Nano and Grain-Orientated Ferroelectric Ceramics Produced by SPS JING LIU 刘景

Nano and Grain-Orientated Ferroelectric
Ceramics Produced by SPS
JING LIU
刘景
Department of Physical, Inorganic and Structural Chemistry
Stockholm University
2007
Doctoral Dissertation 2007
Department of Physical, Inorganic and Structural Chemistry
Stockholm University
S-10691 Stockholm
Sweden
Faculty opponent:
Professor RNDr. Pavol Sajgalik, DrSc.
Institute of Inorganic Chemistry
Slovak Academy of Sciences
Dubravska cesta 9
SK-845 36 Bratislava 45
Slovak republic
Evaluation committee:
Professor Lennart Bergström, Stockholm University
Professor Karin Larsson, Uppsala University
Doc Sten Eriksson, Chalmers University of Technology
Doc Jinshan Pan, Royal Institute of Technology, KTH
© Jing Liu, pp. 1-73
ISBN: 978-91-7155-444-4
Printed in Sweden by Printcenter, US-AB, Stockholm
Dedicated to
my parents, my husband and my daughter
We can't solve problems by using
the same kind of thinking we used
when we created them.
–Albert Einstein
Abstract
Nano-powders of BaTiO3, SrTiO3, Ba0.6Sr0.4TiO3, a mixture of the
composition (BaTiO3)0.6(SrTiO3)0.4 with particle sizes in the range of 60
to 80 nm, and Bi4Ti3O12 with an average particle size of 100 nm were
consolidated by spark plasma sintering (SPS). The kinetics of reaction,
densification and grain growth were studied. An experimental procedure
is outlined that allows the determination of a “kinetic window” within
which dense nano-sized compacts can be prepared. It is shown that the
sintering behaviour of the five powders varies somewhat, but is generally
speaking fairly similar. However, the types of grain growth behaviour of
these powders are quite different, exemplified by the observation that the
kinetic window for the (BaTiO3)0.6(SrTiO3)0.4 mixture is 125 oC, ~75 oC
for Bi4Ti3O12, ~25oC for BaTiO3 and SrTiO3, while it is hard to observe
an apparent kinetic window for obtaining nano-sized compacts of
Ba0.6Sr0.4TiO3. During the densification of the (BaTiO3)0.6(SrTiO3)0.4
mixture the reaction 0.6BaTiO3+0.4SrTiO3 → Ba0.6Sr0.4TiO3 takes place,
and this reaction is suggested to have a self-pinning effect on the grain
growth, which in turn explains why this powder has a large kinetic
window. Notably, SPS offers a unique opportunity to more preciously
investigate and monitor the sintering kinetics of nano-powders, and it
allows preparation of ceramics with tailored microstructures.
The dielectric properties of selected samples of (Ba, Sr)TiO3 ceramics
have been studied. The results are correlated with the microstructural
features of these samples, e.g. to the grain sizes present in the compacts.
The ceramic with nano-sized microstructure exhibits a diffuse transition
in permittivity and reduced dielectric losses in the vicinity of the Curie
temperature, whereas the more coarse-grained compacts exhibit normal
dielectric properties in the ferroelectric region.
The morphology evolution, with increasing sintering temperature, of
bismuth layer-structured ferroelectric ceramics such as Bi4Ti3O12 (BIT)
and CaBi2Nb2O9 (CBNO) was investigated. The subsequent isothermal
V
sintering experiments revealed that the nano-sized particles of the BIT
precursor powder grew into elongated plate-like grains within a few
minutes, via a dynamic ripening mechanism.
A new processing strategy for obtaining highly textured ceramics is
described. It is based on a directional dynamic ripening mechanism
induced by superplastic deformation. The new strategy makes it possible
to produce a textured microstructure within minutes, and it allows
production of textured ferroelectric ceramics with tailored morphology
and improved physical properties.
The ferroelectric, dielectric, and piezoelectric properties of the textured
bismuth layer-structured ferroelectric ceramics have been studied, and it
was revealed that all textured samples exhibited anisotropic properties
and improved performance. The highly textured Bi4Ti3O12 ceramic
exhibited ferroelectric properties equal to or better than those of
corresponding single crystals, and much better than those previously
reported for grain-orientated Bi4Ti3O12 ceramics. Textured CaBi2Nb2O9
ceramics exhibited a very high Curie temperature, d33-values nearly three
times larger than those of conventionally sintered materials, and a high
thermal depoling temperature indicating that it is a very promising
material for high-temperature piezoelectric applications.
VI
List of publications
This thesis is based on the list papers:
I: Jing Liu, Zhijian Shen, Mats Nygren, Bo Su, and Tim W. Button,
Spark Plasma Sintering Behavior of nano-sized (Ba, Sr)TiO3
powders: determination of sintering parameters yielding
nanostructured ceramics, Journal of the American Ceramic
Society, 89 [9] (2006) 2689–2694.
II: Jing Liu, Zhijian Shen, Mats Nygren, Yanmei Kan, and Peiling
Wang, SPS processing of bismuth-layer structured ferroelectric
Ceramics yielding highly textured microstructures, Journal of
the European Ceramic Society, 26 (2006) 3233-3239.
III: Zhijian Shen, Jing Liu, Jekabs Grins, Mats Nygren, Peiling Wang
and Yanmei Kan, Haixue Yan and Ulrich Sutter, Effective grain
alignment in Bi4Ti3O12 ceramics by superplastic-deformationinduced directional dynamic ripening, Advanced Materials,
17(6) (2005) 676-680.
IV: Jing Liu, Zhijian Shen, and Mats Nygren, Consolidation and
dielectric behaviours of Ba0.6Sr0.4TiO3 ceramics with tailored
microstructures, Ferroelectrics, 319 (2005) 109-116.
V: Haixue Yan, Hongtao Zhang, Rick Ubic, Michael J. Reece, Jing
Liu, Zhijian Shen and Zhen Zhang, A Lead Free High Curie
Point Ferroelectric Ceramic, CaBi2Nb2O9, Advanced Materials,
17 (2005) 1261-1265.
VI: Zhijian Shen, Hong Peng, Jing Liu and Mats Nygren, Conversion
from Nano- to Micron-Sized Structures: Experimental
VII
Observations, Journal of the European Ceramic Society, 24
(2004) 3447–3452.
Papers not included in the thesis:
VII: J Petzelt, T Ostapchuk, I Gregora, P Kuzel, J Liu and Z Shen,
Infrared and Raman studies of the dead grain-boundary layers
in SrTiO3 fine-grain ceramics, Journal of Physics: Condensed
Matter, 19 (2007) 196222-16.
VIII: Haixue Yan, Michael J. Reece, Jing Liu, Zhijian Shen, Yanmei
Kan, Peiling Wang, Effect of texture on dielectric properties and
thermal depoling of Bi4Ti3O12 ferroelectric ceramics, Journal of
applied physics, 100 (2006) 076103-3.
IX: Jan Petzelt, Tetyana Ostapchuk, Ivan Gregora, Maxim Savinov,
Dagmar Chvostova, Jing Liu, and Zhijian Shen, Grain boundary
effects on dielectric, infrared and Raman response of SrTiO3
nanograin ceramics, Journal of the European Ceramic Society, 26
(2006) 2855–2859.
X: Haixue Yan, Hongtao Zhang, Rick Ubic, Mike Reece, Jing Liu,
and Zhijian Shen, Orientation dependence of dielectric and
relaxor behaviour in Aurivillius phase BaBi2Nb2O9 ceramics
prepared by spark plasma sintering, Journal of Materials
Science: Materials in Electronics, 17(9) (2006) 657-661.
XI: Xingyuan Guo, Ping Xiao, Jing Liu, Zhijian Shen, Fabrication of
Nanostructured Hydroxyapatite via hydrothermal synthesis
and spark plasma sintering, Journal of the American Ceramic
Society, 88(4) (2005) 1026-1029.
VIII
XII: B. Su, J. Y. He, B. L. Cheng, T. W. Button, J. Liu, Z. Shen, and M.
Nygren, Dielectric properties of spark plasma sintered (SPS)
barium strontium titanate (BST) ceramics, Integrated
Ferroelectrics 61, (2004) 117-122.
XIII: Xingyuan Guo, Julie E. Gough, Ping Xiao, Jing Liu, Zhijian Shen,
Fabrication of Nanostructured Hydroxyapatite and Analysis of
Human Osteoblastic Cellular Response, Journal of Biomedical
materials research: Part A, in press.
XIV: Jing Liu, Zhijian Shen, Haixue Yan, Mike Reece, Yanmei Kan, and
Peiling Wang, Grain-orientated La0.75Bi3.25Ti3O12 ceramics with
un-normal ferroelectric characteristics, Journal of applied
physics, submitted.
IX
X
Table of contents
ABSTRACT........................................................................................................ V
LIST OF PUBLICATIONS............................................................................VII
TABLE OF CONTENTS................................................................................. XI
1 INTRODUCTION..........................................................................................1
1.1
Ferroelectric ceramics ..........................................................................1
1.1.1
Ferroelectrics.................................................................................1
1.1.2
Ferroelectric properties .................................................................2
1.2
(Ba, Sr)TiO3 ferroelectric ceramics.....................................................5
1.2.1
Structure ........................................................................................5
1.2.2
Ferroelectric properties and applications.......................................8
1.3
Bismuth layer-structured ferroelectric ceramics ...............................9
1.3.1
Structures ......................................................................................9
1.3.2
Ferroelectric properties and application ......................................10
1.4
Relation between microstructural features and ferroelectric
properties ...........................................................................................................11
1.4.1
Grain size ....................................................................................11
1.4.2
Grain orientation and texturing ...................................................13
1.4.3
Inhomogeneous compositions and non-equilibrium processing..14
1.4.4
Internal stress and strain engineering ..........................................15
1.5
Spark Plasma Sintering (SPS) ...........................................................16
1.5.1
SPS equipment ............................................................................16
1.5.2
Features of SPS ...........................................................................17
1.5.3
Temperature distribution .............................................................19
1.5.4
Temperature difference ...............................................................21
1.5.5
SPS parameter effects..................................................................25
1.5.6
Applications ................................................................................26
1.6
Aim of this thesis: Nano and Grain Orientated Ferroelectric
Ceramics Produced by SPS ..............................................................................27
1.6.1
Controlling grain size and morphology.......................................27
1.6.2
Achieving 2-D structures ............................................................28
2
EXPERIMENTAL PROCEDURE ........................................................29
XI
2.1
Starting powders.................................................................................29
2.2
Spark plasma sintering.......................................................................29
2.2.1
Sintering ......................................................................................29
2.2.2
Superplastic deformation.............................................................30
2.3
3
Characterization .................................................................................31
THE CONTROL OF GRAIN SIZE AND MORPHOLOGY ..............33
3.1
Shrinkage and densification process .................................................33
3.1.1
Sintering kinetics.........................................................................33
3.1.2
Densification ...............................................................................33
3.2
Grain growth and microstructure evolution ....................................35
3.2.1
Ferroelectric ceramics based on (Ba, Sr)TiO3 .............................35
3.2.2
Bismuth layer-structured ferroelectric ceramics..........................36
3.3
Phase evolution....................................................................................37
3.4
Kinetics ................................................................................................38
3.4.1
Grain growth and kinetic windows .............................................38
3.4.2
Self-pinning effect.......................................................................41
3.4.3
Thermal activation ......................................................................41
3.5
Ferroelectric properties......................................................................42
3.5.1
Ferroelectric properties of (Ba, Sr)TiO3 ceramics.......................42
3.5.2
Ferroelectric properties of nano-sized Bi4Ti3O12 ceramics..........44
4
ACHIEVING 2-D TYPE MICROSTRUCTURE .................................47
4.1
Superplastic deformation yielding textured microstructures .........47
4.1.1
Superplastic deformation kinetics ...............................................47
4.1.2
Textured microstructures.............................................................49
4.2
X-ray studies of textured samples .....................................................51
4.3
Superplastic deformation induced directional dynamic ripening ..52
4.4
Ferroelectric properties......................................................................53
4.4.1
Ferroelectric properties of textured Bi4Ti3O12 ceramics..............53
4.4.2
Ferroelectric properties of textured CaBi2Nb2O9 ceramics..........56
5
XII
SUMMARY..............................................................................................59
6
FUTURE WORK.....................................................................................61
ACKNOWLEDGMENTS ................................................................................62
REFERENCES..................................................................................................64
XIII
XIV
1 Introduction
1.1 Ferroelectric ceramics
1.1.1 Ferroelectrics
The term ferroelectrics refers to materials with spontaneous dipoles that
undergo reversible changes of polar direction when exposed to an
external electrical field with a strength less than the dielectric breakdown
field of the material itself. Thus, three conditions can be discerned to
classify a material as a ferroelectric: (i) the existence of spontaneous
polarization; (ii) possibility to reorient the polarization; (iii) the capacity
of the material to maintain a remnant polarization after being polarized,
i.e. upon the removal of the external electrical field. Based on their
crystal structures, four types of ferroelectric material groups can be
identified, namely: (i) the tungsten-bronze group; (ii) the perovskite
group; (iii) the pyrochlore group; (iv) the bismuth layer-structured oxides
(1)
. Barium titanate, (BaTiO3) and lead zirconate titanate (PZT), both
having perovskite related structures, have so far dominated the basic and
applied research fields of ferroelectric ceramics. The use of PZT ceramics
is environmentally unfavourable due to their high lead content (2), and
compounds belonging to the fourth group have recently been suggested
as lead-free alternatives for ferroelectric ceramics for use at elevated
temperatures.
The interest in ferroelectric ceramics was born in the early 1940s when
the high dielectric constant of barium titanate ceramics was recognized.
Since then, these ceramics have been the heart and soul of several
multibillion-dollar industries, e.g. industries producing high-dielectricconstant capacitors, piezoelectric transducers, positive-temperaturecoefficient devices, electro-optic light valves, and ferroelectric thin-film
memories (1, 3, 4).
1
1.1.2 Ferroelectric properties
Ferroelectric properties are usually low-temperature properties. The
increase of thermal motion at high temperature tends to randomize the
atomic displacements that give rise to the ferroelectric properties. The
temperature at which breakdown occurs is defined as the ferroelectric
Curie Temperature (TC) (5). Upon cooling, the ferroelectric material
undergoes a phase transition at TC and a non-centrosymmetric
(ferroelectric) structure is formed from a centrosymmetric (paraelectric)
one, as exemplified by the temperature dependence of the permittivity of
a BaTiO3 ceramic shown in Fig.1-1. As mentioned above, the
ferroelectric phase transition is associated with small displacements of the
ions from their centrosymmetric positions, or by similar ordering
processes that create a net dipole in the material.
