Remarkably Enhanced Dielectric Stability and Energy Storage Properties in BNT-BST Relaxor Ceramics by A-Site Defect Engineering for Pulsed Power Applications

Lead-free bulk ceramics for advanced pulsed power capacitors show relatively low recoverable energy storage density (W rec ) especially at low electric eld condition. To address this challenge, we proposed an A-site defect engineering to optimize the electric polarization behavior by disrupting the orderly arrangement of A-site ions, in which Ba 0.105 Na (BNS 0.245−1.5x ð 0.5x B 0.325+x T, x = 0, 0.02, 0.04, 0.06, 0.08) lead-free ceramics were selected as the representative. The BNS 0.245−1.5x ð 0.5x B 0.325+x T ceramics were prepared by using pressureless solid state sintering and achieved large W rec (1.8 J/cm 3 ) at a low electric eld (@110 kV/cm) when x = 0.06. The value of 1.8 J/cm 3 is super high as compared to all other W rec in lead-free bulk ceramics under a relatively low electric eld (< 160 kV/cm). Furthermore, a high dielectric constant of 2930 ± 15% in a wide temperature range of 40 ~ 350°C was also obtained in BNS 0.245−1.5x ð 0.5x B 0.325+x T (x = 0.06) ceramics. The excellent performances can be attributed to the A-site defect engineering, which can reduce P r and improve the thermal evolution of polar nanoregions (PNRs). This work conrms that the BNS 0.245−1.5x ð 0.5x B 0.325+x T (x = 0.06) ceramics are desirable for advanced pulsed power capacitors, and will push the development of a series of BNT-based ceramics with high W rec and high temperature stability.


Introduction
Dielectric capacitor is an indispensable component in contemporary electronic devices, which ful lls different functions such as dc blocking, coupling, ltering, and pulse discharge [1,2]. Considering the complicated working environment, especially high temperature (150 ~ 200 o C, even to 300 o C), ceramic dielectrics would be more suitable for energy storage candidates than other polymer materials [3].
Generally, the energy storage properties of ceramic dielectrics can be evaluated by the following equations [4]: where W, W rec , h, E, dP, P max , and P r denote total energy storage density, recoverable energy storage density, energy e ciency, applied external electric eld, polarization increment at E, maximum polarization, and remnant polarization, respectively. Therefore, high P max , low P r, and high breakdown strengthen E b are important factors to achieve high W rec [5,6]. However, high applied electric eld may limit its application in integrated electronic circuits, as well as in wearable or implantable devices requiring low electric eld. Pb-based relaxor ferroelectric and antiferroelectric ceramics had been considered as potential candidates, while the toxic of Pb limits its application [7,8]. It is urgent to design and develop new Pb-free systems with high W rec especially under relatively low electric eld.
Bismuth sodium titanate (Bi 0.5 Na 0.5 TiO 3 , BNT) possesses characteristics of high P max and complicated phase structure, and hence is considered as a promising potential energy storage ferroelectric material [9][10][11]. Noted that a high P r and poor sintering behavior of pure BNT ceramic result in a low W rec . Different methods, therefore, are utilized to improve W rec such as chemical doping, glass modi cation, multilayer structure design, advanced sintering technology [12][13][14][15][16][17]. For chemical doping, it can be divided into chemical equivalent and aliovalent doping. Especially, the chemical aliovalent doping includes "donor" and "acceptor" doping, which can induce different types of defects and improve the properties of materials  [19]. It can be observed that defect engineering can effectively improve energy storage properties by forming defect dipoles. However, it should be pointed out that a few reports on the chemical defects by adjusting element ratio in composition to optimize the energy storage properties of BNT-based ceramics.
