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RESEARCH
RESEARCH: Original Research
Enhanced electric-field-induced strains in
(K,Na)NbO3 piezoelectrics from
heterogeneous structures
Mao-Hua Zhang 1,2,†, Qinghua Zhang 3,†, Ting-Ting Yu 4,†, Geng Li 1,
Hao-Cheng Thong 1, Li-Ying Peng 4, Lisha Liu 1,5, Jing Ma 1, Yang Shen 1,
Zhijian Shen 1, John Daniels 5, Lin Gu 3, Bing Han 4,⇑, Long-Qing Chen 6,
Jing-Feng Li 1, Fei Li 6,7,⇑, Ke Wang 1,⇑
1
State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua University, Beijing, PR China
Department of Materials and Earth Sciences, Nonmetallic Inorganic Materials, Technical University of Darmstadt, Darmstadt, Germany
3
Institute of Physics, Chinese Academy of Sciences, Beijing, PR China
4
Department of Orthodontics, Peking University School and Hospital of Stomatology, Beijing, PR China
5
School of Materials Science and Engineering, UNSW Australia, Sydney, Australia
6
Materials Research Institute, The Pennsylvania State University, University Park, PA, USA
7
Electronic Materials Research Laboratory, Key Laboratory of the Ministry of Education, Xi’an Jiaotong University, Xi’an, PR China
2
Piezoelectrics exhibit mechanical strain in response to electrical stimuli and vice versa. A high level of
electric-field-induced strain with minimal hysteresis is desired for piezoelectric materials when used as
actuators. The past two decades have seen extensive research into lead-free piezoelectrics to replace
Pb(Zr,Ti)O3 and compositional engineering has been demonstrated to be an effective method to tailor
their functional properties. Doped (K,Na)NbO3 (KNN) compositions with elaborate compositional
tuning can exhibit enhanced electromechanical properties. However, a balance between enhanced
properties and non-toxicity of the dopants should be considered. In this work, we propose to use
microstructural engineering to enhance the properties. Based on phase-field simulations, we propose to
take advantage of depolarization energies generated by polar-nonpolar interfaces, to increase the
contribution of domain wall motion to electric-field-induced strain. Heterogeneous ferroelectricparaelectric microstructures were introduced into a KNN ceramic via a two-step sintering process. Their
presence was characterized by high-resolution transmission electron microscopy. Enhanced reversible
domain wall motion was verified by in situ high-energy X-ray diffraction. Electric-field-induced strain
is enhanced by 62% and 200% at 25 °C and 150 °C, respectively. Considering lead-free piezoelectrics
also represent an emerging class of biomaterials for medical technology, the non-toxicity and
biocompatibility of the investigated compositions are examined by in vitro cell viability assays. Our
⇑ Corresponding authors.
†
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E-mail address: Wang, K. (wang-ke@tsinghua.edu.cn)
These authors contributed equally to this work.
1369-7021/Ó 2021 Elsevier Ltd. All rights reserved. https://doi.org/10.1016/j.mattod.2021.02.002
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results demonstrate that microstructural engineering is a promising alternative approach to enhance
the electric-field-induced strain of lead-free piezoelectrics while maintaining biocompatibility
Introduction
Shape-memory alloys [1,2], magnetostrictive materials [3] and
piezoelectrics [4,5] exhibit mechanical strain in response to
external electric stimuli. Among these, piezoelectrics are characterized by highly accurate and rapid strain response with large
blocking force, and thus are widely used in ultrasonic transducers, actuators, sensors, and mechanical energy harvesters [4,6].
Electric-field-induced strain is an important figure of merit for
piezoelectrics when used as an actuator. Typical application scenarios include jet dispensers [7], automotive diesel engine injectors [8], medical endoscopes for autofocus and auto zoom [9] and
high precision cell microinjection systems [10]. Motion of defectpinned domains [11], nonpolar-to-polar phase transitions [12–
16] and nanostructure control [17–19] have been proposed to
enhance the electric-field-induced strain for piezoelectrics. However, a critical issue that may be associated with these approaches
is the significantly increased hysteresis. A high degree of hysteresis not only affects positioning accuracy, but also generates heat
that can result in undesirable temperature rise when piezoelectrics are in service. Thus, alternative approaches to effectively
enhance the electric-field-induced strain without increasing the
hysteresis are greatly needed.
For the last 60 years, perovskite lead–zirconate–titanate (PZT)
ceramics have been the most technologically and commercially
important piezoelectric materials because they possess large
electric-field-induced strain, small hysteresis loss, and excellent
temperature stability [4]. Lead pollution, however, raises
increased environmental concerns during the mining of leadcontaining ores as well as the processing and disposal of PZT
products. In addition, Pb exposure can cause severe chronic poisoning [20,21]. There have been wide-spread scientific activities
in the quest to replace PZT with lead-free alternatives in the past
two decades and considerable attention has been focused on (K,
Na)NbO3 (KNN)-based piezoelectric materials [22–26].