Fig.1-1 Temperature dependence of the permittivity of a BaTiO3 ceramic (6).
2
Fig. 1-2 The hysteresis loops of single crystals of Bi4Ti3O12 (7).
When one increases the voltage applied across a dielectric substance, an
increase of the induced polarization, P, occurs. The polarization observed
on increasing the voltage is not reproduced when the voltage is
decreased, implying that the polarization versus voltage curve exhibits a
hysteresis loop such as that shown in Fig.1-2. Ferroelectrics thus exhibit a
spontaneous polarization (Ps), and a remnant polarization (Pr). The
spontaneous polarization Ps (µC/cm2) of the ferroelectric is characterized
by the net dipole moment density in comparison with the paraelectric
state (8).
The ferroelectricity of ceramics is fundamentally associated with their
domain structures and domain motions (9). When an external electric field
is applied to a ferroelectric, the switching of the adjacent domains of
dipoles along the direction of the applied external field is referred to as
domain reorientation or switching. Domain walls exist between domains
with different polarization orientation. In general, 90° (strain-induced)
and 180° (not strain-induced) domains exist in tetragonal materials,
whereas strain-induced entities of 71°, 109° and 180° domains are
dominant in rhombohedral materials. Macroscopic changes in dimensions
occur when strain-induced domains are switched (4, 8).
3
In general, ferroelectric ceramics are characterized by their remnant
polarization, dielectric and piezoelectric properties. The application of an
electric field across a dielectric leads to a polarization of charge, although
long-range motion of ions or electrons should not occur. For
ferroelectrics such a polarization does not decrease to zero after removing
the electric field, implying the presence of a residual polarization.
Dielectric properties can be defined from the behaviour of the material in
a parallel-plate capacitor:
ε=
C ×d
e0 × A
where ε is the dielectric constant or relative permittivity, e0 is the
permittivity of free space, 8.854 × 10-12 F m-1, C is the capacitance, and
A and d are the area and distance between the plates of the capacitor. ε
depends on the degree of polarization and/or charge displacement that
can occur in the ferroelectrics (5). Under the action of an applied
mechanical stress, piezoelectric crystals polarize and develop electrical
charges on opposite crystal faces. In principle, two effects are thus
operative in piezoelectric ceramics: (i) the direct effect (designated as a
generator) is identified as the electric charge (polarization) generated
from a mechanical stress; (ii) the inverse effect (designated as a motor) is
associated with the mechanical movement generated by the application of
an electric field (1). The polarization, P, and stress, σ, are related to the
piezoelectric coefficient, d, by (5)
P = dσ
where d (often denoted d33) thus is the piezoelectric coefficient (1) . The
direct and inverse effects of piezoelectricity are illustrated in Fig.1-3.
4
Fig.1-3 Direct and inverse effects of piezoelectricity (d33) (10). The figure shows that d33
relates to the generation (direct effect) of a polarization response (Pr) parallel to the
direction of the mechanically applied stress (σ), where Q3 means charges. U3 represents
the elongation of the sample when an external voltage (field) E3 is applied.
1.2 (Ba, Sr)TiO3 ferroelectric ceramics
1.2.1 Structure
BaTiO3 has the perovskite (ABO3) structure, and its unit cell is shown in
Fig.1-4. The structure consists of a corner-linked network of oxygen
octahedra, with Ti4+ ions occupying the B sites within the octahedral cage
and the Ba2+ ions are located at the interstitial A positions created by the
linked Ti4+– oxygen octahedra (1, 11, 12).
Fig.1-4 Unit cell of BaTiO3.
5
Fig.1-5 Structural modifications and associated phase transition temperatures of BaTiO3.
The structural modifications and associated phase transition temperatures
of BaTiO3 are illustrated in Fig.1-5. The symmetry of BaTiO3 changes
from rhombohedral to orthorhombic to tetragonal to cubic with increasing
temperature (12-15). The cubic phase has the space group Pm-3m and a unit
cell parameter of 4.0340 Å(16), while the space group of the tetragonal
phase is P4 mm with lattice parameters a = 3.9945 and c = 4.0335 Å (12) .
The phase transition from orthorhombic to tetragonal occurs at -5oC, and
the tetragonal phase changes to cubic at 127 oC (6, 12-15).
Many ions, e.g. Sr2+ and Pb2+, can substitute for the Ba2+ ions, giving
rise to modifications of the lattice parameters and phase transition
temperatures of Ba1-x (Sr/Pb)xTiO3 (2). Fig.1-6 reveals a tendency of the
lattice shrinking with the doping of Sr2+ in Ba1-xSrxTiO3. Thus at room
temperature the tetragonal phase remains stable when x < 0.4, while for
x > 0.4 the cubic phase is stable, implying that TC is shifted towards
lower temperatures with increasing doping levels.
6
4.04
4.02
Cubic
a
Lattice parameter (Å)
4.00
3.98
3.96
3.94
Tetragonal
c
a
3.92
3.90
3.88
3.86
3.84
0.0
0.2
0.4
0.6
0.8
1.0
x in Ba1-xSrxTiO3
Fig.1-6 Unit cell parameters of Ba1-xSrxTiO3 plotted versus x.
Polarization Up
Polarization Down
Fig.1-7 Two out of the six possible polarization states produced by
displacement of the central cation.
7
Fig.1-8 A SEM micrograph of the domain structure of BaTiO3 (17).
In BaTiO3 the polarization is ascribed to the displacement of the central
Ti4+ ions, as shown in Fig.1-7, where displacement of Ti4+ ions along the
c axis in the tetragonal unit cell is illustrated. Other possible displacement
directions could be along a or b axes in the orthorhombic structure. The
views of “polarization up” and “polarization down” (representing 180°
polarization reversal) seen in Fig.1-7 thus represent two of the six
possible polarization directions. Domains are formed during the
paraelectric–ferroelectric phase transition of BaTiO3 ceramics to relax the
stresses induced by the phase transformation. Often, BaTiO3 ceramics
have two types of domains: (i) herringbone 90° domains (ii) and square
180° domains, as illustrated in Fig.1-8 (9).
1.2.2 Ferroelectric properties and applications
BaTiO3 is well known for its remarkable dielectric properties, with
permittivity values larger than 4000 in the vicinity of TC (127oC) (6).
Normally, ferroelectric materials have high specific electrical resistivity
13
(>10 Ω·cm) (1), d33-values (piezoelectric constant) around 180 (pC/N) (18),
and the hysteresis loops exhibit almost linear parts (1). (Ba, Sr)TiO3
ceramics were the first piezoelectric ceramic transducers ever developed,
8
and (Ba, Sr)TiO3 ceramics have been widely used in recent years as
capacitors, thermostats and electro-optic devices (1, 19).
1.3 Bismuth layer-structured ferroelectric ceramics
1.3.1 Structures
The family of bismuth layer-structured compounds with perovskite
related structures are known as Aurivillius phases. The general formula of
these phases is (Bi2O2)2+(Am-1BmO3m+1)2-, where A is a mono-, di- or
trivalent element (alone or in combination) with cuboctahedral
coordination, B is a transition element octahedrally coordinated to six
oxygen ions, and m is the number of layers of octahedra in the perovskite
slab. The m value can vary from 1 to 6 (20). The structures of most of
Aurivillus phases, except for Bi4Ti3O12, have orthorhombic symmetry (21).
The crystal structure of Bi4Ti3O12, which is the most carefully
investigated compound in this family, is characterized by the occurrence
of pseudo-perovskite layers (Am-1BmO3m+1)2- stacked between (Bi2O2)2+
layers as shown in Fig.1-9 (22). Bi3+ is located on the A-site and Ti4+ on Bsite in the perovskite subcell, and m is equal to 3. The tetragonal
paraelectric phase of Bi4Ti3O12 transforms into the monoclinic
ferroelectric phase with the space group of P1a1 at 675 oC (21). This phase
transformation results in a number of different non-180o domains, which
can be displaced by an external field and thus contribute to the
piezoelectric performance of the material (23).
9
Fig. 1-9 Crystal structure of Bi4Ti3O12.
1.3.2 Ferroelectric properties and application
Many compounds in the family of Aurivillius phases exhibit high Curie
temperatures. Due to their layered structure feature, the crystals of this
family of compounds have highly anisotropic properties (24, 25). The
polarization occurs along the a-b plane (26) and the electrical conductivity
is also higher along this plane (27). Calculations have shown that the
polarization of Bi4Ti3O12 along the a axis (28µC/cm2) is much higher than
that along the c axis (5µC/cm2) (26, 28), whereas the polarization along the
b axis is suppressed due to the presence of glide planes perpendicular to
this axis according to powder XRD research (28). Typical ferroelectric data
of Bi4Ti3O12 are given in Fig.1-10 and Fig.1-2 (7, 23, 29).
Bismuth layer-structured ferroelectrics have been proposed for hightemperature piezoelectric applications and as high temperature sensors in
automotive, aerospace, and power generating devices, as well as in
connection with chemical and materials processing (30). Recently,
10
compounds of this type have been suggested as candidate materials for
non-volatile ferroelectric random-access memories (FeRAMs) (4, 31-33).
Fig. 1-10 Temperature dependence of the permittivity of a single crystal of Bi4Ti3O12 and
of a polycrystalline sample (29).
1.4 Relation between microstructural features and
ferroelectric properties
1.4.1 Grain size
It has been verified that the permittivity of BaTiO3 based ferroelectric
ceramics strongly depends on the grain size. Previous work is
summarized in Fig.1-11 (17, 34, 35). Commonly achieved experiences
include:
(i)
The permittivity of ceramics containing grains larger than 10
µm is almost grain size independent, e.g. the coarse-grained
(20 ~ 50 µm) ceramics of pure BaTiO3 show permittivity
values in the range 1500~2000 at room temperature (34).
11
(ii)
(iii)
(iv)
When the grain size is reduced to a few micrometers, the
permittivity at room temperature increases notably (17).
The highest values of permittivity at room temperature were
observed in ceramics with an average grain size of 0.7~1 µm
(34)
.
Decreasing the grain size even further yields an almost
temperature independent permittivity (17, 34, 36, 37).
6000
o
measured at 25 C
o
measured at 70 C
5000
εr
4000
3000
2000
1000
0
0,1
1
10
a (µm)
100
Fig. 1-11 Dielectric constants of BaTiO3 ceramics measured at 25 and 70 oC, plotted
versus the average diameter of the grains (a) expressed in µm (34).
The grain size dependence of the permittivity of BaTiO3 based ceramics
is strongly related to their domain structures. The equilibrium width of
90° domains is almost constant when the grains are larger than 10 µm,
while it narrows when the size of grains is decreased (34). In other words,
the observation that the room temperature permittivity has a pronounced
maximum at a grain size of 0.7~1 µm is attributed to an increase of
domain wall mobility (34). It has been suggested that the reduction of
permittivity with further decreasing grain size possibly results from
atomic (dipolar) structural changes (17, 34, 38). It is also possible that such
reduction is due to the presence of large amounts of defects
(ferroelectrically dead layers) at the grain boundaries in nano-sized
ferroelectric compacts (37, 39).
12
It is well established that coarse-grained ceramics with an average grain
size of 10~50 µm can be produced by conventional sintering methods,
e.g. pressureless sintering (PLS) (34, 40). Fine-grained BaTiO3 ceramics, i.e.
compacts with grain sizes in the range 1~10 µm, can be preferably
obtained by the hot isostatic pressing (HIP) and hot pressing (HP)
processes (41-44). The real challenge is, however, to prepare truly nanosized bulk ceramics. In this context, it has been reported that nano-sized
dense (Ba, Sr)TiO3 ceramics can be prepared by cold isostatic pressing
followed by hot-pressing (34), and by Spark Plasma Sintering (SPS) (36, 37,
45, 46)
. In particular, SPS has been recognized as a versatile tool that
provides unique possibilities to engineer the sintering kinetics so as to
yield microstructures with tailored grain sizes.
1.4.2 Grain orientation and texturing
Polycrystalline ceramics often exhibit isotropic physical properties even
if the individual grains that constitute the compact are anisotropic,
because the grains usually are randomly orientated with respect to their
crystallographic symmetry. The layered crystal structural feature of
bismuth layer-structured ceramics promotes preferential growth of grains
in directions perpendicular to the stacking axis of the layers, yielding
crystals with plate-shaped habits (28). It has been demonstrated that
ceramics with grain-orientated microstructures exhibit anisotropic
properties.
The aim of aligning anisotropic grains in ceramics is to mimic the
properties of anisotropic single-crystal components, and two approaches
have been implemented to produce textured or grain-orientated
microstructures: (i) Preparation of powders preferably with
needle/platelet morphology and aligning the grains via a shear flow
process followed by pressure-less sintering and/or hot-forming processes
for extended periods of time at high temperatures, to allow the growth of
aligned grains according to the Ostwald ripening mechanism (47-49) ; (ii) In
13
order to improve the grain alignment even further, a small number of
well-developed large (2-5 µm) needles/platelets are added as grain
growth templates (50, 51) . However, with these established techniques few
have succeeded to align grains smaller than a few micrometers (49, 50, 52).
1.4.3 Inhomogeneous compositions and non-equilibrium
processing
(Ba, Sr)TiO3 ferroelectrics exhibiting a flat temperature dependence of
high permittivity around room temperature are desirable for applications
such as capacitors (1). It has been demonstrated that compositionally
graded (Ba, Sr)TiO3 materials can yield an expanded ferroelectricparaelectric transition temperature range.
The ferroelectric properties of ferroelectric ceramics are also strongly
related to the composition. Thus, it has been reported that BaxSr1-xTiO3
capacitors composed of phases with different x-values have several
temperature dependent dielectric peaks, or an almost flat temperature
dependence of high permittivity (53, 54). Pb(Zr0.53Ti0.47)O3 of the
morphotropic boundary composition, i.e. containing both the
rhombohedral and tetragonal phases, show the highest piezoelectric d33
constants, dielectric constants and electromechanical coupling factors (8).
