Journal of Materials Science & Technology  2020 , 44 (0): 54-61 https://doi.org/10.1016/j.jmst.2019.11.012

Research Article

Hole-pinned defect-dipoles induced colossal permittivity in Bi doped SrTiO3 ceramics with Sr deficiency

Yulong Qiaoa, Weili Liab*, Yulei Zhanga, Lu Jinga, Chang Gaoa, Wenping Caoc, Dan Xud, Weidong Feiae*

a School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China
b National Key Laboratory of Science and Technology on Precision Heat Processing of Metals, Harbin Institute of Technology, Harbin, 150001, China
c School of Light Industry, Harbin University of Commerce, Harbin, 150028, China
d Key Laboratory of Quantum Manipulation & Control of Heilongjiang Province, Harbin University of Science and Technology, Harbin, 150080, China
e State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin, 150001, China

Corresponding authors:   * School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China. E-mail addresses: wlli@hit.edu.cn (W. Li),wdfei@hit.edu.cn (W. Fei).

Received: 2019-09-2

Revised:  2019-10-28

Accepted:  2019-11-4

Online:  2020-05-01

Copyright:  2020 Editorial board of Journal of Materials Science & Technology Copyright reserved, Editorial board of Journal of Materials Science & Technology

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Abstract

Bi doped SrTiO3 ceramics with Sr deficiency, i.e. Sr1-1.5xBixTiO3 (x=0, 0.01, 0.05, 0.1), were prepared via conventional solid-state reaction route. A colossal permittivity (CP) over 104 with low dielectric loss less than 0.05 was obtained in x=0.05 Sr1-1.5xBixTiO3 ceramics. In addition, the dielectric constant is maintained at a value greater than 104 in the range of 102-105 Hz and almost frequency independent. Phase structure analysis and density functional theory calculations suggest that the BiSr· - VSr" - BiSr· defect complex with hole-pinned defect-dipoles maybe responsible for the high-performance CP properties. This work gives a new way to achieve high performance CP materials in ABO3 perovskite ceramics.

Keywords: Bi doped SrTiO3 ; Sr deficiency ; Colossal permittivity ; Low dielectric loss

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Yulong Qiao, Weili Li, Yulei Zhang, Lu Jing, Chang Gao, Wenping Cao, Dan Xu, Weidong Fei. Hole-pinned defect-dipoles induced colossal permittivity in Bi doped SrTiO3 ceramics with Sr deficiency[J]. Journal of Materials Science & Technology, 2020, 44(0): 54-61 https://doi.org/10.1016/j.jmst.2019.11.012

1. Introduction

Colossal permittivity (CP) materials have attracted much attention from the community of microelectronics due to their potential applications for device miniaturization [[1], [2], [3], [4], [5], [6]]. High-performance CP materials need to exhibit temperature- and frequency-stable colossal permittivity (CP > 1000) as well as sufficiently low dielectric loss. For practical applications, CP materials should also have large breakdown electric field, excellent voltage independent permittivity and low dielectric loss.

At present, many candidate colossal permittivity materials have been investigated, such as CaCu3Ti4O12 (CCTO) [7,8], doped NiO2 [9] and doped TiO2 [2,[10], [11], [12], [13]]. Especially, the modified TiO2 ceramics have a CP property [[14], [15], [16]]. Among these materials, the origin of colossal permittivity for CCTO and doped NiO2 ceramics are interface effects including surface barrier layer capacitors (SBLC) and internal barrier layer capacitance (IBLC). The dielectric constants of CCTO and doped NiO2 materials can reach about 105 at room temperature. Unfortunately, many experiments have reported that dielectric materials with interface effects had relatively high dielectric loss. For example, dielectric loss of CCTO usually stays over 0.2 [7,8,17,18]. This disadvantage limited the practical application of these CP materials.

