2(JK L"JM). However, when ‘t’ deviates from
its ideal value, ‘a’ changes due to change in ionic sizes compared to the ideally reported
value for the formation of cubic perovskite structure. When ‘t’ is larger, $N"OJ3 "L "JMP is
overestimated and 2(JK L"JM) is underestimated and when ‘t’ is smaller, $N"OJ3 "L "JMP is
underestimated and 2( ! "# $) is overestimated. This accounts for error in the
estimated‘a’ value. Lattice constant obtained from the empirical formula was reported to
show an average relative absolute error, (|%error|) which is the difference between the
theoretically and the experimentally obtained value, of 0.63%. The six-fold assumption is
valid for the B site cation but is incorrect for the A site cation and the anion O, the actual
coordination number being 12 and 2 respectively. Rick Ubic [2007] proposed another
empirical equation (eqn. 4.4) relating the ionic radii and the cubic or pc lattice parameters
to predict the lattice constants with an average error of 0.60%. The average |%error| of
0.63% and 0.60% using eqn. 4.3 and 4.4 was calculated for the 132 known perovskites.
92
Table 4.5. Experimentally deduced lattice parameter of the pc Ba1-xDxTiO3 powders
synthesized using 1:3 sol: support solvent ratio with 5% span80, calcined at 750°C
compared with that derived theoretically using the empirical formulae with (a) and
without (a ) assuming the six-fold coordination for the ions.
Ba1-xDxTiO3
apc
±0.0006
(A°)
aJiang
(A°) |%error|
aRick
(A°) |%error|
a'Jiang
(A°) |%error|
a'Rick
(A°) |%error|
BaTiO3 4.0171 4.0153 0.0458 4.0070 0.2519 4.0712 1.3287 4.0454 0.6998
Ba0.9099Sr0.001TiO3 4.0086 4.0152 0.1640 4.0069 0.0415 4.0711 1.5356 4.0453 0.9082
Ba0.99Sr0.01TiO3 4.0067 4.0144 0.1928 4.0055 0.0296 4.0703 1.5637 4.0442 0.9286
Ba0.95Sr0.05TiO3 4.0014 4.0108 0.2357 3.9992 0.0534 4.0666 1.6050 4.0393 0.9392
Ba0.9Sr0.1TiO3 4.0007 4.0064 0.1414 3.9914 0.2320 4.0621 1.5107 4.0332 0.8053
Ba0.999La0.001TiO3 4.0150 4.0151 0.3078 4.0069 0.5142 4.0711 1.0711 4.0453 0.4409
Ba0.99La0.01TiO3 4.0162 4.0136 0.0644 4.0055 0.2679 4.0699 1.3187 4.0442 0.6925
Ba0.95La0.05TiO3 4.0151 4.0069 0.2033 3.9992 0.3963 4.0645 1.2155 4.0393 0.5997
Ba0.9La0.1TiO3 3.9961 3.9986 0.0611 3.9914 0.1178 4.0578 1.5184 4.0332 0.9183
Ba0.999Ce0.001TiO3 4.0283 4.0151 0.3286 4.0069 0.5350 4.0711 1.0506 4.0453 0.4203
Ba0.99Ce0.01TiO3 4.0156 4.0135 0.0531 4.0055 0.2536 4.0698 1.3301 4.0442 0.7066
Ba0.95Ce0.05TiO3 4.0187 4.0064 0.3089 3.9992 0.4876 4.0640 1.1125 4.0393 0.5093
Ba0.9Ce0.1TiO3 3.9867 3.9974 0.2694 3.9914 0.1197 4.0567 1.7259 4.0332 1.1534
Ba0.999Mg0.001TiO3 4.0142 4.0150 0.0183 4.0067 0.1872
Ba0.99Mg0.01TiO3 4.0137 4.0120 0.0417 4.0040 0.2424
Ba0.95Mg0.05TiO3 4.0170 3.9987 0.4572 3.9916 0.6373
Ba0.9Mg0.1TiO3 3.9961 3.9822 0.3480 3.9761 0.5009
Ba0.999Li0.001TiO3 4.0093 4.0150 0.1418 4.0068 0.0635
Ba0.99Li0.01TiO3 4.0089 4.0122 0.0823 4.0041 0.1186
Ba0.95Li0.05TiO3 4.0159 3.9998 0.4033 3.9926 0.5850
Ba0.9Li0.1TiO3 3.9957 3.9843 0.2867 3.9781 0.4429
Ba0.999K0.001TiO3 4.0156 4.0153 0.0079 4.0069 0.2140 4.0712 1.3660 4.0453 0.7373
Ba0.99K0.01TiO3 4.0139 4.0155 0.0376 4.0055 0.2113 4.0714 1.4109 4.0442 0.7485
Ba0.95K0.05TiO3 4.0190 4.0161 0.0729 3.9992 0.4943 4.0720 1.3022 4.0393 0.5026
Ba0.9K0.1TiO3 3.9954 4.0169 0.5347 3.9914 0.0991 4.0728 1.9016 4.0332 0.9369
93
Table 4.6. Experimentally deduced lattice parameter of the pc Ba1-xDxTiO3 powders
synthesized using 1:3 sol: support solvent ratio with 5% span20, calcined at 750°C
compared with that derived theoretically using the empirical formulae with (a) and
without (a ) assuming the six-fold coordination for the ions.
Ba1-xDxTiO3
apc
±0.0006
(A°)
aJiang
(A°) |%error|
aRick
(A°) |%error|
a'Jiang
(A°) |%error|
a'Rick
(A°) |%error|
BaTiO3 4.0194 4.0153 0.1019 4.0070 0.3082 4.0712 1.2733 4.0454 0.6441
Ba0.999Sr0.001TiO3 4.0204 4.0152 0.1306 4.0069 0.3367 4.0711 1.2450 4.0453 0.6158
Ba0.99Sr0.01TiO3 4.0151 4.0144 0.0168 4.0055 0.2396 4.0703 1.3570 4.0442 0.7205
Ba0.95Sr0.05TiO3 4.0184 4.0108 0.1898 3.9992 0.4801 4.0666 1.1854 4.0393 0.5167
Ba0.9Sr0.1TiO3 4.0075 4.0064 0.0291 3.9914 0.4031 4.0621 1.3426 4.0332 0.6360
Ba0.999La0.001TiO3 3.9958 4.0151 0.4817 4.0069 0.2769 4.0711 1.8497 4.0453 1.2244
Ba0.99La0.01TiO3 4.0168 4.0136 0.0782 4.0055 0.2817 4.0699 1.3051 4.0442 0.6788
Ba0.95La0.05TiO3 4.0025 4.0069 0.1114 3.9992 0.0810 4.0645 1.5257 4.0393 0.9119
Ba0.9La0.1TiO3 4.0140 3.9986 0.3844 3.9914 0.5641 4.0578 1.0794 4.0332 0.4766
Ba0.999Ce0.001TiO3 4.0196 4.0151 0.1110 4.0069 0.3169 4.0711 1.2653 4.0453 0.6363
Ba0.99Ce0.01TiO3 4.0190 4.0135 0.1373 4.0055 0.3380 4.0698 1.2471 4.0442 0.6230
Ba0.95Ce0.05TiO3 4.0170 4.0064 0.2666 3.9992 0.4451 4.0640 1.1543 4.0393 0.5513
Ba0.9Ce0.1TiO3 4.0216 3.9974 0.6034 3.9914 0.7543 4.0567 0.8659 4.0332 0.2883
Ba0.999Mg0.001TiO3 4.0182 4.0150 0.0797 4.0067 0.2854
Ba0.99Mg0.01TiO3 4.0196 4.0120 0.1891 4.0040 0.3902
Ba0.95Mg0.05TiO3 4.0207 3.9987 0.5491 3.9916 0.7293
Ba0.9Mg0.1TiO3 4.0193 3.9822 0.9314 3.9761 1.0852
Ba0.999Li0.001TiO3 4.0170 4.0150 0.0508 4.0068 0.2565
Ba0.99Li0.01TiO3 4.0182 4.0122 0.1488 4.0041 0.3501
Ba0.95Li0.05TiO3 4.0142 3.9998 0.3609 3.9926 0.5425
Ba0.9Li0.1TiO3 4.0168 3.9843 0.8145 3.9781 0.9716
Ba0.999K0.001TiO3 4.0156 4.0153 0.0082 4.0069 0.2142 4.0712 1.3658 4.0453 0.7371
Ba0.99K0.01TiO3 4.0179 4.0155 0.0605 4.0055 0.3096 4.0714 1.3142 4.0442 0.6511
Ba0.95K0.05TiO3 4.0151 4.0161 0.0248 3.9992 0.3963 4.0720 1.3985 4.0393 0.5997
Ba0.9K0.1TiO3 4.0184 4.0169 0.0393 3.9914 0.6766 4.0728 1.3356 4.0332 0.3653
94
The average |%error| associated with the empirical formula devised by Rick Ubic
decreased as the changes in the ionic radii were directly incorporated in the estimation of
‘a’ [Rick, 2007]. In this case also, six-fold coordination assumption produced lower
average error values as compared to the lattice constant obtained using ionic radii values
corresponding to the actual coordination number. As a predictive tool, assumption of six-
fold coordination in the empirical relations has proved to be more accurate. The reason
for this was considered to be the lower accuracy of the measured values of ionic radii for
the actual coordination [Rick, 2007; Moreira and Dias, 2007]. The |%error| calculated for
BaTiO3 structure with experimental lattice parameter value of 4.012 A° using eqn. 4.3
was 0.075 and using eqn. 4.4 was 0.125. This indicates a better fit to be achieved for
BaTiO3 with eqn. 4.3, i.e. with the empirical formula devised by Jiang et al. [2006].
%&'()* + ,-../0#1 ! "# $2 " #,-3.4.# 5 67#8#69:;#16<#8#692= > #,-?@0? ---(4.3)
%ABCD + @-@0E3, " #@-34@F?#1 G#"# $2 " #,-?4?,?#1 !#"# $2 ---(4.4)
Theoretical lattice parameters reported in Tables 4.5 and 4.6 were calculated with the
help of (eqn. 4.3) and (eqn. 4.4) using the ionic radii values reported in Table 4.2
assuming six-fold coordination for all ions and is reported as ‘a’ with subscript
specifying the author. For Ba1-xDxTiO3, two cations namely ‘Ba’ and the dopant ‘D’
occupy the A-site position. Hence, G#for the doped powders was calculated as the
weighted average, of the ionic radii of the two A-site cations, based on their
concentration. The lattice parameters were also evaluated using the ionic radii values for
the actual coordination of 12 and 2 respectively for A site cation and anion O. These
values are reported as ‘a ’ with subscript specifying the author. Experimentally obtained
lattice parameter values for pure BaTiO3 in the present work fit better with the empirical
formula devised by Jiang et al. [2006] having minimum absolute relative error (|%error|)
of 0.046% (Table 4.5) assuming six-fold coordination which is lower than that of 0.075%
reported by Jiang et al. [2006]. Also, the average |%error| calculated for the pure and
95
doped BaTiO3 powders synthesized using span 80 for the lattice parameters devised by
Jiang et al. was 0.1923%. This may be attributed to the fact that in case of nanoparticles,
crystal structures are known to deviate from that of bulk structure [Poole, 2003]. This
possibly is a contributing factor for better match with six fold coordination values in the
present case. Moreover, the modeling was done for cubic lattice structure. The
nanoparticles with pc lattice structure may not confirm to the coordination of the cubic
lattice exactly. The absolute error calculated using the actual coordination values was
higher than that obtained by assuming six-fold coordination, same as that reported by
Rick Ubic [2007]. The absolute error obtained was also higher using eqn 4.4 as compared
to eqn 4.3. All dopants used in the present work except for Mg and Li exist in the 12-fold
coordination state. The maximum coordination number that Mg and Li take is 8, due to
which lattice parameters for these dopants with only six-fold coordination was evaluated
[David, 2005]. Even for the doped samples the absolute errors obtained assuming six-fold
coordination were <1% thereby proving the usefulness of the empirical relations devised
by Jiang et al. and Rick Ubic. From these tables, it is seen that the pc lattice parameter
reported in Table 4.6 were on an average higher than that reported in Table 4.5. This may
be due to the presence of higher % tetragonality in most cases for the powders
synthesized using span 20 for all six dopants and all different concentrations used as
compared to those synthesized using span 80. However, the average |%error| calculated
with respect to the lattice parameter devised by Jiang et al. [2006] was also higher in
powders synthesized using span 20 (0.2360%) than in those synthesized using span 80
(0.1923%). The lattice parameter values, similar to those obtained for Span 80, had a
better fit with the equation devised by Jiang et al.
4.3 FTIR Analysis
Figure 4.10 shows the FTIR pattern for pure BaTiO3 powder synthesized using 1:3 sol:
support solvent with 5% span80 calcined at 500°C, 750°C and 1000°C. The spectra of the
powder calcined at 500°C showed a broad absorption band from 300-900 cm-1attributed
96
Fig
4.1
0. B
aTiO
3 po
wde
r sy
nthe
size
d us
ing
5% s
pan
80 w
ith
vary
ing
calc
inat
ions
tem
pera
ture
at
500 o
C, 7
50 o
C a
nd 1
000 o
C.
4000
3500
3000
2500
2000
1500
1000
500
0
10
20
30
40
50
60
70
80
CO
3
2-
CO
3
2-CO
2
O-H
O-H
O-H
O-H
CO
3
2-
CO
3
2-
CO
3
2- Ti-
O
Transmittance (%)
Waven
um
ber
(cm
-1)
50
00C
75
00C
10
00
0C
97
to the Ti-O vibration, the carbonate peaks (CO32-) at 858 cm-1, 1398 cm-1, 1447 cm-1,
1749 cm-1, and 2453cm-1 and the hydroxyl (O-H) group at 1060 cm-1, 1627 cm-1, 3167
cm-1, 3360 cm-1 and 3600 cm-1. Bands corresponding to the carbonate and hydroxyl
group were noted due to the presence of intermediate BaCO3 phase and water. Peaks
corresponding to carbonate (CO32-) bands have been noted at 856cm-1 (out of plane
deformation), 1382cm-1(vibration), 1430 cm-1 (symmetric stretching), 1450cm-
1(asymmetric stretching), 1750cm-1(symmetric stretching) and 2450cm-1 (combined
symmetric and asymmetric stretching) [Wei et al., 2008; Takeuchi et al., 1997; López et
al., 1999; Kuo et al., 1996; Durán et al., 2001; Un-Yeon et al., 2004; Song et al., 2000].
