Superior Energy Density Achieved in Unfilled Tungsten Bronze Ferroelectrics via Multiscale Regulation Strategy

Abstract The most promising candidates for energy storage capacitor application are relaxor ferroelectrics, among which, the perovskite structure ferroelectric ceramics have witnessed great development progress. However, less attention has been paid on tetragonal tungsten bronze structure (TTBS) ceramics because of their lower breakdown strength and polarization. Herein, a multiscale regulation strategy is proposed to tune the energy storage performances (ESP) of TTBS ceramics from grain, domain, and macroscopic scale. The enhanced relaxor behavior with dynamic polar nanodomains guarantees low remanent polarization, while the refined grains and enlarged bandgap ensure increased breakdown strength. Hence, excellent ESP is realized in unfilled TTBS Sr0.425La0.1□0.05Ba0.425Nb1.4Ta0.6O6 (SLBNT) ceramics with an ultrahigh recoverable energy density of 5.895 J cm−3 and a high efficiency of 85.37%. This achievement notably surpasses previous studies in TTBS ceramics and is comparable to that of perovskite components. Meanwhile, the energy density exhibits a wide temperature, frequency, and cycling fatigue stability. In addition, high power density (257.89 MW cm−3), especially the ultrafast discharge time (t 0.9 = 16.4 ns) are achieved. The multiscale regulation strategy unlocks the energy storage potential of TTBS ceramics and thus highlights TTBS ceramics as promising candidates for energy storage, like perovskite structured ceramics.


Introduction
Dielectric capacitors have received increasing attention as the heart of pulse power systems due to their fast charge-discharge capability and high power density. Nonetheless, dielectric capacitors take up over 25% of the volume and weight of pulse power systems, [1] which contradicts the development trend toward miniaturization and integration for modern electrical and electronic devices. Consequently, achieving high energy storage properties in dielectric materials is identified as the major obstacle to the development of high-end pulse power capacitors. And the energy storage performance (ESP) can be determined by integrating the polarization-electric field (P-E) hysteresis loops: W rec = ∫ P m P r EdP, × 100%, where W rec , W total , , P m , P r , and E represent recoverable energy storage density, total energy storage density, energy storage efficiency, maximum polarization, remanent polarization, as well as the applied electric field, correspondingly. [2,3] Hence, a high P m , low P r , and a high dielectric breakdown strength (BDS) are critical factors for achieving high W rec and . In the meantime, rapid response to an applied electric field, delayed saturation polarization, uniform www.advancedsciencenews.com www.advancedscience.com and dense microstructure, refined grains, and a large bandgap for carrier transfer are required in high-performance dielectric ceramics. [4] Dielectric ceramics can be divided into four typical types: linear dielectric materials (LD), normal ferroelectrics (FEs), antiferroelectric (AFEs), and relaxor ferroelectrics (RFEs). [5] RFEs with high dynamic and weak coupling polar nanoregions (PNRs) seem to be the optimal material for energy storage. And the weakly coupled PNRs are electric field-sensitive and respond quickly to the variation of external electric field, thus gaining a high P m and a negligible P r , [6,7] which will result in a large ΔP (P m -P r ) and ensure the achievements of great W rec and in RFEs. Over the past