However, undoped KNN materials show intrinsically low
piezoresponse (100 pC/N) [27], when compared to that of
PZT (550 pC/N) [28]. Since enhanced piezoelectric properties
are usually achieved at the boundary region between different
phases [29], the most effective approach to improve the piezoresponse for KNN materials is phase boundary engineering [30],
i.e., shifting the polymorphic phase boundary (PPB) to room
temperature. Hence, chemical compositions of KNN materials
have been elaborately tailored to construct the desired phase
boundary [31–32] and hence, enhanced piezoelectric properties
of 500–650 pC/N have been achieved [33–35]. Lead-free piezoceramics are attractive not only because of the enhanced properties. They also represent an emerging class of biomaterials for
medical technology [36–38], wherein PZT materials are excluded
due to the toxicity of lead. For example, lead-free piezoelectrics
can be used as active scaffolds for actuators and sensors for
implantable biomedical devices [39] as well as electroactive bioceramics for hard tissue implantation [40,41]. To this end, not
only high piezoresponse is required, but also biocompatibility
with living cells is important and needs to be evaluated.
Although KNN piezoceramics are featured with a combination
of antibacterial effects and biocompatibility [36,42–44], undoped
or lightly doped KNN compositions were mostly employed in the
existing investigations. Doped KNN piezoceramics typically
exhibit enhanced properties but there are, however, concerns
over the toxicity of some dopant elements [45,46]. Thus, biocompatibility evaluation of the doped compositions exhibiting
high piezoresponse remains to be evaluated and non-toxic
KNN materials exhibiting enhanced piezoelectric properties are
highly desired.
In this work, to enhance the electric-field-induced strain, a
ferroelectric-paraelectric heterogeneous system is conceived and
validated with the help of phase-field simulations. We propose
to utilize the depolarization electrostatic energies generated by
the ferroelectric–paraelectric interfaces to produce the desired
domain structures and to enhance the contribution of non-180°
domain wall motion to the electric-field-induced strain. Our
proposed idea is experimentally implemented in a doped KNN
material. Desired ferroelectric-paraelectric (polar-nonpolar)
microstructures were introduced via a two-step sintering process
and their presence was characterized by high-resolution transmission electron microscopy. Biocompatibility of the investigated
composition is evaluated by in vitro cell viability assays.
Materials and methods
Phase-field simulations
The temporal evolution of the polarization field is described by
the time-dependent Ginzburg-Landau (TDGL) equation,
@Pi ðr; tÞ
@F
¼ L
; ði ¼ 1; 2; 3Þ;
@t
@P i ðr; tÞ
ð1Þ
where L is the kinetic coefficient, F is the total free energy of the
system, r is the space position, and P i ðr; tÞ is the polarization.
oF/oPi(r,t) is the thermal dynamic driving force for the spatial
and temporal evolution of oPi(r,t).
The total free energy of the system is expressed by Eq. (2),
which includes the bulk free energy, elastic energy, electrostatic
energy, and gradient energy:
Z
F ¼ ½f bulk þ f elas þ f elec þ f grad dV
ð2Þ
V
where V is the system volume, fbulk is the Landau bulk free energy
density, felas is the elastic energy density, felec is the electrostatic
energy density and fgrad is the gradient energy density.
The bulk free energy density is expressed by Landau theory,
i.e.,
f bulk ¼ a1 ðP21 þ P22 þ P 23 Þ þ a11 ðP41 þ P 42 þ P43 Þ
þa12 ðP21 P 22 þ P22 P23 þ P23 P21 Þ
þa111 ðP 61 þ P62 þ P 63 Þ þ a112 ½P41 ðP22 þ P 23 Þ
ð3Þ
þ P 42 ðP21 þ P23 Þ þ P43 ðP 21 þ P22 Þ þ a123 P21 P22 P23
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RESEARCH: Original Research
Keywords: Phase-field simulations; Lead-free Piezoelectrics; Field-induced strain; Biocompatibility
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Materials Today
RESEARCH: Original Research
where a1, a11, a12, a111, a112 and a123 are Landau energy coefficients. The values of these coefficients determine the thermodynamic behavior of the bulk phases (paraelectric or ferroelectric).