With increasing calcination temperature from 500°C to 750°C to 1000°C, the 538 cm-1
Ti-O absorption peak grew narrower and sharper. The intensity of carbonate and
hydroxyl peaks also reduced with increasing calcinations temperature, indicating a
decrease in the impurities. The bands in the TiO6 octahedra represent the O-Ti-O bending
(390 cm-1) and Ti-O stretching (538 cm-1, 630cm-1) [Jin et al., 2009; Arya et al., 2003;
Durán et al., 2001; Takeuchi et al., 1997; Wei et al., 2008; Sun et al., 2007; Un-Yeon et
al., 2004]. Tian et al. [2000] also observed narrowing and sharpening of the Ti-O bands
with increasing temperature which they associated with the decreasing number of
vibrational molecules as the amorphous TiO2 transformed into the TiO6 octahedra.
Similar behavior of decreasing carbonate, hydroxyl peak intensities and increasing
sharpness of Ti-O peaks at higher temperature was also achieved for powders synthesized
using span 20. CO2 adsorbed on a metal cation is known to produce a peak at 2340 cm-1
[López et al., 1999]. Hydroxyl (O-H) bands are associated to the peaks at 1050 cm-1
(rotatory mode of H2O), 1600 cm-1 (bending of H2O) and also to the bands in the range of
3220-3097 cm-1 (stretching vibration of OH) and 3400-3600 cm-1 (stretching vibration of
OH and H2O) [Wei et al., 2008; Takeuchi et al., 1997; Asiaie et al., 1996; Mohammad et
al., 2009; López et al., 1999; Arya et al., 2003; Radonji! et al., 2008; Mao et al., 2007;
Polotai et al., 2004].
Figure 4.11 shows the FTIR spectra for doped BaTiO3 with varying dopant concentration
in powders synthesized using span 80 calcined at 750°C. The spectra observed for doped
98
Fig
4.1
1. F
TIR
spe
ctru
m o
f B
a 1-x
DxT
iO3
synt
hesi
zed
usin
g 5%
spa
n 80
for
the
pow
der
calc
ined
at
750 o
C w
ith
vary
ing
x as
0 (
),
0.0
01 (
)
, 0.0
1 (
)
and
0.1
(
).
4000
3500
3000
2500
2000
1500
1000
500
0
20
40
60
80
100
120
140
160
CO
2
K
O-H
O-H
CO
3
2-
CO
3
2-
O-H
CO
32-
O-H
CO
3
2- T
i-O
Wave n
um
ber
(cm
-1)
Transmittance (%)
4000
3500
3000
2500
2000
1500
1000
500
020406080100
120
140
160
CO
2
Li
O-H
O-H
CO
32-
CO
32-O-H
CO
32-
O-H
CO
32-
Ti-
O
Transmittance (%)
Wav
enu
mb
er (
cm-1)
4000
3500
3000
2500
2000
1500
1000
500
020406080100
120
140
160
CO
2
Mg
CO
32-
CO
32-
CO
3
2-
O-H
O-H
O-H
O-H
CO
3
2-
Ti-
O
Transmittance (%)
Wave
nu
mb
er
(cm
-1)
4000
3500
3000
2500
2000
1500
1000
500
0
20
40
60
80
100
120
140
160
Ce
O-H
O-H
CO
3
2-
CO
2
O-H
CO
3
2-
CO
3
2-
O-H
CO
3
2-
Ti-
O
Transmittance (%)
Wave
nu
mb
er
(cm
-1)
4000
3500
3000
2500
2000
1500
1000
500
0
20
40
60
80
100
120
140
160
La
CO
3
2-
O-H
O-H
O-H
O-H
CO
3
2-
CO
3
2-O
-H
CO
3
2- Ti-
O
% Transmittance
Wave n
um
ber
(cm
-1)
4000
3500
3000
2500
200
01500
100
0500
020406080
100
120
140
160
Sr
CO
32-
O-H
O-H
CO
2
CO
3
2-C
O3
2-
O-H
Ti-
O
CO
3
2-
O-H
Transmittance (%)
Wav
e n
um
ber
(cm
-1)
99
BaTiO3 was in general similar to that of undoped BaTiO3 (Figure 4.10). The Ti-O
characteristic absorption peak shifted towards higher wave number from the value noted
for pure BaTiO3for all dopant concentrations with the maximum shift noted at 547cm-1 in
Sr doped BaTiO3 for x=0.01, at 551cm-1 for x=0.01 in La doped BaTiO3 and at 547cm-1
for x=0.001 in Li doped BaTiO3. However, for Ce, Mg and K doped BaTiO3 with low
doping concentration of x=0.001, the Ti-O absorption peak shifted to a higher wave
number noted at 543cm-1 for Ce, Mg and at 547cm-1 for K as the dopant, which then
remained unchanged with increasing dopant concentration. It is known that dopant
addition in pure BaTiO3 changes the cell size, crystal structure and binding energy, which
strongly correlate with the spectral vibrations. When a smaller size ion replaces Ba2+ the
cell size decreases, i.e. a decrease in lattice spacing results in a decrease of the Ti-O bond
length [Jin et al., 2009; Sun et al., 2007]. Reduction of the Ti-O distance increases the
force constant between these ions increasing the strength of the Ti-O bond. That is, with a
decrease in the cell size, the binding energy increases, shifting the wave number of the
absorption peak to the larger value [Jin et al., 2009; Sun et al., 2007; Radonji! et al.,
2008]. With a change in the valence of the dopant ion replacing Ba2+, the electric charge
balance needs to be maintained. For example, with Li1+ and K1+ doping Ba2+, one oxygen
vacancy is produced to maintain the electric charge balance [Sun et al., 2007]. Dopant
ions used in the present work produces powders with reduced lattice spacing as well as
increased cubicity with increasing dopant concentration as discussed in section 4.2. These
ions, being smaller in size than Ba as seen from Table 4.2 except for K which is almost
the same as that of Ba, are expected to produce a higher wave number shift which is
observed in Figure 4.11. It is known that when dopant ion replaces Ti in the BaTiO3
crystal, wave number of the absorption peak shifts towards lower value due to easy
distortion of the TiO6 octahedra [Sun et al., 2007]. This was not observed in the present
study suggesting that the dopant ion does not replace Ti ion.
Figure 4.12 shows the FTIR spectra for doped BaTiO3 with varying dopant concentration
in powders synthesized using span 20, calcined at 750°C. The Ti-O peak shift was in
general towards higher wave number in doped BaTiO3 with respect to the undoped
100
Fig
4.1
2. F
TIR
spe
ctru
m o
f B
a 1-x
DxT
iO3
synt
hesi
zed
usin
g 5%
spa
n 20
for
the
pow
der
calc
ined
at
750 o
C w
ith
vary
ing
x as
0 (
),
0.0
01 (
)
, 0.0
1 (
)
and
0.1
(
).
4000
3500
3000
2500
2000
1500
1000
500
0
20
40
60
80
100
120
140
160
180
200
O-H
CO
3
2-
KC
O2
O-H
O-H
O-H
CO
3
2-
CO
3
2-
CO
3
2- Ti-
O
Transmittance (%)
Waven
um
ber
(cm
-1)
4000
3500
3000
2500
2000
1500
1000
500
0
20
40
60
80
100
120
140
160
180
200
CO
3
2-
Li
O-H
O-H
O-H
CO
2
O-H
CO
3
2-
CO
3
2-
CO
3
2- Ti-
O
Waven
um
ber
(cm
-1)
Transmittance (%)
4000
3500
3000
2500
2000
1500
1000
500
0
50
100
150
200
250
O-H
CO
3
2-
Mg
O-H
O-H
O-H
CO
2
CO
3
2-
CO
3
2-
CO
3
2- Ti-
O
Transmittance (%)
Waven
um
ber
(cm
-1)
4000
3500
3000
2500
2000
1500
1000
500
0
50
100
150
200
250
O-H
O-H
CO
3
2-
Ce
O-H
O-H
CO
2
CO
3
2-
CO
3
2-C
O3
2- Ti-
O
Transmittance (%)
Wave
nu
mb
er
(cm
-1)
4000
3500
3000
2500
2000
1500
1000
500
0
20
40
60
80
100
120
140
160
180
200
La
CO
3
2-O
-H
O-H
CO
3
2-
Ti-
O
CO
3
2-
CO
3
2-
O-H
CO
2
O-H
Transmittance (%)
Wave
nu
mb
er
(cm
-1)
4000
3500
3000
2500
2000
1500
1000
500
0
20
40
60
80
100
120
140
160
180
200
Sr
CO
3
2-
O-H
O-H
CO
2
O-H
O-H
CO
3
2- C
O3
2-
CO
3
2- T
i-O
Transmittance (%)
Waven
um
ber
(cm
-1)
101
BaTiO3 for all dopants. The wave number values for this peak for x=0.1 were noted at
556cm-1 in Sr doped BaTiO3, 555cm-1 in La and Ce doped BaTiO3, 559cm-1 in Mg doped
BaTiO3 and 543cm-1 in K doped BaTiO3. Though the powders synthesized using span 20
facilitated in retaining tetragonality (section 4.2), the decreased lattice spacing along with
an increase in cubicity with increasing dopant concentration must have caused the higher
wave number shift of the Ti-O absorption peak. However, the lattice spacing in Ce doped
BaTiO3 powders synthesized using span 20 showed a small increase in lattice spacing
(table 4.4) which may be due to existence of some amount of Ce4+ replacing Ti leading to
c-axis elongation. Nevertheless, since the Ti-O absorption peak did not produce a
significant shift towards lower wave number, it is possible that the amount of Ce4+ that
replaces Ti must be very small.
4.4 TEM, SEM and EDS Analysis
Figure 4.13 shows TEM micrographs of the powders synthesized using span 80 and span
20 calcined at 750°C for different surfactant concentrations. For powders synthesized
using span 80, with surfactant concentration as 5vol% (Figure 4.13a), 10vol% (Figure
4.13b), 15vol% (Figure 4.13c) and 20vol% (Figure 4.13d), the particle were noted to
have irregular shapes, making it difficult to assign a specific shape to them. For powders
synthesized using 20vol% span 20 (Figure 4.13e), the rod shape formation was observed
along with the presence of irregularly shaped particles. Average particle size calculated
was 57nm for powders synthesized using 5 vol% span 80, 97nm with 10 vol% span 80,
91nm with 15 vol% span 80, 77nm with 20 vol% span 80 and 66nm with 20 vol% span
20.
Figure 4.14 shows the SEM micrographs of the powders calcined at 750oC, for the two
surfactants with varying concentrations. For Span 80, with the concentration varying as
5vol% (Figure 4.14a), 10vol% (Figure 4.14b) and 15vol% (Figure 4.14c), the particle
shapes changed from spherical to less spherical in nature with some irregular structures.
102
Fig 4.13. TEM-BaTiO3 morphology for powder calcined at 750°C using, a) 5% span 80,
b) 10% span 80, c) 15% span 80, d) 20% span 80, and e) 20% span 20.
103
Fig 4.14. SEM-BaTiO3 morphology for powder calcined at 750°C using, a) 5% span 80,
b) 10% span 80, c) 15% span 80, d) 20% span 80, and e) 20% span 20.
a b
cd
e
104
For 20vol% (Figure 4.14d) concentration, very small regular spherical particles were
dispersed with more irregular shaped agglomerated particles. The SEM micrograph of
powders prepared with span 20 (Figure 4.14e) showed presence of elongated structures
instead of spheres. These particles were either rectangular or rod like in nature. Since the
number of grains used to calculate the average particle size was very small in the TEM
micrographs, the average particle size used in the analysis during the present work was
estimated from the SEM micrographs where a larger numbers of grains could be counted
for better averaging.
Micelle shape depends on the concentration of the surfactant used and its HLB value
gives information about its solubility (section 2.7) [Schramm, 2000]. For non-ionic
surfactants, HLB value is also related to the CMC and the critical packing parameter
(CPP) [Zhang et al., 2003; Inderjit, 2004; Schmidts et al., 2009]. CMC increased and CPP
decreased with increasing HLB value. CPP value of span 80 is greater than 1 [Schmidts
et al., 2009]. Span 20 having a higher HLB value (8.6) than span 80 (4.3), will have a
correspondingly smaller CPP value. For a normal micelle the structure is reported to
change from spherical to lamellar with increasing CPP, however, for the reverse micelle,
a higher CPP promotes the formation of spherical structure [Zhang et al., 2003]. Hence,
with comparatively lower CPP than span 80, span 20 surfactant concentration used is
most probably not enough to form spherical structure and hence formed the rod-like
micelles, whereas the concentration of span 80 used is enough for the spherical reverse
micelle formation [Inderjit, 2004; Schramm, 2000; Zhang et al., 2003; Schmidts et al.,
2009]. These micelles act as templates for BaTiO3 particles that take the same shape.
