Ferroelectric and Relaxor-Ferroelectric Phases Coexisting Boosts Energy Storage Performance in (Bi0.5Na0.5)TiO3-Based Ceramics

With the intensification of the energy crisis, it is urgent to vigorously develop new environment-friendly energy storage materials. In this work, coexisting ferroelectric and relaxor-ferroelectric phases at a nanoscale were constructed in Sr(Zn1/3Nb2/3)O3 (SZN)-modified (Bi0.5Na0.5)0.94Ba0.06TiO3 (BNBT) ceramics, simultaneously contributing to large polarization and breakdown electric field and giving rise to a superior energy storage performance. Herein, a high recoverable energy density (Wrec) of 5.0 J/cm3 with a conversion efficiency of 82% at 370 kV/cm, a practical discharged energy density (Wd) of 1.74 J/cm3 at 230 kV/cm, a large power density (PD) of 157.84 MW/cm3, and an ultrafast discharge speed (t0.9) of 40 ns were achieved in the 0.85BNBT-0.15SZN ceramics characterized by the coexistence of a rhombohedral-tetragonal phase (ferroelectric state) and a pseudo-cubic phase (relaxor-ferroelectric state). Furthermore, the 0.85BNBT-0.15SZN ceramics also exhibited excellent temperature stability (25–120 °C) and cycling stability (104 cycles) of their energy storage properties. These results demonstrate the great application potential of 0.85BNBT-0.15SZN ceramics in capacitive pulse energy storage devices.


Introduction
With the growing demand for advanced electrical power systems, dielectric capacitors, as essential elements for electrostatic energy storage, play a decisive role in high-power applications and pulsed power technologies, owing to their unique advantages of fast charge-and-discharge speed, ultrahigh power density, and excellent stability and reliability [1][2][3].Compared to polymer-based capacitors [4], the ceramic-based capacitors possess higher capacitance and better temperature stability, and thus, have received increasing attention both in academic research and commercial applications.However, the relatively low recoverable energy density (W rec ) caused by the low breakdown electric field (E b ) in dielectric ceramics limits their energy storage applications, for which device miniaturization and system intellectualization are necessary [5,6].Therefore, a great deal of research has been carried out to explore high-performance dielectric ceramics with high W rec and energy efficiency (η) as well as excellent temperature and cycling stability.In general, the total energy density (W tot ), the recoverable energy density (W rec ), and the efficiency (η) of ceramic-based dielectric capacitors can be obtained by the Equations ( 1)-(3) [7,8]: Molecules 2024, 29 where P max and P r present the maximum polarization and remanent polarization, respectively; E is the applied external electric field.Equations ( 1)- (3) indicate that a high W rec can be achieved through improving P max , reducing P r , and simultaneously enhancing E b .Different from linear dielectrics (LDs) and anti-ferroelectrics (AFEs), relaxor ferroelectrics (RFEs) have the combined advantages of a high difference (∆P = P max − P r ) and a relatively large breakdown electric field (E b ), and thus, they possess a pronounced energy storage performance [7,[9][10][11][12][13][14].It is recognized that this pronounced energy storage performance for RFEs is generally strongly related to polar nanoregions (PNRs), which are usually driven by random fields caused by composition disorder disrupting ranged ferroelectric domains [15][16][17][18].In the nonergodic relaxor (NR) state, the PNRs freeze and the ferroelectric domains with micro-scale generate strong hysteresis, a large P r and inferior energy storage performance.By increasing the content of the composition and/or temperatures, the power and size of the PNRs increases considerably and decreases rapidly, respectively, and quick merging is achieved, resulting in a large reduction in P r [19].In the dominated ergodic relaxor (ER) state, PNRs with random dipole orientation are easily deflected and ergodic, and the polarization direction and intensity of them can alter with the application of an external electric field [20,21].Despite the transformation from the RFE to the ferroelectric state with the help of random fields, the long-range ferroelectric domains are disrupted, thus producing disordered PNRs with the removal of electric field [22,23]; that is, P r is significantly suppressed and P max remains relatively high.Accordingly, tailoring RFEs to the dominated ER state using PNRs that generate a small P r , a large P max , reduced hysteresis, and enhanced thermal stability is the main collaborative optimization strategy (domain engineering), and enhances the energy storage performance, as shown by the slim polarization versus electric field (P (E)) loops with a significantly delayed polarization saturation.Moreover, a higher electric field can drive a larger P max , which supports the idea that improving E b by multiple effects is crucial for achieving a high energy density of dielectrics.As a result, simultaneous optimization of the external (sample thickness, electrode shape/size, etc.) and internal (e.g., grain size, defects, and grain orientation) factors has been performed to enhance the E b [24,25].For instance, a large E b value of 420 kV/cm was obtained in BiFeO 3 -BaTiO 3 -NaNbO 3 ceramics via reducing the grain size [25].Hence, tailoring the grain size, which aims to enhance E b , was also adopted as a collaborative optimization strategy (i.e. grain size engineering).The decreased grain size means the increased fraction of the grain boundary, which can increase the depletion space charge layers and hence, offers higher potential barriers for charge carriers [26].
It is well known the W rec and η of the typical ferroelectric system (Bi 0.5 Na 0.5 ) 0.94 Ba 0.06 TiO 3 are too small to satisfy the requirements of the system's practical application, as a result of premature polarization saturation and a large P r .In this work, Sr(Zn 1/3 Nb 2/3 )O 3 , selected as an endmember, was introduced to (Bi 0.5 Na 0.5 ) 0.94 Ba 0.06 TiO 3 to partially reduce the ferroelectricity and enhance the relaxation behavior, and eventually, the coexisting ferroelectric and relaxor-ferroelectric phases were constructed, which gave rise to superior energy storage properties (W rec = 5.0 J/cm 3 , η = 82%) in the 0.85BNBT-0.15SZNceramics.