In our simulation, the difference between the ferroelectric (core)
and paraelectric (shell) regions is determined by the energy fbulk,
and different Landau coefficients were used for the ferroelectric
and paraelectric regions. We employed 2D 128 128 discrete grid
points, periodic and stress-free boundary conditions. The grid
space in real space is Dx ¼ Dy ¼ Dz ¼ 1nm. The ferroelectric region
is set to be a circle with a diameter of 90 nm in the 2D simulation
(or a sphere in the 3D simulation). The other regions are set to be
in paraelectric phase. For the ferroelectric region, the Landau
free energy parameters at room temperature (300 K) are set
to be: a1 = 7.04 103 C2m2N, a11 = 2.67 108 C4m6N,
a111 = 3.33 109 C6m10N,
a12 = 2.67 108 C4m6N,
9 6 10
a112 = 3.33 10 C m N and a123 = 3.0 1010 C6m10N. The
ferroelectric region is in tetragonal phase with a spontaneous
polarization of 0.3C m2 at 300 K. For paraelectric region, the
Landau free energy parameters at room temperature (300 K) are
set to be a1 = 2.2 103 C2m2N, a11 = 2.67 108 C4m6N,
a111 = 3.33 109 C6m10N,
a12 = 2.67 108 C4m6N,
9 6 10
a112 = 3.33 10 C m N and a123 = 3.0 1010 C6m10N.
The gradient energy density, which is associated with the formation and evolution of domain walls, can be expressed as:
f grad ¼ 12 G11 ðP 21;1 þ P 22;2 þ P 23;3 Þ þ G12 ðP1;1 P2;2 þ P2;2 P 3;3 þ P 1;1 P3;3 Þ
þ 12 G44 ½ðP1;2 þ P2;1 Þ2 þ ðP2;3 þ P3;2 Þ2 þ ðP1;3 þ P3;1 Þ2 þ 12 G044 ½ðP1;2 P2;1 Þ2 þ ðP2;3 P3;2 Þ2 þ ðP1;3 P3;1 Þ2 ð4Þ
where Gij are gradient energy coefficients. Pi,j denote oPi/orj. The
gradient energy coefficients are chosen to be G11/G110 = 1.5,
0
G44/G110 = 0.75 and G44 /G110 = G44/G110 = 0.75, where G110 = 7.04 1011 C2m4N. According to these parameters, the simulated width of the domain walls is approximately 1–2 nm, which
is consistent with the existing experimental measurements.
The corresponding elastic energy density can be expressed as:
f elas ¼
1
1
cijkl eij ekl ¼ cijkl ðeij e0ij Þðekl e0kl Þ
2
2
ð5Þ
where cijkl is the elastic stiffness tensor, eij is the total strain, ande0kl
is the electrostrictive stress-free strain, i.e., e0kl ¼Q ij kl Pk Pl . The elas-
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and CZ4, respectively). Raw chemicals Li2CO3 (98.0%), Na2CO3
(99.8%), K2CO3 (99.0%), Nb2O5 (99.99%), Ta2O5 (99.9%),
CaCO3 (99.5%) and ZrO2 (99.0 %, all chemicals from Sinopharm, China) were weighed stoichiometrically and were ball
milled for 24 h at 250 min1 in ethanol. The mixed powder
was calcined at 900 °C for 4 h and ball milled again for 24 h
at 250 min1. MnO2 (97.5%, Sinopharm, China) was added
after the calcination as the sintering aid. For preparation of
NSed samples, the powder was pressed into disks 10 mm in
diameter and sintered at 1120 °C. For preparation of SPSed samples, the same powder was pressed into disks 16 mm in diameter, followed by sintering at 1100 °C for 2 h. The sintered pellets
were then crushed into powder, followed by a second-step spark
plasma sintering (SPS). For SPS, the crushed powder was placed
into a cylindrical graphite die with an inner diameter of 10 mm.
After the SPS chamber was vacuumed, the temperature was
raised to 1050 °C at a rate of 100 °C min1 and held for
3 min. A pressure of 50 MPa was applied along the z-axis during
the entire process. SPS was accomplished in an SPS apparatus
(Dr. Sinter 1020 SPS, Sumitomo Coal Mining Co. Ltd., Kawasaki, Japan). After SPS, the sintered samples were annealed in
air at 850 °C for 12 h.
In situ synchrotron X-ray diffraction (XRD)
In situ synchrotron XRD experiments were carried out at beamline ID15A of the European Synchrotron Radiation Facility. A
beam energy of 71.5 keV was used. The sample was placed in
a specifically designed electric field chamber that allows for
application of a field with the beam transmitting through the
sample [47]. Diffraction images were collected in the forward
direction using a Pilatus 2 M CdTe area detector. In this scattering geometry, each diffraction image contains full orientationdependent data of the scattering vector q and the angle with
respect to the applied electric field vector E. Segments of the
measured images were then integrated into sequential onedimensional diffraction patterns. Peak fitting was done sequentially for further interpretation in Igor Pro 7.0. Simultaneously,
in situ macroscopic strain was performed using an optical displacement sensor coupled to the top surfaces of the sample.