The size of particles as a function of surfactant and its concentration was calculated from
Figure 4.14. The values reported here were calculated using 50 grains which included the
smallest and the largest particles observed in each micrograph. For 5vol% span 80 most
of the particles were found to be in 30 to 70nm size range with a few in the range of 90-
100nm, the average of which was calculated to be 57nm. Similarly, for 10vol% span 80
most of the particles were in the 60-80nm range with a few about 120nm range having an
105
average of 92nm. 15vol% span 80 particles were equally distributed from 40nm to 140nm
having an average size of 94nm. Size range of particles decreased when span 80
concentration increased to 20vol%, and was found to be distributed from 15nm to 80nm
averaging out as 42nm. However, few particles in this composition were found to be
above 100nm which when considered increased the average particle size of the powders
to 70nm. Hence 5% surfactant was chosen as the surfactant concentration for the
synthesis of doped powder. With 20vol% span 20 (Figure 4.14e), the particle size ranged
from 18 to 125nm measured along its width averaging to 66nm. Figure 4.15a and b show
the SEM micrographs of BaTiO3 pellets sintered at 750oC with a soaking time of 1 hr and
12 hr respectively. Figure 4.15c shows the micrograph of the pellet prepared from the
same powder but sintered at 1200oC for 2hr. The micrographs showed grain growth and
formation of neck between adjacent grains, as expected with increase in sintering
temperature and soaking time [Kim, 2002]. The EDS spectrum in Figure 4.15d,
confirmed the presence of Ba, Ti and O in the synthesized material. The morphology of
doped powders was similar to that of the undoped, without showing any significant
change. Hence, the characteristic micrographs are only shown here. Table 4.7 compares
the particle size with the crystallite size of pure BaTiO3 powders synthesized using span
80, calcined at 750°C.
4.5 Resistivity Analysis
Figure 4.16 shows positive temperature coefficient of resistivity (PTCR) effect for
BaTiO3 pellets synthesized using 5% span 80 and 5% span 20, sintered at 750oC.
Resistivity measurement was carried out at ambient pressure of 1 atm. during the heating
and cooling cycle on the sintered pellets. Figure 4.16 shows the Tc of the pellets to be
approximately 75ûC, as calculated from the variation of slope with temperature (first
maximum in slope) of the resistivity vs. temperature curve [Zubeda and Sutapa, 2012].
Resistivity values changed from ~ 108 "cm to 1010 "cm for the powders synthesized
using span 80 producing a PTCR jump of 2 orders. PTCR jump is defined as
106
Fig 4.15. Particle morphology of the pellets (5% span 80 powder) sintered at (a) 750°C
for 1h soaking, (b) 750°C for 12h soaking and (c) 1200°C for 2h soaking (d) EDS
spectrum (750°C for 1h soaking).
Table 4.7. Average particle size via SEM and crystallite size via XRD of 750°C calcined
pure BaTiO3 powders synthesized as a function of span 80 concentration.
Span 80
Concentration
Particle size
(nm)
Crystallite size
(nm)
5% 57 27.93
10% 92 29.85
15% 94 30.37
20% 70 33.03
d
a b
c
107
Fig
4.1
6. P
TC
R e
ffec
t in
BaT
iO3 he
at t
reat
ed a
t 75
0°C
syn
thes
ized
usi
ng 5
% (
a) s
pan
80 a
nd (
b) s
pan
20 d
urin
g th
e he
atin
g
and
cool
ing
cycl
e.
3050
7090
110
130
150
170
3.0x
107
2.0x
109
4.0x
109
6.0x
109
8.0x
109
1.0x
1010
1.2x
1010
1.4x
1010
1.0x
1010
1.5x
1010
2.0x
1010
2.5x
1010
3.0x
1010
3.5x
1010
4.0x
1010
!C
("cm
)!
H
("cm
)
Tem
per
atu
re (
0 C)
3050
7090
110
130
150
170
1.0x
106
2.0x
109
4.0x
109
6.0x
109
8.0x
109
1.0x
1010
1.2x
1010
1.4x
1010
9.0x
109
1.2x
1010
1.5x
1010
1.8x
1010
2.1x
1010
2.4x
1010
2.7x
1010
!H
("cm
)
!C
("cm
)
Tem
per
atu
re (
0 C)
108
R = #max/#min [Viviani et al., 2004]. For powders synthesized using span 20 the resistivity
values changed from ~ 107 "cm to 1010 "cm with a 3 order PTCR jump. Though the
resistivity value of undoped BaTiO3 decreased from its otherwise reported value of >109
"cm, the variation of surfactant did not produce a remarkable change in the resistivity
values and were observed to be in the similar range. Values recorded during cooling
cycle had resistivity values higher than that recorded during heating cycle (Figure 4.16), a
trend similar to that observed by J. -G. Kim. The onset of decrease in resistivity at high
temperatures upon heating was considered to be due to the decrease in potential barrier
height that resulted from desorption of chemisorbed oxygen atoms at the grain
boundaries. This corresponds to a peak in the resistivity measurement. Higher resistivity
during cooling cycle compared to the heating cycle was associated with the increase in
potential barrier due to the decrease in number of conduction electrons because of
adsorbed oxygen at grain boundaries [Kim, 2002; Huybrechts et al., 1995]. Shift of
resistivity maximum towards lower temperature during the cooling cycle indicated
hysteresis of the material. This is due to the time taken for the conduction electrons to be
released so as to reduce the potential barrier height. More oxygen adsorb at the grain
boundaries up to the point at which the acceptor level reach the Fermi level releasing the
conduction electrons lowering the potential barrier height [Kim, 2002; Huybrechts et al.,
1995].
Pure BaTiO3 in its bulk form is usually an insulator with resistivity >109"cm with no
PTCR effect showing almost temperature independent resistivity except a small change
in the resistivity noted at the TC [Masó et al., 2008; Huybrechts et al., 1995; Kim, 2002;
Panwar and Semwal, 1991; Heywang, 1971; Darko et al., 2003; Gheno et al., 2007].
Goodman observed a single crystal Sm-doped BaTiO3 to show temperature independent
resistivity with a small increase only at the TC, an effect similar to that observed for
undoped BaTiO3 [Goodman, 1963]. However, when the same material was prepared in
the powdered polycrystalline form, it exhibited PTCR behavior. PTCR effect is therefore
known to be a grain boundary phenomenon, where the grain boundary can be considered
as a low-permittivity non-ferroelectric phase [Goodman, 1963; Huybrechts et al., 1995;
109
Amin, 1994; Panwar and Semwal, 1991; Gheno et al., 2007; Darko et al., 2003; Xue et
al., 2007]. Materials exhibiting PTCR effect shows an abrupt increase in resistivity on
reaching its Curie temperature (Tc). PTCR effect in BaTiO3 is explained using the widely
accepted Heywang model above Tc and Jonker model below Tc, discussed below
[Heywang, 1971; Huybrechts et al., 1995; Huybrechts et al., 1993; Huybrechts B. et al.,
1993; Panwar and Semwal, 1991; Al-Allak et al., 1988; Hishita et al., 1990; Kim, 2002].
In the ferroelectric phase below Tc BaTiO3 possess spontaneous polarization [Makoto and
kouichi, 2003; Moulson and Herbert, 2003]. Neighboring grains have different crystal
orientation due to which the polarization direction differ from one grain to another
[Jonker, 1983; Huybrechts et al., 1995; Al-Allak et al., 1988; Hishita et al., 1990; Panwar
and Semwal, 1991; Amin, 1994]. This results in a net polarization perpendicular to the
grain boundaries, at which surface charges are created. Jonker roughly estimated half of
the grain boundaries to be positively charged and half to be negatively charged. Grain
boundary area with negative surface charges will either have a partially or a completely
filled depletion layer, decreasing the potential barrier. Positively charged grain boundary
area increases the potential barrier, but is insignificant as the conduction electrons always
follow the lowest barrier path [Jonker, 1983; Huybrechts et al., 1995; Al-Allak et al.,
1988; Hishita et al., 1990; Panwar and Semwal, 1991; Amin, 1994]. Hence a decrease of
resistivity with increasing temperature is observed in the ferroelectric phase.
Acceptor states along the grain boundaries act as electron traps that attract electrons from
the bulk creating an electron depletion layer within the bulk of the grain which results in
the grain boundary barrier (HI) given as [Heywang, 1961; Huybrechts et al., 1995;
Huybrechts et al., 1993; Kim, 2002; Amin, 1994; Huybrechts B. et al., 1993; Chatterjee
et al., 1999; Srimala et al., 2008; Pavlovi! et al., 2002]
HI +# J#KLM1N2O#PQPRS1N2KT
---(4.5)
where, e is the electron charge, Ns is the density of trapped electrons at the grain
boundaries, Nd the charge carrier concentration, $o permittivity of free space and $gb
110
relative permittivity of the grain boundary region. At Tc electron traps are below Fermi
level and the electron trap density equals the density of trapped electrons at the grain
boundary, because all electron traps are filled. Just above Tc BaTiO3 changes to
paraelectric phase, where $gb decreases. Decrease in $gb increases HI which results in a
rapid increase of resistivity (U) given as [Heywang, 1961; Huybrechts et al., 1995;
Huybrechts et al., 1993; Hishita et al., 1990; Al-Allak et al., 1988; Panwar and Semwal,
1991; Amin, 1994; Huybrechts B. et al., 1993]
U + V#WXY ZJ#[QDN \ ---(4.6)
where, A is a constant, T the temperature and ] the Boltzmann constant. Increase in HI
prevails until the acceptor states reaches the Fermi level at which the density of trapped
electrons decreases due to reemission of electrons in to the bulk, resulting in the
maximum value of resistivity [Heywang, 1971]. The resistivity then starts decreasing,
resulting in the negative temperature coefficient of (NTC) resistivity [Huybrechts B. et
al., 1993; Huybrechts et al., 1993; Panwar and Semwal, 1991].
BaTiO3 deviates from its insulating behavior to semiconducting, either when (1) it is
doped with donors, (2) when gases are adsorbed on the surface of the grains or (3) when
cation vacancies are created [Huybrechts et al., 1995]. 3d-element doping is also known
to influence the PTCR effect [Huybrechts et al., 1995; Huybrechts et al., 1993]. A
decrease in Tc from the value of 120oC is also observed for BaTiO3 crystals with a
decrease in particle size [Kareiva et al., 1999]. Annealing and cooling in an oxidative
atmosphere results in oxidation of the GB without oxidizing the grains [Huybrechts et al.,
1995; Huybrechts B. et al., 1993]. Oxidation of GB leads to higher electron trap densities
which produce a notable change in the PTCR effect by formation of either (1) barium
vacancies (VýBa), (2) oxidizing 3d-elements to a higher oxidation state, or (3) oxygen
adsorption [Huybrechts et al., 1995; Masó et al., 2008; Kim, 2002]. VýBa produce a small
change in the PTCR properties as compared to the larger change produced by oxygen
vacancies (VúúO). This is due to the slower diffusion rate of VýBa from the grain boundary
towards the grain and higher diffusion rates of VúúO from the grain towards the grain
boundary [Huybrechts et al., 1995; Huybrechts B. et al., 1993]. Moreover, the size range
111
of the particle makes the material predominantly cubic at room temperature. This also
contributes towards the PTCR behavior as it is associated with the presence of
paraelectric cubic phase [Hishita et al., 1990; Yoon et al., 2003; Al-Allak et al., 1988].
Cation vacancies if present during sintering can also be a contributing factor towards the
observed PTCR characteristics [Masó et al., 2008]. However, this is most likely not the
case as the diffusion of VýBa below 800°C is negligible [Huybrechts et al., 1995]. In the
present work, the PTCR effect of the undoped BaTiO3 can be associated with oxygen
adsorption and the smaller grain size which increase the GB area.
Figures 4.17 and 4.18 shows resistivity behavior of Ba1-xDxTiO3 powders (D = Sr, La,
Ce, Mg, Li and K) with varying dopant concentration synthesized using span 80 and span
20 respectively. Measurement conditions were same as that maintained during the
analysis of undoped BaTiO3. Tables 4.8-4.11 reports the parameters deduced from the
resistivity-temperature profile for these powders. Similar to undoped BaTiO3, these
powders also showed higher resistivity during cooling cycle than that measured during
heating cycle irrespective of the dopant used.
To induce semiconductivity and maximize PTCR effect, BaTiO3 is specifically doped
with donors (higher valence ion) e.g. Ba2+ with La3+ or Ti4+ with Nb5+ whereas to shift the
Curie temperature towards the higher or lower value it is doped with isovalent ions e.g.
Ba2+ with Sr 2+ or Ti4+ with Zr4+ [Heywang, 1971; Huybrechts et al., 1995; Amin, 1994;
Darko et al., 2003; Kim, 2002; Panwar and Semwal, 1991; Slipenyuk et al., 2003]. Each
PTCR BaTiO3 grain possesses heterogeneous electrical properties due to the presence of
a core of pure ferroelectric BaTiO3 surrounded by a paraelectric shell containing the
dopant [Fiorenza et al., 2009, Yoon et al., 2003]. However, in the present work, PTCR
effect is not only obtained for doped BaTiO3 but also for the undoped BaTiO3. Here the
sintering was carried out at ambient pressure in atmospheric air. Doping modified the
crystal structure due to difference in the ionic radii of the dopant and the ion being
replaced as seen from the XRD studies (section 4.2).
112
Fig
4.1
7.a
. Res
isti
vity
pro
file
of
Ba 1
-xS
r xT
iO3
synt
hesi
zed
usin
g 5%
spa
n80
for
the
pow
der
calc
ined
at
750 o
C w
ith
vary
ing
x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
2.0
x10
7
3.0
x10
9
6.0
x10
9
9.0
x10
9
1.2
x10
10
1.5
x10
10
1.8
x10
10
1.0
x10
10
6.0
x10
10
1.1
x10
11
1.6
x10
11
2.1
x10
11
2.6
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
2.0
x10
7
3.0
x10
9
6.0
x10
9
9.0
x10
9
1.2
x10
10
1.5
x10
10
1.8
x10
10
8.0
x10
9
2.8
x10
10
4.8
x10
10
6.8
x10
10
8.8
x10
10
Tem
pera
ture
(0C
)
!H
("cm
)!