Structural/compositional inhomogeneities can be obtained on different
length scales: at an atomic level, i.e. materials containing different
structural entities; at a micron-sized level, i.e. materials containing coreshell structured grains; at a macro-sized level, i.e. laminated materials
gradients formed by, for instance, casting suspensions of different
compositions/grain sizes.
It has been verified that an appropriate lamination of the selected
components having various phase transition temperatures could yield a
desirable flat temperature permittivity profile. Thus, laminated Ba(Zr1xTix)O3 ceramics with x values varying from one end to the other of a
14
cylindrical pellet have been prepared and found to yield a flat temperature
dependence of permittivity in the range of 20-120oC (55, 56).
Recently, the possibility of achieving materials with structural/
compositional inhomogeneities has been extended by the use of new
sintering processes such as SPS technique. In this context, it has been
reported that dense (Ba, Sr)TiO3 ceramics containing several predesigned phase compositions can be prepared by spark plasma sintering
(54)
.
1.4.4 Internal stress and strain engineering
It is well known that strains have a profound influence on the
ferroelectric properties of ceramics. The existence of internal residual
stress in fine-grained ferroelectric ceramics, which is generated by the
absence of 90o twinning within the grains, gives rise to an increase of the
permittivity as the ceramic cools below the Curie temperature (57). For
example, fine-grained BaTiO3 ceramics have a very high permittivity of
~4000 at room temperature (34, 35, 57, 58). Recently, novel ferroelectric
properties have been achieved in (Ba, Sr)TiO3 ferroelectric thin films by
introducing pre-designed strains into their structures, yielding an increase
of the ferroelectric transition temperature (TC) up to nearly 500oC, and
achieving a remnant polarization approximately 250% higher than that
observed for BaTiO3 single crystals (59, 60). However, one should be aware
that very high strains give rise to crack formation (19, 61). On the other
hand, the low permittivity of BaTiO3 bulk ceramics consisting of truly
nano-sized grains has partly been ascribed to strain, as discussed above
(17, 38, 39)
.
An important issue is to modify the performance of ferroelectrics by
strain engineering. In connection with phase transitions, strains are
introduced in ferroelectric compacts, and to the extent that the compacts
contain phases of different structures or phases of the same structure but
different compositions, strains are introduced due to the mismatch of the
15
thermal expansion coefficients. These types of strain can all be
engineered. In the case of BaTiO3 based ceramics, the formation of
compositional core–shell grain structures during sintering through a
dissolution-precipitation process has been used to introduce strains in
fully dense compacts (53, 62-64). Recently, it has also been demonstrated
that strains can be introduced into thin films of BaTiO3 by growing these
thin films on single-crystal substrates that have lattice parameters
somewhat different from those of BaTiO3, e.g. GdScO3, DyScO3 (59, 60).
1.5 Spark Plasma Sintering (SPS)
1.5.1 SPS equipment
Spark Plasma Sintering (abbreviated SPS) is a comparatively new
sintering process that was developed from the mid-1980s to the early
1990s and currently attracts growing attention among productions
engineers as well as materials researchers (65). The basic configuration of
an SPS unit is shown in Fig.1-12. It consists of a uniaxial pressure device,
in which the water-cooled punches serve also as electrodes, a watercooled reaction chamber that can be evacuated, a pulsed DC generator,
and a position- and temperature-regulating system.
Fig. 1-12 Basic configuration of a typical SPS set-up (66).
16
1.5.2 Features of SPS
SPS resembles the hot pressing (HP) process in several respects, i.e. the
precursor powder (green body) is loaded in a die, and a uniaxial pressure
is applied during the sintering process. However, instead of using an
external heat source, a pulsed direct current is allowed to pass through the
electrically conducting pressure die and, in appropriate cases, also
through the sample. This implies that the die also acts as a heat source
and that the sample is heated from both outside and inside.
Due to the high conductivity of the pressure dies, low voltages (max. 15
V) and strong currents (max. 5500A) are used in our set-up (Dr. Sinter
2050 SPS). The use of pulsed direct current also implies that the samples
are exposed to a pulsed electric field during the sintering process. The
process inventors originally claimed that the pulses generate spark
discharges and even that plasma is created between the powder particles,
which explains why the process is named spark plasma sintering.
However, today most researchers do not believe in the occurrence of
plasma, and the presence of sparks is still under debate.
Sintering temperatures up 2200oC can be used in the SPS unit, but in
the high-temperature region we can not prepare samples with diameters
larger than 20 mm, due to current limitation: larger dies require stronger
currents than small ones. In most cases the temperature was measured
with a thermocouple inserted into the graphite die, but for sintering
temperatures exceeding 1000oC the temperature was recorded by an
optical pyrometer focused on the surface of the die. Temperature and
temperature distribution plays a very important role in the SPS process,
and it is accordingly discussed in some detail below.
One of the advantages of the SPS set-up is that we can apply a much
higher uniaxial pressure than that possible in most HP units. A maximum
sintering pressure of 200 KN can be applied in our unit.
17
A unique feature of the SPS process is the possibility of using very fast
heating rates (up to 600oC min-1) and very short holding times (minutes)
still obtaining fully dense sample. Again the maximum heating rate that
one can use depends on the size of the pressure dies, i.e. it is difficult to
use heating rates higher than 200oC min-1 for a die with an inner diameter
of 50 mm, whereas it is possible to use a heating rate of 600oC min-1 for a
die with an inner diameter of 12 mm. However, too fast heating rates will
strongly affect the temperature distribution (see below). In this study
heating rates in the range of 50-200oC min-1 have been used.
The pulses have a duration of 3.3 ms, and a pulse sequence consisting
of twelve pulses followed by a period of 6.6 ms of zero current (12:2)
was used in this study. Such an on/off pulse sequence is shown in Fig.113. The on/off pulse sequence can be adjusted from 99:1 to 1:9.
Figure 1-13 The on/off pulsed sequence used in this study.
Thus, the consolidation rate in SPS is greatly enhanced, and the sintering
temperature can be a few hundred degrees lower than that typically used
in conventional sintering processes such as HP. Four factors that
contribute to the enhanced densification rate can be discerned: (i) the use
of rapid heating and cooling rates; (ii) the rapid transfer of heat because
the die itself acts as a heating element; (iii) the application of a
mechanical pressure exceeding that used in the conventional hot-pressing
18
process; (iv) the use of a pulsed DC current to heat the sample, implying
that samples are also exposed to a pulsed electric field during sintering.
1.5.3 Temperature distribution
It has been a very hot topic to study the temperature distribution in the
SPS samples and dies. The temperature distribution is strongly dependent
on parameters such as size and shape of the pressure die and associated
punches, their thermal and electrical conductivities, the pressure and
heating rate used, the contact resistance between the punches and the die,
the electrical and thermal properties and size of the sample to be
compacted, etc.
In most modelling experiments it is assumed that the geometry of the
die is centric with an appropriate size ratio between dies and punches, and
it is assumed that the thermal and electrical conductivities of the die and
punches are homogenous, so as to avoid a temperature gradient within the
sample region. According to our experience, the temperature distribution
in a small die is more homogeneous than in a large die, i.e. we can easily
obtain fully homogeneous, transparent nano-sized hydroxyapatite ceramic
plates with a diameter of 12 mm and a thickness of 2-3 mm, whereas it is
very difficult to sinter fully transparent plates with a diameter of 30 mm
and the same thickness. The experimental experience we have suggests
that the temperature in the central part of large die is lower than in the
vicinity of the die walls. Recently, we have also found that the contact
resistance between the punches and the die is of great importance. Higher
contact resistance results in an inhomogeneous temperature distribution.
It has also been verified that the electrical properties of the specimens
have an important influence on the temperature distributions inside the
graphite die as well as in the specimens. Thus, in a non-conducting
sample (Y-doped ZrO2) one has observed larger thermal gradients than in
the case of an electrically conductive sample (TiN) (67), indicating that the
19
temperature distribution within the non-conducting sample is not as
homogeneous as within a conducting sample.
According to our experience, the temperature distribution in thin
samples is also better than in thicker ones, especially in the hightemperature region (T > 1000oC). However, in the low temperature
region (T < 1000oC), the temperature distribution within thick samples
(4-6 mm) is quite homogeneous, especially for samples having diameters
in the range of 8-12 mm, and when low heating rates are used (20-40oC
min-1).
Experiments and models have shown that high heating rates also
strongly affect the temperature distribution, especially when nonconducting materials are compacted. A high heating rate (> 100oC min-1)
implies that the final temperature overshoots the pre-set, typically by 2575oC depending on the size of the die. Larger dies yield higher overshoot,
smaller ones lower. An example is that the temperature overshoot is
approximately 50oC when applying the heating rate of 100oC min-1 for a
die with outer and inner diameters of 35 mm and 15 mm, respectively (68).
A more homogeneous temperature distribution is obtained within a
minute or two, depending on the size of the sample (die). Anyhow, the
temperature gradient associated with the use of high heating rates in
combination with short dwell times can give rise to non-uniform
microstructures, and accordingly inhomogeneous mechanical properties
(67)
. The problem with temperature overshoot can be avoided by changing
the heating rate from high to low some 50oC below the preset maximum
temperature. Doing so, the temperature distribution at the preset
temperature will be much more homogeneous, while the dwell time will
be approximately half a minute longer than preset one.
The holding time also influences the temperature distribution within
samples. Experiments verified that increasing holding times yield a more
homogeneous temperature distribution within the sample, especially for
non-conduction materials. Here the thermal conductivity of the sample
20
plays an important role. In this connection it can be noted that the thermal
conductivity varies strongly with the density of the sample.
The mechanical pressure also affects the temperature distribution
within samples. Experiments showed that a lower pressure yielded a more
homogeneous temperature distribution within the sample, especially for
non-conducting materials.
In general, the presence of a temperature distribution (gradient) within
the sample promotes the thermal diffusion processes, i.e. it promotes
densification, sintering and grain growth processes. It goes without
saying that the presence of temperature gradients also can give rise to
non-uniform microstructures. Thus, densification experiments with nanosized SrTiO3 powders have proved that fully densified nano- structured
SrTiO3 ceramics are obtained using a heating rate of a 100oC min-1.
However, when a heating rate of 200oC min-1 was applied the obtained
ceramics exhibited inhomogeneous microstructures containing abnormal
micro-sized grains embedded in micron sized ones, i.e. a wider
temperature distribution (steeper gradient) in combination with an
increased temperature overshoot yields inhomogeneous microstructures.
In most of my densification experiments I have used a graphite die with
inner and outer diameters of 12 and 40 mm, respectively, and a height of
60 mm, a heating rate of 100 oC min-1, a pressure of 50 MPa, and similar
amounts of powders in order to obtain comparable sintering curves.
1.5.4 Temperature difference
It is established that the temperature at the surface of the pressure die is
lower than that within the sample. This implies that the recorded
(measured) temperature is lower than the “experimental” (actual)
temperature. Fig.1-14 illustrates the position of seven temperature sensors
used in an early work (69). The surface temperature is normally recorded
21
by an optical pyrometer while the other positions are normally furnished
with thermocouples.
Fig. 1-14 Schematic description of the positions of seven temperature sensors. The 1, 2, 3
die positions, the centre position, border position and bottom position are normally
furnished with thermocouples, while the temperature at the surface-position is normally
recorded with an optical pyrometer focused onto the surface position (adopted from
reference (69)).
In one experiment, a die with inner and outer diameters of 10 and 40 mm,
respectively, and a height of 60 mm was used. The dies were filled with
titanium and alumina powders, respectively, for the purpose of yielding
fully dense tablets with a diameter of 10 mm and a height of 10 mm. A
pressure of 37.5 MPa was used, and a square pulsed current with an
on/off pattern of 1:1 and a pulse-discharge time and -cut time of 100 ms
was used to heat the samples. The temperatures were recorded at the
centre and the 1-3 die positions (see Fig. 1-14). In the case of titanium no
temperature difference was recorded up to 1000oC (recorded at centre
position), but when the temperature reached to 1300oC (recorded at centre
position) in the alumina compact, the temperature at this position was
22
slightly lower than that of the 1-3 die positions; no temperature difference
between the 1-3 positions was discerned, however (69). The experimental
conditions used in this study are quite different from those used in my
experiments, implying that these findings are difficult to compare
adequately.
In another experiment the temperature at the bottom position was
compared with that of the surface position, see Fig.1-14, using the same
SPS unit as ours, and applying a pulse sequence of 12:2, a pressure of 50
MPa and a heating rate of 200oC min-1. In this case a die with an inner
diameter of 19 mm, an outer one of 45 mm, and a height of 38 mm was
used. Alumina was used, and the final height of the fully dense tablet was
3 mm. The temperature at the surface of the pressure die was monitored
by the optical pyrometer, while the one at the bottom position was
monitored by a thermocouple. At low temperatures (around 600oC, the
lower limit for our optical pyrometer), the difference between at bottom
position and surface position was negligible, but with increasing
temperature the difference increased, and the temperature at the bottom
position was about 150oC higher than that of the surface position at
1350oC (recorded at the surface position) (70, 71).
In another SPS experiment the temperature difference between the
border and centre position was studied, using the same SPS unit as ours
(72)
. Graphite dies with inner and outer diameters of 40 mm and 90 mm,
respectively, and a height of 60 mm were filled with TiB2 and BN
powders, and the samples were sintered using a heating rate of 170oC
min-1. The temperature difference increased with increasing temperature,
and at 1700oC (recorded at the border position) the temperature at this
position was 450oC higher than that of the centre position (72). This
suggests that large samples in combination with high heating rates yield
large temperature differences.
In another experiment, the temperature difference between surface and
sample interior was determined by melting silicon and lithium silicate
23
powder. In this case a die with inner and outer diameters of 19 and 44.6
mm, respectively, and a height of 38.1 mm was used. A heating rate of
15oC min-1 and a pressure of 15MPa were applied, using a pulse on/off
pattern of 12:2. The experimental temperatures (surface) were monitored
up to 1100oC via a K-type platinum thermocouple attached to the die
surface, while a single-colour pyrometer was used for higher
temperatures. The result was that the actual temperature exceeded the
“experimental” (surface) temperature, and that the difference increased
considerably from 65oC at 650oC (focus on the surface position), 170oC at
1030oC (focus on the surface position) to 240oC at 1180oC (focus on the
surface position) (73). In the same article, Zavaliangos et al. state that the
temperature difference between the surface position and the centre
position in electrically conducting materials, e.g. graphite, is about 1015% less than in non-conductive materials, e.g. alumina (73).