decade, multiple strategies have been investigated to strengthen the ESP of perovskite structure RFEs, including nanodomain design/engineering, [5,8,9] superparaelectric design, [10][11][12] and defect engineering [13] et al. Hence, high W rec (typically higher than 4 J cm −3 ) and (no less than 80%) are achieved in all types of perovskite components, such as BaTiO 3based (0.6BaTiO 3 -0.4Bi(Mg 0.5 Ti 0.5 )O 3 , 4.49 J cm −3 and 93%), [14] NaNbO 3 -based (0.68NaNbO 3 -0.32(Bi 0.5 Li 0.5 )TiO 3 , 8.73 J cm −3 and 80.1%), [15] Bi 0.5 Na 0.5 TiO 3 -based (0.65(0.84Bi 0.5 Na 0.5 TiO 3 -0.16K 0.5 Bi 0.5 TiO 3 )-0.35(Bi 0.2 Sr 0.7 TiO 3 ), 4.06 J cm −3 and 87.3%), [16] BiFeO 3 -based (0.57BiFeO 3 -0.33BaTiO 3 -0.1NaNbO 3 , 8.12 J cm −3 and 90%), [5] and K 0.5 Na 0.5 NbO 3 -based (0.85K 0.5 Na 0.5 NbO 3 -0.15Bi(Ni 0.5 Zr 0.5 )O 3 , 8.09 J cm −3 and 88.46%) ceramics. [17] Tetragonal tungsten bronze structure (TTBS) ferroelectric ceramics are well-known as the second largest family of ferroelectrics after the perovskite ferroelectrics. However, TTBS ceramics have received much less attention for energy storage study due to their low saturated polarization and poor BDS related to abnormal grain growth. [18] The connections between the BO 6 octahedron's common vertices form the framework of TTBS ceramics and three types of interstices: square A1, pentagonal A2, and triangular C sites. [19] And subsequently, TTBS ceramics are categorized into three groups on the basis of the ion occupancy situation: fully filled (A and C sites are all occupied), filled (all A sites are occupied, C sites are empty), and unfilled (A sites are partially occupied, C sites are empty). [20,21] Some investigations into TTBS ceramics for energy storage have been conducted, but the W rec is frequently unsatisfactory (less than 4 J cm −3 ). For instance, Sr 0.5 Na 0.5 Nb 2 O 5 -based ((Sr 2 NaNb 3.5 Ta 1.5 O 6 (W rec = 3.99 J cm −3 and = 91.7%)), [22] Sr 0.5 K 0.5 Nb 2 O 5 -based (Sr 2 KNb 4.2 Ta 0.8 O 15 (W rec = 3.84 J cm −3 and = 93.2%)), [23] and Sr 0.5 Ba 5 Nb 2 O 6 -based (0.7Sr 0.75 Ba 0.25 Nb 2 O 6 -0.3PbZr 0.52 Ti 0.48 O 3 (W rec = 3.0 J cm −3 and = 81.5%) [24] ). Prior research has established that the ion doping modulated regularity of ferroelectric and relaxor behavior in TTBS ceramics, which is primarily a result of cation occupancy and structural distortion. [25][26][27][28][29][30][31] In accordance with the crystalchemical framework, a small A-site average ion radius causes a reduction in O-B-O bond length and benefits the improvement of relaxor properties. [32] Meanwhile, the weakened B-O bonding, local structural fluctuation or distortion, and order-disorder distribution all support the modification of the B-site cation to improve relaxor behavior. [29,33,34] Additionally, the introduction of A-site vacancies can induce extra changes in the local coordination environment, structure relaxation, and charge/structure disorder. [26,[35][36][37] Consequently, various methods can be adopted to adjust the relaxor behaviors of TTBS ceramics, and their flexible structures and abundant properties tunability provide a suitable realm for the effective design of high energy storage performance. [38] Hence, the TTBS ceramic family merits more attention and investigation as an energy storage material. The innovative design of high ESP TTBS ceramics would assist in further comprehending the regulation mechanism between composition, structure, and properties, and may open a new window for future development of high performance energy storage materials.