C
("cm
)
30
50
70
90
110
130
150
170
1.0
x10
7
2.1
x10
8
4.1
x10
8
6.1
x10
8
8.1
x10
8
1.0
x10
9
1.2
x10
9
1.0
x10
9
1.5
x10
9
2.0
x10
9
2.5
x10
9
3.0
x10
9
3.5
x10
9
!H
("cm
)
Tem
pera
ture
(0C
)
!C
("cm
)
30
50
70
90
110
130
150
170
1x10
7
1x10
8
2x10
8
3x10
8
4x10
8
1x10
9
2x10
9
3x10
9
4x10
9
5x10
9
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
113
Fig
4.1
7.b
. Res
isti
vity
pro
file
of
Ba 1
-xL
a xT
iO3
synt
hesi
zed
usin
g 5%
spa
n80
for
the
pow
der
calc
ined
at
750 o
C w
ith
vary
ing
x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
1.0
x10
8
3.1
x10
9
6.1
x10
9
9.1
x10
9
1.2
x10
10
1.5
x10
10
1.8
x10
10
1.5
x10
10
2.0
x10
10
2.5
x10
10
3.0
x10
10
3.5
x10
10
4.0
x10
10
Tem
pera
ture
(0C
)
!H
("cm
)!
C
("cm
)
30
50
70
90
110
130
150
6x10
7
1x10
10
2x10
10
3x10
10
4x10
10
1.0
x10
10
3.0
x10
10
5.0
x10
10
7.0
x10
10
9.0
x10
10
1.1
x10
11
1.3
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
6x10
8
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
6x10
10
1x10
10
1x10
11
2x10
11
3x10
11
4x10
11
5x10
11
6x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1x10
7
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
1x10
11
2x10
11
3x10
11
4x10
11
5x10
11
6x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
114
Fig
4.1
7.c
. Res
isti
vity
pro
file
of
Ba 1
-xC
e xT
iO3
synt
hesi
zed
usin
g 5%
spa
n80
for
the
pow
der
calc
ined
at
750 o
C w
ith
vary
ing
x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
2.0
x10
8
2.0
x10
10
4.0
x10
10
6.0
x10
10
8.0
x10
10
1.0
x10
11
1.2
x10
11
1.4
x10
11
5.0
x10
10
2.5
x10
11
4.5
x10
11
6.5
x10
11
8.5
x10
11
1.1
x10
12
1.3
x10
12
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
3.0
x10
8
2.0
x10
10
4.0
x10
10
6.0
x10
10
8.0
x10
10
5.0
x10
10
1.0
x10
11
1.5
x10
11
2.0
x10
11
2.5
x10
11
3.0
x10
11
3.5
x10
11
4.0
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1.0
x10
8
5.1
x10
9
1.0
x10
10
1.5
x10
10
2.0
x10
10
2.5
x10
10
3.0
x10
10
3.0
x10
10
8.0
x10
10
1.3
x10
11
1.8
x10
11
2.3
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
3x10
7
2x10
9
4x10
9
6x10
9
8x10
9
1x10
10
5x10
9
3x10
10
5x10
10
7x10
10
9x10
10
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
115
Fig
4.1
7.d
. Res
isti
vity
pro
file
of
Ba 1
-xM
g xT
iO3
synt
hesi
zed
usin
g 5%
spa
n80
for
the
pow
der
calc
ined
at
750 o
C w
ith
vary
ing
x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
3x10
7
1x10
10
2x10
10
3x10
10
4x10
10
2.0
x10
10
7.0
x10
10
1.2
x10
11
1.7
x10
11
2.2
x10
11
2.7
x10
11
3.2
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
5x10
7
1x10
10
2x10
10
3x10
10
4x10
10
2.0
x10
10
7.0
x10
10
1.2
x10
11
1.7
x10
11
2.2
x10
11
2.7
x10
11
!C
("cm
)
!H
("cm
)
Tem
pera
ture
(0C
)
30
50
70
90
110
130
150
170
8.0
x10
7
2.0
x10
10
4.0
x10
10
6.0
x10
10
8.0
x10
10
4x10
10
1x10
11
2x10
11
3x10
11
4x10
11
!C
("cm
)
Tem
pera
ture
(0C
)
!H
("cm
)
30
50
70
90
110
130
150
170
7.0
x10
7
2.0
x10
10
4.0
x10
10
6.0
x10
10
8.0
x10
10
1.0
x10
11
1.2
x10
11
1.4
x10
11
5x10
10
2x10
11
3x10
11
4x10
11
5x10
11
6x10
11
7x10
11
8x10
11
Tem
pera
ture
(0C
)
!C
("cm
)
!H
("cm
)
116
Fig
4.1
7.e
. Res
isti
vity
pro
file
of
Ba 1
-xL
i xT
iO3
synt
hesi
zed
usin
g 5%
spa
n80
for
the
pow
der
calc
ined
at
750 o
C w
ith
vary
ing
x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
9x10
6
1x10
9
2x10
9
3x10
9
4x10
9
5x10
9
1.0
x10
9
3.0
x10
9
5.0
x10
9
7.0
x10
9
9.0
x10
9
1.1
x10
10
1.3
x10
10
1.5
x10
10
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1.0
x10
7
5.0
x10
9
1.0
x10
10
1.5
x10
10
2.0
x10
10
2.5
x10
10
3.0
x10
10
3.5
x10
10
2.0
x10
10
7.0
x10
10
1.2
x10
11
1.7
x10
11
2.2
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
6.0
x10
7
5.0
x10
10
1.0
x10
11
1.5
x10
11
2.0
x10
11
5.0
x10
10
2.5
x10
11
4.5
x10
11
6.5
x10
11
8.5
x10
11
1.1
x10
12
1.3
x10
12
1.5
x10
12
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
3x10
8
1x10
11
2x10
11
3x10
11
4x10
11
5x10
11
6x10
11
1.0
x10
11
6.0
x10
11
1.1
x10
12
1.6
x10
12
2.1
x10
12
2.6
x10
12
3.1
x10
12
Tem
pera
ture
(0C
)
!C
("cm
)!
H
("cm
)
117
Fig
4.1
7.f
. R
esis
tivi
ty p
rofi
le o
f B
a 1-x
KxT
iO3
synt
hesi
zed
usin
g 5%
spa
n80
for
the
pow
der
calc
ined
at
750 o
C w
ith
vary
ing
x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
30
50
70
90
110
130
150
170
2x10
7
1x10
9
2x10
9
3x10
9
4x10
9
5x10
9
6x10
9
3x10
9
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
6x10
10
!C
("cm
)
!
("cm
)
Tem
pera
ture
(0C
)
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
1x10
7
1x10
8
2x10
8
3x10
8
4x10
8
5x10
8
6x10
8
5x10
8
2x10
9
3x10
9
4x10
9
5x10
9
6x10
9
!
("cm
)
Tem
pera
ture
(0C
)
!C
("cm
)
30
50
70
90
110
130
150
170
5.0
x10
7
2.0
x10
10
4.0
x10
10
6.0
x10
10
8.0
x10
10
1.0
x10
11
1.2
x10
11
4x10
10
1x10
11
2x10
11
3x10
11
4x10
11
5x10
11
6x10
11
7x10
11
8x10
11
!
("cm
)
Tem
pera
ture
(0C
)
!C
("cm
)
30
50
70
90
110
130
150
170
1.0
x10
8
5.0
x10
10
1.0
x10
11
1.5
x10
11
2.0
x10
11
7.0
x10
10
2.7
x10
11
4.7
x10
11
6.7
x10
11
8.7
x10
11
1.1
x10
12
1.3
x10
12
!
("cm
)
Tem
pera
ture
(0C
)
!C
("cm
)
118
Table 4.8. Resistivity values of Ba1-xDxTiO3 pellets during the heating cycle with varying
D and x in powders synthesized using span 80.
D x 0.001 0.01 0.05 0.1
Sr
#30° ("-cm) 1.70x107 1.58x107 2.59x107 2.32x107
TC (°C) 65 55 75 75
#max ("-cm) 7.80x108 1.21x109 1.72x1010 1.79x1010
Tmax (°C) 170 170 145 150
PTCR jump 1 2 3 3
La
#30° ("-cm) 3.06x107 6.56x108 6.42x107 3.36x107
TC (°C) 85 70 70 75
#max ("-cm) 5.42x1010 5.60x1010 4.48x1010 1.69x1010
Tmax (°C) 170 120 110 140
PTCR jump 3 2 3 3
Ce
#30° ("-cm) 4.25x107 1.19x108 3.36x108 2.44x108
TC (°C) 85 75 75 80
#max ("-cm) 1.07x1010 3.10x1010 8.25x1010 1.32x1011
Tmax (°C) 155 170 130 155
PTCR jump 3 2 2 3
Mg
#30° ("-cm) 9.03x107 9.17x107 6.46x107 3.86x107
TC (°C) 95 85 55 60
#max ("-cm) 1.39x1011 9.33x1010 4.39x1010 4.13x1010
Tmax (°C) 130 130 120 125
PTCR jump 4 3 3 3
Li
#30° ("-cm) 3.10x108 7.12x107 1.62x107 8.89x106
TC (°C) 70 90 80 85
#max ("-cm) 6.67x1011 2.27x1011 3.67x1010 4.72x109
Tmax (°C) 130 125 145 170
PTCR jump 3 4 3 3
K
#30° ("-cm) 1.32x108 7.19x107 2.04x107 1.05x107
TC (°C) 80 75 70 60
#max ("-cm) 1.97x1011 1.37x1011 6.13x109 6.28x108
Tmax (°C) 125 125 135 170
PTCR jump 3 4 2 1
119
Table 4.9. Resistivity values of Ba1-xDxTiO3 pellets during the cooling cycle with varying
D and x in powders synthesized using span 80.
D x 0.001 0.01 0.05 0.1
Sr
#max ("-cm) 4.86x109 3.41x109 9.69x1010 2.66x1011
Tmax (°C) 110 120 80 80
La
#max ("-cm) 5.23x1011 5.90x1011 1.20x1011 3.88x1010
Tmax (°C) 135 75 80 100
Ce
#max ("-cm) 9.05x1010 2.33x1011 3.97x1011 1.21x1012
Tmax (°C) 90 90 90 95
Mg
#max ("-cm) 7.17x1011 4.60x1011 2.86x1011 3.13x1011
Tmax (°C) 95 95 70 75
Li
#max ("-cm) 3.12x1012 1.33x1012 2.27x1011 1.38x1010
Tmax (°C) 95 95 95 115
K
#max ("-cm) 1.21x1012 8.29x1011 5.63x1010 5.29x109
Tmax (°C) 95 110 85 90
120
Fig
4.1
8.a
. R
esis
tivi
ty p
rofi
le o
f B
a 1-x
Sr x
TiO
3 sy
nthe
size
d us
ing
5% s
pan2
0 fo
r th
e po
wde
r ca
lcin
ed a
t 75
0 oC
wit
h va
ryin
g x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
1.0
x10
7
2.0
x10
10
4.0
x10
10
6.0
x10
10
8.0
x10
10
5.0
x10
10
2.5
x10
11
4.5
x10
11
6.5
x10
11
8.5
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1.0
x10
7
5.0
x10
9
1.0
x10
10
1.5
x10
10
2.0
x10
10
2.5
x10
10
3.0
x10
10
6.0
x10
10
9.0
x10
10
1.2
x10
11
1.5
x10
11
1.8
x10
11
2.1
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1x10
7
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
6x10
10
4x10
10
1x10
11
2x10
11
3x10
11
4x10
11
5x10
11
Tem
pera
ture
(0C
)
!H
("cm
)!
C
("cm
)
30
50
70
90
110
130
150
170
3.0
x10
7
5.0
x10
9
1.0
x10
10
1.5
x10
10
2.0
x10
10
2.5
x10
10
3.0
x10
10
3.5
x10
10
4.0
x10
10
4.5
x10
10
2.0
x10
10
7.0
x10
10
1.2
x10
11
1.7
x10
11
2.2
x10
11
2.7
x10
11
3.2
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
121
Fig
4.1
8.b
. R
esis
tivi
ty p
rofi
le o
f B
a 1-x
La x
TiO
3 sy
nthe
size
d us
ing
5% s
pan2
0 fo
r th
e po
wde
r ca
lcin
ed a
t 75
0 oC
wit
h va
ryin
g x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
5x10
7
1x10
10
2x10
10
3x10
10
4x10
10
5.0
x10
10
1.0
x10
11
1.5
x10
11
2.0
x10
11
2.5
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1x10
7
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
3.0
x10
10
6.0
x10
10
9.0
x10
10
1.2
x10
11
1.5
x10
11
1.8
x10
11
2.1
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
3.0
x10
7
2.0
x10
10
4.0
x10
10
6.0
x10
10
8.0
x10
10
1.0
x10
11
1.2
x10
11
6x10
10
2x10
11
3x10
11
4x10
11
5x10
11
6x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1x10
7
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
6x10
10
7x10
10
3x10
10
1x10
11
2x10
11
3x10
11
4x10
11
5x10
11
6x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
122
Fig
4.1
8.c
. R
esis
tivi
ty p
rofi
le o
f B
a 1-x
Ce x
TiO
3 sy
nthe
size
d us
ing
5% s
pan2
0 fo
r th
e po
wde
r ca
lcin
ed a
t 75
0 oC
wit
h va
ryin
g x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
1x10
7
1x10
10
2x10
10
3x10
10
4x10
10
5.0
x10
10
1.0
x10
11
1.5
x10
11
2.0
x10
11
2.5
x10
11
3.0
x10
11
3.5
x10
11
4.0
x10
11
!H
("cm
)
Tem
pera
ture
(0C
)
!C
("cm
)
30
50
70
90
110
130
150
170
3x10
7
2x10
10
4x10
10
6x10
10
8x10
10
1x10
11
5x10
10
3x10
11
5x10
11
7x10
11
9x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
2x10
7
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
5x10
10
2x10
11
3x10
11
4x10
11
5x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1x10
8
2x10
10
4x10
10
6x10
10
8x10
10
01x10
11
2x10
11
3x10
11
4x10
11
5x10
11
6x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
123
Fig
4.1
8.d
. R
esis
tivi
ty p
rofi
le o
f B
a 1-x
Mg x
TiO
3 sy
nthe
size
d us
ing
5% s
pan2
0 fo
r th
e po
wde
r ca
lcin
ed a
t 75
0 oC
wit
h va
ryin
g x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
7x10
7
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
6x10
10
1x10
10
1x10
11
2x10
11
3x10
11
4x10
11
5x10
11
6x10
11
7x10
11
Tem
pera
ture
(0C
)
!C
("cm
)
!H
("cm
)
30
50
70
90
110
130
150
170
2.0
x10
7
2.0
x10
9
4.0
x10
9
6.0
x10
9
8.0
x10
9
1.0
x10
10
1.2
x10
10
1.4
x10
10
1.6
x10
10
1.8
x10
10
7.0
x10
9
5.7
x10
10
1.1
x10
11
1.6
x10
11
2.1
x10
11
2.6
x10
11
3.1
x10
11
3.6
x10
11
Tem
pera
ture
(0C
)
!C
("cm
)
!H
("cm
)
30
50
70
90
110
130
150
170
5x10
7
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
5.0
x10
9
2.0
x10
11
4.0
x10
11
6.0
x10
11
8.1
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1.0
x10
8
2.0
x10
10
4.0
x10
10
6.0
x10
10
8.0
x10
10
2.0
x10
10
2.2
x10
11
4.2
x10
11
6.2
x10
11
8.2
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
124
Fig
4.1
8.e
. R
esis
tivi
ty p
rofi
le o
f B
a 1-x
Li x
TiO
3 sy
nthe
size
d us
ing
5% s
pan2
0 fo
r th
e po
wde
r ca
lcin
ed a
t 75
0 oC
wit
h va
ryin
g x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
1x10
7
1x10
9
2x10
9
3x10
9
4x10
9
5x10
9
4x10
9
1x10
10
2x10
10
3x10
10
4x10
10
5x10
10
6x10
10
7x10
10
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
5.0
x10
7
3.0
x10
9
6.0
x10
9
9.1
x10
9
1.2
x10
10
1.5
x10
10
1.0
x10
10
3.0
x10
10
5.0
x10
10
7.0
x10
10
9.0
x10
10
1.1
x10
11
Tem
pera
ture
(0C
)
!H
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
2.0
x10
7
5.0
x10
9
1.0
x10
10
1.5
x10
10
2.0
x10
10
1.0
x10
10
4.0
x10
10
7.0
x10
10
1.0
x10
11
1.3
x10
11
1.6
x10
11
1.9
x10
11
2.2
x10
11
Tem
pera
ture
(0C
)
!C
("cm
)!