Based on our previous work, a binary solid solution of (Bi 0.5 Na 0.5 ) 0.65 (Ba 0.3 Sr 0.7 ) 0.35 TiO 3 (BNT-BST) is considered as a good energy storage material due to "clamped" behavior in P-E loop and high dielectric constant e r (~4000) at room temperature [20,21]. However, its relatively high P r and poor dielectric temperature stability make it hard to obtain high W rec . Based on the above considerations, we proposed The phase structure evolution was identi ed using X-ray diffraction (XRD, D8-Advance, Bruker, Germany) with Cu K α radiation and Raman spectroscopy (LabRAM HR800, HORIBA). The microstructure features of the ceramic sample were observed by a scanning electron microscope (SEM, JSM-6700F, JEOL, Japan) after polishing and thermally-etching. For electrical performance testing, ceramic samples were polished to smooth and parallel in both surfaces and then painted Ag electrode on both sides. The dielectric properties of ceramics were measured using a precision impedance analyzer (HP4294A, Agilent) over a temperature range from − 100 o C to 400 o C. For P-E loop and charge-discharge measurements, the ceramic samples were polished and covered a central electrode with margin blank on both sides. The P-E hysteresis loops and I-E curves were examined using a ferroelectric analyzer at 10 Hz (Trek model 609B) based on a standard Sawyer-Tower circuit. The temperature dependent charge-discharge capability was tested using a designed RLC circuit (CFD-003, Tongguo Technology, Shanghai, China) being connected to a temperature controlled chamber. It is well known that Raman spectroscopy is an effective tool to investigate the crystalline structure information and phase transition. Figure 2(a) shows Raman spectra of BNS 0.245−1.5x ð 0.5x B 0.325+x T ceramics at room temperature. Generally, BNT material possesses 16 active phonon modes, and the irreducible representation is Γ Raman = 4A 1 + 1B 1 + 3B 2 + 8E based on the group theory [22,23]. Figure 2 Fig. 2(d), which should be owing to the increase of Sr vacancy. Meanwhile, the v3 phonon mode's wave number shows a shifting upward as well, which due to the tilt or twist of the Ti-O bond. This may be ascribed to the increase of Sr vacancies and empty in lattices. Furthermore, the wave number and signal intensity of the v6 phonon mode remain basically unchanged, while the wave number of the v5 phonon mode decreases slowly, as shown in Fig. 2(d). These changes could be due to structural transition, possibly due to TiO 6 octahedral distortion caused by the increase of Sr vacancy and empty in lattices with x. grains and extremely high density. As x value increases, the grain size gradually increases (see Fig. 3f), which should be attributed to the following two reasons: the one is that Bi 3+ replaces the A-site ion, the lattice shrinkage will cause stress, and the other is that Bi 3+ replaces Sr 2+ , in order to maintain charge balance, defects such as Sr vacancy and empty lattices will be formed, and stress will also be generated, which accelerates the mass transfer rate between particles and weakens the competition with the adjacent grains, leading to accelerated grain growth and densi cation promotion of ceramics during the sintering process [27]. where ε r , T m , C are the relative dielectric constant, the absolute temperature corresponding to the maximum dielectric constant ε m , the Curie constant, respectively. The diffuseness factor γ decreases with decreasing relaxor degree, varies between 1 (for normal ferroelectric) and 2 (for ideal relaxor ferroelectric) [29]. With increasing x value, γ value increases from 1.71 to 2.00, indicating relaxor characteristic is effectively enhanced, for pure BNT-BST ceramics, one dielectric peak at T m and obvious relaxor behavior below T m can be observed. After composition modi cation, wave-like double peaks can be found in dielectric spectra, and corresponding temperature is named as T m1 and T m2 , respectively. Noted that dielectric relaxor behavior only exists at a temperature below T m1 , and thus dielectric peak at T m1 is similar to that at T m . Thus, a new dielectric peak at T m2 should be induced by Bi-Sr ratio change.