Nowadays, Hu et al. [19] found that (Nb + In) co-doped rutile TiO2 (NITO) ceramics exhibit colossal permittivity as well as low dielectric loss (~0.05) over a temperature range from 20 to 200 °C. The origin of high CP performance in this material is attributed to the highly localized electron-pinned defect-dipoles. The triangular 2In3+-V-Ti3+ complexs assist in binding electrons, while the diamond-shaped Nb5+-2Ti3+-ATi complexs provide delocalized electrons that bind together, resulting in large defect dipole clusters. However, the colossal dielectric permittivity and low dielectric loss obtained in (Nb + In) co-doped TiO2 ceramics strongly depend on the synthesis conditions. In the follow-up work of Hu et al., the colossal permittivity in (Nb + Al) co-doped TiO2 ceramics is closely related to the distribution of Ti3+ [20]. The defect-complex structure with localized charges in co-doped TiO2 is too complex and sensitive to the formation conditions. This makes the donor acceptor co-doped TiO2 found by some researchers have no CP properties [21]. Even some co-doped TiO2 ceramics can have very large dielectric constant, their dielectric loss is much bigger than 0.1 since charges are not localized perfectly in the defect-complex structures and the presence of interface effect [22].

If a simple way could be created to construct defect-complex structures with more stable dielectric performance, it would be much easier to obtain CP materials. Some studies have reported that many different defect complexes can be easily formed by single doping or donor-acceptor co-doping in ABO3 perovskite [[23], [24], [25], [26]]. SrTiO3 (STO) is an ABO3 perovskite incipient ferroelectric with low dielectric loss, but remains in a paraelectric phase down to near 0 K [[27], [28], [29], [30]]. This makes STO have a good dielectric temperature stability, since STO does not undergo a ferroelectric phase transition near room temperature. However, as the electrical properties of STO is very sensitive to doping elements and preparation conditions, a crossover occurs from an insulating to a semiconducting, then a metallic behaviour occurs when different modified strategies are used in STO [28]. Thus, in many reports, the dielectric properties of STO vary greatly and can be regulated by means of doping. The CP property of doped STO ceramics has been reported by some authors [[31], [32], [33], [34], [35]], but the dielectric loss of some doped STO ceramics are very high. Although the authors speculated or simulated possible electronic-pinned defect complexes to explain the CP properties in STO ceramics, the large dielectric loss indicates that the electrons are clearly not “pinned” perfectly by these defect structures. According to defect chemistry, two kinds of defects may exist in pure STO ceramics. They are oxygen vacancy (VO··) and Sr vacancy (VSr"). The trivalent impurity metal cation defect (ASr·) is introduced into the STO crystal when a trivalent metal cation substituted the A-site in STO. In order to keep electrical neutrality of STO systems, more VSr" defects would be produced or a part of Ti4+ change to Ti3+. In fact, the defect of ASr· is a positive electricity centre with a weak binding effect on electrons which can provide delocalized electrons or localized hole charges. Donor doping inhibits the generation of VO·· defects. Although VO·· defects can still exist in doped STO because of the high-temperature sinter process [28], it is more difficult to form a defect complex consisting of VO·· and Ti3+. Other defect complexes consisting of ASr· and Ti3+ cannot limit the migration of electrons, too. Therefore, only VSr" which acts as an acceptor and provide hole charges can help to form localized charges defect complex with trivalent impurity metal cation defect. Unlike VO··, metal cation vacancies are more stable and less sensitive to formation conditions. If ASr· defect and VSr" can form defect complex, STO may have a potential to achieve CP property.

In order to obtain CP and low dielectric loss characteristics in STO ceramics, Bi doped STO with pre-designed Sr deficiency has been prepared using a special heat treatment process (sintered at 1500 °C) and the dielectric properties of the doped ceramics were investigated in this study. CP properties with a low dielectric loss and a wide frequency stability range are obtained in the Sr1-1.5xBixTiO3 ceramics. First-principle calculations have been carried out to investigate the defect configurations.

2. Experimental and method

The Sr1-1.5xBixTiO3 ceramics were prepared by solid-state reaction route. The raw materials Bi2O3 (99.9%), SrCO3 (99.95%), TiO2 (99.8%) powders were weighed according to the chemical formula and milled in ethanol using zirconia balls for 12 h with 300 rpm. The suspensions were then dried at 100 °C for 8 h. Following this, the powders were calcined at 900 °C for 4 h in an alumina crucible and ball-milled again by the same method. After drying, the powders were mixed with polyvinyl alcohol (PVA) and pressed into a 10 mm diameter pellet. Slow heating at 600 °C for 1 h burned out the binder. The compacted pellets were covered by pure STO powders at the surface and bottom then sintered at 1500 °C for 2.5 h in air. Silver electrodes were formed on both surfaces of the disk and annealed at 600 °C for 30 min. Then some samples were kept in a drying oven at 80 °C for 48 h. In order to analyse the contribution of different electrodes to permittivity, Cu, Al and Ag electrodes were coated on both surfaces of the polished pellet by vacuum evaporation systems and fired at 200 °C for 30 min for dielectric properties testing.