H
("cm
)
30
50
70
90
110
130
150
170
4.0
x10
7
5.0
x10
9
1.0
x10
10
1.5
x10
10
2.0
x10
10
2.5
x10
10
3.0
x10
10
3.5
x10
10
2.0
x10
10
7.0
x10
10
1.2
x10
11
1.7
x10
11
2.2
x10
11
Tem
pera
ture
(0C
)
!C
("cm
)
!H
("cm
)
125
Fig
4.1
8.f
. R
esis
tivi
ty p
rofi
le o
f B
a 1-x
KxT
iO3
synt
hesi
zed
usin
g 5%
spa
n20
for
the
pow
der
calc
ined
at
750 o
C w
ith
vary
ing
x
mea
sure
d du
ring
hea
ting
(H
) an
d co
olin
g (C
) cy
cle.
x =
0.1
x =
0.0
5
x =
0.0
1x =
0.0
01
30
50
70
90
110
130
150
170
2.0
x10
7
2.2
x10
8
4.2
x10
8
6.2
x10
8
8.2
x10
8
1.0
x10
9
6.0
x10
8
5.6
x10
9
1.1
x10
10
1.6
x10
10
2.1
x10
10
2.6
x10
10
Tem
pera
ture
(0C
)
!
("cm
)
!C
("cm
)
30
50
70
90
110
130
150
170
1x10
7
1x10
8
2x10
8
3x10
8
4x10
8
5x10
8
4.0
x10
8
2.4
x10
9
4.4
x10
9
6.4
x10
9
8.4
x10
9
1.0
x10
10
1.2
x10
10
1.4
x10
10
!
("cm
)
Tem
pera
ture
(0C
)
!C
("cm
)
30
50
70
90
110
130
150
170
2.0
x10
7
3.0
x10
9
6.0
x10
9
9.0
x10
9
1.2
x10
10
1.5
x10
10
1.0
x10
10
4.0
x10
10
7.0
x10
10
1.0
x10
11
1.3
x10
11
1.6
x10
11
!
("cm
)
Tem
pera
ture
(0C
)
!C
("cm
)
30
50
70
90
110
130
150
170
1x10
7
1x10
10
2x10
10
3x10
10
4x10
10
2.0
x10
10
9.0
x10
10
1.6
x10
11
2.3
x10
11
3.0
x10
11
3.7
x10
11
!C
("cm
)
!
("cm
)
Tem
pera
ture
(0C
)
126
Table 4.10. Resistivity values of Ba1-xDxTiO3 pellets during the heating cycle with
varying D and x in powders synthesized using span 20.
D x 0.001 0.01 0.05 0.1
Sr
#30° ("-cm) 3.67x107 8.96x107 3.58x107 3.53x107
TC (°C) 75 75 75 70
#max ("-cm) 4.65x1010 6.81x1010 3.02x1010 9.38x1010
Tmax (°C) 125 140 150 135
PTCR jump 3 3 3 3
La
#30° ("-cm) 1.94x107 3.44x107 4.06x107 5.69x107
TC (°C) 95 85 90 90
#max ("-cm) 6.80x1010 1.15x1011 5.23x1010 4.17x1010
Tmax (°C) 135 135 150 150
PTCR jump 3 4 3 3
Ce
#30° ("-cm) 1.31x108 2.54x107 3.81x107 1.18x107
TC (°C) 80 90 90 80
#max ("-cm) 8.76x1010 5.29x1010 9.87x1010 4.16x1010
Tmax (°C) 130 145 135 145
PTCR jump 2 3 3 3
Mg
#30° ("-cm) 1.11x108 5.91x107 3.26x107 8.99x107
TC (°C) 70 80 90 70
#max ("-cm) 9.17x1010 5.34x1010 1.76x1010 5.62x1010
Tmax (°C) 115 105 115 115
PTCR jump 2 3 3 3
Li
#30° ("-cm) 4.12x107 2.41x107 5.44x107 1.82x107
TC (°C) 75 75 75 70
#max ("-cm) 3.71x1010 1.96x1010 1.64x1010 5.47x109
Tmax (°C) 135 145 140 140
PTCR jump 3 3 3 2
K
#30° ("-cm) 6.26x107 2.04x107 1.48x107 2.18x107
TC (°C) 80 75 65 70
#max ("-cm) 4.01x1010 1.45x1010 5.59x108 1.03x109
Tmax (°C) 140 160 140 130
PTCR jump 3 3 1 2
127
Table 4.11. Resistivity values of Ba1-xDxTiO3 pellets during the cooling cycle with
varying D and x in powders synthesized using span 20.
D x 0.001 0.01 0.05 0.1
Sr
#max ("-cm) 3.17x1011 5.16x1011 1.98x1011 8.09x1011
Tmax (°C) 90 100 100 95
La
#max ("-cm) 5.62x1011 5.19x1011 1.94x1011 2.42x1011
Tmax (°C) 95 105 120 100
Ce
#max ("-cm) 6.00x1011 4.65x1011 8.94x1011 3.81x1011
Tmax (°C) 100 100 100 100
Mg
#max ("-cm) 9.11x1011 7.88x1011 3.31x1011 6.76x1011
Tmax (°C) 85 75 80 85
Li
#max ("-cm) 2.42x1011 2.05x1011 1.11x1011 6.64x1010
Tmax (°C) 90 95 105 90
K
#max ("-cm) 3.57x1011 1.70x1011 1.35x1010 2.55x1010
Tmax (°C) 95 90 70 80
128
In powders synthesized with span 80, PTCR jump measured during the heating cycle was
1 order for x=0.001, 2 order for x=0.01 and 3 order for x=0.05 and x=0.1 Sr
concentration. In powders synthesized using span 20 however, a 3 order PTCR jump is
obtained for all concentrations of Sr. Resistivity values noted at 30°C and at Tmax were
higher in powders synthesized using span 20 than in span 80 for Sr doped BaTiO3. It was
clearly seen that resistivity maxima during heating cycle was above 120°C, but during
cooling cycle this value appeared at lower temperature indicating hysteresis of the
material. Similar to that noted for pure BaTiO3, increase in resistivity during cooling
cycle results from the adsorbed oxygen at the grain boundaries and the time taken for the
conduction electrons to be released in to the material results in the hysteresis of the
material. PTCR jump measured in La doped BaTiO3 powders synthesized using span 80
and span 20 was 3 orders for all concentrations except for x=0.01for which it showed a 2
order (span 80) and a 4 order (span 20) PTCR jump. Resistivity values noted at 30°C was
higher in powders synthesized using span 80 than in span 20 for all La concentrations
except for x=0.1. But at Tmax, the resistivity values noted were higher in powders
synthesized using span 20 than span 80 for all La concentrations. For La concentrations
at which the resistivity maxima during heating cycle was close to 135°C, during cooling
cycle this value appeared closer to 100°C and for those concentrations which showed the
resistivity maxima below 130°C, during cooling cycle it was noted below 90°C. Ce
doped BaTiO3 produced powders with 2-3 order PTCR jump. In powders synthesized
using span 80, this jump was of 2 order for x=0.01 and x=0.05, and for those using span
20, it was 2 order for x=0.001 and 3 order for all other Ce concentrations. At 30°C
resistivity values noted in powders synthesized using span 80 was higher than that in span
20 except for x=0.001 Ce concentration. At Tmax, the resistivity values noted were higher
in powders synthesized using span 20 than span 80 for all Ce concentrations except for
x=0.1. Mg doped BaTiO3 powders synthesized using span 80 and span 20 measured a 3
order PTCR jump for all Mg concentrations except for x=0.001. The PTCR jump for
x=0.001 was 4 order and 2 order for powder synthesized using span 80 and span 20
respectively. The resistivity values noted at 30°C were higher in powders synthesized
using span 80 than span 20 for Mg concentration of x=0.01 and x=0.05 for the other
129
concentrations the opposite effect was noted. However, at Tmax the resistivity values were
higher in powders synthesized using span 80 except for x=0.1 where the powders
synthesized using span 20 produced higher resistivity values. Similar to that observed
with other dopants, when the Tmax value was closer to 130°C during heating cycle, that
noted during cooling cycle was closer to 100°C, and for those below 130°C during
heating cycle, the noted Tmax during cooling cycle was below 90°C. Li doped BaTiO3
also produced powders with a PTCR jump of 3 order. The exception to this was for Li
concentration of x= 0.1 and x=0.01 powders synthesized using span 80 and span 20
respectively. For x=0.1 (span 80) a 4 order PTCR jump was noted and for x=0.01 (span
20) a 2 order PTCR jump was noted. For low Li concentration of x=0.001 and x=0.01,
the resistivity values noted at 30°C were higher for powders synthesized using span 80
than those using span 20, and for higher concentrations of x=0.05 and x=0.1 the powders
synthesized using span 20 were higher than those using span 80. At Tmax, the resistivity
value noted was higher for powder synthesized using span 80 for Li concentration of
x=0.1. For all other concentrations, higher resistivity values were noted for powders
synthesized using span 20. Tmax noted during the cooling cycle was closer to 100°C
irrespective of that noted during heating cycle. K doped BaTiO3 powders showed a
variation in the PTCR jump from 1 order to 4 orders. In the powders synthesized using
span 80 with increasing K concentration from x=0.001, x=0.01, x=0.05 and x=0.1 the
PTCR jump noted was 3order, 4 order, 2 order and 1 order respectively. In powders
synthesized using span 20 however, a 3 order PTCR jump was noted for low K
concentration of x=0.001 and x=0.01, which decreased to a 1 order PTCR jump for
x=0.05 and 2 order jump for x=0.1. Resistivity values noted at 30°C and at Tmax for
powders synthesized using span 80 was lower for x=0.1 than span 20 and was higher for
all other K concentrations. Tmax noted during the heating cycle was above 120°C and that
noted during cooling cycle was closer to 100 except for higher K concentration powders
synthesized using span 20. Powders synthesized using span 20 were elongated in
structure with an average higher % tetragonality as compared to those synthesized using
span 80. Higher tetragonality leads to higher resistivity in BaTiO3 as the ferroelectric
BaTiO3 is an insulator. Hence, relatively higher resistivity obtained in powders
130
synthesized using span 20 than in span 80 is justified. Also, the variation of barrier height
is reported to be due to the variation of charge impurity concentration at the surface
[Panwar and Semwal, 1991]. Barrier height and the density of charge carriers change
according to the impurity concentration at the grain boundary that ultimately controls the
PTCR jump [Panwar and Semwal, 1991]. Hence, higher amount of grain boundaries exist
in those powders that produce a higher PTCR jump as these allow more oxygen to be
adsorbed which results in a higher potential barrier.
Similar to that noted for pure and Sr doped BaTiO3, the resistivity measured during
cooling cycle was higher than that measured during heating cycle for all other dopants the
reason for which is mentioned earlier. No drastic change in the resistivity of the
synthesized BaTiO3 was observed due to doping. Even the range of resistivity noted for
the doped and the pure BaTiO3 was similar. This may be due to oxygen adsorption and
the smaller crystallite size rather than doping, playing the dominant role in inducing the
PTCR effect in the present study. However, the changes noted with dopant concentration
may be attributed to the charge carrier concentrations that may differ with addition of
different dopants.