Results And Discussion
Temperature stability of dielectric constant (TCC) is a crucial factor to in uence its application scenes. In general, it can be calculated by the equation as follows: where e base denote dielectric constant at a based temperature, other is accorded to above mentioned. As x value increases, the dielectric peak at T m1 is suppressed, while the temperature difference ΔT (T m2 -T m1 ) is gradually enhanced, which is bene cial for enhancing temperature stability. Fig. 4(f) shows TCC of BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics with different x values chosen 150 o C as based temperature. With increasing x value, dielectric temperature stability is effectively enhanced especially at high temperature, which is attributed to the appearance of a new dielectric peak at T m2 . For x = 0.06 composition, a wide temperature range of TCC at ± 15% corresponds to 40~350 °C.
In order to further explore the reason for generating dielectric peak at T m2 , a variable Raman spectrum is used. Fig. 5 shows Raman spectra and wave number of Raman peaks of BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics with x = 0.06 over a temperature range from room temperature to 275 o C. As temperature increases, the signals of modes v2 and v6 gradually present a disappeared tendency, while other modes change slightly, as shown in Fig. 5(a). All modes display a decreased tendency, while no abruptly change in wave number (Fig. 5b). The result illustrates that mode v2 and v6 is sensitive to phase structure of BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics with x = 0.06. The slow transition of modes v2 and v6 both demonstrate phase structure of x = 0.06 composition just have a slight change. Therefore, it can be concluded that the new dielectric peak at T m2 should be attributed to the thermal evolution of PNRs affected by the concentration of Sr vacancies.
Figure 6(a) shows P-E loops of BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics at 60 kV/cm and room temperature. As x value increases, P-E loops gradually go slim, which is bene cial for improving energy storage properties, the maximum polarization at a given electric eld slightly decreases with the increase of x value. Combined with I-E loops at the same conditions, current peaks gradually diffuse, and corresponding intensity decreases, indicating the relaxor characteristic is enhanced [30]. In addition, P max , P r and P max -P r versus x value for BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics at 60 kV/cm is exhibited in Fig.   6(b). A relatively high P max -P r of 27.52 μC/cm 2 can be obtained at x = 0.06 composition due to a rapid decrease in P r , because the PNRs are dynamic sensitive to external electric eld stimuli. Figs. 6(c)-6(d) show energy e ciency h and recoverable energy density W rec of BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics at different electric elds, respectively. As x value increases, h presents an increased tendency obtaining a high value for x = 0.06 composition, and then decreases again with further increasing x value. Meanwhile, BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics with x = 0.06 possess a maximum W rec of 1.8 J/cm 3 only at a low electric eld of 110 kV/cm, as shown in Fig. 6(d).
In order to investigate the working stability at different external elds, temperature, frequency and electric fatigue dependent energy storage properties of BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics with x = 0.06 have been examined. Fig. 7(a) shows P-E loops of x = 0.06 ceramics over a temperature range of 30~150 o C at Page 7/18 60 kV/cm and 10 Hz. As temperature increases, P-E loops gradually go slim, and keep a high P s and low P r . Therefore, W rec and h of x = 0.06 ceramics possess good temperature stability, as exhibited in Fig.   7(b). Meanwhile, frequency dependent P-E loops of x = 0.06 ceramics at 60 kV/cm are displayed in Fig.  7(c). It can be seen that energy loss has a slight increase tendency during discharge process, which may be related to vacancy defect pin domain wall. W rec and h of x = 0.06 ceramics, therefore, show a slight decrease in value at frequency of 1~100 Hz, as shown in Fig. 7(d). Finally, P-E loops as functions of cycle numbers and corresponding W rec and h for x = 0.06 ceramics are illustrated in Figs. 7(e)-(f), respectively.
Noted that polarization of x = 0.06 ceramics keeps a stable value at 10 Hz after 10 5 electric cycles.
Obviously, x = 0.06 ceramics possess a good fatigue endurance, and W rec and h as functions of cycles are illustrated in Fig. 7(f).