The displacement of the sample surface was used to calculate
the macroscopic strain.
tic constants and electrostrictive coefficients are set to be the same
for the ferroelectric and paraelectric regions, that are SD
11 =
12
SD
m2/N,
and
SD
20 1012 m2/N,
12 = 7.520 10
44 =
12
2
2 4
2 4
m /N, and Q11 = 0.089 C m , Q12 = 0.030 C m ,
20 10
and Q44 = 0.034 C2m4.
The electrostatic energy density is given by:
1
f eles ¼ Ein
Pi Eex
i Pi
2 i
ð6Þ
where Ein
i is the electric field induced by the dipole moments in
the specimen and Eex
i is an applied external electric field.
Solid-state synthesis
Spark plasma sintered (SPSed) and normal sintered (NSed) ceramic samples share the same nominal composition (1 x)
(Na0.49K0.49Li0.02)(Nb0.8Ta0.2)O3-xCaZrO3 (hereafter abbreviated
CZ100x, x = 0, 0.03, 0.035, and 0.04, i.e., CZ0, CZ3, CZ3.5
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Transmission electron microscope (TEM) characterization
TEM specimens were obtained by mechanically polishing to
approximately 20 lm, followed by argon-ion milling (PIPS691,
Gatan, Pleasanton, United States) to a thickness of 50–100 nm.
Microstructure and elemental distributions were investigated
using ARM-200CF (JEOL, Tokyo, Japan) transmission electron
microscope operated at 200 keV and equipped with double
spherical aberration (Cs) correctors. All STEM images were filtered using HREM-Filters Pro/Lite released by HREM Research
Inc. Atomic positions were determined by fitting with Moment
Method & Contour using the CalAtom Software developed by
Prof. Fang Lin.
Other experimental details regarding electrical measurements
of electrical properties, microstructural characterizations, and
biocompatibility are found in Supplementary data.
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Results and discussion
Materials design
The proposed structure to take advantage of the ferroelectricparaelectric interfaces to produce desired domain structures is
shown in Fig. 1. Using phase-field simulations, we adopted a
core–shell system, where the core and shell are in tetragonal ferroelectric and paraelectric phases, respectively. After poling along
the [010] direction (y-direction), a homogeneous ferroelectric
system is expected to approach a single-domain state (Fig. 1b),
while the core–shell system is in a multi-domain state with considerable non-180° domain walls (Fig. 1a). This is because in the
core–shell system, depolarization electrical energy, which is associated with polarization discontinuity, is very high if the ferroelectric core is in the single-domain state. To minimize the free
energy of the whole system, ferroelectric domains with polarizations not only in the y-direction are present. A schematic description of the core–shell system is given in Fig. S1.
Fig. 1c shows simulated microstructural evolution of the core–
shell system under an applied electric field along the y-direction.
With increasing electric field, a gradual growth of the ferroelectric domains with polarization along the y-direction and a
shrinkage of the domains with polarization along other directions are observed. As a result of the domain wall motion, a
highly enhanced electric-field-induced strain is observed for the
ferroelectric-paraelectric heterogeneous system (the core–shell
system) when compared to that of the homogeneous ferroelectric system (Fig. 1d). More importantly, only a minimal increase
in the hysteresis is observed for the strain hysteresis loop of the
ferroelectric-paraelectric system. The low hysteresis is due to
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reversible and low-hysteretic domain wall motion, which can
be explained as follows. In the ferroelectric–paraelectric heterogeneous system, depolarization electric field, which provides
the restoring forces to revert the domain wall motion, depends
on the magnitude of the applied electric field. The domain structure of the ferroelectric core is expected to approach a single
domain state with increasing electric field; thus, the polarization
discontinuity around the ferroelectric-paraelectric interfaces and
the associated depolarization electric field are increased. On the
other hand, the depolarization electric field reduces with decreasing electric field. Because of this characteristic, depolarization
electric field plays a role in avoiding abrupt and/or discontinuous
variation in the domain wall motion during this process. Therefore, although the electric-field-induced strain is enhanced, minimal hysteresis can be maintained for the ferroelectricparaelectric system. Since the piezoresponse of a polycrystalline
material is an ensemble average of the property over all grains
with various crystallographic orientations, the case when the
applied electric field is along the [110] direction was simulated,
as shown in Fig. S2. Despite the smaller strain enhancement in
this case, the enhanced strain also observed for the
ferroelectric-paraelectric system.