The temperature distribution at a fixed temperature has also been
studied by modelling calculations. The theoretical analysis of the
temperature distribution was based on Fourier’s law and Ohm’s laws.
These calculations show that the heat transfer between the graphite parts
and the samples depends on the presence of contact resistances and the
properties of sample, i.e. whether the sample is a good electrical and/or
thermal conductor. Accordingly, the calculated temperature distribution
for good conductors such as titanium is different from that of nonconductors, e.g. alumina and zirconia (67, 69, 70, 73). Including a
consideration of the influence of thermal and electrical contact
resistances, the temperature distribution calculated by modelling matches
the results obtained by experiment much better than without considering
the contact resistances (67, 73).
Dr. Salamon in our group has also estimated the temperature difference
between the centre of the sample and the surface of the die. The die
surface temperatures were recorded with an optical pyrometer focused on
the surface of a graphite die with inner and outer diameters of 15 mm and
35 mm, respectively, and a height of 30 mm. The sample consisted of two
24
pre-sintered BN-Si3N4 plates with a small cavity in the middle. Gold,
silver and platinum wires were placed in this cavity. The samples were
heated to temperatures in the range 900-1700oC using a heating rate of
100oC min-1, a mechanical pressure of 50 MPa and a holding time of 3
min. The melting temperatures of these metals were used as fixed points.
The outcome of these experiments were as follows: During the heating
part of the experiment the temperature at the centre is estimated to be 120
degrees higher than at the surface of the die, and 60 degrees higher under
isothermal heat treatment conditions (68). These results are similar to
previous findings by Anselmi-Tamburini, U et al. (70, 71).
Generally speaking, both experimental findings and calculations seem
to indicate that the temperature difference between the surface of the die
and the sample interior is comparatively small at low temperatures,
whereas the difference increases with increasing temperature. At high
temperatures (> 1100oC), the temperatures at the border position or
bottom position (see Fig 1-14) are higher than at the surface position, and
also higher than at the centre position. However, the values of those
temperature differences vary depending on the size and shape of the die
and punches, the thermal and electrical properties of the material to be
sintered, and if the die is or is not properly thermal isolated, etc. Also,
using a graphite blanket wrap on the die surface during SPS high
temperature sintering can efficiently diminish temperature differences.
Thus, the calculated temperature difference could not always match with
the experimental results.
1.5.5 SPS parameter effects
During an SPS process the temperature is always of great importance,
whether the studies concern densification and/or sintering rates, grain
growth, plastic deformation and/or solid-state reactions, since all these
processes are thermally activated.
25
At temperatures below an onset temperature (Ton), it is very difficult to
obtain fully dense ceramic samples, even if one applies longer holding
times or high pressures. The effect of the holding time for the
densification process is positive at temperatures exceeding Ton. The grain
growth rate is substantially much higher at temperatures above a certain
critical temperature (Tg) than at T < Tg (46).
It is well known that the application of mechanical pressure promotes
the removal of pores and enhances diffusion, implying that increasing
pressures may enhance the densification process. At temperatures above
Ton, the effect of the pressure for the densification process is efficient,
whereas at < Ton this effect is negligible (36, 45, 46, 71).
The effect of different pulse patterns has not been studied extensively.
It has, however, been reported that increasing the on/off ratio yielded
smaller grain sizes above Tg but not below Tg, and that the densification
rate curve is shifted towards higher temperatures as the pulse on/off ratio
is increased (46, 71, 74).
It has been suggested that enhanced densification efficiency in the SPS
process can in part be explained by the application of a pulsed electric
field. The grain-boundary diffusion and grain-boundary migration
processes should thus be enhanced by the presence of a pulsed electric
field (46). It has been shown that sintering in the presence of a liquid phase
is greatly enhanced by a pulsed electric field, and so is the deformation
processes at elevated temperatures. Sialon samples are thus
superplastically deformed around 1500oC at a rate of ~10-3 s-1 in the SPS
unit, while identical experiments in a HP set-up yielded compressive
strain rates on the order of 10-5 s-1 at several hundreds degrees higher
temperatures (46, 75).
1.5.6 Applications
26
SPS has been successfully applied to preparing many kinds of materials,
e.g. advanced alloy materials, functional graded materials (FGM), finegrained ceramics, amorphous materials, target materials, thermoelectric
materials, nano-composites (46, 65, 76-82). More recently, SPS has been used
more extensively to sinter ferroelectric ceramics (36, 37, 40, 45, 54, 83-85).
SPS can also be applied for investigation of the sintering kinetics and
for rapid densification of various materials paired with minimized grain
growth (36, 45, 76, 79-82, 86-88). By controlling the sintering kinetics, tailored
microstructures can thus be produced, and accordingly materials with
tailored properties.
1.6 Aim of this thesis: Nano and Grain Orientated
Ferroelectric Ceramics Produced by SPS
The main work of this thesis has been focused on fabrication of nano and
grain orientated ferroelectric ceramics by SPS, and also their ferroelectric
properties.
1.6.1 Controlling grain size and morphology
One of the main purposes of the present work is controlling the grain
size, from nano to micro, and the grain morphology of ferroelectric
ceramics. We have carried out spark plasma sintering of five nano-sized
powders, namely Bi4Ti3O12, BaTiO3, SrTiO3, Ba0.6Sr0.4TiO3, and a powder
mixture of BaTiO3 and SrTiO3 having a Ba/Sr mole ratio of 6/4. The
latter mixture is used to study the influence of the solid-state reaction that
leads to the formation of Ba0.6Sr0.4TiO3 on the kinetics of densification
and grain growth. The aim is to outline a procedure that allows us to
define a “kinetic window” within which it is possible to densify nanopowders of these compounds into fully dense compacts containing grains
of almost the same size as the staring powders.
27
1.6.2 Achieving 2-D structures
Another purpose of this work is to establish a new advanced processing
route to effectively achieving anisotropic 2-D structures with tailored
grain morphology by SPS, i.e. producing grain-aligned lead-free
ferroelectric ceramics with improved performance properties. This work
was initially motivated by the idea of rationally combining our two recent
findings that: (i) Rapid anisotropic grain growth occurs above the onset
temperature (Tg) of grain growth when high heating rates are applied, via
a dynamic ripening mechanism (76); (ii) Nano-ceramics can easily undergo
rapid superplastic deformation in the presence of a pulsed electrical
current/field (75).
28
2 Experimental procedure
2.1 Starting powders
Commercially available monophasic nano-crystalline powders of BaTiO3
(abbreviated BT below), SrTiO3 (ST), Ba0.6Sr0.4TiO3 (BST64) produced
by TPL Inc., Albuquerque, NM, USA, were used in this study. The
particle sizes of the powders are ca. 60~80 nm. The starting powders
were used as received. A powder mixture of BT and ST with a Ba/Sr
mole ratio of 6/4 was produced by ball milling for 24 h, using zirconia
balls and 2-propanol as milling media. The mixed powder was dried at
40°C. This powder mixture is abbreviated MBST64 below (54).
The nanocrystalline Bi4Ti3O12 (abbreviated BIT below) powder was
prepared by a hydrolysis technique. Bi(NO3)3⋅5H2O and Ti(OC4H9)4 were
used as starting materials. Powders calcined at 600°C proved to be well
crystallized single-phase with a BET surface area of 12 m2/g and an
average particle size of 100 nm (52).
The CaBi2Nb2O9 (CBNO) powder was prepared by heat treatment of
appropriate mixtures of Bi2O3 (99.975%), CaCO3 (99%), Nb2O5 (99.9%)
at 950oC for 4 h (40). The grain size of the resulting powder was in the
range of 0.7 to 1 µm.
2.2 Spark plasma sintering
2.2.1 Sintering
Samples of 12 mm in diameter and 2 mm in thickness were prepared in
vacuum under a uniaxial mechanical pressure of 50 MPa. In some cases a
pressure of 75 MPa or 100 MPa was applied to insure densification at the
lowest sintering temperatures. The pressure was applied at room
29
temperature and held constant until the end of the sintering cycle. The
peak current and the voltage reached in the present cases are ca. 2000 A
and 4 V, respectively. In most cases the temperature was measured with a
thermocouple inserted into the graphite die at a position 2 mm from
surface into the die along Y axis (see Fig.1-14), but when the sintering
temperature exceeded 1000oC the temperature was recorded by an optical
pyrometer focused on the surface of the die. The set-up is provided with a
dilatometer for recording the shrinkage and shrinkage rate, and these data
were stored on a computer. The following parameters were also recorded:
temperature, pressure, current and voltage. The linear shrinkage and
shrinkage rate discussed below are defined as -∆L/L0 and -d(∆L/ L0)/dt,
respectively, with L0 being the thickness of the sample at room
temperature with pressure applied. The ∆L-values were corrected for the
contribution related to the expansion of the die. Since the mass and
diameter of the sample are constant during the SPS process, the linear
shrinkage defined above is also a measure of the volume shrinkage.
Initially the sintering behaviour of each kind of powder in the SPS unit
was tested by consolidating a sample of 12 mm in diameter and 5 mm in
thickness, using a heating rate of 100 oC min-1 and a pressure of 50 MPa,
in order to determine the onset temperature of densification (Ton) and the
temperature at which the sample had achieved its final density (Tfin).
Thereafter, a series of samples were prepared by selecting isothermal
sintering temperatures (Tiso) within the temperature interval that ranges
from slightly below Ton to well above Tfin, in order to investigate the
evolution of microstructure with temperature and to find out the
temperature where substantial grain growth occurs. The cooling rate from
Tfin down to 500oC was approximately 350oC min-1.
2.2.2 Superplastic deformation
Compressive deformation tests were performed in the SPS apparatus.
Fully dense cylindrical compacts of Bi-based ceramics containing
equiaxed nano-sized grains were prepared by SPS. Such a compact,
30
having a diameter of 12 mm and a height of ∼6 mm, was loaded into a
graphite die having an inner diameter of 20 mm. This die was heated at a
rate of 100oC min-1 under a constant uniaxial compressive load that
corresponded to an initial compressive stress of 40 MPa, applied via the
punches of the graphite die, implying that the compressive stress level
was decreased to 20 MPa at 50% strain. The compressive deformation
strain and stain rate discussed below are defined as -∆Ld/ L0, and -d(∆Ld/
L0)/dt, respectively, where ∆Ld and L0 represent the shrinkage of sample
height and the original height of the sample before deformation,
respectively. Both of the ∆Ld and L0 values were corrected for the
contribution related to the expansion of the die.
All sintered and deformed samples were annealed in a muffle furnace in
air in the temperature range 650-1000oC for 200-300 min to remove the
surface graphite and ensure full oxidation before any further
characterization.
2.3 Characterization
The bulk densities were measured according to Archimedes’ principle.
The microstructures of the samples were evaluated from micrographs of
fractured or thermally etched surfaces, recorded in a scanning electron
microscope (SEM; Model 880, JEOL, Tokyo, Japan). The average grain
sizes were determined by an image analysis program (Image tool,
UTHSCSA) through investigating more than 100 grains in SEM images.
Selected samples were also examined by transmission electron
microscopy (TEM). The crystalline phases were characterized by X-ray
powder diffraction (XRD) studies, using a focusing Guinier-Hägg camera
with CuKαl radiation and silicon as an internal standard. Additional X-ray
powder diffraction data were collected on a STOE SCADIP powder
diffractometer (CuKαl radiation, Ge monochromator) equipped with a
linear PSD detector, using a rotating sample in symmetric transmission
31
mode. The unit cell parameters were calculated with the program TREOR
(89)
. The refinement was made using GSAS and 109 reflections for d ≥
1.46 Å. The Lotgering orientation factor f =
(p − p 0 )
∑ I 00l
with p =
(1 − p 0 )
∑ I hkl
and P0 for a random-orientation powder pattern was calculated using
intensities for 120 reflections with 2θ ≤ 69.5°. The preferred orientation
was modelled in the refinement by the March-(Dollase) function (90),
O ph
⎛
sin 2 A ⎞
⎟
= ⎜⎜ R 02 cos 2 A +
R 0 ⎟⎠
⎝
−3 / 2
where A is the angle between the preferred orientation axis, (00l), and the
reflection vector.
The dielectric properties were measured at different frequencies, using
a HP 4194A impedance analyzer within the temperature range –70 to
70°C, and using an Agilent 4284A LCR meter in the high-temperature
region, e.g. from room temperature to 1100°C. Ag was used as electrode
material. The piezoelectric constant (d33) was measured by a quasi-static
d33 meter (ZJ-3B, CAS). The ferroelectric polarization hysteresis loops
were recorded by a ferroelectric hysteresis measurement unit (NPL, UK).
The thermal depoling experiments were conducted by annealing poled
samples for 2 h at various temperatures up to 1000oC and then
performing the d33 measurements at room temperature.
32
3 The control of grain size and morphology
3.1 Shrinkage and densification process
3.1.1 Sintering kinetics
The normalized shrinkage of the BaTiO3 (BT), SrTiO3 (ST), Bi4Ti3O12
(BIT) and CaBi2Nb2O9 (CBNO) samples are plotted versus temperature in
Fig.3-1. The densification process of all samples progresses very rapidly.
Once activated, the densification is completed within 2 min within a very
narrow temperature interval ranging from 650oC to 950oC. The nanostructured Bi4Ti3O12 (BIT) powder shows a very low densification onset
temperature (Ton) with a maximum shrinkage rate of 8.9x10-3 s-1 at 790oC.
Although Ton of this powder is as low as 625oC, the main part of the
densification does not occur until above 750oC. The nano-powders of
BaTiO3 (BT), SrTiO3 (ST), the BaTiO3/SrTiO3 powder mixture
(MBST64) and the Ba0.6Sr0.4TiO3 powder (BST64) exhibit higher Ton and
similar values of maximum shrinkage rate, e.g. 1.2x10-2 s-1 at 930oC,
5.6x10-3 s-1 at 890oC, 9.7x10-3 s-1 at 940oC and 6.3x10-3 s-1 at 910oC,
respectively. The densification curves of these powders are similar, and
the main part of the densification occurs at T > 800oC, although Ton
values are as low as 625oC (BT), 750oC (ST), 700oC (MBST64) and
700oC (BST64), respectively. The nano-sized powder of BIT shows a
lower densification onset temperature (~625oC) and higher maximum
shrinkage rate (8.9x10-3 s-1 at 790oC) than the micron-sized powder of
CBNO (~825oC and 4.5x10-3 s-1 at 950oC, respectively).