In this work, we proposed a multiscale regulation strategy to improve the energy storage performance of Sr 0.5 Ba 0.5 Nb 2 O 5 unfilled tungsten bronze structure ceramics (Figure 1). Via La 3+ and Ta 5+ doping, the designed Sr 0.425 La 0.1 □ 0.05 Ba 0.425 Nb 1.4 Ta 0.6 O 6 (SLBNT) ceramics achieved enhanced relaxor behavior by destroying the ferroelectric long-range order and inducing polar nanoregions at domain scale. Extra vacancies were also introduced to satisfy electrical neutrality and further facilitate structural relaxation and disorder. In addition, the substitution of large bandgap oxides Ta 2 O 5 (≈4 eV) and La 2 O 3 (≈4.3 eV) can enlarge the bandgap. Meanwhile, the strong refractory nature of Ta 5+ can significantly inhibit grain growth and reduce the average grain size. The refined grains and enlarged bandgap from the grain scale and macroscopic scale, guarantee together an increased E b . Ultimately, the SLBNT ceramics demonstrated excellent energy storage performance (W rec = 5.895 J cm −3 and = 85.37%), which was superior to all previous reports on TTBS ceramics. And the high power density (257.89 MW cm −3 ) and ultrafast discharge time (t 0.9 = 16.4 ns) further incorporated the potential of SLBNT ceramics for application. This work developed a novel lead-free preferential material for pulse power applications and highlighted the research prospects of TTBS ceramics for energy storage. With the variation in composition, the P m decreases monotonously, as also can be seen in Figure 2d and the inset of Figure 2a, which demonstrates the composition dependent P-E loops at fixed E of 300 kV cm −1 . The BDS of the samples, however, increases evidently, from ≈300 kV cm −1 for SBN to ≈400 kV cm −1 for SLBN, and ultimately reaches ≈574 kV cm −1 for SLBNT ( Figure 2d) with the introduction of La 3+ and Ta 5+ . Results from the codoping strategy, a steadily enhanced relaxor character with slimmer P-E loops and lower P r has also been realized, as revealed from the E dependent P-E loops in Figure 2b. The ESPs of three compositions are shown in Figure 2d. SBN ceramics present a poor ESP (W rec = 2.592 J cm −3 and = 76.18%) due to the comparatively large P r and low E b . For SLBN, despite the fact that an enhancement of 33.01% was realized in E b , the W rec is nevertheless lower than 4 J cm −3 . Finally, the designed SLBNT ceramics achieve an ultrahigh E b of ≈574 kV cm −1 with a further increase of 43.5%. As a result, W rec increases to 5.895 J cm −3 with an enhancement of 50.69%, and is also improved to 85.37%. The achievement of optimized composition of Sr 0.425 La 0.1 □ 0.05 Ba 0.425 Nb 1.4 Ta 0.6 O 6 can also be supported by a systematical investigation on the   [17,18,[22][23][24][27][28][29][38][39][40][41][42][43][44][45][46][47][48][49][50][51][52][53] Adv. Sci. 2023, 10, 2300227 Figure 3. a) XRD patterns of the SBN, SLBN, and SLBNT ceramics at room temperature. b) Enlarged views of (311), (002), and (620) peak. c) Raman spectra of SBN, SLBN, and SLBNT ceramics at room temperature. d) Raman shift of 2 and 5 modes and e) FWHM of 2 mode as a function of composition. SEM images and grain size distributions of f) SBN, g) SLBN, and h) SLBNT ceramics. i) Impedance spectra of SBN, SLBN, and SLBNT ceramics at 500°C. j) The optical dielectric constant as a function of phonon energy. k) the experimental ellipsometric (dots) and the best-fitting (solid lines) spectra of SLBNT ceramics (the inset shows the E g of all components). structure and properties of Sr 0.425 La 0.1 □ 0.05 Ba 0.425 Nb 2−x Ta x O 6 ceramics with different Ta 5+ concentrations, which is summarized in Figures S1-S4 (Supporting Information). A comparison of the ESP of our SLBNT ceramics and previously reported lead-free bulk ceramics are illustrated in Figure 2e. It is evident that the excellent overall performance of SLBNT ceramics leads the field in tetragonal tungsten bronze structure