To implement the proposed ferroelectric-paraelectric heterogeneous structure in a ceramic material, a two-step sintering process, including both normal sintering (NS) and spark plasma
sintering (SPS) as the first and second step, respectively, was
employed. First, normal sintered (NSed) samples were prepared
by conventional solid-state reaction, and the obtained
microstructure is homogeneous and without the observation of
FIGURE 1
Simulated microstructural evolution and electric-field-induced strain for a ferroelectric-paraelectric system under an electric field along the [010] direction. (a)
A core–shell system after poling along the [010] direction. (b) A homogeneous ferroelectric system after poling along the [010] direction. (c) Microstructural
evolution of the core–shell system with an increasing electric field along the [010] direction. The x- and y-axes represent the [100] and [010] directions,
respectively. Simulated unipolar electric-field-induced strains for the ferroelectric-paraelectric system and a mono-domain ferroelectric system are compared
in (d). The ferroelectric-paraelectric system exhibits an enhanced strain with minimal increase in the hysteresis. The arrows denote the polar vectors of each
grid. Three-dimensional (3D) microstructural evolution of the ferroelectric-paraelectric system with an increasing electric field along the [010] direction is
shown in Fig. S3.
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core–shell structures. Second, NSed samples were crushed into
powder, followed by SPS, during which heterogeneous structures
are induced. The spark plasma sintered (SPSed) and NSed samples
represent the ferroelectric-paraelectric system (Fig. 1a) and the
homogeneous ferroelectric system (Fig. 1b) in the simulations,
respectively. Possible reasons for the appearance of the heterogeneous structures may be related to the different particle size of
the powder for SPS when compared to that of the powder for
NS, as well as the experimental conditions of SPS (externally
applied uniaxial stress and vacuum environment), which are discussed in the Supplementary Data (Fig. S4).
Heterogeneous structures of the SPSed sample were examined
by backscattered electron image, as shown in Fig. 2(a). Core-shell
structures were observed for the SPSed sample while were not
observed for the NSed sample. The presence of the heterogeneous structures is further examined by dielectric characterizations. Temperature-dependent dielectric permittivity for the
SPSed samples with different amounts of CaZrO3 dopants is
shown Fig. 2(b). Curie temperature decreases with increasing
CaZrO3 content, accompanied by a broadening in the dielectric
peak and a decrease in the peak value. The broadening in the
dielectric peak is usually considered to be resulted from compositional heterogeneity [48–49], which leads to local deviations in
the unit-cell distortion and thus a distribution of transition temperatures [50]. A deviation from the Curie-Weiss law above Tm is
observed for all the investigated SPSed samples, as shown in Fig .
S5. Here, a modified Curie-Weiss equation is employed to quantify the broadening in the dielectric peak for different samples, as
shown below [48,51]:
1 1
ðT T m Þc
¼
e em
C
ð1Þ
The characteristic parameter c denotes the degree of diffuseness as well as the degree of compositional heterogeneity. For
ideal ferroelectrics with no heterogeneity, the value of c is 1,
while for ideal relaxor ferroelectrics with significant structural
heterogeneity, the value is 2 [48,51]. For the investigated compositions, it was observed that an increase in c is related to an
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increase in the CaZrO3 content, indicating a higher level of
heterogeneity is present for samples with higher CaZrO3 content
(Fig. 2c). As the CaZrO3 content increases, the parameter c eventually saturates to a maximum value of 2.01 ± 0.01 for the SPSed
sample. This indicates that a considerable degree of heterogeneity is reached; therefore, CZ3.5 composition is employed for further investigation.
Microstructural characterization
The presence of local structural heterogeneity was observed
directly by transmission electron microscopy (TEM). A lowmagnification high-angle annular dark-field (HAADF) image of
the SPSed sample is given (Fig. 3a). A contrast between the core
and shell regions within one individual grain is observed for
the multiple grains in the investigated area. A magnified image
focused on an individual grain is shown in Fig. 3b, where
domain patterns are observed in the core region while are missing in the shell region. The atomic structures within the core
and shell regions of the grain are characterized by highresolution HAADF images, and the positions of the A-site (K/
Na/Ca) and B-site (Nb/Ta) atomic columns are given in Fig. 3c
and Fig. 3d, respectively. According to the positions, atomic displacements are presented as vectors pointing from the center of a
B-site cation to the center of its four nearest neighboring A-site
cations. These atomic displacements represent the magnitudes
and directions of the polar vectors for each unit-cell column
[52,53]. For the core region, the averaged displacements along
the [0k0] and [00l] directions are 0.184 Å and 0.026 Å, respectively and thus, the averaged polarization is close to the [0k0]
direction. Although there are some fluctuations present in individual polar vectors (Fig. 3c), the majority of the polar vectors
align themselves along the [0k0] direction, which indicates the
core is in tetragonal ferroelectric phase. On the other hand, polar
vectors in the shell region are of smaller magnitudes (0.080 Å and
0.044 Å along the [0k0] and [00l] directions, respectively), as
shown in Fig. 3d. Above observations indicate the core region
is in the polar state with large atomic displacement and ferroelectric domains, while the shell is close to the non-polar state with
FIGURE 2
Microstructural and dielectric evidence of the structural heterogeneity. Backscattered electron (BSE) images of (a) the SPSed sample. (b) Temperaturedependent relative dielectric permittivity of the SPSed samples with different amounts of CaZrO3 dopants (measured at 1 kHz). (c) lg(1/e 1/em) versus lg
(T Tm) at temperatures above Tm (which represents the temperature where the permittivity maximizes in the temperature-dependent permittivity curves)
to obtain the parameter c, which is generally recognized as the degree of diffuseness and can also be regarded as the degree of compositional heterogeneity
in this case. A large c value indicates a high level of heterogeneity. Temperature-dependent loss factor of the SPSed samples with different amounts of
CaZrO3 dopants is shown in Fig. S6.