Charge-discharge characteristic is an essential factor for dielectric materials to evaluate its energy storage capabilities, and thus charge-discharge measurement is ful lled at a speci ed circuit. Generally, discharge energy density W d can be calculated by the equation as following [31]: where R is load resistance (100 W), i is the maximum discharge current, and V is the effective volume of ceramic between two electrodes. Figs. 8(a)-(b) show room temperature underdamped discharge waveform and corresponding W d of BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics with x = 0.06 ceramics at different electric elds, respectively. As the electric eld increases, the maximum discharge current I max and W d both gradually enhances. It should be mentioned that W d is less than W rec at the same electric eld for x = 0.06 compositions. This may be attributed to the following two reasons [32]: the one is that the domain can not switch quickly to respond to the external electric eld, and the other is that equivalent series resistance (ESR) generate joule heat during charge-discharge process. The discharge rate is characterized by evaluating factor t 0.9 (dashed line in Fig. 8b), which represents the time needed for releasing 90% of all stored energy [33]. Fig. 8(b) shows that t 0.9 is about 0.1 µs for x = 0.06 ceramics at room temperature, which illustrates energy can be released by a pulse current way in a short time. Variable temperature discharge current curves as function time for BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics with x = 0.06 ceramics are displayed in Fig. 8(c). As temperature increases, the maximum discharge current I max basically keeps a stable value while W d possesses an obvious enhancement as shown in Fig. 8(d). The discharge capability of x = 0.06 ceramics possesses good temperature stability, which is bene cial for the application in high temperature environment.
To better evaluate the energy storage properties of BNS 0.245-1.5x ð 0.5x B 0.325+x T ceramics, we compare W rec and E max of x= 0.06 and 0.08 composition with some lead-free ceramic bulks reported previously [34][35][36][37][38][39][40][41][42][43][44]. It can be seen from Fig. 9(a) that a large W rec (>1.5 J/cm 3 ) usually requires a high E b (>160 kV/cm) to produce high polarization, especially for some BT-based and KNN-based materials. In this work, a high W rec can be achieved under a relatively low electric eld, which exceeds other BNT-based energy storage ceramics at the same electric eld, even other lead-free systems. With further comparing W rec and h of different compositions, as shown in Fig. 9(b). It should be pointed that high W rec and h is hard to obtain simultaneously in one system in uenced by heat loss at electric eld. Note that BNS 0.1555 ð 0.03 B 0.385 T (x = 0.06) ceramics possess a relatively high W rec (>1.5 J/cm 3 ), together with high h (>70%) under relatively low electric eld (<160 kV/cm), demonstrating it is potential to obtain both high W rec and h, which should be a promising candidate for power ceramic capacitors application.

Conclusions
In this work, A-site defect engineering was proposed to improve the energy storage performance of BNS 0.245−1.5x ð 0.5x B 0.325+x T lead-free ceramics. High recoverable energy density of 1.8 J/cm 3 under low electric eld (@110 kV/cm) and energy e ciency of 72% are achieved simultaneously in the sample with x = 0.06. This good energy storage performance is attributed to the A-site defect engineering that can reduce P r . The ceramic also exhibits satisfactory thermal, frequency, and cycling stabilities as well as a high charge-discharge rate. BNS 0.245−1.5x ð 0.5x B 0.325+x T (x = 0.06) ceramics show a high dielectric constant of 2930 ± 15% in a wide temperature range of 40 ~ 350°C. This high-temperature stability is attributed to the A-site defect engineering, which can improve the thermal evolution of PNRs. All these advantages indicate that BNS 0.245−1.5x ð 0.5x B 0.325+x T ceramics is suitable for solid state pulse power ceramic capacitors, and A-site defect engineering is a robust strategy to improve the W rec and hightemperature stability of lead-free ceramics.           P-E loops as a function of (a) temperature, (c) frequency, and (e) cycle numbers for BNS0.245-1.5xð0.5xB0.325+xT ceramics with x = 0.06. Wrec and η as a function of (b) temperature, (d) frequency, and (f) cycle numbers for BNS0.245-1.5xð0.5xB0.325+xT ceramics with x = 0.06.