The first-principle calculations based on density functional theory (DFT) were performed using the materials studio CASTEP. The exchange-correlation interactions are approximated by generalized gradient approximation (GGA) according to Perdew, Burkeand Ernzerh (PBE) method. The ionic potentials were described by Ultrasoft-pseudopotentials plane wave. A 2 × 2 × 2 Monkhorst Pack k-point mesh for the primitive cell was used for Brillouin zone integration. A plane wave cutoff energy of 500 eV was applied for all calculations. The calculations were completed using cubic symmetry supercell. This supercell is made by 3 × 3 × 3 cells with two Bi ions and one Sr vacancy in each supercell. The different configurations of Bi doped STO 3 × 3 × 3 supercells are shown in Fig. 6.

The phase purity and crystal structure were determined using X-ray diffraction (XRD) on a Philips X’Pert PRO diffractometer with Cu Kα radiation. The 2θ result of fine scan XRD was obtained by Jade 6.0 fit peak profile analyse. The surface microstructure of the samples after sintering was collected on a Zeiss Merlin Compact scanning electron microscope (SEM). The dielectric properties were measured using an Agilent 4294A impedance analyser. The temperature dependent dielectric properties were tested by software control of the connected TZDM-RT-600 Four-channel signal acquisition and Agilent 4294A impedance analyser instruments. Low temperature dielectric constant was measured by Physical Property Measurement System.

3. Results and discussion

Fig. 1 shows the frequency dependent dielectric constant and loss measured at room temperature for Bi doped and undoped STO ceramics. The dielectric constant of pure STO is about 250. As the Bi content increases from 0% to 10%, the dielectric constant first increases and then decreased to 102 magnitudes. Colossal permittivity about 104 is obtained in Bi doped STO ceramics with x = 1% and 5%. The best performance of dielectric properties is obtained in the Sr0.925Bi0.05TiO3 ceramics with a CP about 14000 and a low dielectric loss less than 0.05 over a broad frequency range. In addition, the dielectric constant and loss of 5% sample are almost independent to frequency from 102 Hz to 105 Hz.

Fig. 1.   Frequency dependent dielectric properties. (a) Frequency dependent dielectric constant of Bi doped and undoped STO ceramics. (b) Frequency dependent dielectric loss of Bi doped and undoped STO ceramics.

The temperature dependent dielectric constant and loss are shown in Figs. 2 and 3 . The dielectric constant of ceramics with x = 1% and x = 5% undergoes a slight decrease from 30 to 100 °C, then a weak dielectric relaxation occurs between 100-250 °C, at last rises slightly after 250 °C. Comparing to x = 0% and x = 10% samples, ceramics with x = 1% and 5% have stable dielectric constant maintaining at about 104 in the testing temperature range from room temperature to 400 °C. The x = 5% sample exhibits relatively stable temperature dependent colossal permittivity as well as low dielectric loss (<0.1) in the range of 30-130 °C.

Fig. 2.   Temperature dependent dielectric constant of Bi doped STO ceramics. (a) Temperature dependent dielectric constant of undoped STO. (b) Temperature dependent dielectric constant of x = 1% sample. (c) Temperature dependent dielectric constant of x = 5% sample. (d) Temperature dependent dielectric constant of x = 10% sample.

Fig. 3.   Temperature dependent dielectric loss of Bi doped STO ceramics. (a) Temperature dependent dielectric loss of undoped STO. (b) Temperature dependent dielectric loss of x = 1% sample. (c) Temperature dependent dielectric loss of x = 5% sample. (d) Temperature dependent dielectric loss of x = 10% sample.

Surface morphology of BixSr1-1.5xTiO3 ceramics with different x are displayed in Fig. 4. From the figure, undoped STO has spherical grains with non-uniformly distributed sizes about 5-20 μm and possesses some pores. Different from the undoped STO, the BixSr1-1.5xTiO3 ceramics with x = 1% and 5% show irregular flaky polygon grains and dense microstructures without any pores. The grain size of Bi doped STO is larger than that of undoped STO. But the grain sizes have no significant difference between x = 1% and x = 5% samples. Ceramics with the concentration of x = 10% have rough surface microstructure without visible grain boundaries. Bi contributes to the sintering of STO ceramics, making the crystal grains grow more easily at higher temperatures. So that the grain size after Bi doping is larger than that of pure STO. However, when the doping amount is too high, the increased content of Bi lead to remelting of the ceramic surface.