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small atomic displacement. High-magnification HAADF image
and EDS mapping of the grain are examined, as shown in Fig. 3e.
Distinct element segregations observed in the EDS mapping are
in good agreement with the contrast in the HAADF image. In
the core region where striped domains are present, a compositional depletion of tantalum and an enrichment of niobium
and potassium are observed. It was shown that Ta doping can significantly decrease the Curie temperature [54–55] and thus, Tariched regions are expected to be close to the paraelectric state.
A comprehensive characterization of the element distribution
in multiple grains is shown in Fig. S7, where tantalum and niobium segregations are the most significant. This may be due to
Ta5+ (CN = 6, 0.64 Å) and Nb5+ (CN = 6, 0.64 Å) ions have the
lowest ionic radii and are most likely to diffuse. Potassium segregation is more obvious than sodium segregation, which may be
explained by the fact that potassium is more volatile than
sodium [56] despite the larger ionic radii of potassium. The presence of stripe domains in the core region indicates the presence
of ferroelectric polar state, while the regions without domains are
expected to be close to the non-polar state. Both the dielectric
(Fig. 2) and TEM characterizations (Fig. 3) confirm successful
introduction of local heterogeneous structures (i.e., coexistence
of ferroelectric and paraelectric regions) into the SPSed sample.
Electric-field-induced strains
Unipolar electric-field-induced strains of the SPSed and NSed
samples were characterized and compared, as shown in Fig. 4.
At room temperature, a significant increase of 52–62% in the
unipolar strains was observed for the SPSed sample when com-
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pared to that for the NSed sample. The degree of hysteresis,
which is defined as the strain difference DS at half of the maximum electric field divided by the maximum strain (Fig. S8), is
13% for the SPSed sample at 4 kV/mm, which is similar to that
for the NSed counterpart (11% at 4 kV/mm, see Table S2).
According to previous phase-field simulations, the enhanced
strain with minimal hysteresis increase observed for the SPSed
sample is resulted from enhanced domain wall motion due to
the presence of ferroelectric-paraelectric heterogeneous structures, which were experimentally confirmed by the TEM characterization. Note that the observed strain enhancement for the
SPSed sample is not as significant as that in the phase-field simulations when the electric field is along the [010] direction
(Fig. 1d). This is because the strain enhancement for a real polycrystalline piezoceramic is expected to be a value between the
simulated value along the [010] direction and that along the
[110] direction.
Temperature-dependent electric-field-induced strains were
examined and the strain values were also presented in the form
of normalized piezoelectric constant d33* (strain/electric field),
as shown in Fig. 4b and Fig. 4c, respectively. At 25 °C, a d33* of
500 pm V1 is observed for the SPSed sample, which is 62%
higher than that for the NSed sample. At elevated temperatures,
d33* increases with increasing temperature between 25 °C and
85 °C and becomes stable between 85 °C and 150 °C. At
120 °C, d33* reaches a peak value of 605 pm V1, which is
160% higher than that of the NSed sample at the same temperature. At 150 °C, d33* of the SPSed sample is 200% higher than
that of the NSed sample. Temperature-dependent polarization
FIGURE 3
TEM characterization of the local heterogeneous structures and element segregations in the SPSed sample. (a) Low-magnification HAADF image of the grains.
A contrast between the core and the shell is present for multiple grains. (b) High-magnification angular bright-field (ABF) image of an individual grain along
the [100] zone axis. Regular striped domains are present in the core while are absent in the shell of the grain. Atomic-resolution HAADF image for the (c) core
and (d) shell, recorded along the crystallographic [100] direction. The polar vectors (arrows) are given for each unit-cell column in the atomic-resolution
image. (e) Energy-dispersive X-ray spectroscopy (EDS) mapping results within an individual grain.
