Simultaneously high-energy storage density and responsivity in quasi-hysteresis-free Mn-doped Bi0.5Na0.5TiO3-BaTiO3-(Sr0.7Bi0.2□0.1)TiO3 ergodic relaxor ceramics

ABSTRACT High-energy storage density, responsivity and efficiency, i.e. WR = 1.07 J/cm3, ξ = 119 J/(kV m2) and η = 92%, were simultaneously obtained in Mn-doped 0.62Bi0.5Na0.5TiO3-0.06BaTiO3-0.32(Sr0.7Bi0.2□0.1)TiO3 ergodic relaxor ceramics. Appropriate Mn doping was beneficial to enhance breakdown field strength. Moreover, temperature and different atmosphere-dependent impedance spectroscopy results indicated that oxygen vacancies were the conductivity mechanism for all samples. The valence state of Mn together with the conjugation between Mn ion and oxygen vacancies was confirmed by X-ray photoelectron spectra and electric paramagnetic resonance. The above results indicate that quasi-hysteresis-free loops with high-energy storage performances can be obtained by the induced defect complex. IMPACT STATEMENT High-energy storage density WR = 1.07 J/cm3, responsivity ξ = 119 J/(kV m2) and efficiency η = 92% were simultaneously obtained in quasi-hysteresis-free ceramics by introducing defect complex. GRAPHICAL ABSTRACT


Introduction
High performances of dielectric materials are desired to meet the ever-increasing consumer and industrial requirements for electronic components. The ubiquitous lead zirconium titanate (PZT)-based piezoelectric ceramics, though dominating the electronic markets, are gradually abandoned due to its contamination to the environment. Recently, lead-free perovskite-type ceramics have received a flurry of interests and much progress has been made [1,2]. Among them, lead-free candidate to the application in energy storage capacitors [4]. Later, much attention has been paid to the energy storage properties in BNT-based ceramics [5][6][7][8]. As a matter of fact, apart from the advantage of the solidstate dielectric bulks that can operate at high temperature, BNT-based ceramics can also exhibit a good temperature stability because of their diffuse phase transition behavior due to the transformation between R3c and P4bm polar nanoregions (PNRs) [9,10].
The widespread application of power electronics demands high-energy storage/responsivity and efficiency material, the W R and efficiency in BNT-based ceramics is limited yet so far. With respect to enhancing W R ; generally, two routes are usually adopted to in the BNT-based material.
(i) The coexistence of nonergodic and ergodic phase will result in large maximum polarization (P max ) and low remnant polarization (P r ), thus enlarging the gap between P max and P r , which is beneficial to obtain large W R (usually with pinched P-E loops). This can be usually realized via composition and temperature modulation [5,7]. (ii) Another approach to enhance W R is improving the breakdown field strength (BDS). The addition of 'hard' dopant (Fe, Mn, etc.) and complex-ion (such as (Ni 1/3 Nb 2/3 ) 4+ ) is believed to be an effective method to improve BDS [11,12].
For example, significant results have been obtained in (1−x)Bi 0.48 La 0.02 Na 0.48 Li 0.02 Ti 0.98 Zr 0.02 O 3−x Na 0.73 Bi 0.09 NbO 3 ceramics reported by Yang et al. The maximum energy storage density can be obtained for the sample with x = 0.10 at room temperature (RT), with an energy storage density of 2.04 J/cm 3 at 178 kV/cm, the performance of which is outstanding in lead-free ceramics. However, the energy storage efficiency is only around 55% [13]. Balancing of energy storage density and efficiency in the BNT-material has been a difficult problem. The transformation between the relaxor phase and long ferroelectric order is always accompanied by giant strain with large hysteresis in BNT-based materials. The charge-discharge time is within few micro-seconds, the exceptionally large strain in BNT-based ceramics will lead to large localized stress and may easily crack during repeated charge-discharge cycles, especially when high electric field is applied [5]. Moreover, due to hysteresis, a large portion of the energy will dissipate as heat during the discharging process, which inevitably results in high energy loss and reduce service life for electronic components. It is worthy to note that the pure ergodic relaxor phase can get rid of the above two shortcomings. However, the P max for the pure ergodic phase is usually low when far away from the nonergodic-ergodic phase region due to the enhanced local random field that impedes the appearance of long-range ferroelectric order, then low W R will be obtained. So we envisaged whether a low hysteresis loops can be designed and large energy storage density/responsivity and efficiency can be obtained simultaneously under moderate electric field in BNT-based ceramics?