4.6 Impedance Spectroscope Analysis (IS)
Figure 4.19 a shows the dielectric constant vs. temperature behavior of pure BaTiO3
pellets synthesized using span 80 for different frequencies. Figure 4.19b and 4.19c shows
the dielectric constant ($ r ) and dielectric loss (tan%) for the same pellet at a frequency of
1 kHz and 3 MHz respectively. The pellet was sintered at 750°C for a soaking time of
12hrs. For a frequency of 1 kHz the permittivity maximum was observed at 75oC which
was same as the Tc noted from the resistivity curves. With an increase in the frequency to
3 MHz the Tc shifted to 80°C. Maximum value of permittivity obtained for 1 kHz
frequency at Tc (75°C) was 9000, the loss (tan%) at this frequency was 3 times $ r . $ r
decreased with increasing frequency to as low as 33 for 3 MHz frequency at which tan%
131
Fig
4.1
9. B
aTiO
3 di
elec
tric
res
pons
e fo
r 5%
spa
n80
sam
ple
sint
ered
at
750°
C (
12 h
ours
soa
king
), m
easu
red
at (
a)
fixe
d fr
eque
ncie
s,
(b)
1 kH
z fr
eque
ncy
and
(c)
3 M
Hz
freq
uenc
y.
25
50
75
100
125
150
175
200
05
10
15
20
25
30
35
0.0
0.3
0.6
0.9
1.2
1.5
#r'
Tem
pera
ture
(0C
)
3 M
Hz
tan
$
25
50
75
100
125
150
175
200
0
1000
2000
3000
4000
5000
6000
7000
8000
9000
0.0
0.5
1.0
1.5
2.0
2.5
3.0
#r'
Tem
pera
ture
(0C
)
1 k
Hz
tan
$
cb
a
25
50
75
100
125
150
175
200
0
2000
4000
6000
8000
10000
Tem
peratu
re (
0C
)
#r'
KH
z1
KH
z1
0
KH
z1
00
KH
z5
00
MH
z3
132
in the material was 1.6 times $ r . Similar trend was observed for BaTiO3 powders
synthesized using span 20 as seen in Figure 4.20 for the pellet sintered at 750°C 1 hr
soaking time. Figure 4.20b and 20c shows the $ r and tan% for the same pellet at a
frequency of 1 kHz and 3 MHz respectively. Similar shift in value of Tc with changes in
frequency has been noted by several researchers for doped BaTiO3 ceramics which they
associated with the increase in the relaxor behaviour of the material [Joshi et al., 2006;
Elsayed and Haffz, 2005; Zhu et al., 2007; Chou et al., 2007; Aparna et al., 2001].
Lattice defects and strains force the structure to become cubic and decrease the
polarization in BaTiO3 as the particle size reduces from micrometer to nanometer size.
That is, the dipole interaction with each other forms the basis for the observed
temperature dependent properties near the transition temperature. If the coupling strength
of the bulk is larger than the surface, the surface will be relatively disordered [Asiaie et
al., 1996]. This disordered surface will tend to disorder the bulk. With decreasing particle
size, more and more of the bulk gets disordered due to the disordered surface. At very
small particle size, the particle will be completely disordered and the polar state will be
destroyed. Hence there exists a critical particle size below which the ferroelectricity is not
sustained [Asiaie et al., 1996]. The synthesized powders have average particle size in the
nanometer range and pseudocubic structure indicating the presence of predominantly
cubic and small amounts of tetragonal phases. The particle size most probably is small
enough to possess a disordered surface. This may be the reason for the huge decrease of
the dielectric constant with increasing frequency. The shift of the transition temperature
towards the lower temperature from the standard value of 120°C may correspond to this
disordered phase of the surface of the particles [Asiaie et al., 1996].
Relaxor behavior in ferroelectrics is characterized by frequency-dispersive dielectric
parameters ($ r and tan%), slim polarization-electric field loops and polar nanodomains
[Zhang et al., 2007; Chou et al., 2007; Xiong et al., 2007; Bokov and Ye, 2006; Delgado,
2005; Dash, 2009; Liu et al., 2007; Kerfah et al., 2010; Shvartsman and Lupascu, 2012].
Materials that exhibit this behaviour are termed as relaxor ferroelectrics or relaxors, and
133
Fig
4
.20
. B
aTiO
3 di
elec
tric
res
pons
e fo
r 5%
spa
n20
sam
ple
sint
ered
at
750°
C (
1 ho
ur s
oaki
ng),
mea
sure
d at
(a
) fi
xed
freq
uenc
ies,
(b
) 1
kHz
freq
uenc
y an
d (c
) 3
MH
z fr
eque
ncy.
cb
a
25
50
75
100
125
150
175
200
140
150
160
170
180
190
0.2
6
0.3
1
0.3
6
0.4
1
0.4
6
3 M
Hz
#r'
Tem
pera
ture
(0C
)
tan
$
25
50
75
100
125
150
175
200
200
400
600
800
1000
1200
1400
1600
1800
2000
0123456789
#r'
Tem
pera
ture
(0C
)
1 k
Hz
tan
$
25
50
75
100
125
150
175
200
1.0
x10
2
6.0
x10
2
1.1
x10
3
1.6
x10
3
2.1
x10
3
#r'
Tem
peratu
re (
0C
)
1 K
Hz
10
KH
z
10
0K
Hz
50
0 K
Hz
3M
Hz
134
are often referred to as “ferroelectrics with diffuse phase transition” [Bokov and Ye,
2006; Kerfah et al., 2010; Alexei and Yuji, 1998; Shvartsman and Lupascu, 2012]. The
terminology “diffuse phase transition” appears due to the broadening of the dielectric
peak in the temperature dependent dielectric constant profile [Zhang et al., 2007; Xiong
et al., 2007; Bokov and Ye, 2006; Zhu et al., 2007; Shvartsman and Lupascu, 2012]. With
increasing frequency both, a shift in Tm (temperature corresponding to maximum
dielectric constant) towards higher temperature and a decrease in maximum dielectric
constant as well as a deviation from Curie-Weiss law in the vicinity of Tm, are associated
with relaxors [Chou et al., 2007; Xiong et al., 2007; Shvartsman and Lupascu, 2012].
Relaxors also show an increase in the dielectric loss and the temperature of the maximum
dielectric loss with increasing frequency [Chou et al., 2007; Kerfah et al., 2010]. From
Figure 4.19 and 4.20, the higher dielectric loss obtained for the synthesized BaTiO3
powders may be attributed to the relaxor behaviour of these powders. Also, the
semiconducting nature increases the losses in the material because the increase in
semiconductivity implies loss of ferroelectricicity resulting in the loss of polar domains in
the material [Masó et al., 2008]. Values of dielectric constant from 600 to 12000 has been
reported for BaTiO3 prepared by various synthesis methods and having various particle
size ranging between 120nm to 5.5&m [Takeuchi et al., 1997; Gorelov et al., 2011; Joshi
et al., 2006; Huarui and Lian, 2003; Masó et al., 2008]. Xu and Gao, however, noted
70nm particles with 80% tetragonality to possess a dielectric constant value of 6900 at
room temperature contradicting the otherwise accepted relation between particle size and
the powder phase [Huarui and Lian, 2003]. Structure and dielectric properties of BaTiO3
are known to be dependent on its particle size [Gorelov et al., 2011; Frey and Payne,
1996; Takeuchi et al., 1997; Xue et al., 2007; Fong et al., 2004]. Particle size reduction
causes a deviation of these parameters from those of bulk ceramics or single crystals
[Frey and Payne, 1996]. This deviation is referred to as the size effect of BaTiO3 which
occurs at a critical size [Frey and Payne, 1996]. Critical size is the grain size at which
BaTiO3 transforms from tetragonal to cubic structure [Joshi et al., 2006; Blake, 2004;
Pithan et al., 2006; Huang et al., 2006; Fong et al., 2004]. Different synthesis methods
135
have produced BaTiO3 with various critical sizes ranging from 25 to 200nm, which is
dependent on the processing paramaters such as heat treatment temperature and cycle,
impurities in the material, etc. [Joshi et al., 2006; Takeuchi et al., 1997; Huarui and Lian,
2003; Vijatovi! et al., 2008]. With decreasing particle size, decrease in the dielectric
constant occurs, for example for a particle size of 5.5&m having dielectric constant value
of 10,000 the decrease in particle size to 1&m results in the decrease in dielectric constant
value to 6000 [Masó et al., 2008; Joshi et al., 2006]. As seen from Figure 4.19 and 4.20,
a decrease in dielectric constant value and a shift of Tc was observed for the powders
synthesized using span 80 and span 20 in the present work, as reported for relaxors.
However, since the phase transition was not diffused they cannot be called as proper
relaxor ferroelectric. Dielectric constant ($r) of a normal ferroelectric is known to
decrease with increasing frequency as it becomes difficult for the domains to change their
orientation at the same rate as that of the electric field [Elissalde and Ravez, 2001;
Garbarz-Glos et al., 2009]. The ferroelectric phase, irrespective of the amount present,
follows the field at lower frequencies by orienting the dipoles in the direction of the
applied field. At higher frequencies these ferroelectric domains may not be able to follow
the applied field at the same speed and hence lag behind. Moreover the non-ferroelectric
phases present may be the reason for its relaxor type frequency-dispersion behavior
observed in the present study.
Figure 4.21 shows the dielectric constant vs. temperature behavior of doped BaTiO3
pellets synthesized using span 80 for 3 MHz frequency with varying dopant
concentration. At low concentration of Sr, x=0.001 and x=0.01, powders synthesized
using span 80 produced a single transition peak (Tc) at 50°C and 60°C respectively.
Along with this transition peak these powders showed presence of a broad hump at
120°C. With increasing Sr concentration x=0.05 and x=0.1, a broad peak was obtained
with the Tc at 135° C and 110°C respectively. A similar two transition peak behavior was
also observed by Zhou et al. [2006] which they attributed to the core-shell structure of
doped BaTiO3. They associated the dielectric peak with the ferroelectric-paraelectric
phase transition of the unreacted pure BaTiO3 grain core at 120°C, and the peak closer to
136
Fig
4.2
1.
Die
lect
ric
resp
onse
of
Ba 1
-xD
xT
iO3 p
elle
ts s
ynth
esiz
ed u
sing 5
% s
pan
80 a
t 3 M
Hz
freq
uen
cy w
ith v
aryin
g D
and x
as 0
.00
1 (
'),
0.0
1 (
), 0
.05 (!
) an
d 0
.1(
).
40
60
80
100
120
140
160
180
200
100
110
120
130
140
150
160
170
180
190
K
'
Tem
pera
ture
(0C
)
40
60
80
100
120
140
160
180
200
80
85
90
95
100
105
110
Li
'
Tem
pera
ture
(0C
)
40
60
80
100
120
140
160
180
200
40
60
80
100
120
140
Mg
'
Tem
pera
ture
(0C
)
25
50
75
100
125
150
175
200
4.0
4.5
5.0
5.5
100
200
300
400
Ce
'
Tem
pera
tur
(0C
)
40
60
80
100
120
140
160
180
200
0
20
40
60
80
100
La
'
Tem
pera
ture
(0C
)
40
60
80
100
120
140
160
180
200
0
50
100
150
200
250
300
350
Sr
'
Tem
pera
tur
(0C
)
137
room temperature was associated with the grain shell. Hence, the peak above 120°C may
be due to the ferroelectric core and that at around 110°C due to the shell with higher Sr
content in the Sr doped BaTiO3 powders. Also, similar behavior of the transition peaks
merging showing one peak with increasing Nb incorporation in BaTiO3 was observed by
Maso et al. [2008] which they associated with the quasi-ferroelectric behavior of the
material. Broadening of the peak was indicative of the ferroelectric to paraelectric phase
transition to be diffusive in nature [Zhang et al., 2007; Masó et al., 2008]. This behaviour
of diffusive phase transition with higher values of dopant concentration was also
observed by Zhang, Li and Viehland for 3% - 6% La doping in BaTiO3 [Zhang et al.,
2007]. Sr doping in BaTiO3 is known to decrease the Tc towards lower temperature from
the normally reported Tc of 120°C. In the present work the Tc in the undoped BaTiO3 was
observed at 80°C at 3MHz frequency and the material exhibited a relaxor-type
ferroelectric behavior. Similar relaxor type behavior with the TC reported to be lower than
the ideally reported value of 120°C was also observed with Sr doped BaTiO3 synthesized
using span 80. The dielectric values with Sr doping at 3 MHz frequency were higher
(>50) than those of undoped BaTiO3 (33) powders synthesized using span 80.
Transition peak was observed below 100°C for all La concentrations of powders
synthesized using span 80. For x=0.001 La concentration, TC value was lower than that of
undoped BaTiO3 (Figure 4.19.c). However, with increasing La concentration the shift in
the peak was noted towards higher temperature and the dielectric constant values were
higher than that of undoped BaTiO3. La doping in BaTiO3 was known to decrease the Tc
towards room temperature from the normally reported Tc of 120°C and also increase its
dielectric constant when La ion replaces the Ba ion in the crystal [Da-Yong et al., 2006;
Chunlin et al., 2012; Morrison et al., 1998, 1999]. Though the increase in dielectric
constant value noted was as expected, the shift in the transition temperature with
increasing La concentration was contrary to the reported trend. This was seen from the
higher temperature shift of the transition temperature noted at ~50°C for x=0.001 to
~80°C for x=0.1 as compared to the value noted for pure BaTiO3 at 75°C. This may be
due to the suppressed tetragonality of the synthesized powders as seen from Figure 4.8
138
and Table 4.3.
Similar to that observed with La doping, Ce doped BaTiO3 powders showed presence of
a single transition peak below 100°C for low concentrations of Ce (x=0.001 and x=0.01).
Dielectric constant values of these powders increased with increasing Ce concentrations.
A lower temperature shift of TC has been reported with Ce doping at Ti site [Da-Yong et
al., 2006; Makovec and Kolar, 1997; Zhi et al., 2002]. However, the shift in TC noted
towards higher value suggests that Ce does not replae Ti in the synthesized powders.