Results and Discussion
Figure 1a displays the XRD patterns of (1 − x)BNBT-xSZN (x = 0-0.18)ceramics with different doping concentrations of SZN.All samples exhibit a typical perovskite structure, and not any impurity phase was detected, indicating that the SZN have completely diffused into the BNBT matrix crystal lattice and formed new solid solutions.The enlarged XRD patterns in the 2θ range of 46-47 • are depicted in Figure 1b.It is evident that the (200) diffraction peaks gradually move to lower degrees with increasing the SZN content, indicating a lattice expansion, which is due to the substitutions of Sr 2+ with a relative large ionic radius (R Sr 2+ = 1.44 Å) for Na + (R Na + = 1.39 Å) and Bi 3+ (R Bi 3+ = 1.38 Å) on the A-sites, and the substitutions of Zn 2+ (R Zn 2+ = 0.74 Å) and Nb 5+ (R Nb 5+ = 0.64 Å) for Ti 4+ (R Ti 4+ = 0.605 Å) on the B-sites [27].The rietveld refinements of the XRD patterns for (1 − x)BNBT-xSZN ceramics, and the TEM selected area electron diffraction (SAED) for representative 0.85BNBT-0.15SZNceramics are given in Figure 2 to further identify the phase structures.The low values of all the reliability factor of patterns (R p < 7%), the reliability factor of weighted patterns (R wp < 8%), and the goodness-of-fit indicator (χ 2 < 2%), as shown in Figure 2a-e, indicate that the structural model is valid and the refinement results are consistent with the experimental data.Therefore, we can infer that the crystal structures evolve from R and T phases for the ceramics with x = 0 to coexisting R, T, and C phases for the ceramics with x = 0.08-0.The 1/2(ooe) and 1/2(ooo) superlattice spots (where o and e stand for odd and even Miller indices, respectively) were observed in the sample, demonstrating the existence of T and R phases in a pseudo-cubic matrix, which is consistent with the results of the Rietveld refinement of XRD pattern (Figure 2d).It should be pointed out that these superlattice diffractions are not observed in the XRD patterns, as these superlattice diffractions are weak.Combined with the results of XRD refinements and SAED patterns, we can infer that R, T, and pseudo-cubic phases coexist in the x = 0.15 sample.It is worth noting that the desired coexistence of a ferroelectric state (R and T phases) and a relaxor-ferroelectric state (pseudo-cubic phase) constructed by introducing SZN into BNBT ceramics is conducive to enhancing the P max and E b values simultaneously, and thus, results in an excellent energy storage density.