3.1.2 Densification
33
A pressure of 50 MPa and a holding time of only 2 min was used in the
isothermal densification experiments, and these experiments revealed that
densities higher than 95% could be achieved at T ≥ 900oC for the BIT
powder, and at T ≥ 925oC for the ST, BST64 and MBST64 powders. The
ST powder can be compacted to a similar level of density by prolonging
the heating time at 900oC and using a pressure of 100 MPa, but using the
same sintering parameters did not yield any improvement of the density
for the BST64 powder. BIT samples with densities equal to or exceeding
97% TD were obtained at T > 850oC using a pressure of 50 MPa and zero
holding time, while for the samples prepared at 800 ≤ T ≤ 850oC a
pressure of 75 MPa and a holding time of 3 min was used, and finally,
dense BIT samples could be prepared at 775oC using a pressure of 100
MPa and a holding time of only 3 minutes. CBNO compacts with
densities of 95% TD or more can be prepared at T ≥ 925oC using a
pressure of 100 MPa and a holding time of 3 min.
100
BT
ST
-∆L/∆Lmax (%)
80
60
40
20
0
500
600
700
800
900
o
Temperature ( C)
(a)
34
1000
1100
-∆L/∆Lmax %
100
BIT
CBNO
80
60
40
20
0
500
600
700
800
900
o
Temperature ( C)
1000
1100
(b)
Fig. 3-1 Normalized shrinkage of BaTiO3 (BT) and SrTiO3 (ST) (a) and Bi4Ti3O12 (BIT)
and CaBi2Nb2O9 (CBNO) (b) plotted versus temperature. A heating rate of 100oC/min and
a pressure of 50 MPa were used.
3.2 Grain growth and microstructure evolution
3.2.1 Ferroelectric ceramics based on (Ba, Sr)TiO3
The microstructures of ST compacts densified at different temperatures
are shown in Fig.3-2. The ST compact prepared at 900oC contains
homogeneous nano-sized grains of similar size as the starting powder,
and distinctly bimodal microstructures are formed in samples prepared at
925 and 950oC, i.e. the microstructures contain large grains as well as
fine ones, while large equi-axed grains are formed at 1000oC. This
suggests that the onset temperature for grain growth, Tg, is ~925oC. The
temperature dependence of the microstructures of the BT, BST64 and
MBST64 exhibit the same trend but with different onset temperatures for
grain growth, see below.
35
Fig. 3-2 SEM micrographs of ST samples sintered at (a) 900oC, (b) 925oC, (c) 950oC, (d)
1000oC. A holding time of 2 min and a pressure of 50 MPa were used at T ≥ 925oC,
whereas a pressure of 100 MPa was used at 900 oC.
3.2.2 Bismuth layer-structured ferroelectric ceramics
The microstructures of the BIT compounds are given in Fig.3-3. It is
obvious that the grain growth mechanism is different in comparison with
the (Ba, Sr)TiO3 materials discussed above. Nano-sized grains are thus
found in samples prepared below 850oC (Tg), while elongated plate-like
grains are formed at high temperatures, and no bimodal features can be
discerned.
The microstructure of the CBNO compact prepared at 925oC reveals the
presence of small platelet grains of the same size as in the starting
powder, while the same grain growth trend at higher temperatures as seen
for BIT is observed, see below.
36
Fig. 3-3 SEM micrographs of BIT samples sintered at (a) 800oC, (b) 850oC, (c) 950oC, (d)
1000oC, using zero holding time and a pressure of 50 MPa for samples prepared at
T ≥ 850oC, while for the samples prepared below 850oC a pressure of 75 MPa, and a
holding time of 3 min were used.
3.3
Phase evolution
The XRD profiles around the {200} peak of samples of BST64 and
MBST64 compositions sintered at different temperatures and of the
starting powders are shown in Fig.3-4. The BST64 sample patterns can
all be refined with cubic symmetry. The MBST64 samples prepared at
925oC or below are composed of two phases, however, whereas patterns
of samples prepared above this temperature can basically be refined with
cubic symmetry.
The X-ray studies verify that all of the BT, ST and BIT samples are
monophasic, including the BIT ceramics sintered at 1000oC. All ST
sample patterns can be refined with cubic symmetry, but BT samples
37
O
1000 C
O
O
O
925 C
O
900 C
original powders
30
40
2θ /( )
*
C
950
O
C
925
O
C
900
O
C
original powders
50
o
Silicon
C
O
1000
950 C
C ∗
O
1050
*
(111)
MBST64
1100
(111)
(110)
(200)
(110)
∗ Silicon
BST64
(200)
prepared above Tg are tetragonal, while the X-ray powder patterns of the
nano-sized ones, i.e. the ones prepared below Tg, contain peaks that are
heavily broadened, but they all seem to have cubic symmetry. The origin
of the broad peaks in the X-ray pattern of the MBST64 sample prepared
at 900oC might in part be ascribed to the presence of various cubic phases
with slightly different chemical compositions.
30
40
50
o
2θ/( )
(a)
(b)
Fig. 3-4 X-ray diffraction patterns of the starting powders of BST64 and MBST64 and
of samples sintered at different temperatures: (a) BST64, (b) MBST64.
3.4 Kinetics
3.4.1 Grain growth and kinetic windows
The grain sizes and density data obtained at the various densification
temperatures, Tiso, are plotted versus Tiso in Fig.3-5, and a “kinetic
window” is revealed within which full densification is achieved paired
with a very limited grain growth. It is evident that the investigated
powders can be differentiated into four groups: (i) The BST64 powder
38
does not have any apparent kinetic window; (ii) The BT and ST powders
have a narrow kinetic window of ~25oC; (iii) The BIT powder has a
kinetic window of ~75oC; (iv) The MBST64 powder mixture has a broad
kinetic window of ~125oC.
Below a certain critical temperature, Tg, the grain growth progresses
very slowly, but above Tg the grain growth takes place dramatically fast.
Thus, homogeneous nano-grained microstructures (see Fig.3-2.a and
Fig.3-3.a) are observed in samples prepared below Tg for both (Ba,
Sr)TiO3 ceramics and BIT ferroelectrics, while distinct bimodal
microstructures (see Fig.3-2.d) are formed in (Ba, Sr)TiO3 samples
prepared above Tg, i.e. the microstructures of the latter contain large
grains that are embedded in a matrix composed of fine ones. The BT, ST,
and BST64 powders have a Tg value of 925oC, while corresponding value
for the MBST64 powder is 1050oC. It can be noted that the grains larger
than 10 µm are observed in samples prepared above Tg, and they grew
from a size of ~200 nm to ~10 µm within 1-2 min.
When the BIT samples were sintered at temperatures above Tg
(> 850oC), the nano-sized particles with an average length of 0.15 µm
grew to elongated plate-like grains with an average length of 2.58 µm
along the a-b direction, while the grain growth along c direction was
much more restricted, i.e. the average thickness increased from 0.15 µm
to 0.52 µm (see Fig. 3-3.d), yielding a microstructure containing almost
randomly orientated elongated platelet grains that formed via a dynamic
ripening mechanism (76). The microstructure of the CBNO compact
prepared at 925oC reveals the presence of equiaxed grains of similar size
(~1 micron) as in the starting material, see below (Fig. 4-4.c). The
microstructures of samples prepared above 925oC showed these compacts
to consist mostly of randomly orientated plate-shaped grains, but the
length-to-width ratios of the latter are substantially less than those
observed in the BIT compacts. Some large equiaxed grains could also be
found.
39
4.0
3.5
Length
96
3.0
Thickness
Grain Size (µm)
Relative Density (%)
100
2.5
92
Tg
88
Kinetic Window
2.0
1.5
1.0
84
80
700
0.5
750
800
850
900
950
1000
0.0
1050
o
Temperature ( C)
(a)
24
20
BT
ST
Density
80
Grain Size (µm)
Relative Density (%)
100
16
Tg
60
12
40
20
0
8
4
Kinetic Window
850
900
950
1000
o
Temperature ( C)
(b)
24
Density
20
80
16
60
12
40
Tg
Tg
8
20
MBST64
BST64
Kinetic Window
0
900
950
1000
1050
1100
Grain Size (µm)
Relative Density (%)
100
4
0
o
Temperature ( C)
(c)
Fig. 3-5 Relative densities and grain sizes plotted versus sintering temperature (Tiso).
The kinetic windows within which fully dense nano-sized ceramics were obtained are
marked. (a) BIT, (b) BT and ST; (c) BST64 and MBST64
40
The fact that we are able to define a kinetic window within which it is
possible to densify nano-sized powders paired with a very limited grain
growth opens up new possibilities for optimizing microstructures.
3.4.2 Self-pinning effect
It can thus be noted that the MBST64 powder exhibits a kinetic window
as broad as 125oC, while that of the BST64 powder is almost nonexistent. The densification of the former powder is accompanied with a
solid-state reaction, i.e. BaTiO3 and SrTiO3 react to form Ba0.6Sr0.4TiO3
during the compaction. Apparently, the ongoing solid-state reaction
retards the grain growth. It is well known that the grain growth process
can be retarded by addition of second phase in connection with the
sintering of ceramics. The added particles have a pinning effect on the
grain growth process. In our case, and in analogy with the pinning effect,
the ongoing reaction gives rise to a self-pinning effect on the grain
growth process. This mechanism could potentially be applied for
preparation of other types of ceramics with nano-grained microstructures.
3.4.3 Thermal activation
As shown in Figs.3-1 to 3-4, the densification and grain growth
mechanisms are clearly thermally activated, and the maximum in the
shrinkage rate curve occurs at a relative density of ~ 0.8, which is in good
agreement with previous findings from pressureless sintering (PLS) and
hot-pressing (HP) experiments (91-94). It is commonly observed that rapid
grain growth takes place during the final stage of sintering micron-sized
powders. However, in the case of nano-sized powders it seems that the
grain growth takes place also in the beginning of the sintering process (93,
95)
. In the nano-sized porous green bodies the grain growth seems initially
to be driven by surface diffusion, but when the relative density has
reached a value around 0.8, the particle/grain growth of nano particles is
mainly driven by the size difference between particles/grains present,
which apparently also is a thermally activated process, and this process is
41
activated at a lower temperature in the case of nano-sized particles than
the grain growth process of micron-sized particles (96).
3.5 Ferroelectric properties
3.5.1 Ferroelectric properties of (Ba, Sr)TiO3 ceramics
The temperature dependence of the dielectric constant and loss tangent of
MBST64 and BST64 compacts recorded at a frequency of 10 kHz are
shown in Fig.3-6. The samples prepared above Tg, having bimodal
microstructures similar to those depicted in Fig.3-2, exhibit a maximum
in the permittivity around the Curie temperature (∼0oC). The Curie
temperature observed here is in agreement with previous findings for
BST ceramics of the same composition but prepared by conventional
pressureless sintering (97). The sample prepared below Tg possessing a
nano-sized microstructure exhibited low permittivity values, and the
permittivity is almost independent of temperature, indicating that the
transition in permittivity is very diffuse, and the dielectric losses in the
ferroelectric region decrease. The abnormally high temperature
dependence of the dielectric loss (dissipation) of the BST64 sample
consolidated at 900oC is ascribed to its low density, while the temperature
dependence of the dielectric loss of the other samples is less evident, see
Fig. 3-6.b.
42
7000
O
MBST64(950 C)
O
MBST64(1100 C)
O
BST64(950 C)
O
BST64(900 C)
10 KHz
6000
5000
εr
4000
3000
2000
1000
-80
-60
-40
-20
0
20
40
60
80
Temperature (oC)
(a)
0.12
10 KHz
0.10
O
MBST64(950 C)
O
MBST64(1100 C)
O
BST64(950 C)
O
BST64(900 C)
D
0.08
0.06
0.04
0.02
0.00
-80
-60
-40
-20
0
20
40
60
80
Temperature (oC)
(b)
Fig. 3-6 Temperature dependence of the dielectric constants (εr) (a) and loss tangents (D)
(b) at 10 KHz of the BST64 and MBST64 samples having nano-sized microstructures
(BST64 (900oC), MBST64 (950oC)) and micron-sized microstructures (BST64 (950oC),
MBST64 (1100oC)).
It has been well established that grain size has a profound influence on
the dielectric properties of ferroelectric ceramics through the interaction
of domain structures and grain boundaries; i.e. nano-sized ferroelectric
compacts exhibit lower permittivity values than micron-sized compacts
43
of the same composition. This is in agreement with the findings that the
MBST64 and BST64 compacts prepared below Tg exhibit substantially
lower permittivity values than the ones prepared above Tg, especially
around the Curie temperature where the contribution of domain walls to
the permittivity is high (17, 34). As the temperature dependences of the
permittivity of the BST64 and MBST64 ceramics prepared below Tg are
very similar, this suggests that it is unlikely that the diffuse phase
transition occurring in both samples originates from local compositional
fluctuation in MBST64. The origin of the decrease of permittivity with
the decrease of grain size has been explored in connection with the recent
successes in preparing dielectric film and bulk ceramics with truly nanosized grain structures (37, 39, 98). The present work provides additional
experimental data to the on-going debate concerning the contribution of
distorted (“dead”) grain boundary layers to the total permittivity (37, 39).
The dielectric losses are less dependent on grain size. As can be seen
from Fig.3-6.b, the MBST64 ceramic consolidated at 950oC exhibits
dielectric loss tangent (dissipation) < 0.02 within the whole temperature
range (from –70 to 70oC), whereas a decline of dielectric loss slightly
below the Curie temperature is observed for the MBST64 compact
consolidated at 1100oC. This decline of dielectric loss originates mainly
from the domain wall motion in large grains (99).
3.5.2 Ferroelectric properties of nano-sized Bi4Ti3O12
ceramics
The temperature dependence of the dielectric constant and loss tangent of
the nano-structured BIT sample, depicted in Fig. 3-3.a, and measured at 1
MHz, is given in Fig.3-7. The sample exhibits a normal ferroelectric
temperature dependence of the dielectric constant and loss tangent with a
Curie temperature of ∼675oC. The Curie temperature is in good
agreement with previous findings (7, 100). The nano-sized sample exhibited
slightly lower dielectric constant values compared with micro-sized BIT
samples (7, 30, 101). It has been reported that it is difficult to polarize pure
44
BIT ceramics because of their high conductivity (100). We were able to
polarize this material at 190oC using a DC field of 4.2 KV/mm for 5 min,
and we obtained a piezoelectric constant (d33) of 6.0 pC/N. The room
temperature P-E hysteresis loop of the nano-structured BIT ceramic at 1
Hz is given in Fig.3-8. Saturation polarization was never reached, and a
comparatively low remnant polarization (Pr) value (2.5µC/cm2) was
obtained.