ferroelectric ceramics and can even be comparable with some state-of-the-art perovskite structure ceramics. [17,18,[22][23][24][27][28][29][38][39][40][41][42][43][44][45][46][47][48][49][50][51][52][53] Figure 3a shows the X-ray diffractometer (XRD) patterns of SBN, SLBN, and SLBNT ceramics. All the patterns reveal pure tungsten bronze structure, indicating that the La 3+ and Ta 5+ successfully entered the SBN host lattice. The magnified patterns around 32°and 46°are presented in Figure 3b. Due to the smaller ionic radius and higher electronegativity, La 3+ (1.36 Å (12 coordination), 1.1) occupies the A1 site like Sr 2+ (1.44 Å (12 coordination), 0.95). As a direct consequence, the unit cell shrinks, and the lattice parameters decline as a result of decreasing ion radius, let alone the formation of extra A-site vacancies due to alio-valent ion substitution. The impact of extra vacancy is concluded as an important local distortion: a Nb/Ta cube around a vacancy is compressed in c axis, which means a larger tilts of octahedra. [54] Accordingly, the diffraction peaks slightly shift to higher 2 values from SBN to SLBN. On the contrary, the diffraction peaks of SLBNT show a pronounced shift toward lower 2 value although Ta 5+ (0.64 Å (12 coordination)) and Nb 5+ (0.64 Å (12 coordination)) have same ion radii. This shift in diffraction peaks could be attributed to the difference in electronegativity between Ta 5+ (1.51) and Nb 5+ (1.59), which leads to a weaker covalent bond. [33] According to the second-order Jahn-Teller effect associated with the electron orbital hybridization between B-site cations and oxygen, the ferroelectricity of TTB originates from the displacement of the B-site ion along the center of the octahedron. [25] While the different electron configurations of Ta 5+ can directly change the displacement of polar cations, resulting in the tilting of the oxygen octahedron and enhanced relaxor behaviors. Moreover, the diffraction peak around 46°steadily split into two peaks (the lowangle (002) peak and the high-angle (620) peak). This splitting www.advancedsciencenews.com www.advancedscience.com of diffraction peaks is ascribed to the reduced tetragonality (c/a). The more accurate Rietveld refinement is illustrated in Figure  S5 (Supporting Information), and the crystallographic data are summarized in Table S1 (Supporting Information). It shows that all compositions belong to the tetragonal crystal system with a space group of P4bm, and the tetragonality (c/a) decreases with the ions doping. Meanwhile, compression of structural deformation along the polar c axis (a decreasing lattice parameter c) points to the suppressed ferroelectric nature, which means that the decreased tetragonality would be related to the relaxor behavior. The decreased tetragonality disturbs the displacement of the B-site ferroelectric active ion along the center of the octahedron, resulting in strong relaxation of SLBNT. [25,27,36] The local structures of SBN, SLBN, and SLBNT ceramics were also examined by using Raman spectroscopy.  [49] Smaller-ion-radius La 3+ can increase stretching vibrations while lessening the interaction between metal ions and oxygen. [55] Moreover, the addition of Ta 5+ can also weaken the strong interaction because of its different electronic configuration from Nb 5+ . [52] Consequently, the 2 and 5 modes exhibit a redshift attributable to the weakened covalency of the B-O bond and the expansion of the BO 6 octahedron, as displayed in Figure 3d. [51] The increased full width at half maximum (FWHM) of 2 mode in Figure 3e also indicates the enhanced disorder in short-range structure and the increasing distortion degree of the BO 6 octahedron. [49] In a word, these variations and distortions of the BO 6 polar unit are affected by La 3+ and Ta 5+ in different ways, whereas, both contribute to local polarization and enhanced relaxor characteristic.