Results and Discussion
Figure 1a displays the XRD patterns of (1 − x)BNBT-xSZN (x = 0-0.18)ceramics with different doping concentrations of SZN.All samples exhibit a typical perovskite structure, and not any impurity phase was detected, indicating that the SZN have completely diffused into the BNBT matrix crystal lattice and formed new solid solutions.The enlarged XRD patterns in the 2θ range of 46-47° are depicted in Figure 1b.It is evident that the (200) diffraction peaks gradually move to lower degrees with increasing the SZN content, indicating a lattice expansion, which is due to the substitutions of Sr 2+ with a relative large ionic radius (RSr 2+ = 1.44 Å) for Na + (RNa + = 1.39 Å) and Bi 3+ (RBi 3+ = 1.38 Å) on the A-sites, and the substitutions of Zn 2+ (RZn 2+ = 0.74 Å) and Nb 5+ (RNb 5+ = 0.64 Å) for Ti 4+ (RTi 4+ = 0.605 Å) on the B-sites [27].The rietveld refinements of the XRD patterns for (1 − x)BNBT-xSZN ceramics, and the TEM selected area electron diffraction (SAED) for representative 0.85BNBT-0.15SZNceramics are given in Figure 2 to further identify the phase structures.The low values of all the reliability factor of patterns (Rp < 7%), the reliability factor of weighted patterns (Rwp < 8%), and the goodness-of-fit indicator (χ 2 < 2%), as shown in Figure 2a-e, indicate that the structural model is valid and the refinement results are consistent with the experimental data.Therefore, we can infer that the crystal structures evolve from R and T phases for the ceramics with x = 0 to coexisting R, T, and C phases for the ceramics with x = 0.08-0.15,and eventually, to the C phase for the ceramics with x = 0.18.The detailed phase volume fractions of the (1 − x)BNBT-xSZN (x = 0-0.18)ceramics are displayed in the insets of Figure 2a-e.Figure 2f and Figure 2g show the SAED patterns along [001]c and [110]c directions, respectively, for the representative sample x = 0.15.The 1/2(ooe) and 1/2(ooo) superlattice spots (where o and e stand for odd and even Miller indices, respectively) were observed in the sample, demonstrating the existence of T and R phases in a pseudo-cubic matrix, which is consistent with the results of the Rietveld refinement of XRD pattern (Figure 2d).It should be pointed out that these superlattice diffractions are not observed in the XRD patterns, as these superlattice diffractions are weak.Combined with the results of XRD refinements and SAED patterns, we can infer that R, T, and pseudo-cubic phases coexist in the x = 0.15 sample.It is worth noting that the desired coexistence of a ferroelectric state (R and T phases) and a relaxor-ferroelectric state (pseudo-cubic phase) constructed by introducing SZN into BNBT ceramics is conducive to enhancing the Pmax and Eb values simultaneously, and thus, results in an excellent energy storage density.The SEM images of surface morphologies and the corresponding grain size distributions of the (1 − x)BNBT-xSZN ceramics are shown in Figure 3a-e.It can be seen that all the samples present a compact microstructure and clear grain boundary, suggesting good sintering behavior.The average grain size and relative density as a function of SZN doping content are presented in Figure 3f.It is evident that the grain size first increases slightly and then decreases with increasing SZN content.The slight increase in grain size for the ceramics x ≤ 0.10 may be related to the fact that a trace amount of SZN dopant acts as a sintering assistant and raises the kinetic energy of grain growth.However, the consistent reduction of grain size with further increase of SZN doping content (x > 0.10) can be explained as follows.The substitution of larger Sr 2+ ions for smaller Na + and Bi 3+ on the A-sites, and the large Zn 2+ and Nb 5+ ions for smaller Ti 4+ ions on the B-sites increases the lattice strain energy and thus, the grain boundary mobility is inhibited [28].It is worth noting that the relative density of all samples is larger than 97.5%, demonstrating a dense microstructure.It is well known that a dense microstructure and refined grains can substantially increase the breakdown electric field (E b ) [29,30].3f.It is evident that the grain size first increases slightly and then decreases with increasing SZN content.The slight increase in grain size for the ceramics x ≤ 0.10 may be related to the fact that a trace amount of SZN dopant acts as a sintering assistant and raises the kinetic energy of grain growth.However, the consistent sites, and the large Zn 2+ and Nb 5+ ions for smaller Ti 4+ ions on the B-sites increases the lattice strain energy and thus, the grain boundary mobility is inhibited [28].It is worth noting that the relative density of all samples is larger than 97.5%, demonstrating a dense microstructure.It is well known that a dense microstructure and refined grains can substantially increase the breakdown electric field (Eb) [29,30].The temperature dependence of permittivity curves is a powerful tool for revealing the composition-induced transition between the ferroelectric and relaxor-ferroelectric states.Figure 4a,b display the temperature dependences of the permittivity (ε) and loss tangent (tanδ) measured at different frequencies for two representative compositions x = 0 and x = 0.15, respectively.It is evident that two dielectric peaks (labeled as Tm and Ts) exist in the ε-T curves.The dielectric peak Tm is originated from the transition from the ferroelectric to the paraelectric phase, while the dielectric peak Ts is caused by the transformation of ferroelectric rhombohedral to tetragonal PNRs [31].As compared to the ceramic with x = 0, a strong frequency dependence of Tm and an apparent broadening of Tm peaks are observed in the ceramic with x = 0.15, demonstrating frequency dispersion and diffused phase transition behaviors, which are the typical characteristics of relaxor ferroelectrics [32].In order to clearly compare the changes of dielectric properties, the temperature dependence of ε and tanδ curves of the (1 − x)BNBT-xSZN (x = 0-0.18)ceramics at The temperature dependence of permittivity curves is a powerful tool for revealing the composition-induced transition between the ferroelectric and relaxor-ferroelectric states.Figure 4a,b display the temperature dependences of the permittivity (ε) and loss tangent (tanδ) measured at different frequencies for two representative compositions x = 0 and x = 0.15, respectively.It is evident that two dielectric peaks (labeled as T m and T s ) exist in the ε-T curves.The dielectric peak T m is originated from the transition from the ferroelectric to the paraelectric phase, while the dielectric peak T s is caused by the transformation of ferroelectric rhombohedral to tetragonal PNRs [31].As compared to the ceramic with x = 0, a strong frequency dependence of T m and an apparent broadening of T m peaks are observed in the ceramic with x = 0.15, demonstrating frequency dispersion and diffused phase transition behaviors, which are the typical characteristics of relaxor ferroelectrics [32].In order to clearly compare the changes of dielectric properties, the temperature dependence of ε and tanδ curves of the (1 − x)BNBT-xSZN (x = 0-0.18)ceramics at 100 kHz is depicted in Figure 4c.It can be seen that the T m and T s move to a lower temperature and a widened permittivity platform can be found near T m with the increasing of SZN concentration, which may be attributed to the increased disorder degree of A-and B-site ions, the reduced structural stability, and thus, the enhanced relaxor behavior through the introduction of SZN [33][34][35].Furthermore, the values of room temperature ε, ε m (the maximum of ε at T m ) and tanδ exhibit a gradual downward tendency with the increase of SZN content, which may be attributed to the enhanced relaxor behavior, and are conducive to enhancing the E b .In addition, the relaxor behavior of dielectric ceramics can be quantitatively assessed by the diffusion coefficient (γ), which can be calculated through the modified Curie-Weiss law [36]: where ε is the permittivity, ε m is the maximum of ε at T m , and C is the Curie constant.