However, for x = 0.1 Ce concentration, the dielectric constant was almost independent of
temperature. Unlike La and Ce, Mg doped BaTiO3 powders synthesized using span 80
showed the presence of a double peak. Dielectric constant decreased continuously with
increasing Mg content except for x=0.001 Mg concentration. The dielectric constant
became temperature independent for x=0.1 Mg concentration. For Li and K doping in
powders synthesized using span 80 the presence double peaks were noted for all the
dopant concentrations. Hence characteristic curves corresponding to x=0.1 are presented
in Fig. 21. The first peak was noted at 70°C for both dopants and the second peak at
140°C for Li and at150°C for K doped BaTiO3. As mentioned above, TC at higher
temperature was associated with the BaTiO3 core and that at lower temperature with the
paraelectric shell due to doping. Figure 4.22 shows the dielectric constant vs. temperature
behavior of doped BaTiO3 pellets synthesized using span 20 for 3 MHz frequency with
varying dopant concentration. BaTiO3 doped with Sr showed presence of two peaks, one
around 70°C and the other at above 120°C. At higher Sr concentration i.e. x=0.05 and
x=0.1 the two peaks broadened showing the diffusive phase transition similar to that
obtained with powders synthesized using span 80. The dielectric values with Sr doping at
3 MHz frequency were smaller than those of undoped BaTiO3 powders synthesized using
span 20 as seen from Figure 4.20c. For x=0.001 La concentration the TC noted was same
as that for undoped BaTiO3 (Figure 4.20.c). However with increasing La concentration,
no significant trend in the change in TC was noted. For x=0.1 La concentration, the
dielectric constant became temperature independent. Similar temperature independent
dielectric constant with increasing dopant concentration was observed by Maso et al.
139
Fig
4.2
2.
Die
lect
ric
resp
onse
of
Ba 1
-xD
xT
iO3 p
elle
ts s
ynth
esiz
ed u
sing 5
% s
pan
20 a
t 3 M
Hz
freq
uen
cy w
ith v
aryin
g D
and x
as 0
.00
1 ("
), 0
.01 (
), 0
.05 (!
) an
d 0
.1(
).
20
4060
80
100
120
140
160
180
200
30
40
50
60
70
80
90
Mg
Tem
pera
ture
(0 C
)
e'
40
60
80
100
120
140
160
180
200
10
20
30
40
K
'
Tem
pera
ture
(0 C
)
40
60
80
100
120
140
160
180
200
5
10
15
20
25
30
35
Li
'
Tem
pera
ture
(0C
)
4060
80
100
120
140
160
180
200
60
80
100
120
140
160
Ce
Tem
pera
tur
(0 C)
'
40
60
80
100
120
140
160
180
200
0
20
40
60
80
100
120
140
160
180
200
La
'
Tem
pera
ture
(0C
)
40
60
80
100
120
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20
30
40
50
60
70
80
90
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110
Sr
'
Tem
pera
tur
(0C
)
140
[2008] for Nb doped BaTiO3 which they confirmed was typical of a non- ferroelectric
material. Dielectric constant was higher for low Ce concentration compared to that
obtained for x = 0.1 Ce concentration. Similar to that observed with Sr and La doped
BaTiO3, two peaks were noted in Ce doped BaTiO3 powders. The lower temperature
peak at ~ 70°C was dominant over that obtained at around140°C. Hence suggesting that
the dopant ion does not dominate the variation in the TC and the dielectric constant
obtained. For Mg, Li and K doping in powders synthesized using span 20 the presence of
single peak was noted for all dopant concentrations. Hence characteristic curves
corresponding to x=0.1 are presented in Figure 4.22. These dopants showed presence of
only one transition peak at ~ 70°C for Mg and Li dopant and at 80°C for K dopant.
However, the dielectric transition in these samples was neither very sharp nor did it show
a diffused phase transition. This implies that no core-shell type structure formation took
place for these dopants. As mentioned above, the lower amount of tetragonal phase led to
the presence of the phase transition and the presence of the cubic phase resulted in a
decrease of the dielectric constant.
141
Chapter 5 Conclusion and Future Scope of Work
5.1 Conclusion
The present work was aimed at synthesizing pure and doped barium titanate nanoparticles
using sol-gel emulsion technique. The objective was to study the effect of surfactants
(span 80 and span 20) and dopants on the particle size, structural and electrical properties
of the synthesized powders. Pure and doped BaTiO3 particles were prepared with the
standardized protocol to produce powders in the nano range. The properties of the
synthesized material were characterized with respect to calcination temperature,
crystalline behavior, resistivity and impedance parameters and surface morphologies.
Emulsion used in the present study was w/o type. The small sized micelles generated
acted as templates during synthesis. The optimized synthesis method involved the
addition of a translucent BaTiO3 sol (water medium) to a support solvent made up of the
oil medium and surfactant to form the emulsion. From the SEM micrographs it was noted
that the powders synthesized using span 80 produced mostly spherical shaped particles
and those synthesized using span 20 produced predominantly rod-like particles. This
shape dependence could be justified from the values of HLB, CPP and the relation
governing the two. Particles with an average size of 57nm for particles synthesized using
5% span80, 70nm with 20% span 80 and 66nm using 20% span 20 were obtained. Use of
5% span 80 produced particles with spherical shape which distorted with increasing
surfactant concentration to 20%. Onset of BaTiO3 started from 600°C with the absence of
any distinct splitting of the (200) peak at ~45° up to a calcination temperature of 900°C
indicating predominance of cubic phase in the synthesized powders. A shoulder (002)
peak was present around that of (200) peak for powders calcined at a temperature of
750°C. However, for the powders calcined at 1000°C distinct presence of (200) and (002)
peak was obtained which confirmed an increase in the % tetragonality. In pure BaTiO3,
synthesized using span 80 and span 20, < 50% tetragonality was achieved. Particles
142
formed were of nanometer size with the presence of both the tetragonal (smaller
component) and the cubic phases (larger component) that led to pseudocubic symmetry
of the crystal structure. Experimentally obtained lattice parameters showed a better fit
with those calculated using the empirical formula devised by Jiang et. al for powders
synthesized using span 80 as well as span 20. The |%error| for powders synthesized using
span 80 was 0.046% and that using span 20 was 0.102% in comparison to the 0.075%
reported by Jiang et. al. The better fit was attributed to the smaller particle size which
better approximated to the six-fold coordination than bulk material. FTIR analysis of pure
BaTiO3 synthesized using span 80 showed the Ti-O absorption peak to become narrower,
sharper and the intensity of the carboxyl, hydroxyl peaks to decrease with increasing
calcination temperature from 500°C to 1000°C. Similar behavior was also achieved for
pure BaTiO3 synthesized using span 20. PTCR effect was noted for the synthesized pure
BaTiO3 powder using span 80 as well as that using span 20. Pure BaTiO3 is traditionally
known to be an insulator. It deviates from its insulating behavior either when it is doped
with donors, when gases are adsorbed on the grain surfaces, or when cation vacancies are
created. PTCR effect is known to be grain boundary phenomena. Increase in resistivity of
BaTiO3 is associated to the increase in potential barrier height made up of acceptor states
at the grain boundary. Oxidation of grain boundary either with barium vacancies,
oxidation of 3d-elements to a higher oxidation state, or due to adsorbed oxygen results in
higher electron trap density. For pure BaTiO3 synthesized using span 80, the room
temperature resistivity obtained was~ 108 #cm and the same for span 20 was ~ 10
7 #cm
which was lower than the standard value of 109 #cm - 10
12 #cm reported for pure
BaTiO3. The value increased to ! 1010
#cm with increase in temperature. Increase in
resistivity starts at the TC, which was evaluated as 75°C in the synthesized powders,
confirming that the TC of pure BaTiO3 synthesized in the present work decreased from
the standard value of 120°C. This was attributed to the adsorbed oxygen and the
predominance of cubic structure of the crystals due to small crystallite size. Resistivity
measured during cooling cycle in the ambient atmosphere was higher than that measured
during heating cycle because of the increased potential due to the decrease in the number
of conduction electrons as a result of oxygen adsorption at the grain boundaries.
143
Impedance spectroscopy measurements showed the TC of pure BaTiO3 synthesized using
span 80 to be 75°C and that using span 20 to be 80°C for 1 kHz frequency. With
increasing frequency, the dielectric constant decreased and showed a relaxor-type
ferroelectric behavior which resulted from the smaller particle size with disordered
surface. Dielectric loss noted in these powders was higher than that reported in the
literature. This was attributed to the relaxor behavior and the semiconducting nature of
the synthesized powders due to the loss of polar domains as a result of lower
ferroelectricity of the powders.
Doping help in tailoring the properties of BaTiO3 as the dopant ion strongly affect its
structure and properties. Sr, La, Ce, Mg, Li and K were the chosen dopants in the present
study. Sr doped BaTiO3 is known to decrease TC towards room temperature and possess
non-linear dielectric behavior, having various application as electromechanical sensors,
phase shifters, tunable filters, hig-Q resonators, etc. La being a donor induces
semiconductivity giving rise to PTCR effect in BaTiO3 to be used as a PTCR thermistor.
La and Ce doping is also known to decrease TC towards room temperature. The dielectric
properties of these dopants have led to its use in capacitors. Mg as an acceptor was used
to modify the PTCR effect in BaTiO3 and also reduce TC towards room temperature. Li
and K doping in BaTiO3 were only studied for its structural properties and in modified
BaTiO3 for its PTCR effect. Hence, the present work aimed at studying the effect of these
dopants on the structural and electrical properties of nano sized BaTiO3 particles
synthesizrd via sol-gel emulsion technique. The trends observed in the structural and
electrical properties of doped BaTiO3 were similar to that noted for pure BaTiO3,
irrespective of the type of dopant (isovalent/ aliovalent) used.
Doping BaTiO3 with Sr decreased the cell volume and pseudocubic lattice parameter for
Ba0.9Sr0.1TiO3 powder synthesized using both span 80 and span 20. An overall shift of the
diffraction peak towards the higher angle was obtained, corresponding to the decrease in
lattice spacing as well as an increase in the cubic phase of these powders as compared to
that of pure BaTiO3. Such behavior has also been reported in Sr doped BaTiO3, and has
144
been attributed to the smaller ionic size Sr replacing Ba. Dominance of cubic phase in the
synthesized powders was also confirmed with the presence of a single (200) peak in
powders synthesized using span 80 and low % tetragonality (18.29%, x=0.1) calculated
for those synthesized using span 20. The pc lattice parameter value of BaTiO3 decreased
with Sr doping from 4.0171Aû (undoped) to 4.0007Aû for x=0.1 Sr concentration in
powders synthesized using span 80. Similarly in powders synthesized using span 20 the
pc lattice parameter value decreased from 4.0194Aû (undoped) to 4.0075Aû for Sr
concentration corresponding to x=0.1. From the FTIR analysis, the Ti-O characteristic
peak in BaTiO3 shifted towards higher wave number from 538cm-1
noted for pure
BaTiO3 for all Sr concentration in the powders synthesized using span 80 and span 20.
This was assigned to the decrease in lattice spacing along with the increase in cubic phase
noted in the synthesized powders. The maximum shift of the Ti-O absorption peak was
9cm-1
in powder synthesized using span 80 and 18cm-1
in powder synthesized using span
20. Doping changes the cell size, crystal structure and binding energy, due to which
replacing Ba with a smaller ion such as Sr has been reported to shift the wave number
towards a higher value. PTCR effect observed in Sr doped BaTiO3 powders synthesized
using both span 80 and span 20 measured higher resistivity values during cooling cycle
than heating cycle similar to that observed in pure BaTiO3. The range of resistivity value
noted was not highly altered with Sr addition and appeared from 107#cm to 10
10#cm in
powders synthesized using span 80 and span 20. Sr doping in BaTiO3 has been reported
to produce a shift in TC from the standard value of 120°C towards room temperature.
However, the TC in the Sr doped BaTiO3 powder was noted in the same range as that of
the synthesized pure BaTiO3. This is possibly due to the dominant role played by the
smaller particle size and the adsorbed oxygen at the grain boundaries than the dopant ion
in lowering the TC. Powders synthesized using span 80 showed presence of two peaks in
the dielectric measurements for lower Sr concentration (x=0.001 and 0.01) which merged
to produced single broad dielectric peak indicting its diffusive behavior for higher
concentration of Sr (x=0.05 and 0.1). The powders synthesized using span 20 also
showed presence of two transition peaks, one at low temperature around 70°C and the
other around 120°C. Similar behavior of two transition peak presence was reported for
145
Nb doped BaTiO3, which was attributed to the quasi-ferroelectric behavior. Hence, a
similar quasi-ferroelectric nature of the synthesized powders due to the smaller size and
the incorporation of Sr must have led to the presence of two peaks.
La3+
doping in BaTiO3 conformed to the smaller radii dopant ion trend towards forming
the predominantly cubic structure and produced smaller cell volume and pseudocubic
lattice parameter for Ba0.9La0.1TiO3 irrespective of the surfactant used during synthesis.
Higher angle shift of (110) diffraction peak in these powders similar to that observed in
Sr doped BaTiO3 was assigned to the smaller lattice spacing along with an increase in
cubicity of the crystal structure. This was also observed with the presence of a single
(200) peak in La doped BaTiO3 (x=0.1) powders synthesized using span 80 and low %
tetragonality (19.64%, x=0.1) calculated for those synthesized using span 20. A decrease
in the pc lattice parameter value was noted from 4.0171Aû for undoped to 3.9961Aû for
x=0.1 La doped BaTiO3 in powders synthesized using span 80. A small decrease in the pc
lattice parameter value in powders synthesized using span 20 was noted from 4.0194Aû
for undoped to 4.0140Aû for x=0.1 La concentration in BaTiO3. This showed that there
was a decrease in cell volume with La doping in BaTiO3. FTIR results obtained with La
doping showed higher wave number shift of the Ti-O absorption peak in the powders
synthesized using span 80 as well as span 20 due to the decrease in lattice spacing as well
as the increase in cubicity with increasing dopant concentration. The maximum shift of
the Ti-O absorption peak was 13cm-1
in powder synthesized using span 80 and 17cm-1
in
powder synthesized using span 20. La doping like any other donor dopant has been
reported to induce semiconductivity in BaTiO3. La doping gives rise to the PTCR effect
due to electronic conduction resulting from charge compensation. However, in the
present study, the PTCR jump and also the room temperature resistivity value was of the
same order as that obtained for the undoped BaTiO3, indicating that the La doping did not
play a dominant role in the PTCR effect observed in the synthesized powders. Hence, the
adsorbed oxygen at the grain boundaries and the particle size achieved plays the decisive
role in producing semiconducting material. The resistivity range obtained with La doping
was 107#cm to 10
11#cm. TC noted in these powders was closer to that of the synthesized
146
undoped BaTiO3 and varied from 70°C - 95°C. TC measured was below 100°C for La
doped BaTiO3 powders synthesized using span 80 and span 20. In powders synthesized
using span 80, with increasing La concentration from x=0.001 to 0.1, an increase in TC
was noted which was not observed in those synthesized using span 20. Increase in
dielectric constant and shift of TC towards room temperature is reported for La doped
BaTiO3. This was however not observed in the synthesized powder which is attributed to
the suppressed tetragonality and the presence of the non-ferroelectric phase in the
synthesized powders.