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FIGURE 4
Electric-field-induced strain and temperature stability of the SPSed and NSed samples. (a) Room-temperature unipolar electric-field-induced strains under
different electric fields at 0.1 Hz. (b) Temperature-dependent strains under 4 kV mm1 at 0.1 Hz. (c) Temperature-dependent normalized piezoelectric
constant d33* under different electric fields. (d) Comparison in the temperature dependences of d33* with PZT4 and PZT5H materials, normalized to their
room temperature value d33*RT. Data for PZT4 and PZT5H are taken from Ref. [28].
hysteresis loops of the SPSed sample from 20 °C to 150 °C is
shown in Fig. S9. The temperature-dependence of the strains
(d33*) of the SPSed and NSed samples was compared to that of
two commercial PZT materials, PZT4 and PZT5H, as shown in
Fig. 4d. Among all the PZT materials, PZT4 material is known
for a combination of good thermal stability and high piezoresponse. It can be seen that d33* for the SPSed sample varies less
than 20% between 20 °C and 150 °C, exhibiting good temperature stability that is comparable to that for the PZT4 sample.
The good temperature stability of the SPSed sample is evaluated
from the perspective of temperature-dependent polarization values, as shown in Fig. 10. In contrast, d33* for both the NSed sample and PZT5H material show strong temperature dependence.
This indicates that the SPSed material is a promising candidate
for actuator applications that require thermally stable strain
performance.
To demonstrate that the high strain in the SPSed sample arises
from reversible domain wall motion, in situ synchrotron X-ray
diffraction (XRD) experiments were performed. The length scale
of the core–shell structures produces a very high degree of particle size and strain broadening in XRD patterns. Additionally, in
this case, any ferroelectric/ferroelastic distortion occurring is
small, with expected c/a ratio of 1.004. Thus, the high-energy
XRD results cannot conclusively resolve tetragonal ferroelectric
in the core region from tetragonal or pseudo-cubic paraelectric
50
shell regions of the material. Since the presence of heterogeneous
structures was already confirmed by TEM, high-energy XRD will
be focused on characterization of domain switching in the SPSed
sample.
Using in situ synchrotron XRD experiments, the contributions
of electric-field-induced lattice variation and domain switching
to the strain for the SPSed sample can be quantitatively characterized by analyzing the variations in the (200)pc, (111)pc and
(222)pc type reflections. Diffraction profiles for the (111)pc,
(002)pc/(200)pc and (222)pc peaks during the application of a
unipolar electric field cycle (Emax = 3 kV mm1) and contour
plots of the (002)pc/(200)pc reflections are shown in Fig. 5(a)
and Fig. 5(b), respectively. From the diffraction patterns, changes
in the relative intensities of the (002)pc/(200)pc peaks and variation in the peak positions are visible. The total electric-fieldinduced strain in perovskite ceramics results from the combined
contributions of intrinsic lattice strain and extrinsic strain due to
non-180° ferroelectric domain switching. [57] Since the in situ
field-dependent X-ray data collected here include the full orientation dependence of domain switching magnitudes, it is possible to calculate the domain switching strain. Detailed
calculations of the electric-field-induced extrinsic strains as well
as the proportion of reversible domain switching are provided
in Fig. S11. Due to the very small ferroelastic distortion of this
system, it is not possible to calculate an intrinsic strain contribu-
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FIGURE 5
In situ synchrotron XRD experiments. Evolution of in situ synchrotron X-ray diffraction profiles for the SPSed sample with externally applied fields. (a) The peak
profiles of (111)pc, (002)pc/(200)pc and (222)pc reflections at selected states during the application of unipolar electric field: remanent state after the first
unipolar cycle (E0), at 3 kV mm1 (Emax) and after the removal of the electric field (Erem). (b) Contour plots of (002)pc/(200)pc reflections during application of
the electric field. The profiles are indexed based on the tetragonal symmetry.
FIGURE 6
Biocompatibility of the investigated KNN compositions and PZT. (a) SEM images of L929 cells cultured on the surfaces of PZT, CZ3.5, and CZ4. Scale bar
is100 lm. (b) Quantitative evaluations of cells density were analyzed based on the SEM images. Error bars represent the standard deviation based on 10
repeated measurements (n = 5). N, number of cells. *P < 0.05, ***P < 0.001. (c) Cell proliferation characterization via CCK-8 assay of L929 cells on the surfaces
of PZT, CZ3.5 and CZ4 after 1, 3, and 5 days of incubation. Error bars represent the standard deviation based on 10 repeated measurements (n = 3). *P < 0.05,
**P < 0.01.
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tion from well resolved peaks as demonstrated previously by Pramanick et al. [58]. Quantification of the extrinsic strain shows
that non-180° ferroelectric domain switching accounts for 50%
of the apparent macroscopic strain for the SPSed sample. The
domain switching fraction of 50% is similar to that of a soft
PZT [58] or other microstructurally inhomogeneous systems
[59], which accounts for the enhanced electric-field-induced
strains of the SPSed sample.