The γ value can be obtained from the slope of the linearly fitted curve of ln(1/ε − 1/ε m ) as a function of ln(T − T m ), as shown in Figure 4d; the γ increases as the doping content increases and is determined to be 1.96 for the sample with x = 0.15, suggesting a strong relaxor behavior.
100 kHz is depicted in Figure 4c.It can be seen that the Tm and Ts move to a lower temper-ature and a widened permittivity platform can be found near Tm with the increasing of SZN concentration, which may be attributed to the increased disorder degree of A-and Bsite ions, the reduced structural stability, and thus, the enhanced relaxor behavior through the introduction of SZN [33][34][35].Furthermore, the values of room temperature ε, εm (the maximum of ε at Tm) and tanδ exhibit a gradual downward tendency with the increase of SZN content, which may be attributed to the enhanced relaxor behavior, and are conducive to enhancing the Eb.In addition, the relaxor behavior of dielectric ceramics can be quantitatively assessed by the diffusion coefficient (γ), which can be calculated through the modified Curie-Weiss law [36]: 1 ε where ε is the permittivity, εm is the maximum of ε at Tm, and C is the Curie constant.The γ value can be obtained from the slope of the linearly fitted curve of ln(1/ε − 1/εm) as a function of ln(T − Tm), as shown in Figure 4d; the γ increases as the doping content increases and is determined to be 1.96 for the sample with x = 0.15, suggesting a strong relaxor behavior.Figure 5a displays the bipolar P(E) loops of (1 − x)BNBT-xSZN ceramics measured at 120 kV/cm.The (Bi0.5Na0.5)0.94Ba0.06TiO3ceramic (i.e., x = 0) presents a square-like P(E) hysteresis loop with high values of both Pmax and Pr, indicating a strong ferroelectric characteristic.With the increase of SZN concentration, both the Pmax and Pr values decrease and the P(E) loops become slim, implying the long-range order in the ferroelectric state is gradually disrupted and transformed into a relaxor-ferroelectric state [37].Moreover, the Eb values are summarized in Figure 5b, and the Weibull distribution plots prove the reliability of the test results of Eb, as all the shape parameters β (the slope of the fitted lines) are Figure 5a displays the bipolar P(E) loops of (1 − x)BNBT-xSZN ceramics measured at 120 kV/cm.The (Bi 0.5 Na 0.5 ) 0.94 Ba 0.06 TiO 3 ceramic (i.e., x = 0) presents a square-like P(E) hysteresis loop with high values of both P max and P r , indicating a strong ferroelectric characteristic.With the increase of SZN concentration, both the P max and P r values decrease and the P(E) loops become slim, implying the long-range order in the ferroelectric state is gradually disrupted and transformed into a relaxor-ferroelectric state [37].Moreover, the E b values are summarized in Figure 5b, and the Weibull distribution plots prove the reliability of the test results of E b , as all the shape parameters β (the slope of the fitted lines) are larger than 11 [38,39].Figure 5c shows the unipolar P(E) loops of the (1 − x)BNBT-xSZN ceramics measured at E b ; the corresponding energy storage properties (W rec and η) are calculated using Equations ( 1)-( 3) and displayed in Figure 5d.It is worth noting that the optimal energy storage properties (W rec ~5.0 J/cm 3 , η ~82%) are obtained in the ceramics with x = 0.15 under an electric field of 370 kV/cm.The ferroelectric properties and energy storage performance of (1 − x)BNBT-xSZN ceramics are summarized in Table 1.The comparison of W rec and η between the sample with x = 0.15 obtained in this work and some other representative BNT-based ceramics is presented in Figure 5e [40][41][42][43][44][45][46][47].One can see that the sample with x = 0.15 exhibits a satisfactory energy storage performance.larger than 11 [38,39].Figure 5c shows the unipolar P(E) loops of the (1 − x)BNBT-xSZN ceramics measured at Eb; the corresponding energy storage properties (Wrec and η) are calculated using Equations ( 1)-( 3) and displayed in Figure 5d.It is worth noting that the optimal energy storage properties (Wrec ~ 5.0 J/cm 3 , η ~ 82%) are obtained in the ceramics with x = 0.15 under an electric field of 370 kV/cm.The ferroelectric properties and energy storage performance of (1 − x)BNBT-xSZN ceramics are summarized in Table 1.The comparison of Wrec and η between the sample with x = 0.15 obtained in this work and some other representative BNT-based ceramics is presented in Figure 5e [40][41][42][43][44][45][46][47].One can see that the sample with x = 0.15 exhibits a satisfactory energy storage performance.The prominent advantages of pulsed power devices are their exceptional power density and fast discharge speed, which exceed those of other conventional power supplies.Therefore, in order to accurately evaluate the practical pulsed charge-and-discharge properties, the under-damped and over-damped discharge current curves of 0.85BNBT-0.15SZNceramics under various electric fields were measured, and the results are shown in Figure 6.It can be seen from Figure 6a that the peak current (Imax) increases significantly with the increment of an applied electric field.The current density (CD) and power density  The prominent advantages of pulsed power devices are their exceptional power density and fast discharge speed, which exceed those of other conventional power supplies.Therefore, in order to accurately evaluate the practical pulsed charge-and-discharge properties, the under-damped and over-damped discharge current curves of 0.85BNBT-0.15SZNceramics under various electric fields were measured, and the results are shown in Figure 6.It can be seen from Figure 6a that the peak current (I max ) increases significantly with the increment of an applied electric field.The current density (C D ) and power density (P D ) can be obtained from the under-damped discharge current curves using Equations ( 5) and ( 6), respectively: where S is the electrode area, E is the applied electric field, and I max is the peak current [48].Figure 6b gives the C D and P D values of the 0.85BNBT-0.15SZNceramics as a function of electric field.Obviously, the C D and P D values gradually increase with increasing electric field, and the maximum values of 1434.94A/cm 2 and 157.84 MW/cm 3 , respectively, are obtained at 220 kV/cm.Figure 6c shows the over-damped discharge current curves (I(t)) of the 0.85BNBT-0.15SZNceramics under various electric fields.Generally, the discharge energy density (W d ) can be obtained by integrating the I(t) curves according to Equation ( 7): where R is the load resistance (200 Ω in this work), I is the discharge current, t is the discharge time, and V is the volume of sample.It is obvious that the W d values gradually increase from 0.67 J/cm 3 to 1.74 J/cm 3 as the electric field increases from 130 kV/cm to 230 kV/cm.It should be pointed out that the recoverable energy density (W rec ) obtained by integrating the P(E) loop is usually higher than the discharged energy density (W d ) measured by the discharge current curve under the same electric field.This phenomenon can be attributed to different testing mechanisms.The P(E) loop test is almost in the millisecond level (i.e., 1-100 Hz), while the charge-and-discharge test is mostly in the order of microseconds or nanoseconds, and the hysteresis effect caused by fast domain switching is more significant [49].Therefore, the discharged energy density is lower than the recoverable energy storage density.In addition, the discharge speed is also a crucial parameter for the application of pulse power devices.Generally, the discharge speed is determined by the time (t 0.9 ) required to release 90% of the total energy density (W d ) [50].