Ce3+
doping in BaTiO3 synthesized using span 80 and span 20 also resulted in the shifting
of (110) diffraction peak towards higher angle compared to that noted for pure BaTiO3.
Ce doping showed both (200) and (002) peaks which confirmed the presence of
tetragonal phase in the synthesized powders irrespective of the surfactant used. This may
possibly be due to Ce3+
getting incorporated in the Ba-[TiO6] sub lattice, or presence of
some amount of Ce4+
which may replace Ti causing the c-axis elongation. However, the
overall trend of decreasing lattice spacing along with an increase in cubicity resulted in
the characteristic Ti-O absorption peak to shift towards higher wave number in FTIR
absorption spectra of powders synthesized using span 80 and span 20. The maximum
shift of the Ti-O absorption peak was 5cm-1
in powder synthesized using span 80 and
17cm-1
in powder synthesized using span 20. PTCR effect observed in Ce doped BaTiO3
was similar to that noted for undoped BaTiO3. The resistivity value ranged from 108#cm
to 1011#cm in powders synthesized using span 80 and from 10
8#cm to 1010#cm in those
using span 20. The major characteristics noted like the diffusive relaxor-type behavior
and lower TC value with Ce doping was similar to that noted for other powders. Like La,
Ce doping in BaTiO3 has also been reported to decrease the TC towards room
temperature. This was however not observed in the synthesized powders. Powders
synthesized using span 80 showed presence of one single diffused peak but those
synthesized using span 20 showed presence of two peaks which was similar to that
reported for core-shell structure. The basic reason for this behavior is the particle size
147
which makes the material predominantly cubic having very small amount of tetragonal
structure, due to which the dielectric behavior becomes more diffused in these powders
Mg2+
doping in BaTiO3 synthesized using span 80 and span 20 shifted the (110)
diffraction peak towards higher angle. The presence of (200) and (002) peaks was noted
for all Mg concentrations in powders synthesized using span 80 and span 20. As the ionic
radius of Mg is closer to that of Ti than Ba, some amount of Mg must have replaced Ti
causing c-axis elongation. On the other hand, the (110) diffraction peak shift showed Mg
to have possibly replaced Ba. The cell volume shrinkage in powders synthesized using
span 80 was 1.57% and for those synthesized using span 20, the negligible shrinkage of
0.02% was noted for x=0.1 Mg concentration. From FTIR spectra, a higher wave number
shift of the Ti-O characteristic peak with respect to that of pure BaTiO3was noted on Mg
doping in powders synthesized using span 80 and span 20. The maximum shift of the Ti-
O absorption peak was 5cm-1
in powder synthesized using span 80 and 21cm-1
in powder
synthesized using span 20. This higher wave number shift behavior was attributed to the
decrease in lattice spacing along with the increase in cubic phase of the synthesized
powders with increasing dopant concentration. The resistivity values noted ranged from
107#cm to 10
11#cm for powders synthesized using span 80 and from 107#cm to
1010#cm in those synthesized using span 20. The trend observed in the PTCR behavior of
Mg doped BaTiO3 powders irrespective of the surfactant used during synthesis was
similar to that of undoped BaTiO3. With Mg doping, impedance spectroscopy showed the
presence of double peak with decreasing dielectric constant as the Mg concentration
increased in powders synthesized using span 80. This was attributed to the core-shell type
behavior obtained in these powders. However, in powders synthesized using span 20 only
one peak was noted which does not show the presence of core-shell type behavior.
Li1+
, due to its smaller ionic radius upon doping produced cubic powders with smaller
cell volume and pseudocubic lattice parameter compared to that of undoped BaTiO3 for
x=0.1 concentration. A shift of (110) diffraction peak towards higher angle with respect
to that of undoped BaTiO3 was achieved in Li doped BaTiO3 powders for all Li
148
concentrations synthesized using span 80 and span 20. In powders synthesized using span
80, presence of one single (200) peak for x=0.1 Li concentration indicated predominance
of cubic phase. However in powders synthesized using span 20, both (200) and (002)
peaks were present due to comparatively greater amount of tetragonality. % tetragonality
calculated for powders synthesized using span 20 was 15.57% for x=0.1 Li concentration.
Cell volume reduced from 64.83 A°3 for the undoped to 63.79 A°
3 in powders synthesized
using span 80 and in those using span 20 it reduced from 64.94 A°3 for the undoped
to
64.81 A°3
for x=0.1 Li concentration. FTIR analysis showed higher wave number shift of
the characteristic Ti-O absorption peak in powders synthesized using span 80 and span 20
as seen with other dopants. In powder synthesized using span 80, the maximum shift of
the Ti-O absorption peak was 9cm-1
and in powder synthesized using span 20, it was
17cm-1
. No drastic change in the PTCR properties with Li doping in BaTiO3 was noted.
The resistivity value ranged from 106#cm to 10
11#cm in powders synthesized using span
80 and from 107#cm to 10
10#cm in those synthesized using span 20. Ba0.9Li0.1TiO3
powder synthesized using span 80 was more diffusive in nature with double dielectric
peak presence than that synthesized using span 20 which showed presence of a single
peak. This was similar to that noted for Mg doped BaTiO3 synthesized.
K1+
doping in BaTiO3 powders synthesized using span 80 and span 20 also showed
smaller cell volume and pseudocubic lattice parameter compared to that of undoped
BaTiO3. Higher angle shift of (110) diffraction peak with doping was also observed. In
powders synthesized using span 80, presence of a single (200) peak indicated it to be
predominantly cubic in nature for x=0.1 K concentration. Calculated % tetragonality in
powders synthesized using span 20 was 16.42% which showed the presence of both (200)
and (002) peaks for x=0.1 K concentration. Cell volume reduced by 1.62% in powders
synthesized using span 80 and by 0.08% in those using span 20 for x=0.1 K concentration
compared to those of undoped samples. Higher wave number shift of the characteristic
Ti-O absorption peak with K doping was noted in the synthesized powders irrespective of
the surfactant used. The maximum shift of the Ti-O absorption peak was 9cm-1
in powder
synthesized using span 80 and 5cm-1
in powder synthesized using span 20.PTCR effect in
149
K doped BaTiO3 was similar to that observed with undoped BaTiO3. Resistivity values
noted for powders synthesized using span 80 varied from 107#cm to 10
11#cm and those
using span 20 varied from 107#cm to 10
10#cm. TC noted from the resistivity
measurements was in the range of 60°C to 80°C in K doped powders. Ba0.9K0.1TiO3
powder synthesized using span 80 was more diffusive in nature with the presence of two
dielectric peaks than that synthesized using span 20 which showed a single dielectric
peak presence.
Pure BaTiO3 synthesized in the present study was pseudocubic in nature on calcining at
750°C irrespective of the surfactant used. The % tetragonality obtained in these powders
was small even when calcined at 1000°C. PTCR effect was noted in these powders with a
reduction in the transition temperature. Relaxor type dielectric properties were achieved
as the tetragonality in the synthesized powders was low due to the nano size particles.
Wherein, the normal trend of insulating undoped BaTiO3 synthesized using conventional
methods, calcined above 1000°C are tetragonal in structure. Dopants used in the present
study in general showed the lattice spacing to have reduced as well as the cubic phase to
have increased with increasing dopant concentration in the synthesized powders. This
was confirmed from the XRD and the FTIR results. One of the reasons for this is that the
ions used were smaller than Ba except for K which was almost of the same ionic radius.
When a smaller ion replaces Ba in BaTiO3, it shrinks the unit cell increasing the bond
strength of the ions. However, certain elements such as Ce (Ce4+
) and Mg have been
reported to replace Ti ions in powders synthesized using solid-phase method calcined at
temperatures >1000°C and in hydrothermal reaction calcined at ~220-260°C, which is
also considered to have resulted in the presence of tetragonality in the synthesized
powders. PTCR and dielectric properties noted was similar to that observed for undoped
BaTiO3. This is due to the smaller particle size that plays a dominant role in controlling
its properties.
150
5.2 Future Scope of Work
The present work showed that the properties exhibited by barium titanate nano particles
were different from those reported for bulk material. Since the size of the particles is
related to the calcination temperature, further studies can be concentrated on exploring
modification in the current synthesis route to fabricate fully crystalline BaTiO3 powders
at lower temperatures. Efforts can also be driven toward controlling the various synthesis
parameters to get particles of controlled size and shape which will lead to better
properties. In the present work all the synthesized powders were pseudocubic in nature,
thereby restricting the use of the synthesized material for PTCR purposes. Synthesizing
tetragonal BaTiO3 powders reproducibly with controlled shapes and sizes can be an
important and interesting area to explore further.
151
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LIST OF PUBLICATIONS
Journal
Zubeda B. H. Aga and Sutapa R. R. (2012) Electrical and structural characterization of
PTCR pure BaTiO3 nanopowders synthesized by sol-gel emulsion technique. Journal of
Electroceramics. 28: pg. 109-117.
International Conferences
Zubeda B. H. Aga and Sutapa R. R. (2011) Synthesis of sub-micron size pure and doped
barium titanate powders by sol-gel emulsion technique, In: Third International
Conference on frontiers in nanoscience and technology Cochin Nano – 2011, 14th-17th
August 2011, Cochin, India.
Zubeda B. H. Aga and Sutapa R. R. (2011) Strontium doped barium titanate nano-
powders synthesized by sol-gel emulsion technique, In: International Conference on
nanomaterials and nanotechnology, 18th-21st December 2011, Delhi, India.
National Conferences
Zubeda B. H. Aga and Sutapa R. R. (2011) Synthesis od sub-micron size barium titanate
powders by sol-gel emulsion technique, In: chESA & Department of Chemical
Engineering, Pravara Rural College sponsored 10th National Level Technical Symposium
CHEMSTORM 2011, 31st March 2011, Loni, India.
Zubeda B. H. Aga and Sutapa R. R. (2011) Effect of lanthanum doping on barium
titanate nano-powders synthesized by sol-gel emulsion technique, In: UGC sponsored
National Seminar on Nanomaterials: synthesis characterization and applications, 2nd-3rd
February 2012, Margao, India.
183
The present work was funded by DRDO, Ministry of Defence, Government of India,
New Delhi, under the sanctioned title: “Synthesis of improved ferroelectric materials by
sol-gel emulsion technique” bearing Ref. No.: ERIP/ER/0605047/M/01/941 on May 7,
2007.
I greatfully acknowledge the following people for the assistance received in
characterization, without which this thesis would not have been complete.
TGA, DSC Dr. N. N. Ghosh, Department of Chemistry, BITS, Pilani – K.
K. Birla Goa Campus.
Dr. Bhanudas Naik, Department of Chemistry, BITS, Pilani –
K. K. Birla Goa Campus.
FTIR Dr. Ervin Desa, Department of Physics, Goa University.
Dr. R. B. Tangsali, Department of Physics, Goa University.
Mr. Pranav Naik, Department of Physics, Goa University.
SEM Dr. Rosilda Selvin, Department of Chemical Engineering,
Lunghwa University Taiwan.
Dr. Rajan S., Director, NCAOR, Goa.
Dr. Rahul Mohan, NCAOR, Goa.
Ms. Shahina Gazi, NCAOR, Goa.
TEM SAIF, IIT Bombay.
Impedance Spectroscopy Dr. Ervin Desa, Department of Physics, Goa University.
Dr. R. B. Tangsali, Department of Physics, Goa University.
Mr. Pranav Naik, Department of Physics, Goa University.
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BIOGRAPHY
Ph. D. Student
Ms. Zubeda Bi H. Aga is a Ph.D. student in the Department of Chemical Engineering,
BITS, Pilani - K. K. Birla Goa Campus. She joined as A Junior Research Fellow for a
DRDO sponsored project in the year 2007. The duration of the project was three years.
She registered for Ph.D. in the same year and submitted her research proposal in the year
2012. She received her Masters Degree from the Physics Department, Goa University,
India in the year 2007.
Supervisor
Prof. Sutapa Roy Ramanan is a faculty in the Department of Chemical Engineering, BITS
Pilani – K. K. Birla Goa Campus. She was awarded her Ph. D. in Materials Science in
1995 from Jadavpur University, Kolkata. After completing her Ph.D., she worked as a
research associate for duration of two years at C.G.C.R.I. Calcutta, India. Subsequetly
she joined University of Sains Malaysia as a visiting post-doctoral fellow for one year
and continued to serve the university as an assistant professor at school of materials and
minerals res. Engg. For seven years. Following this, she joined BITS, Pilani - K. K. Birla
Goa Campus in 2005. Till date she has acted as a supervisor for five students for their
Masters Degree projects and one Ph.D. student, and is currently supervising one Ph.D.
student. Presently she is involved in reviewing research papers and acts as a reviewer for
Journal of American Ceramic Society, Indian Journal of Pure and Applied Physics,
Applied surface Science, Inductril and Engineering Chemistry Research, etc.