Results and Discussion
The microstructures of SBN, SLBNT, and SLBNT ceramics are displayed in Figure 3f-h, and the insets show the grain size distribution. SBN presents an inhomogeneous microstructure with anisometric and columnar-shaped grain, which is caused by the faster growth rate along the (001) facet. [34] The amount of columnar-shaped grain decreases noticeably in SLBN and then disappears in SLBNT. The average grain size (G a ) decreases from 3.02 μm (SBN) to 2.07 μm (SLBN) and eventually reduces to 1.61 μm (SLBNT). The refined grains can be ascribed to the reduction of sintering temperature from 1370°C for SBN to 1230 and 1300°C for SLBN and SLBNT, respectively, which might originate from the lattice activation after La 3+ doping. [5] Moreover, the refractory nature of Ta 5+ can significantly inhibit grain growth, which finally reduces the average grain size and increases the proportion of equiaxed grains for SLBNT. [22,23] In accordance with the relationship between G a and E b , described as E B ∝(G a ) −a , [56] the grain boundaries contain depletion layers having higher barriers for charge delivery, [57] thus reducing the overall ceramics' conductivity. Moreover, as seen in Figure 3i; and Figure S6 (Supporting Information), the resistance increases with the decrease of G a from SBN to SLBN, and SLBNT. The impedance spectra with two semicircles represents the impedance response of grains and grain boundaries, respectively. The G a of SLBN and SLBNT decrease obviously, resulting in more depletion regions and higher grain boundary resistivity. It means that the contribution of grain boundary to resistivity is more evident. Consequently, the SLBN and SLBNT show only one semicircle in impedance spectra. [1,58] Figure 3j presents the dielectric functions ( = 1 + 2 ) measured by spectroscopic ellipsometry. The experimental and simulation spectra match well through the Tauc-Lorentz (TL) model as shown in Figure 3k. [59] With the substitution of large bandgap oxide Ta 2 O 5 (≈4 eV) and La 2 O 3 (≈4.3 eV), the calculated E g increased from 3.26 eV of SBN to 3.28 eV of SLBN, and 3.39 eV of SLBNT. The large bandgap makes it more difficult for the electrons to transmit from the valence band to the conduction band, reducing the probability of electron ionization and collision during the breakdown process, which can further boost the BDS. [52,[60][61][62] To sum up, the refined grains and enlarged band gap demonstrate the effectiveness of our regulation strategy of ESP at the grain scale and macroscopic scale.
The temperature dependent dielectric permittivity ( m ) and dielectric loss (tan ) are shown in Figure 4a-c over the range of −120 to 200°C. In distinction to SBN's sharp and frequencyinsensitive dielectric peak, as frequency increases, the dielectric peaks of SLBN and SLBNT are broad and decrease with increasing temperature. Meanwhile, the decrement of T m to below 0°C indicates the appearance of PNRs at room temperature. Consequently, SLBN and SLBNT ceramics present slimmer P-E loops than pure SBN ceramics. The introduction of La 3+ and Ta 5+ causes ionic disorder and compositional fluctuation, destroying the long-range-ordered ferroelectric microdomains and decreasing the size of domains. [21,37,63] Besides, extra structural vacancies at the A-site further weaken the coupling of PNRs, thereby influencing the relaxor behavior. [51] Additionally, the enhanced relaxor behavior is also quantitatively estimated by the calculated diffuseness degree ( ) in Figure S7  , where m is the maximum permittivity and C is a constant). [64] The value of rises dramatically and monotonously from SBN (1.   the PFM signals fade faster with a relaxation time of 3 min. In contrast, the poled region and piezoresponse of SLBNT fade very fast (1 min) owing to the high dynamic PNRs, which is consistent with typical relaxor behavior. [65] In summary, the decreased domain size favors easier polarization reorientation after the external electric field removal and thus benefits for smaller P r and slimmer P-E loops.