The discharge time (t 0.9 ) is determined to be only 40 ns, proving the ultrafast discharge speed of the ceramic x = 0.15.These results suggest that the 0.85BNBT-0.15SZNceramics exhibit great potential for applications in high power pulse systems.To better reveal the microscopic origin of the outstanding energy storage performance of the 0.85BNBT-0.15SZNceramics, the transmission electron microscopy (TEM) measurements were carried out.Figure 7a,b show the bright field TEM (BF-TEM) images of two representative grains in the 0.85BNBT-0.15SZNceramics.Of particular interest is that both lamellar-shaped domains, which are the signature of the ferroelectric phase [51], and blotchy domains, which correspond to the PNRs, are observed in the ceramics with x = 0.15. Figure 7c and 7d show the HR-TEM lattice fringe images obtained from the areas of the lamellar-shaped domains (Figure 7a) and the polar nanoregions (Figure 7b), respectively, indicating the ordered arrangement of atoms and the fine crystalline quality.To further elucidate the local polarization fluctuation behavior of the ferroelectric domains and PNRs, the inverse FFT patterns of the selected polar regions (30 × 30 nm 2 , marked by the red box in Figure 7c,d) are given in Figure 7e and Figure 7f, respectively.The yellow regions in the inverse FFT images represent the local polar regions [52,53].Interestingly, both areas display a remarkable local structural heterogeneity characteristic, which is caused by the random field generating due to the cations with different valences and sizes To better reveal the microscopic origin of the outstanding energy storage performance of the 0.85BNBT-0.15SZNceramics, the transmission electron microscopy (TEM) measurements were carried out.Figure 7a,b show the bright field TEM (BF-TEM) images of two representative grains in the 0.85BNBT-0.15SZNceramics.Of particular interest is that both lamellar-shaped domains, which are the signature of the ferroelectric phase [51], and blotchy domains, which correspond to the PNRs, are observed in the ceramics with x = 0.15. Figure 7c,d shows the HR-TEM lattice fringe images obtained from the areas of the lamellar-shaped domains (Figure 7a) and the polar nanoregions (Figure 7b), respectively, indicating the ordered arrangement of atoms and the fine crystalline quality.To further elucidate the local polarization fluctuation behavior of the ferroelectric domains and PNRs, the inverse FFT patterns of the selected polar regions (30 × 30 nm 2 , marked by the red box in Figure 7c,d) are given in Figures 7e and 7f, respectively.The yellow regions in the inverse FFT images represent the local polar regions [52,53].Interestingly, both areas display a remarkable local structural heterogeneity characteristic, which is caused by the random field generating due to the cations with different valences and sizes occupying the A-(Na + , Bi 3+ , Ba 2+ , and Sr 2+ ) and B-(Ti 4+ , Zn 2+ , and Nb 5+ ) sites in the unit cell [54].It is well known the local heterogeneity can impede the coherent length of dipoles and disrupt the long-range polar correlation [20,55].Therefore, one can see the inverse FFT images (Figure 7e) obtained from the lamellar-shaped domains display larger sized and more tightly connected polar regions as compared with those (Figure 7f) obtained from polar nano-regions, providing solid evidence that SZN doping disrupted the ferroelectric long-range order, reduced the size of local polar regions, and formed PNRs.It is worth noting that the lamellar-shaped domains (i.e., the ferroelectric phase) can help to enhance the P max under an external electric field, while the PNRs (i.e., the relaxor-ferroelectric phase) can delay polarization saturation to realize a high E b in the dielectric ceramics.Hence, we can infer that the coexistence of ferroelectric and relaxor-ferroelectric phases gives rise to a high energy storage density (W rec ).Furthermore, the not-too-high efficiency (η < 90%) may be closely associated with the switching of the residual lamellar-shaped ferroelectric domains under an external electric field, which usually results in a hysteretic P(E) loop.The temperature and cycle stability are crucial for energy storage applications.The temperature and cycle dependent unipolar P(E) loops are given in Figure 8a and 8b, and corresponding Wrec and η values of the x = 0.15 ceramics under an electric field of 200 kV/cm are given in Figure 8c and 8d, respectively.Apparently, all the P(E) loops do not change significantly and the Pr and Pmax values present minimal variation, leading to slight fluctuations of Wrec (<13%) and η (<11%) in the temperature range of 25-120 °C and incredibly small variations of Wrec (<2%) and η (<4%) during 10 4 cycles for the ceramics with x = 0.15, suggesting superior temperature stability and fatigue-resistant behavior of energy storage performance.Figure 9a and 9b show the over-damped discharge current curves and Wd(t) curves at different temperatures under 150 kV/cm for the x = 0.15 ceramics, respectively.Meanwhile, the corresponding Wd and t0.9 values under various temperatures are calculated and summarized in Figure 9c and 9d, respectively.It should be noted here that the Wd and η decrease slightly, from 1.0 J/cm 3 to 0.9 J/cm 3 , and from 52 ns to 41 ns when the temperature rises from 25 °C to 120 °C, exhibiting variations of less than 10% and 21%, respectively, which means the 0.85BNBT-0.15SZNceramics have an excellent temperature stability of charge-and-discharge properties.The outstanding temperature stability may be closely related to the gentle fluctuation of permittivity (ε) with respect to temperature, as shown in Figure 4c.The reliable temperature stability of the energy stor- The temperature and cycle stability are crucial for energy storage applications.The temperature and cycle dependent unipolar P(E) loops are given in Figure 8a,b, and corresponding W rec and η values of the x = 0.15 ceramics under an electric field of 200 kV/cm are given in Figures 8c and 8d, respectively.Apparently, all the P(E) loops do not change significantly and the P r and P max values present minimal variation, leading to slight fluctuations of W rec (<13%) and η (<11%) in the temperature range of 25-120 • C and incredibly small variations of W rec (<2%) and η (<4%) during 10 4 cycles for the ceramics with x = 0.15, suggesting superior temperature stability and fatigue-resistant behavior of energy storage performance.Figure 9a and 9b show the over-damped discharge current curves and W d (t) curves at different temperatures under 150 kV/cm for the x = 0.15 ceramics, respectively.Meanwhile, the corresponding W d and t 0.9 values under various temperatures are calculated and summarized in Figure 9c and 9d, respectively.It should be noted here that the W d and η decrease slightly, from 1.0 J/cm 3 to 0.9 J/cm 3 , and from 52 ns to 41 ns when the temperature rises from 25 • C to 120 • C, exhibiting variations of less than 10% and 21%, respectively, which means the 0.85BNBT-0.15SZNceramics have an excellent temperature stability of charge-and-discharge properties.The outstanding temperature stability may be closely related to the gentle fluctuation of permittivity (ε) with respect to temperature, as shown in Figure 4c.The reliable temperature stability of the energy storage performance widens the temperature usage range of the 0.85BNBT-0.15SZNceramics in dielectric capacitors.