Fig. 4.   Surface morphology of Bi doped STO ceramics with different x. (a) x = 0%; (b) x = 1%; (c) x = 5%; (d): x = 10%.

In order to explore the possible origin of CP phenomena, the complex impedance plots of BixSr1-1.5xTiO3 ceramics are carried out at different temperatures to investigate the interface effect on grain boundary. Fig. 5 displays the impedance spectra plots with different doping concentrations of BixSr1-1.5xTiO3 ceramics. The impedance spectra plots of all the ceramics can be fitted well by a single R-C equivalent circuit model where R and C represent the resistance and corresponding capacitor. No obvious grain/grain boundary effects can be found. These results indicate that the origin of CP for x = 1% and 5% BixSr1-1.5xTiO3 ceramics is not mainly the IBLC effect.

Fig. 5.   Complex impedance plots of BixSr1-1.5xTiO3 ceramics with different x. (a) x = 0%; (b) x = 1%; (c) x = 5%; (d) x = 10%.

To explore the influence of SBLC effect on the ceramics, Cu, Al and Ag electrodes are used to measure the dielectric constant of BixSr1-1.5xTiO3 ceramics with x = 5%, respectively. As these measurements are done by the same sample, the thickness of ceramic might decrease. But very small decrease of thickness has little effect on the calculation of dielectric constant. The geometry data of the ceramic sample is measured repeatedly after removing and coating with new electrodes to ensure the accuracy of the dielectric constant. The complex impedance plots and the dielectric constant measured from different types of electrodes are displayed in Fig. 6(a) and (b). From the figure, different electrodes have little effect on the dielectric constant. And no obvious interface effects can be observed from the complex impedance plots which are measured by different types of electrodes. These results suggest that the SBLC effect is not the major reason for CP property of BixSr1-1.5xTiO3 ceramics. The leakage current result of pure STO and Bi doped STO ceramics with x = 5% is shown in Fig. 6(c). The forward and reverse leakage currents are symmetrical, but Bi doping causes a significant increase in leakage current which may be related to the IBLC effect. Thus, this suggests that interface effect may be related to dielectric property.

Fig. 6.   (a) Frequency dependent dielectric constant of x = 5% BixSr1-1.5xTiO3 ceramics with Ag, Al and Cu electrodes. (b) Complex impedance plots of x = 5% BixSr1-1.5xTiO3 ceramics with different electrodes. (c) Leakage current result of pure STO and x = 5% BixSr1-1.5xTiO3 ceramics.

No IBLC or SBLC effects are obtained in Bi doped STO ceramics. Another possible explanation is electron-pinned defect-dipoles like the In, Nb co-doped TiO2 ceramics. In order to gain further insight into the nature of the CP of BixSr1-1.5xTiO3, XRD patterns of Bi doped STO ceramics are studied. Fig. 7(a) shows XRD patterns of Bi doped STO and pure STO. The major diffraction peaks match well with the pure STO perovskite phase (PDF 73-0661 or 84-0444). Fine scan XRD method has been widely used in analysing the distribution of defects [[36], [37], [38], [39], [40]]. In this study, fine scan XRD patterns of these ceramics were obtained at room temperature as shown in Fig. 7(b). From the figure, all 2θ of the peaks decrease with increasing x from 1% to 10%, which means the lattice parameters rise with the increase of Bi content. Since the ionic radii of Bi3+ (0.117 nm) is smaller than Sr2+ (0.144 nm) [41], the lattice parameters should reduce with the increase of x according to the size of ionic radii, but the XRD shows the opposite result. In the BixSr1-1.5xTiO3 system, VSr" is pre-designed. So that the content of VSr" increases with the increase of x. It has been verified that the radius of VSr" is significantly larger than Sr2+ ion [35,42]. Therefore, the lattice parameters rise with increasing x. From Fig. 7(b), it can also be found that the 2θ of different peaks has different changing trends. In order to compare the change of 2θ, the 2θ of (200) peaks were used to minus 2θ of (110) peaks and (111) peaks respectively, the results are shown in Fig. 7(c) and (d). Fig. 7(e) shows the data that 2θ of (111) peaks minus 2θ of (110) peaks. From these figures, the values of 2θ(200)-2θ(110) and 2θ(200)-2θ(111) increase with x increasing, but the value of 2θ(111)-2θ(110) is almost unchanged. Since all the 2θ decrease with x increasing, the different changing trends of 2θ(200)-2θ(110), 2θ(200)-2θ(111) and 2θ(111)-2θ(110) means that 2θ(200) decreases more slowly than 2θ(110) and 2θ(111). This result indicates that the (200) interplanar spacing (d200) increases more slowly than others with increasing x. STO is a cubic perovskite crystal. In Bi doped STO, many researchers have reported that Bi3+ substitutes Sr2+ and occupies the A-site of STO. So that, the case where Bi3+ occupies the Ti-site is not considered. As the size of Bi3+ is smaller than Sr2+, so the interplanar spacing should reduce if only Bi3+ substitutes Sr2+, which is inconsistent with the experimental results. Since the size of VSr" is much larger than Sr2+, the increase of the interplanar spacing is due to the VSr". Also due to the increase of the interplanar spacing in different directions is different, we have reasons to believe that the interplanar spacing in the [100] direction increase slowly is because the BiSr· and VSr" form defect complex in this direction.