It is extensively reported that appropriate doping of transition metal ions with different valences (e.g. Mn, Fe ions) is beneficial for the enhancement of resistivity and piezoelectric properties [14,15]. In this letter, the energy storage properties and defect chemistry of Mn-doped 0.62Bi 0.5 Na 0.5 TiO 3 -0.06BaTiO 3 -0.32(Sr 0.7 Bi 0.2 0.1 )TiO 3 ergodic relaxor ceramics have been studied in detail. The impedance spectroscopy, X-ray photoelectron spectra (XPS) and electric paramagnetic resonance (EPR) were also investigated to explore the role of the Mn dopant in modifying the defect microstructure and energy storage properties. The results indicated that quasi-hysteresis-free loops with high-energy storage performances can be obtained by inducing defect complex. The experimental details are shown in the supplemental data.

Results and discussions
Figure 1(a) shows the XRD patterns of the 0-2 Mn ceramics. It can be seen that the main perovskite phase can be formed for these ceramics and a high content of Mn addition will result in the appearance of second phase (Sr 6 Bi 2 O 9 ). Moreover, all single diffraction peaks signify the pseudocubic symmetry for the studied samples. The inset of Figure 1(a) shows the SEM microstructure for 1 Mn ceramic morphology as a representative. All compositions are well densified with pore-free microstructure. The Mn addition has a minor effect on the grain size and it is in the vicinity of 1.5 μm.
The distribution of the BDS of ferroelectric ceramics is considered to follow the Weibull function [16]. According to the Weibull distributions, the BDS data of the 0-2 Mn ceramics are plotted in Figure 1(b). Twenty samples for each composition are measured to obtain the average BDS values. The results show that all the plots display a relatively good linearity and the Weibull modulus β is more than 8, indicating the reliability of the fitting [17]. Besides, β is correlated with the range of BDS, and the intercept determines the magnitude of BDS, as shown by the dashed line. Apparently, the BDS first increases from 0 Mn to 1 Mn sample, and gradually decreases with a further addition of the Mn content. It is well known that the BDS depends on a couple of factors, including porosity, grain size, conductivity mechanism and extrinsic parameters such as sample thickness and electrode area. For the 0-2 Mn ceramics, the evolution of RT dielectric loss and the appearance of second phase may be a key factor to determine the BDS value. As shown in Figure 1(c), for low Mn doping with 0-1 Mn ceramics, the dielectric loss is gradually depressed. However, more Mn amount will enhance the dielectric loss, as indicated by the dashed rectangle in Figure 1(c), which may be due to the fact that excessive Mn content provides more pathways for the localized electron hoping process; at the same time, the appearance of the second phase will also make the ceramics leaky, so the ceramics will breakdown more easily with the application of electric field [18]. Therefore, the value of BDS is maximized for 1 Mn ceramic with the low dielectric loss at RT. Moreover, the dielectric constant of all samples exhibits a shoulder at T s and a broad maximum at T m , both of which are closely related with the evolution of PNRs and an indication of relaxor phase transition behavior [19].