2.0
900
1.8
700
1.4
εr
600
εr
1.6
1 MHz
1.2
D
1.0
500
0.8
400
0.6
D
800
0.4
300
0.2
200
0.0
0
100
200
300
400
500
600
700
800
Temperature (oC)
P (µC/cm2)
Fig. 3-7 Temperature dependence of the dielectric constants (εr) and loss tangents (D) of a
nano-structured BIT ceramic. The data are recorded at a frequency of 1MHz.
20
15
10
5
0
-150
-100
-50
0
-5
50
100
150
E (kV/cm)
-10
-15
-20
Fig. 3-8. P-E hysteresis loop of a nano-structured BIT ceramic. The data are recorded at a
frequency of 1 Hz.
45
46
4 Achieving 2-D type microstructure
4.1 Superplastic deformation
microstructures
yielding
textured
4.1.1 Superplastic deformation kinetics
The pre-forms for the deformation experiments were prepared by SPS as
described above. Thus a fully dense cylindrical sample with a diameter of
12 mm and a height of ∼6 mm was placed into a die with 20 mm inner
diameter, and heated at a rate of 100 oC min-1 in vacuum to the preset
deformation temperature (1000oC ~ 1150oC). A constant uniaxial load
corresponding to an initial compressive stress of 40 MPa was applied
either at room temperature or when reaching the preset temperature.
Holding times of 5 min were applied.
The normalized compressive deformation strain curves of the BIT and
CBNO samples are plotted versus temperature in Fig.4-1. The
superplastic deformation process of the BIT sample started at 760oC, and
that of the CBNO sample was activated at 920oC. A strain of ∼57% for
the BIT sample was achieved within a short period of time (1 min), and
the maximum strain rate reached at ∼840oC was as high as 1.1x10-2 s-1. A
strain of ∼67% was obtained for the CBNO sample, with a maximum
strain rate of 1.3×10-2 s-1 at 1020oC. To the best of our knowledge, a
compressive strain rate as high as 10-2 s-1 has not previously been
observed in connection with deformation of oxide ceramics, although
great efforts have been made during the last two decades to improve the
ductility of ceramic materials.
Isothermal deformation compressive strain curves of BIT samples
recorded at different temperatures are plotted versus time in Fig.4-2. In
these cases the pressure was applied when the preset temperature had
47
been reached. It is evident that both the compressive strain and strain rate
increased with increasing deformation temperature.
80
70
-∆Ld/Lo (%)
60
BIT
CBNO
50
40
30
20
10
0
700
800
900
1000
o
Temperature ( C)
1100
1200
Fig.4-1 Normalized compressive deformation strain curves of the BIT and CBNO samples
plotted versus temperature. The curves were recorded using a heating rate of 100oC/min
and an initial stress of 40 MPa.
Fig.4-2 Normalized compressive deformation strain curves of BIT samples deformed at
different temperatures, plotted versus time. A load corresponding to an initial compressive
stress of 40 MPa was applied at the preset temperature.
48
Isothermal deformation strain experiments on a series of BIT samples
with different microstructures, i.e. exhibiting the nano-sized, sub-micron
sized and micron-sized structures depicted in Fig.3-3.a, b, and c, have
also been performed. The deformation experiments were carried out at
850oC and, as above, the pressure was applied when the preset
temperature had been reached. The recorded strain data are plotted versus
time in Fig.4-3. The deformation strain of the nano-sized sample is
obviously higher than that of the sub-micron sized one, and substantially
higher than that of the micron sized one.
50
a
b
c
−∆Ld/L0(%)
40
30
20
10
0
0
2
4
6
8
10
Time (min)
Fig.4-3 Isothermal deformation strain curves of a series BIT compacts with (a) nanosized, (b) submicron-sized, and (c) micro-sized microstructures, deformed at 850oC. The
curves were recorded using an initial stress of 40 MPa.
4.1.2 Textured microstructures
Scanning electron micrographs of polished and thermally etched surfaces
of BIT and CBNO samples, recorded before and after superplastic
deformation, are shown in Fig.4-4. The micrographs of the deformed
samples are recorded perpendicular to the shear flow direction. It clearly
appears that the well facetted, thin elongated grains are fairly randomly
orientated parallel to the shear flow direction, whereas they pile up in the
49
perpendicular direction, yielding a compact with a 2-D microstructure.
Comparing the microstructures of the CBNO and BIT samples, it is
apparent that the degree of grain alignment in the CBNO ceramic is less
pronounced, i.e. the possibility to obtain highly grain-orientated
microstructures containing grains with high aspect ratios is limited by the
fact that the compact to be deformed contains micron-sized grains.
Fig.4-4 SEM micrographs of polished and thermally etched surfaces of the BIT and
CBNO samples before and after deformation. The latter are recorded perpendicular to the
shear flow direction, (a) BIT before deformation, (b) BIT after deformation, (c) CBNO
before deformation, and (d) CBNO after deformation.
Scanning electron micrographs of BIT samples, recorded after
superplastic deformation at 900 and 1000oC are shown in Fig. 4-5. The
micrograph of the deformed sample is recorded parallel and
perpendicular to the shear flow direction. The densities of the ceramics
after deformation are 97% and 92% TD, respectively. It is clearly seen
that both of them have textured 2-D microstructures, and that the sample
deformed at 900oC exhibits smaller grains (∼1 µm) than the one prepared
50
at 1000oC. It is clear from this figure that the aspect ratio of the resulting
grains varies with the temperature.
(1)
(2)
Fig.4-5. SEM micrographs of polished and thermally etched surfaces of the textured BIT
samples recorded parallel (a) and perpendicular (b) to the shear flow direction. The
samples were deformed at (1) 900oC using a holding time of 5 min; (2) 1000oC using a
holding time of 10 minutes. In these experiments an initial stress of 40 MPa was used.
4.2 X-ray studies of textured samples
Fig. 4-6 shows X-ray diffraction (XRD) patterns of the surfaces of BIT
and CBNO samples parallel to the shear flow direction. The X-ray
powder patterns reveal the presence of monophasic materials. The XRD
patterns exhibit strong (00l) diffraction peaks verifying the highly
textured feature of the samples, in agreement with the microstructural
51
studies described above. A Lotgering orientation factor of 99% was found
for the textured BIT sample, and that of the CBNO sample was 70%.
10
30
50
(1119)
(00 20)
(0018)
(0016)
40
(1115)
(0012)
(0010)
20
(1 1 13)
(315) / (135)
(0 2 10)
(0 0 14)
(220)
(028)
(119)
(0 0 10)
BTO
(0014)
(117)
(115)
(200) / (020)
(113)
(008)
(004)
(006)
(006)
(004)
(008)
CBNO
60
2θ/(o)
Fig.4-6 X-ray diffraction (XRD) patterns of the surfaces of the deformed BIT and CBNO
samples, recorded parallel to the shear flow direction. Notice the strong (00l) diffraction
peaks indicating preferred grain-alignment.
4.3 Superplastic deformation induced directional
dynamic ripening
The main features of the isothermal deformation experiments using predensified compacts can be summarized as follows:
(i)
The deformation of equiaxed nano-sized pre-densified
compacts of BIT proceeds very fast, i.e. compressive strain
rates of the order of 10-2 s-1 can easily be achieved, and the
deformation is associated with a transportation of BITspecies as the nano-sized compact is transformed into an
almost 100% grain aligned compact containing elongated
platelet grains with aspect ratios exceeding 10;
(ii)
The aspect ratio of the BIT grains can be tailored by selecting
the temperature for the isothermal deformation experiment;
52
(iii)
(iv)
(v)
The deformation of pre-densified compacts of BIT
containing sub-micron sized equiaxed grains proceeds less
fast than the deformation of the nano-sized compacts. The
deformation of pre-densified compacts of BIT containing
randomly orientated elongated platelet grains is very much
retarded;
The deformation of the micron-sized compacts of CBNO also
proceeds rapidly, but the aspect ratios of the formed grains
and the resulting degree of grain alignment are much smaller
than for the BIT samples;
Apparently, a favourable interaction occurs between
superplastic deformation and grain growth, i.e. the
superplastic deformation process induces a directional
dynamic ripening mechanism that yields anisotropic grains.
Grain growth induced strain hardening, which normally acts
as an obstacle to further deformation, is obviously minimized
by the fact that the growth of grains occurs preferentially
parallel to the shear flow direction.
Thus, the mechanism behind such a rapid grain growth and grain
alignment process has been termed by us: superplastic deformation
induced directional dynamic ripening.
4.4 Ferroelectric properties
4.4.1 Ferroelectric properties of textured Bi4Ti3O12
ceramics
The temperature dependence curves of the dielectric constant and loss
tangent of the BIT sample with the anisotropic microstructure shown in
Fig.4-4.b are given in Fig.4-7, measured at 1 MHz. First of all, the
dielectric constant values parallel (a/b direction, curve [//]) and
53
perpendicular (c direction, curve [⊥]) to the shear flow direction are
comparable to the corresponding single-crystal data, and are much higher
than those previously reported for any grain-orientated stoichiometric
BIT ceramics (7). Although it has been reported that it is difficult to
polarize a pure BIT ceramic because of its high conductivity (100), we
polarized our material in the a/b direction at 175oC using a DC field of
1.9 KV/mm for 5 min, and we obtained a piezoelectric constant (d33) of
11.4 pC/N. The corresponding d33 value in the c direction, using a DC
field of 5.7 KV/mm for 5 min, is 0.4 pC/N, which again verifies the
anisotropic microstructure of this sample. The thermal depoling
experiments revealed that the d33 value of 11.4 pC/N in a/b direction was
retained up to 676oC, the Curie temperature of the material, and then
disappeared.
2000
[//]
1600
εr
1200
800
400
[⊥]
0
100
200
300
400
500
o
Temperature ( C)
(a)
54
600
700
4.0
3.5
[//]
3.0
D
2.5
2.0
1.5
[⊥]
1.0
0.5
0.0
0
100
200
300
400
500
600
700
o
Temperature ( C)
(b)
Fig. 4-7 Dielectric properties of the textured BIT ceramic: (a) the temperature dependence
of the dielectric constants (εr) and (b) the temperature dependence of the loss tangents
2
P (µC/cm )
(D). Measurements were performed parallel [//] and perpendicular [⊥] to the shear flow
direction, using a frequency of 1 MHz.
[//]
[⊥]
40
30
20
10
0
-125 -100 -75 -50 -25 0
-10
25
50
75 100 125
E (kV/cm)
-20
-30
-40
Fig. 4-8. P-E hysteresis loops of the textured BIT ceramic measured parallel [//] and
perpendicular [⊥] to the shear flow direction. The measurements were performed at room
temperature, using a frequency of 1 Hz.
55
The polarization hysteresis loops of the grain-orientated sample,
determined at room temperature at a frequency of 1 Hz, are given in Fig.
4-8. The loops are strongly anisotropic, with much higher remnant
polarization, Pr, in the a/b direction than in the c direction. The
spontaneous polarization Ps in the a/b direction is 27µC/cm2, which is
consistent with the corresponding calculated data for a single crystal in
the direction (28).
We ascribe the improved anisotropic dielectric and piezoelectric
properties described above to the perfect alignment and refinement of the
anisotropic grains. The former promotes the rotation of the spontaneous
polarization as it does in similar type of materials developed by any
already established grain-alignment techniques, while the latter increases
the domain wall density by forming small elongated grains (29), verifying
the pronounced benefit gained by applying our new processing strategy.
The width of the 90o domains present along the a-b direction in this
sample with grain size of ~0.4x1.3x4 micron3 is thus below 1 µm,
compared to 8.5 µm in a millimetre-sized single crystal (101).
4.4.2 Ferroelectric properties of textured CaBi2Nb2O9
ceramics
Fig.4-9 shows the P-E hysteresis loops for the textured CaBi2Nb2O9
(CBNO) sample having the microstructure depicted in Fig.4-4.d The
remnant polarization Pr, parallel to the shear flow direction, Pr [//], is
much higher than that perpendicular to the shear flow direction, Pr [⊥].
Due to the thickness of the sample (0.10~0.15 mm) and experimental
restrictions (maximum voltage equal to 4 kV), it was not possible to
obtain saturated P-E hysteresis loops, but the CBNO sample seems to
have a high value of the coercive field (Ec). Similar problems have
previously been noted for high Tc Bi3NbTiO9 ceramics (102). The average
piezoelectric constant d33 values of the CBNO samples were 19.5±0.3
56
[//]
[⊥]
-300
-200
P (µC/cm2)
([//]) and 0.2±0.1 ([⊥]), and these values are nearly three times as high as
those of conventionally sintered materials (40).
-100
15
10
5
0
-5
0
100
200
300
E (kV/cm)
-10
-15
Fig. 4-9 P-E hysteresis loops of textured CBNO ceramics, recorded at room temperature
using a frequency of 10 Hz. [//]: Parallel to the shear flow direction; [⊥]: Perpendicular to
the shear flow direction.
Fig. 4-10 shows the temperature dependence of the dielectric constant
and loss of CBNO at 1 MHz. The Curie point of the samples both parallel
([//]) and perpendicular ([⊥]) to the shear flow direction is 943±2oC,
which is very close to that of conventionally sintered CBNO ceramics
(940±2oC) (40). The dielectric constant and loss of the samples are higher
parallel to (εr [//]) than perpendicular to (εr [⊥]) the shear flow direction,
in agreement with our findings for textured BIT ceramics, see above.
The thermal depoling behaviour of textured CBNO samples parallel to
[//] the shear flow direction is shown in Fig.4-11, where the piezoelectric
constant, d33, of the annealed samples is plotted against the annealing
temperature. All samples annealed up to 800oC exhibited the same d33
value, which rapidly drops above 900°C.