Fig. 7.   XRD patterns of Bi doped STO ceramics. (a) The XRD patterns of Bi doped STO ceramics. (b) The fine scan XRD patterns of (110), (111), (200) peaks. (c)-(e) The change of 2θ(200)-2θ(110), 2θ(200)-2θ(111) and 2θ(111)-2θ(110) with increased x%.

To further confirm the above analysis, first-principle calculations based on the density functional theory (DFT) were performed on Bi doped STO 3 × 3 × 3 supercells. Detailed calculation steps and methods are shown in the experimental methods section. Four possible configurations with different BiSr· - VSr" defect complexs in STO (001) plane are shown in Fig. 8. As described in the fine scan XRD results analysis, the BiSr· - VSr" defect complexes are easily formed along the [100] direction, so we considered that the BiSr· - VSr" defect complexes are more likely to be parallel to the (001) plane. For the doping supercell itself, in order to ensure the electrical neutrality of the system, each VSr" in the supercell should combine with two Bi3+ impurity ions to ensure the charge balance. From Fig. 8, there are nine Sr2+ positions in the (001) plane of STO 3 × 3 × 3 supercell. The central Sr2+ position is replaced by VSr" and then two other Sr2+ positions are chosen to be substituted by Bi. Finally, four possible configurations can be obtained, and according to the crystal structure symmetry, the BiSr· - VSr" defect complexes in other cases are equivalent to one of the four configurations in the figure. In order to know the change of lattice parameters after Bi doping, undoped STO 3 × 3 × 3 supercell was also calculated to compare lattice parameters as shown in Fig. 9(a). The values of lattice parameters and volume of supercells are shown in Table 1. The results show that the lattice parameters ‘a, b, c’ changed a lot with the different configurations. Bi doping caused the lattice parameters and the volume of supercells increase. This is consistent with the fine scan XRD results. It is worth mentioning that only the lattice parameter ‘a’ of configuration 2 and 4 has a little decrease compared to undoped STO, but the lattice parameter ‘b’ of configuration 2 increase the most, only configuration 4 has the smallest lattice distortion. Therefore, configuration 4 is the most possible defect complex which caused the abnormal variation of d200 in Bi doped STO lattice. The schematic diagram of configuration 4 is shown in Fig. 9(b), in this configuration two BiSr· and one VSr" form the BiSr· - VSr" - BiSr· defect complex in a straight line parallel to [100] direction, the two BiSr· move to VSr" and the lattice parameter ‘a’ decrease, while the lattice parameters b and c are little larger than undoped STO. This result can explain the previous fine scan XRD results well. As the size of VSr" is significantly larger than Bi3+ and Sr2+ ions, the lattice parameters of the doped STO ceramics should rise with the increase of VSr" content. However, when BiSr· and VSr" form the BiSr· - VSr" - BiSr· configuration in [100] direction, the large lattice spacing change caused by VSr" is partially neutralized by the aggregation of Bi3+ with a smaller ionic radius in [100 direction. Therefore, the (200) interplanar spacing (d200) of Bi doped STO increases more slowly. In the BiSr· - VSr" - BiSr· configuration, Sr vacancy provides hole charges and Bi3+ prefers to accept them, the hole charges are pinned in the BiSr· - VSr" - BiSr· defect complex to form hole-pinned defect-dipoles which caused the Bi doped STO ceramics possess CP property with very low loss.