Characterization
The phase structure of (1 − x)BNBT-xSZN ceramics was identified by X-ray diffractometer (XRD, TD-3700, Dandong Tongda Technology Co. Ltd., Dandong, China).Structural refinement was performed using the Rietveld refinement program GSAS.The microstructure was characterized by means of field-emission scanning electron microscope (FE-SEM, Carl Zeiss, Oberkochen, Germany).The grain size distributions of the samples were obtained using Nano Measurer software (version number 1.2).The domain morphology was observed by using transmission electron microscope (TEM, FEI Talos F200X, Waltham, MA, USA) operated at 200 kV.The temperature dependence of dielectric properties was tested by a precision LCR meter (Agilent E4980A, Santa Clara, CA, USA).For energy storage measurements, the samples were polished to a thickness of ~0.06 mm, and then both sides were sputtered gold electrodes with a diameter of 0.5 mm.The polarizationelectric field (P-E) hysteresis loops were measured by a ferroelectric analyzer (RT1-Premier II, Radiant Technologies Inc., Alpharetta, GA, USA) at room temperature in a frequency of 10 Hz.A commercial charge-discharge system (CFD-003, Tongguo Technology, Shanghai, China) was employed to measure the practical energy release performance.

Conclusions
In summary, (1 − x)BNBT-xSZN (x = 0, 0.08, 0.10, 0.15 and 0.18) dielectric ceramics were synthesized by a conventional solid-state reaction method.The optimal energy storage properties, with a high W rec of 5.0 J/cm 3 and an acceptable η of 82% under 370 kV/cm, were obtained for the x = 0.15 ceramics.The XRD refinement and TEM measurements demonstrate that the introduction of SZN disrupted the long-range ferroelectric order, driving the formation of PNRs and eventually resulting in the coexistence of ferroelectric and relaxor-ferroelectric phases in the x = 0.15 ceramics.The coexistence of ferroelectric and relaxor-ferroelectric phases contributes to simultaneously enhanced P max and E b values, and thus, gives rise to excellent energy storage properties.Meanwhile, the 0.85BNBT-0.15SZNceramics exhibited a large power density (P D = 157.84MW/cm 3 under 220 kV/cm) and ultrafast discharge speed (t 0.9 = 40 ns).Moreover, an outstanding temperature stability of energy storage and charge-and-discharge properties, as well as a cycle stability, were also achieved in the 0.85BNBT-0.15SZNceramics, which is of great importance for the operation of energy storage dielectric materials in harsh environments.All the merits demonstrate that the 0.85BNBT-0.15SZNceramic is a promising candidate for high-power energy storage devices.
15, and eventually, to the C phase for the ceramics with x = 0.18.The detailed phase volume fractions of the (1 − x)BNBT-xSZN (x = 0-0.18)ceramics are displayed in the insets of Figure 2a-e.Figures 2f and 2g show the SAED patterns along [001] c and [110] c directions, respectively, for the representative sample x = 0.15.