The temperature stability of energy storage properties plays a vital role in practical applications. Figure 5a displays the P-E loops of SLBNT ceramics over a wide temperature range from −120 to 120°C at 350 kV cm −1 and 10 Hz. the P-E loops of SLBNT ceramics remain slim as the temperature increases. Furthermore, the P m and P r exhibit a small variation of 2.969 and 1.085 μC cm −2 , correspondingly. In consequence, the W rec shows a slight decline from 2.960 J cm −3 at −120°C to 2.405 J cm −3 at 120°C with a variation of less than ±10.76%, and the decreases from 87.15% to 77.55%, indicating the relatively satisfying temperature stability. Local structure changes under temperature variation in SLBNT ceramics are revealed using temperaturedependent Raman spectroscopy. As seen in Figure 5c, the num-ber of Raman peaks is unaffected by changes in temperature, indicating that SLBNT ceramics have a stable phase structure. [66] As the temperature rises from −140 to 300°C, the peak intensity of external vibration modes (associated with the motions of the A site cation) below 200 cm −1 gradually increases. The enhanced motion of A site ions has a huge impact on the tilting of the BO 6 octahedron, which is related to the displacement of the B-site ion along the c axis. [55,67] As shown in Figure 5d, the 2 mode, which originates from the O-B-O bending vibration of the BO 6 octahedron, exhibits negligible red or blueshifts and only fluctuates slightly around ≈608 cm −1 . In addition, its peak intensity exhibits a decreasing tendency and its full width at half maximum (FWHM) gradually increases throughout the test temperature range (Figure 5e). These structural features prove suppressed polarity and a higher disorder degree in local structure, [15,29,68] which further confirms the existence of nanodomains in a wide T range. Figure 5f illustrates the temperature dependent XRD patterns of SLBNT ceramics in the range of −160-290°C. Invariant diffraction peaks are evident which supports the stable structure of SLBNT in the whole T range. The magnified parts around 46°are also given in Figure 5g. The splitting (002) and (620) peaks display no detectable changes from −160 to 290°C, suggesting the invariant P4bm symmetry upon heating and cooling. The temperature stable structure, as shown by the above Raman and XRD analyses, guarantees the reliability of SLBNT to work in harsh environments under high temperatures or belowzero temperatures.
The P-E loops for SLBNT ceramics are also evaluated at 350 kV cm −1 and room temperature under a wide range of cycle numbers and frequencies to further assess the stability of its ESP. Figure  6a,b; and Figure S10(a) (Supporting Information) provide the P-E loops, W rec , , and the corresponding P m , P r , ΔP after 10 6 cycles. The P m and P r slightly decrease at first and subsequently increase with the accumulation of cycles. And the W rec maintains its current value of about 2.630 J cm −3 , accompanied by a negligible change (±0.64%). As displayed in Figure S10(b,c) (Supporting Information); and Figure 6c, the SLBNT ceramics preserve slim P-E loops with low hysteresis loss in the range 10-250 Hz, indicating high frequency stability. The variations of P m and P r are as small as 0.45 and 0.33 μC cm −2 , respectively. The W rec only slightly decreases from 2.657 J cm −3 at 10 Hz to 2.653 J cm −3 at 250 Hz, exhibiting a variation of less than ±0.56%.
The charge-discharge tests can be used to determine the discharge properties of SLBNT ceramics for practical applications. The overdamped discharge current curves at room temperature are displayed in Figure 6d. The discharge energy density (W dis ) is calculated from the equation: . [45] The integral curve is plotted in Figure 6e with a W dis of 2.061 J cm −3 and an ultrafast t 0.9 of 16.4 ns. Figure 6f presents the underdamped discharge waveforms of SLBNT ceramics at room temperature. The value of current peaks increases linearly as electric fields increase linearly. The current density (C D ) and power density (P D ) were calculated according to following equations: C D = I max /S, P D = EI max /2S. [69] As shown in Figure 6g, the C D exhibits the same increasing tendency as I max , reaching the maximum values of 1473.68 A cm −2 and 28.32 A at 350 kV cm −1 . And the P D increases from 22.26 to 257.89 MW cm −3 . Figure 6h,i examines the temperature stability of the discharge properties at 200 kV cm −1 from 20 to 120°C. The discharge curve remains unchanged, with a slight reduction in peak current. Thus, the calculated W dis only exhibits a slight decrease from 0.667 to 0.551 J cm −3 . Besides, t 0.9 is measured to be shorter than 26.4 ns in the whole test temperature range, and the slight increase of t 0.9 is attributed that the reaction speed of PNRs to electric fields is improved by the enhanced random electric field. The excellent pulse discharging properties under different electric fields and temperatures, as well as the outstanding stability, further emphasize the promise of SLBNT for practical pulse power capacitor applications.