57
700
40
600
at 1 MHz
εr [//]
500
εr [⊥]
30
D [//]
D [⊥]
20
D
εr
400
300
10
200
100
0
0
200
400
600
800
Temperature (oC)
1000
Fig. 4-10 Temperature dependence of dielectric constant and loss for textured CBNO
ceramics, using a frquency of 1 MHz. [//]: Parallel to the shear flow direction; [⊥]:
Perpendicular to the shear flow direction.
20
d33 (pC/N)
15
10
5
0
0
200
400
600
800
1000
Annealing Temperature (oC)
Fig. 4-11 Piezoelectric constant, d33, of annealed samples plotted versus the annealing
temperature for textured CBNO ceramics. The data are recorded parallel to the shear flow
direction.
58
5 Summary
The sintering behaviour of nano-powders, such as Bi4Ti3O12 (BIT) with
an average particle size of 100 nm, BaTiO3 (BT), SrTiO3 (ST),
Ba0.6Sr0.4TiO3 (BST64), and a mixture of the composition
(BaTiO3)0.6(SrTiO3)0.4 (MBST64), all with particle sizes in the range of
60 to 80 nm, have been studied by spark plasma sintering (SPS). It was
verified that both densification and grain growth are thermally sensitive
in all nano-powders. The temperature region where the densification
mechanism is activated is lower than that where grain growth occurs,
which allows us to produce ceramics with tailored microstructures, e.g.
ceramics of BIT, ST, BST and MBST with microstructures ranging from
nano-sized to micron-sized, via those that exhibit distinct bimodal
features.
Using the sintering procedure outlined above, we have been able to
establish a “kinetic window” within which it is possible to prepare dense
samples having nano-sized microstructures. The sintering behaviour of all
the nano-sized powders is fairly similar, whereas the widths of the kinetic
window are different. Thus, the mixture of the composition MBST64
exhibits a window as broad as 125oC, while the window of the nanopowder of the same composition Ba0.6Sr0.4TiO3 is reduced to one single
temperature. The densification of the former powder is accompanied by a
solid-state reaction, and this reaction is suggested to have a self-pinning
effect on the grain growth.
The evolution, with increasing sintering temperature, of the
microstructure in bismuth layer-structured ferroelectrics such as
Bi4Ti3O12 (BIT) and CaBi2Nb2O9 (CBNO) has been investigated, and it is
shown that the nano-sized particles in the precursor powder transform
into elongated plate-like (along the a/b direction) grains, whereas no
evident grain growth occurs along the c direction when the sintering
temperature exceeds Tg. Thus, a microstructure containing elongated
platelet grains is formed via a dynamic ripening mechanism.
59
It is demonstrated that the superplastic deformation induced directional
dynamic ripening is a very effective grain alignment process. This new
process makes it possible to produce a textured microstructure within
minutes by using pre-densified nano-sized BIT ceramics or micro-sized
CBNO ceramics. We have recorded compressive strain rates of the order
10-2 s-1, which is approximately ten times faster than the fastest rates
achieved by any other established technique. The combination of
effective grain alignment and grain refinement opens up new possibilities
for developing anisotropic ceramics with unique and/or improved
performance, and the processing concept described here should, in
principle, have broad applicability to the production of a wide range of
ceramics with a 2-D type of tailored microstructure consisting of
anisotropic grains.
It has been well established that the ferroelectricity of ferroelectric
ceramics is profoundly influenced by their microstructural features, such
as: grain size, grain morphology and texture. It is verified that nanostructured BST ceramics exhibit a diffuse transition in permittivity and
reduced dielectric losses, whereas the dielectric constant of BST ceramics
containing micron sized grains or a mixture of nano-sized and micronsized grains exhibits normal permittivity values in the ferroelectric
region. It was found that the textured bismuth-based ceramic
ferroelectrics, obtained by the superplastic deformation induced
directional dynamic ripening process, exhibit highly anisotropic
ferroelectric properties, i.e. substantially better ferroelectric properties
than those of “conventional” grain-orientated BIT ceramics, and equal to
or better than corresponding single-crystal properties. Textured CBNO
exhibits a combination of high Tc, stable d33 and high thermal depoling
temperature, indicating that is a very promising candidate for hightemperature piezoelectric applications.
60
6 Future work
Some suggestions for future interesting investigations are listed below:
I will still focus on developing tailored microstructure in ferroelectric
ceramics, introducing and applying our new non-equilibrium assembling
concept and using this concept to tailor microstructures. Accordingly I
will mix two sub-micron sized Ba1-xSrxTiO3 powders with different x
values and SPS them to full density. By selecting appropriate sintering
parameters the heat treatment will be interrupted before monophasic
samples, i.e. samples with one single x value, are formed; and the
ferroelectric properties of the resulting compacts will be investigated.
Furthermore, I will use the same concept to prepare other ferroelectric
compounds or compositions in order to extend the understanding and
application of this new sintering process.
Mechanical pressure plays an important role also in fabrication of nanostructured ferroelectric ceramics by SPS. I will investigate the influence
of pressure on the microstructure of ceramics, in order to obtain more
ferroelectric ceramics with controlled grain size and morphology by
applying suitable pressures.
61
Acknowledgments
I would like to take this opportunity to express my gratitude and
appreciation to the people who have contributed in different ways
towards the completion of this thesis:
I want to express my deepest gratitude to Professor Zhijian Shen, my
supervisor, for your inspiring guidance and help, and also for supporting
and encouraging me throughout these years. It is a great pleasure to work
and discuss with you and to share in your knowledge and enthusiasm.
Professor Mats Nygren, my co-supervisor, for inviting me to Sweden
and accepting me as a PhD student; for all the support and help from you.
It is a great honour to work with you.
Associate professor Jekabs Grins, my co-supervisor, for your constant
warm-hearted help and fruitful discussions of my work.
Dr. Haixue Yan and Dr. Michael Reece, my collaboration partners, for
all of your support, for discussion and valuable suggestions. I am very
happy to cooperate with both of you.
Professor Peiling Wang, Professor Tinglian Wen, and Professor Lian
Gao, for your collaboration on producing high quality nano-powders, for
good suggestions relating to my work and, of course, for your
consideration to me.
Dr. Bo Su, Dr. Dou Zhang, Professor Tim W. Button, Dr. Xingyuan
Guo and Dr. Ping Xiao, my collaboration partners, for all kinds of
support of my work.
Professor Sven Lidin and Professor Lennart Bergström, for creating a
stimulating and pleasant scientific environment, and of course, for your
interest in my work and thesis and for always being helpful.
62
Dr. Kjell Jansson, for teaching me to use the scanning electronic
microscope (SEM) and for always being available for endless questions
about that technique. Dr. Ulrich Sutter and Dr. Geoff West, for the
piezoresponse force microscopy and transmission electron microscope
studies.
All of my colleagues at the department, for providing such a nice
working atmosphere. It is my pleasure to work with all of you.
To my family and relatives in China. Especially, I heartily thank my
Aunt Fenglan Liu, Uncle Guoyin Chen, Aunt Aiyun Song and Uncle
Guoqiang Yu for all kinds of help, encouragement and consideration to
me and to my parents in China.
To my parents, Xinrong Xu and Dequan Liu: thank you for giving me
life, and always encouraging and supporting me. I also appreciate your
helping me take care of my little daughter during my thesis writing.
To my husband, Wenlong Yao: thank you very much for your love,
your continuous support and encouragement, and also for your many
interesting and valuable ideas relating to my work. Then to my little
daughter, Lucy Lingyan Yao: you don’t know how blessed we feel every
time we look at your cheerful and lovely face. You also don’t know how
deeply we love you. I would say you are the best present I got from our
God. You, father and me, we are a team, who always support each other,
no matter what may come.
63
References
1. Haertling G. H. Ferroelectric ceramics: history and technology.
Journal of the American Ceramic Society, 1999, 82, 797-818.
2. Demartin-Maeder M., Damjanovic D. Piezoelectric Materials and
Devices. In Piezoelectric Materials and Devices, ed. N. Setter. (EPFL
Swiss Federal Institute of Technology), Lausanne, 2002, pp. 389.
3. Maruno S., Kuroiwa T. & Mikami N., et al. Model of leakage
characteristics of (Ba, Sr)TiO3 thin films. Appl. Phys. Lett., 1998, 73,
954-6.
4. Park B. H., Kang B. S., Bu S. D., Noh T. W., Lee J. & Jo W.
Lanthanum-substituted bismuth titanate for use in non-volatile memories.
Nature (London), 1999, 401, 682-4.
5. West A. R. Electrical Properties. In Basic Solid State Chemistry, ed.
A. R. West. John Wiley & Sons, Ltd, Baffins Lane, Chichester, England,
2001, pp. 293-362.
6. West A. R., Adams T. B., Morrison F. D. & Sinclair D. C. Novel high
capacitance materials: BaTiO3:La and CaCu3Ti4O12. Journal of the
European Ceramic Society, 2004, 24, 1439-48.
7. Kim S. K., Miyayama M. & Yanagida H. Electrical anisotropy and a
plausible explanation for dielectric anomaly of Bi4Ti3O12 single crystal.
Mater. Res. Bull., 1996, 31, 121-31.
8. Setter N. ABC of Piezoelectricity and Piezoelectric Materials. In
Piezoelectric Materials in Devices, ed. N. Setter. EPFL Swiss Federal
Institute of Technology, Lausanne, 2002, pp. 1-4.
9. Zivkovic L. M., Stojanovic B. D., Pavlovic V. B., Nikolic Z. S.,
Marinkovic B. A. & Sreckovic T. V. SEM investigation of domain
structure in (Ba, Ca, Pb)TiO3. Journal of the European Ceramic Society,
1999, 19, 1085-7.
64
10. Wersing W. Applications of piezoelectric materials: an introductory
review. In Piezoelectric materials in devices, ed. N. Setter. EPFL Swiss
Federal Institute of Technology, Lausanne, Switzerland, 2002.5, pp. 31.
11. Megaw H. D. Refinement of the structure of BaTiO3 and other
ferroelectrics. Acta Cryst., 1962, 15, 972-3.
12. Evans H. T., An x-ray diffraction study of tetragonal barium titanate.
Acta Cryst., 1961, 14, 1019-26.
13. HARRIS E., KAY H. F. Size distribution of tobacco smoke particles.
Nature (London), 1959, 183, 741-2.
14. Harada J., Axe J. D. & Shirane G. Neutron-scattering study of soft
modes in cubic barium titanate. Physical Review B: Solid State, 1971,
[3]4, 155-62.
15. Kwei G. H., Lawson A. C., Billinge S. J. L. & Cheong S. W.
Structures of the ferroelectric phases of barium titanate. J. Phys. Chem.,
1993, 97, 2368-77.
16. Wittels M. C., Sherrill F. A. Fast-neutron effects in tetragonal barium
titanate. J. Appl. Phys., 1957, 28, 606-9.
17. Frey M. H., Payne D. A. Grain-size effect on structure and phase
transformations for barium titanate. Physical Review B, 1996, 54, 315868.
18. Saito Y., Takao H. & Tani T., et al. Lead-free piezoceramics. Nature
(London), 2004, 432, 84-7.
19. Nowotny J., Rekas M. Dielectric ceramic materials based on alkaline
earth metal titanates. Key Eng Mat, 1992, 66-67, 45-143.
20. Wu Y., Forbess M. & Seraji S., et al. Doping effect in layer
structured SrBi2Nb2O9 ferroelectrics. J. Appl. Phys., 2001, 90 (10),
5296-302.
21. Newnham R. E., Wolfe R. W. & Dorrian J. F. Structural basis of
ferroelectricity in the bismuth titanate family. Mater. Res. Bull., 1971, 6,
1029-39.
65
22. Subbarao E. C. Ferroelectricity in Bi4Ti3O12 and its solid solutions.
Physical Review, 1961, 122, 804-7.
23. Cummins S. E., Cross L. E. Electrical and optical properties of
ferroelectric single crystals. J. Appl. Phys., 1968, 39, 2268-74.
24. Sakata K., Takenaka T. & Shoji K. Hot-forged ferroelectric ceramics
of some bismuth compounds with layer structure. Ferroelectrics, 1978,
22, 825-6.
25. Takenaka T., Nagata H. & Suzuki M. Bismuth layer-structured
ferroelectrics with high curie temperatures. Ceram. Int., 2004, 30, 2053.
26. Shimakawa Y., Kubo Y. & Tauchi Y., et al. Crystal and electronic
structures of Bi4-xLaxTi3O12 ferroelectric materials. Appl. Phys. Lett.,
2001, 79, 2791-3.
27. Demartin M., Damjanovic D. Lead Free Piezoelectric Materials. In
Piezoelectric Materials in Devices, ed. N. Setter. EPFL Swiss Federal
Institute of Technology, Lausanne, 2002, pp. 1-4.
28. Dorrian J. F., Newnham R. E., Smith D. K. & Kay M. I. Crystal
structure of bismuth titanate. Ferroelectrics, 1971, 3, 17-27.
29. Soga M., Noguchi Y., Miyayama M., Okino H. & Yamamoto T.
Domain structure and polarization properties of lanthanum-substituted
bismuth titanate single crystals. Appl. Phys. Lett., 2004, 84, 100-2.
30. Fuierer P. A., Nichtawitz A. Electric field assisted hot forging of
bismuth titanate. ISAF '94, Proceedings of the IEEE International
Symposium on Applications of Ferroelectrics, 9th, University Park, Pa.,
Aug.7-10, 1994, 1994, 126-9.
31. de Keijser M., Dormans G. J. M. Chemical vapor deposition of
electroceramic thin films. MRS Bull, 1996, 21, 37-43.
32. Watanabe T., Funakubo H., Mizuhira M. & Osada M. Site definition
and characterization of La-substituted Bi4Ti3O12 thin films prepared by
metalorganic chemical vapor deposition. J. Appl. Phys., 2001, 90, 65335.
66
33. de Araujo C. A., Cuchiaro J. D., McMillan L. D., Scott M. C. &
Scott J. F. Fatigue-free ferroelectric capacitors with platinum electrodes.
Nature (London), 1995, 374, 627-9.
34. Arlt G., Hennings D. & De With G. Dielectric properties of finegrained barium titanate ceramics. J. Appl. Phys., 1985, 58, 1619-25.
35. Bell A. J., Moulson A. J. & Cross L. E. The effect of grain size on
the permittivity of barium titanate(IV). Ferroelectrics, 1984, 54, 487-90.