Fig. 8.   Different ion position of BiSr· - VSr" defect complex in STO 3 × 3 × 3 supercells (001) plane.

Fig. 9.   (a) Lattice parameters of different BiSr· - VSr" defect complex configurations. (b) Schematic diagram of BiSr· - VSr" - BiSr· defect complex at (001) plane. The diagram at the bottom explains the lattice distortion in BixSr1-1.5xTiO3.

Table 1   Lattice parameters and volume of different supercells.

Configurationa (Å)b (Å)c (Å)Volume (Å3)
STO11.809211.809211.80921647.81
112.051211.777411.75491668.40
211.807112.005111.74951665.41
311.900611.901611.74901664.02
411.802711.818611.81851648.57

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4. Conclusion

In summary, the dielectric properties of Bi doped STO ceramics fabricated by conventional solid-state reaction method were measured. As a result, a frequency-independent dielectric response with a colossal dielectric permittivity ~14000 and a very low dielectric loss (>0.05) which is very suitable for practical application were obtained in Bi doped STO ceramics with x = 5%. Fine scan XRD data and first-principle calculation results suggest that the high-performance CP behaviour of Bi doped STO ceramics maybe attribute to the hole-pinned defect-dipoles composed of VSr" and BiSr· which hopping hole charges are localized by BiSr· - VSr" - BiSr· defect complex. The leakage current result suggests that interface effect may be related to dielectric property. The present study can provide a general approach towards an appropriate compositional design for high performance colossal permittivity in STO based systems.

Acknowledgment

This work was supported by the National Natural Science Foundation of China [Grant Nos. 51677033, 51802061, 51702069].


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The immense potential of colossal permittivity (CP) materials for use in modern microelectronics as well as for high-energy-density storage applications has propelled much recent research and development. Despite the discovery of several new classes of CP materials, the development of such materials with the required high performance is still a highly challenging task. Here, we propose a new electron-pinned, defect-dipole route to ideal CP behaviour, where hopping electrons are localized by designated lattice defect states to generate giant defect-dipoles and result in high-performance CP materials. We present a concrete example, (Nb+In) co-doped TiO₂ rutile, that exhibits a largely temperature- and frequency-independent colossal permittivity (&amp;gt; 10(4)) as well as a low dielectric loss (mostly &amp;lt; 0.05) over a very broad temperature range from 80 to 450 K. A systematic defect analysis coupled with density functional theory modelling suggests that 'triangular' In₂(3+)Vo(••)Ti(3+) and 'diamond' shaped Nb₂(5+)Ti(3+)A(Ti) (A = Ti(3+)/In(3+)/Ti(4+)) defect complexes are strongly correlated, giving rise to large defect-dipole clusters containing highly localized electrons that are together responsible for the excellent CP properties observed in co-doped TiO₂. This combined experimental and theoretical work opens up a promising feasible route to the systematic development of new high-performance CP materials via defect engineering.
[20] W. Hu, K. Lau, Y. Liu, R.L. Withers, H. Chen, L. Fu, B. Gong, W. Hutchison, ACS Chem. Mater. 27(2015) 4934-4942.

[Cited within: 1]     

[21] Y. Yu, L.D. Wang, W.L. Li, Y.L. Qiao, Y. Zhao, Y. Feng, T.D. Zhang, R.X. Song, W.D. Fei, Acta Mater. 150(2018) 173-181.

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[22] Y.Q. Wu, X. Zhao, J.L. Zhang, W.B. Su, J. Liu,Appl. Phys. Lett. 107(2015), 242904.

[Cited within: 1]     

[23] T.D. Zhang, W.L. Li, Y. Zhao, Y. Yu, W.D. Fei,Adv. Funct. Mater. 28(2018), 1706211.

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[24] Y.L. Zhang, W.L. Li, Y.L. Qiao, Y. Zhao, Z.Y. Wang, Y. Yu, H.T. Xia, Z. Li, W.D. Fei,Appl. Phys. Lett. 112(2018), 093902.