Figure 2 .
Figure 2. Rietveld refinement of XRD patterns of (1 − x)BNBT-xSZN ceramics with (a) x = 0, (b) x = 0.08, (c) x = 0.10, (d) x = 0.15, and (e) x = 0.18; SAED patterns along [001]c (f) and [110]c (g) directions for the sample with x = 0.15.The SEM images of surface morphologies and the corresponding grain size distributions of the (1 − x)BNBT-xSZN ceramics are shown in Figure 3a-e.It can be seen that all the samples present a compact microstructure and clear grain boundary, suggesting good sintering behavior.The average grain size and relative density as a function of SZN doping content are presented in Figure3f.It is evident that the grain size first increases slightly and then decreases with increasing SZN content.The slight increase in grain size for the ceramics x ≤ 0.10 may be related to the fact that a trace amount of SZN dopant acts as a sintering assistant and raises the kinetic energy of grain growth.However, the consistent

Figure 5 .
Figure 5. (a) Bipolar P(E) hysteresis loops measured at 120 kV/cm and 10 Hz.(b) Weibull distribution of breakdown electric fields.(c) Unipolar P(E) hysteresis loops measured at E b and 10 Hz.(d) W rec and η of (1 − x)BNBT-xSZN (x = 0-0.18)ceramics calculated at their breakdown electric fields.(e) Comparison of W rec and η between the 0.85BNBT-0.15SZNsample and some other representative BNT-based lead-free ceramics.