36. Liu J., Shen Z., Nygren M., Su B. & Button T. W. Spark plasma
sintering behavior of nano-sized (Ba, Sr)TiO3 powders: determination of
sintering parameters yielding nanostructured ceramics. Journal of the
American Ceramic Society, 2006, 89, 2689-94.
37. Zhao Z., Buscaglia V. & Viviani M., et al. Grain-size effects on the
ferroelectric behavior of dense nanocrystalline BaTiO3 ceramics. Physical
Review B: Condensed Matter and Materials Physics, 2004, 70,
024107/1,024107/8.
38. Tsunekawa S., Fukuda T., Ozaki T. & Yoneda Y. Study on
ferroelectric domains in BaTiO3 crystalline films and bulk crystals by
atomic force and scanning electron microscopies. J. Appl. Phys., 1998,
84, 999-1002.
39. Sinnamon L. J., Saad M. M., Bowman R. M. & Gregg J. M.
Exploring grain size as a cause for "dead-layer" effects in thin film
capacitors. Appl. Phys. Lett., 2002, 81, 703-5.
40. Yan H., Zhang H. & Ubic R., et al. A lead-free high-curie-point
ferroelectric ceramic, CaBi2Nb2O9. Advanced Materials (Weinheim,
Germany), 2005, 17, 1261-5.
41. Brandmary R. J., Brown A. E. & Dunlap A. M. Annealing effects on
microstructure and dielectric properties of hot-pressed, ultrafine-grained.
U.S.A ECOM Technical Report, 1965, 2614.
42. Xue L. A., Chen Y., Gilbart E. & Brook R. J. The kinetics of hot
pressing for undoped and donor-doped barium titanate (BaTiO3)
ceramics. J. Mater. Sci., 1990, 25, 1423-8.
67
43. Gao L., Hong J. S., Miyamoto H. & Torre S. D. D. L. Bending
strength and microstructure of Al2O3 ceramics densified by spark plasma
sintering. Journal of the European Ceramic Society, 2000, 20, 2149-52.
44. Hirata Y., Nitta A., Sameshima S. & Kamino Y. Dielectric properties
of barium titanate prepared by hot isostatic pressing. Mater Lett, 1996,
29, 229-34.
45. Liu J., Shen Z. & Nygren M. Consolidation and dielectric behaviours
of Ba0.6Sr0.4TiO3 ceramics with tailored microstructures. Ferroelectrics,
2005, 319, 335-42.
46. Shen Z., Johnsson M., Zhao Z. & Nygren M. Spark plasma sintering
of alumina. Journal of the American Ceramic Society, 2002, 85, 1921-7.
47. Holmes M., Newnham R. E. & Cross L. E. Grain-oriented
ferroelectric ceramics. American Ceramic Society Bulletin, 1979, 58,
872.
48. Kimura T., Yoshimoto T., Iida N., Fujita Y. & Yamaguchi T.
Mechanism of grain orientation during hot-pressing of bismuth titanate.
Journal of the American Ceramic Society, 1989, 72, 85-9.
49. Takenaka T., Sakata K. Grain orientation and electrical properties of
hot-forged bismuth titanium oxide (Bi4Ti3O12) ceramics. Japanese
Journal of Applied Physics, 1980, 19, 31-9.
50. Horn J. A., Zhang S. C., Selvaraj U., Messing G. L. & TrolierMcKinstry S. Templated grain growth of textured bismuth titanate.
Journal of the American Ceramic Society, 1999, 82, 921-6.
51. Watanabe H., Kimura T. & Yamaguchi T. Particle orientation during
tape casting in the fabrication of grain-oriented bismuth titanate. Journal
of the American Ceramic Society, 1989, 72, 289-93.
52. Kan Y., Wang P., Li Y., Cheng Y. & Yan D. Fabrication of textured
bismuth titanate by templated grain growth using aqueous tape casting.
Journal of the European Ceramic Society, 2003, 23, 2163-9.
53. Hennings D., Rosenstein G. Temperature-stable dielectrics based on
chemically inhomogeneous barium titanate(IV). Journal of the American
Ceramic Society, 1984, 67, 249-54.
68
54. Su B., He J. Y. & Cheng B. L., et al. Dielectric properties of spark
plasma sintered (SPS) barium strontium titanate (BST) ceramics.
Integrated Ferroelectr., 2004, 61, 117-22.
55. Fujii T., Nakamura R. & Ito S. Fabrication of functional gradient
Ba(Zr,Ti)O3 ferroelectrics ceramics by hot isostatic pressing. Funtai
oyobi Funmatsu Yakin, 2000, 47, 1210-5.
56. Weber U., Greuel G., Boettger U., Weber S., Hennings D. & Waser
R. Dielectric properties of Ba(Zr,Ti)O3-based ferroelectrics for capacitor
applications. Journal of the American Ceramic Society, 2001, 84, 759-66.
57. Buessem W. R., Cross L. E. & Goswami A. K. Effect of twodimensional pressure on the permittivity of fine-and coarse-grained
barium titanate. Journal of the American Ceramic Society, 1966, 49, 369.
58. Buessem W. R., Cross L. E. & Goswami A. K. Phenomenological
theory of high permittivity in fine-grained barium titanate. Journal of the
American Ceramic Society, 1966, 49, 33-6.
59. Lee H. N., Christen H. M., Chisholm M. F., Rouleau C. M. &
Lowndes D. H. Strong polarization enhancement in asymmetric threecomponent ferroelectric superlattices. Nature (London), 2005, 433, 3959.
60. Choi K. J., Biegalski M. & Li Y. L., et al. Enhancement of
ferroelectricity in strained BaTiO3 thin films. Science, 2004, 306, 1005-9.
61. Ota T., Tani M., Hikichi Y., Unuma H., Takahashi M. & Suzuki H.
Dielectric properties of BaTiO3-based ceramics with gradient
compositions. Ceramic Transactions, 1999, 100, 51-60.
62. Kim J., Kang S. L. Formation of core-shell structure in the BaTiO3SrTiO3 system. Journal of the American Ceramic Society, 1999, 82,
1085-8.
63. Song X. Y., Chen D. R. & Yin Z. W. Grain core-grain shell structure
in Nb-doped BaTiO3 capacitor ceramics. Ceramic Transactions, 1994, 41,
129-34.
69
64. Chazono H., Kishi H. Sintering characteristics in the BaTiO3-Nb2O5Co3O4 ternary system: II, Stability of so-called "core-shell" structure.
Journal of the American Ceramic Society, 2000, 83, 101-6.
65. Tokita M. Trends in advanced spark plasma sintering systems and
technology. Functionally gradient materials and unique synthetic
processing methods from next generation of powder technology. Funtai
Kogaku Kaishi, 1993, 30, 790-804.
66. Sumitomo Heavy Industries website. http://www.shi.co.jp/sps/.
67. Vanmeensel K., Laptev A., Hennicke J., Vleugels J. & Van der Biest
O. Modelling of the temperature distribution during field assisted
sintering. Acta Materialia, 2005, 53, 4379-88.
68. Salamon D., Shen Z. & Sajgalık P. Rapid formation of a-sialon
during spark plasma sintering: Its origin and implications. Journal of the
European Ceramic Society, 2007, 27, 2541-7.
69. Matsugi K., Kuramoto H., Hatayama T. & Yanagisawa O.
Temperature distribution at steady state under constant current discharge
in spark sintering process of Ti and Al2O3 powders. J. Mater. Process.
Technol., 2003, 134(2), 225-232.
70. Anselmi-Tamburini U., Gennari S., Garay J. E. & Munir Z. A.
Fundamental investigations on the spark plasma sintering/synthesis
process. II: Modeling of current and temperature distributions. Materials
Science & Engineering, A: Structural Materials: Properties,
Microstructure and Processing, 2005, A394, 139-48.
71. Anselmi-Tamburini U., Garay J. E., Munir Z. A., Tacca A., Maglia
F. & Spinolo G. Spark plasma sintering and characterization of bulk
nanostructured fully stabilized zirconia: Part I. Densification studies. J.
Mater. Res., 2004, 19, 3255-62.
72. Wang Y., Fu Z. Study of temperature field in spark plasma sintering.
Materials Science & Engineering, B: Solid-State Materials for Advanced
Technology, 2002, B90, 34-7.
73. Zavaliangos A., Zhang J., Krammer M. & Groza J. R. Temperature
evolution during field activated sintering. Materials Science &
Engineering, A: Structural Materials: Properties, Microstructure and
Processing, 2004, A379, 218-28.
70
74. Chen W., Anselmi-Tamburini U., Garay J. E., Groza J. R. & Munir
Z. A. Fundamental investigations on the spark plasma sintering/synthesis
process. Materials Science & Engineering, A: Structural Materials:
Properties, Microstructure and Processing, 2005, A394, 132-8.
75. Shen Z., Peng H. & Nygren M. Formidable increase in the
superplasticity of ceramics in the presence of an electric field. Advanced
Materials (Weinheim, Germany), 2003, 15, 1006-9.
76. Shen Z., Zhao Z., Peng H. & Nygren M. Formation of tough
interlocking microstructures in silicon nitride ceramics by dynamic
ripening. Nature (London), 2002, 417, 266-9.
77. Tokita M. Present situation and future prospects of spark plasma
sintering (SPS) system. Shinsozai, 1996, 7, 19-28.
78. Chaim R., Shen Z. & Nygren M. Transparent nanocrystalline MgO
by rapid and low-temperature spark plasma sintering. J. Mater. Res.,
2004, 19, 2527-31.
79. Jones M. I., Hirao K., Hyuga H., Yamauchi Y., Shen Z. & Nygren
M. Wear properties of self-reinforced a-SiAlON ceramics produced by
spark plasma sintering. Wear, 2004, 257, 292-6.
80. Huang Z., Shen Z., Lin L., Nygren M. & Jiang D. Spark-plasmasintering consolidation of SiC-whisker-reinforced mullite composites.
Journal of the American Ceramic Society, 2004, 87, 42-6.
81. Okamoto M., Akimune Y., Furuya K., Hatano M., Yamanaka M. &
Uchiyama M. Phase transition and electrical conductivity of scandiastabilized zirconia prepared by spark plasma sintering process. Solid
State Ionics, 2005, 176, 675-80.
82. Yamaguchi N., Isobe R. & Ohashi O. A novel fabrication technique
of porous hydroxyapatite ceramics. Transactions of the Materials
Research Society of Japan, 2004, 29, 2659-62.
83. Takeuchi T., Capiglia C., Balakrishnan N., Takeda Y. & Kageyama
H. Preparation of fine-grained BaTiO3 ceramics by spark plasma
sintering. J. Mater. Res., 2002, 17, 575-81.
84. Takeuchi T., Betourne E. & Tabuchi M., et al. Dielectric properties
of spark-plasma-sintered BaTiO3. J. Mater. Sci., 1999, 34, 917-24.
71
85. Liu J., Shen Z., Nygren M., Kan Y. & Wang P. SPS processing of
bismuth-layer structured ferroelectric ceramics yielding highly textured
microstructures. Journal of the European Ceramic Society, 2006, 26,
3233-9.
86. Shen Z., Peng H., Liu J. & Nygren M. Conversion from nano- to
micron-sized structures: experimental observations. Journal of the
European Ceramic Society, 2004, 24, 3447-52.
87. Guo X., Xiao P., Liu J. & Shen Z. Fabrication of nanostructured
hydroxyapatite via hydrothermal synthesis and spark plasma sintering.
Journal of the American Ceramic Society, 2005, 88, 1026-9.
88. Chaim R., Shen Z. & Nygren M. Transparent nanocrystalline MgO
by rapid and low-temperature spark plasma sintering. J. Mater. Res.,
2004, 19, 2527-31.
89. Werner P. E., Eriksson L. & Westdahl M. TREOR, a semiexhaustive trial-and-error powder indexing program for all symmetries.
Journal of Applied Crystallography, 1985, 18, 367-70.
90. Dollase W. A. Correction of intensities for preferred orientation in
powder diffractometry: application of the March model. Journal of
Applied Crystallography, 1986, 19, 267-72.
91. Lange F. F. Powder processing science and technology for increased
reliability. Journal of the American Ceramic Society, 1989, 72, 3-15.
92. Shi J. L. Relations between coarsening and densification and mass
transport path in solid-state sintering of ceramics: model analysis. J.
Mater. Res., 1999, 14, 1378-88.
93. Shi J. L. Relation between coarsening and densification in solid-state
sintering of ceramics: experimental test on superfine zirconia powder
compacts. J. Mater. Res., 1999, 14, 1389-97.
94. Shi J. L. Thermodynamics and densification kinetics in solid-state
sintering of ceramics. J. Mater. Res., 1999, 14, 1398-408.
95. Mayo M. J. Processing of nanocrystalline ceramics from ultrafine
particles. International Materials Reviews, 1996, 41, 85-115.
72
96. Shen Z., Nygren M. Microstructural prototyping of ceramics by
kinetic engineering: applications of spark plasma sintering. Chem Rec,
2005, 5, 173-84.
97. Su B., Holmes J. E., Meggs C. & Button T. W. Dielectric and
microwave properties of barium strontium titanate (BST) thick films on
alumina substrates. Journal of the European Ceramic Society, 2003, 23,
2699-703.
98. Luan W., Gao L., Kawaoka H., Sekino T. & Niihara K. Fabrication
and characteristics of fine-grained BaTiO3 ceramics by spark plasma
sintering. Ceram. Int., 2004, 30, 405-10.
99. Cheng B. L., Su B., Holmes J. E., Button T. W., Gabbay M. &
Fantozzi G. Dielectric and mechanical losses in (Ba,Sr) TiO3 systems.
Journal of Electroceramics, 2002, 9, 17-23.
100. Shulman H. S., Testorf M., Damjanovic D. & Setter N.
Microstructure, electrical conductivity, and piezoelectric properties of
bismuth titanate. Journal of the American Ceramic Society, 1996, 79,
3124-8.
101. Shen Z., Liu J. & Grins J., et al. Effective grain alignment in
Bi4Ti3O12 ceramics by superplastic-deformation-induced directional
dynamic ripening. Advanced Materials (Weinheim, Germany), 2005, 17,
676-80.
102. Pardo L., Castro A., Millan P., Alemany C., Jimenez R. & Jimenez
B. (Bi3TiNbO9)(x)(SrBi2Nb2O9)(1-x) Aurivillius type structure piezoelectric
ceramics obtained from mechanochemically activated oxides. Acta
Materialia, 2000, 48, 2421-8.
73