Conclusion
In conclusion, we proposed a multiscale regulation strategy for attaining excellent energy storage performance in unfilled TTBS relaxor ferroelectrics. Benefiting from the induced polar nanoregions, refined grains, as well as enlarged bandgap via La 3+ and Ta 5+ doping, the Sr 0.425 La 0.1 □ 0.05 Ba 0.425 Nb 1.4 Ta 0.6 O 6 ceramics simultaneously realize a record-high W rec among TTBS ceramics of 5.895 J cm −3 and a high of 85.37% accompanied with an ultrahigh E b of ≈574 kV cm −1 . Meanwhile, excellent frequency stability (10-250 Hz), cycling fatigue stability (up to 10 6 times), wide temperature stability (−120-120°C), high power density (P D = 257.89 MW cm −3 ), and ultrafast discharge rate (t 0.9 = 16.4 ns) are likewise demonstrated in the designed ceramics. The overall energy storage performance can also be comparable with the extensively-studied perovskite structured ferroelectric materials. This work not only provides a promising candidate for highperformance pulse power capacitors but also underlines the energy storage potential of tetragonal tungsten bronze structure relaxor ferroelectric materials. After uniform mixing via ball milling for 4 h with ethyl alcohol and zirconium balls, these slurries were dried and calcined at 1200°C for 3 h in an air atmosphere. Then, the powders were ball milled again with smaller zirconium balls (1 mm in diameter) to obtain smaller particle sizes. After that, these slurries were dried, mixed with solution of poly(vinyl alcohol) (PVA), and pressed into disks with a diameter of 13 mm. Finally, after removing the binder at 800°C for 2 h, the disks were sintered at 1230-1370°C for 4 h.

Experimental Section
Structural Characterization: XRD (Bruker, D8 ADVANCE for variable temperature tests and D2 PHASER for room temperature tests) with Cu Ka radiation ( = 1.5406 Å) and Raman spectroscopy (Renishaw inVia reflex) were used to determine the crystal structure. The microstructure was recorded by a field emission scanning electron microscope (FE-SEM, S-4800, Hitachi, Tokyo, Japan). Spectroscopic ellipsometry (SE) measurements was used to measure the optical dielectric constant in the photon energy range of 1.24-4.96 eV (1000-250 nm) (V-VASE by J. A. Woollam Co., Inc.), and a three-layer structure model (air, rough layer, and ceramics) is used to fit the spectra. The response of domains were studied by a commercial AFM systen (Jupiter XR, Oxford, UK). The selected area electron diffraction and domain morphology observation were performed using a JEM-2100F filed-emission transmission electron microscope (FE-TEM, JEOL, Japan).
Properties Measurement: The P-E loops were measured by a ferroelectric measuring system (aixACCT TF Analyzer 2000E) with the sample size of 0.08 mm (thickness) × 0.785 mm 2 (Au electrode area). The temperature-and frequency-dependent dielectric properties and complex impedance were measured by a broad frequency/temperature dielectric spectrometer (Novocontrol GmbH, Concept 80). The charge-discharge test under overdamped (with a load resistance of 100 Ω) and underdamped conditions were measured by a commercial charge-discharge platform (CFD-001, Gogo Instruments Technology, Shanghai, China). The samples for the charge-discharge test were 0.1 mm in thickness with an Au electrode of 1.5 mm in diameter.

Supporting Information
Supporting Information is available from the Wiley Online Library or from the author.