Molecules 2024 , 15 Figure 6 .
Figure 6.(a) Under-damped pulse discharge current curves of 0.85BNBT-0.15SZNceramics at various electric fields.(b) Variation of CD and PD as a function of the applied electric field.(c) Overdamped pulse discharge current curves of 0.85BNBT-0.15SZNceramics at various electric fields.(d) Discharge energy density (Wd) as a function of time for 0.85BNBT-0.15SZNceramics.

Figure 6 .
Figure 6.(a) Under-damped pulse discharge current curves of 0.85BNBT-0.15SZNceramics at various electric fields.(b) Variation of C D and P D as a function of the applied electric field.(c) Over-damped pulse discharge current curves of 0.85BNBT-0.15SZNceramics at various electric fields.(d) Discharge energy density (W d ) as a function of time for 0.85BNBT-0.15SZNceramics.
Molecules 2024, 29, x FOR PEER REVIEW 10 of 15 shaped ferroelectric domains under an external electric field, which usually results in a hysteretic P(E) loop.

Figure 7 .
Figure 7. (a,b) TEM bright-field images of two representative grains for the sample with x = 0.15.(c,d) HR-TEM lattice fringe images.(e,f) Inverse fast Fourier transform (IFFT) images converted from the areas marked with the red square lines in Figure 7c and Figure 7d, respectively, for 0.85BNBT-0.15SZNceramics.

Figure 7 .
Figure 7. (a,b) TEM bright-field images of two representative grains for the sample with x = 0.15.(c,d) HR-TEM lattice fringe images.(e,f) Inverse fast Fourier transform (IFFT) images converted from the areas marked with the red square lines in Figures7c and 7d, respectively, for 0.85BNBT-0.15SZNceramics.

Molecules 2024 , 15 Figure 8 .
Figure 8. (a,b) Unipolar P-E loops and (c,d) corresponding Wrec and η values as functions of temperature and cycles for the ceramics with x = 0.15 under an electric field of 200 kV/cm.

Figure 9 .
Figure 9. (a) Over-damped discharge current curves, (b) Wd dependence of time measured at different temperatures, and variations of (c) Wd and (d) t0.9 as a function of temperature for the ceramics with x = 0.15 under an electric field of 150 kV/cm.

Figure 8 . 15 Figure 8 .
Figure 8. (a,b) Unipolar P-E loops and (c,d) corresponding W rec and η values as functions of temperature and cycles for the ceramics with x = 0.15 under an electric field of 200 kV/cm.

Figure 9 .
Figure 9. (a) Over-damped discharge current curves, (b) Wd dependence of time measured at different temperatures, and variations of (c) Wd and (d) t0.9 as a function of temperature for the ceramics with x = 0.15 under an electric field of 150 kV/cm.

Figure 9 .
Figure 9. (a) Over-damped discharge current curves, (b) W d dependence of time measured at different temperatures, and variations of (c) W d and (d) t 0.9 as a function of temperature for the ceramics with x = 0.15 under an electric field of 150 kV/cm.

Table 1 .
Summary of ferroelectric properties and energy storage performance of (1 − x)BNBT-xSZN ceramics.

Table 1 .
Summary of ferroelectric properties and energy storage performance of (1 − x)BNBT-xSZN ceramics.