Finally, the safety and biocompatibility of the investigated
lead-free compositions (CZ3.5 and CZ4) were examined and
compared with that of PZT ceramics. Herein, we cultured L929
cells on the substrates of lead-free compositions (CZ3.5 and
CZ4) with a comparison of PZT ceramics to demonstrate the
in vitro biocompatibility of the investigated lead-free compositions. Cells viability and proliferation were assessed by SEM
and cholecystokinin octapeptide (CCK-8) analyses on the surfaces of PZT, CZ3.5, and CZ4. SEM images characterize distinct
cell densities after L929 cells were cultured on different modified
interfaces on days 1, 3, and 5, as shown in Fig. 6a. Statistically,
the number of cells were increased on each of the surfaces over
time, and the cells packed more density on surface of CZ3.5
and CZ4 than that of PZT on day 3 and 5, indicating that the surfaces of CZ3.5 and CZ4 facilitate cells proliferation (Fig. 6b). In
cell viability studies with the CCK-8 assay (Fig. 6c), lower values
indicate increased cytotoxicity. In this experiment, we set coverglass as the control group, the optical density (OD) values of this
group were consistently higher than those of the other groups.
The least amount of viable cells was found on the surface of
PZT during all the test periods, especially on days 3 and 5. With
increasing incubating time, CZ3.5 and CZ4 group displayed similar cell viability, while the OD values of them were significantly
elevated when compared with that of the PZT group, indicating
their better biocompatibility. This result was further supported
by quantitative evaluations of L929 cells according to the SEM
images at the same time points (Fig. 6b). Our results indicate
the biocompatibility reported for undoped and/or slightly doped
KNN materials [36,45,42,60] are maintained for the investigated
KNN compositions, which are doped with lithium, tantalum, calcium and zirconium. The investigated CZ3.5 composition exhibits a combination of enhanced piezoresponse and good
biocompatibility.
Conclusions
In summary, based on phase-field simulations and a series of
experimental characterizations, we demonstrated the possibility
to enhance electric-field induced strain via the introduction of
ferroelectric-paraelectric heterogeneous structures into a doped
KNN composition using a two-step sintering process. Increased
depolarization energies arising from ferroelectric-paraelectric
interfaces enhance domain wall motion and hence, enhance
the magnitude of electric-field-induced strain. The non-toxicity
and biocompatibility of the investigated compositions are consistent with previous reports on undoped and/or lightly doped
KNN materials. This work shows local structural engineering
could be used to tailor the functional properties of piezoceramics.
The investigated KNN materials are non-toxic and with
enhanced electric-field-induced strain, and thus are potential
candidates for implantable biomedical devices
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CRediT authorship contribution statement
Mao-Hua Zhang: data curation, formal analysis, investigation, writing-original draft, writing-review & editing. Qinghua
Zhang: data curation, formal analysis, software. Ting-Ting
Zhu: data curation, formal analysis, methodology. Geng Li:
investigation. Hao-Cheng Thong: investigation. Li-Ying
Peng: investigation. Lisha Liu: data curation, formal analysis,
software. Jing Ma: writing-review & editing. Yang Shen:
writing-review & editing, supervision. Zhijian Shen: writingreview & editing, supervision. John Daniels: supervision, validation, writing-review & editing. Lin Gu: supervision. Bing
Han: conceptulization, funding acquisition, supervision.
Long-Qing Chen: supervision, funding acquisition. JingFeng Li: supervision. Fei Li: conceptulization, funding aquisiion, supervision, writing-review & editing. Ke Wang: conceptulization, funding aquisiion, supervision, writing-review &
editing, project administration.
Declaration of Competing Interest
The authors declare that they have no known competing
financial interests or personal relationships that could have
appeared to influence the work reported in this paper.
Acknowledgements
K. W. acknowledges the support of the National Nature
Science Foundation of China (Grant Nos. 51822206, 52032005,
51761135118). F.L. acknowledges the support of the National
Nature Science Foundation of China (Grant Nos. 51572214 and
51761145024). The effort at Penn State is supported by U.S.
Department of Energy, Office of Basic Energy Sciences, Division
of Materials Sciences and Engineering under Award DE-FG0207ER46417.
Appendix A. Supplementary data
Supplementary data to this article can be found online at
https://doi.org/10.1016/j.mattod.2021.02.002.
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