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[25] W.P. Cao, W.L. Li, Y. Feng, T. Bai, Y.L. Qiao, Y.F. Hou, T.D. Zhang, Y. Yu, W.D. Fei,Appl. Phys. Lett. 108(2016), 202902.

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[26] W.P. Cao, W.L. Li, X.F. Dai, T.D. Zhang, J. Sheng, Y.F. Hou, W.D. Fei, J. Eur. Ceram. Soc. 36(2016) 593-600.

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DOI      URL      PMID      [Cited within: 1]      Abstract

Piezoelectricity is inherent only in noncentrosymmetric materials, but a piezoelectric response can also be obtained in centrosymmetric crystals if subjected to inhomogeneous deformation. This phenomenon, known as flexoelectricity, can significantly affect the functional properties of insulators, particularly thin films of high permittivity materials. We have measured strain-gradient-induced polarization in single crystals of paraelectric SrTiO3 as a function of temperature and orientation down to and below the 105 K phase transition. Estimates were obtained for all the components of the flexoelectric tensor, and calculations based on these indicate that local polarization around defects in SrTiO3 may exceed the largest ferroelectric polarizations. A sign reversal of the flexoelectric response detected below the phase transition suggests that the ferroelastic domain walls of SrTiO3 may be polar.
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[30] J. Liu, Q. Liu, Z. Nie, S. Nie, D. Lu, P. Zhu, Ceram. Int. 45(2019) 10334-10341.

DOI      URL      [Cited within: 1]     

[31] K. Tsuji, W.T. Chen, H. Guo, X.M. Chen, T.K. Lee, W.H. Lee, C.A. Randall, RSC Adv. 6(2016) 92127-92133.

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[32] L. Tong, D. Zhang, H. Wang, Q.J. Li, Y. Yu, Y.D. Li, S.G. Huang, Y.M. Guo, C.C. Wang, Mater. Lett. 180(2016) 256-259.

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[33] Z. Wang, M. Cao, Q. Zhang, H. Hao, Z. Yao, Z. Wang, Z. Song, Y. Zhang, W. Hu, H. Liu, J. Am. Ceram. Soc. 98(2015) 476-482.

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[34] Z. Wang, M. Cao, Z. Yao, Q. Zhang, Z. Song, W. Hu, Wei, Q. Xu, H.Hao, H. Liu, Z. Yu, J. Eur. Ceram. Soc. 34(2014) 1755-1760.

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[35] A. Tkach, O. Okhay, A. Almeida, P.M. Vilarinho, Acta Mater. 130(2017) 249-260.

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[36] Y. Feng, W.L. Li, D. Xu, Y.L. Qiao, Y. Yu, Y. Zhao, W.D. Fei, ACS Appl. Mater. Interfaces 8 (2016) 9231-9241.

DOI      URL      PMID      [Cited within: 1]      Abstract

The high piezoelectricity of ABO3-type lead-free piezoelectric materials can be achieved with the help of either morphotropic phase boundary (MPB) or polymorphic phase transition (PPT). Here, we propose a new defect engineering route to the excellent piezoelectric properties, in which doped smaller acceptor and donor ions substituting bivalent A-sites are utilized to bring local lattice distortion and lower symmetry. A concrete paradigm is presented, (Li-Al) codoped BaTiO3 perovskite, that exhibits a largely thermo-stable piezoelectric constant (&amp;gt;300 pC/N) and huge mechanical quality factor (&amp;gt;2000). A systematic analysis including theoretical analysis and simulation results indicates that the Li(+) and Al(3+) ions are inclined to occupy the neighboring A-sites in the lattice and constitute a defect dipole (ionic pairs). The defect dipoles possess a kind of dipole moment which tends to align directionally after thermo-electric treatment. A mechanism related to the defect symmetry principle, phase transition, and defect migration is proposed to explain the outstanding piezoelectric properties. The present study opens a new development window for excellent piezoelectricity and provides a promising route to the potential utilization of lead-free piezoelectrics in high power applications.
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[40] Y.L. Zhang, W.L. Li, S.C. Xu, Z.Y. Wang, Y. Zhao, J. Li, W.D. Fei, J. Mater. Chem. A 6 (2018) 24550-24559.

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[42] A. Tkach, A. Almeida, J. Agostinho Moreira, T.M. Correia, M.R. Chaves, O. Okhay, P.M. Vilarinho, I. Gregora, J. Petzelt,Appl. Phys. Lett. 98(2011), 052903.

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