In order to ascertain the conduction mechanism in the 0-2 Mn ceramics, impedance spectroscopy is adopted as it is known to be an effective technique to characterize the contribution of the grain, grain boundary and electrode interface to the electrical conduction behavior in ferroelectric ceramics [20]. Complex impedance spectra are measured in the selected temperature range of 470-580°C for 0-2 Mn ceramics, as shown in Figure 2(a-e) (dots). It is worth noting that the semicircles do not appear at a lower temperature because of the high resistivity. A single semicircle is observed from the complex impedance results indicating that a single relaxation mechanism dominates the impedance. The value of Z decreases with an increase in the measuring temperature, demonstrating that all samples exhibit a negative temperature coefficient of resistance behavior that is analogous to the semiconductor [21]. Due to the non-Debye relaxation behaviors, a parallel R-C circuit is simulated to fit the impedance data, as shown in Figure 2(f). By using the ZView software, the resistance (R b ) and capacitance (C b ) are simulated based on the fitting results and a close agreement between the experimental (dots) and fitting results (line) is obtained. The fitting results are listed in the supplementary material ( Table 1). The conductivity of the bulk σ b and activation energy E a can be calculated by the formula [22]: where h is the sample thickness, S is the electrode area, R b is the bulk resistance, E a is the activation energy of conduction, k B is the Boltzmann constant and T is the absolute temperature. The variation of bulk conductivity with the temperature is plotted in Figure 2(g). It can be seen that the experimental data obey the Arrhenius law according to the fitting results. As seen from Figure 2  materials [24]. The E a refers to the dissociation of defect dipoles and corresponds to the ability of transition from constrained oxygen vacancies to the freely moved ones. The gradual enhancement of E a with Mn addition is believed to be related to the reduction of free oxygen vacancies [25]. Therefore, appropriate Mn doping can hinder the migration of oxygen vacancies and the local electron hopping and then effectively enhance the BDS. Figure 3(a,b) shows the Z (f ) and M (f ) spectra for x = 1 Mn ceramic over 470-580°C as a representative. Impedance spectroscopy with a single peak indicates the absence of the grain-boundary component in x = 1 Mn ceramic as there is no additional arcs at a lower frequency range [26], which is also in agreement with the above discussions. Complex impedance plots measured in nitrogen, air and oxygen atmosphere are shown in Figure 3(c). It can be seen that the impedance plots in air and oxygen are nearly the same, which predicts that the resistivity remains unchanged when is 1 Mn ceramic is exposed to air and oxygen atmospheres. However, when the 1 Mn ceramic is exposed to nitrogen atmosphere, the resistivity is drastically reduced. A similar phenomenon has also been witnessed in other Mn-doped BNT-BT-SBT ceramics. The obvious variation of resistivity in the inert nitrogen atmosphere indicates that ceramics exhibits an n-type semiconductor behavior, which is probably due to the loss of oxygen according to previous studies [27,28]. The normalized activation energies E a for three different atmospheres are displayed in Figure 3(d). It can be obviously seen that the E a increases with an increase in the oxygen partial pressure. It is known that exposure to nitrogen atmosphere for these ceramics is easily to be induced the oxygen vacancies. The enhancement of E a indicates a decrease in the concentration of oxygen vacancies. Figure S1 (a1-a5) shows the XPS profile of O 1 s of the 0-2 Mn ceramics. It is found that the spectra are composed of two peaks by using Gaussian fitting. The first peak located at the low binding energy around 529.6 eV is assigned to be lattice oxygen. Another peak in the vicinity of 531.4 eV is attributed to the presence of oxygen vacancies in the 0-2 Mn ceramics, which confirms the existence of oxygen vacancies. Ti 2p 3/2 and Ti 2p 1/2 XPS peaks are displayed in Figure S1 (b1-b5). According to the fitting results, the peak position has not changed for the all compositions within the resolution of XPS  (shown by the dashed line), probably indicating the positive charge V ·· 0 is not compensated by the reduction of Ti 4+ to Ti 3+ state [29]. However, further systematic research is needed to clarify the valence of Ti under this condition.
To further give an insight into the role of Mn element in the BNT-BT-SBT ceramics, X-band EPR spectra are characterized for the all compositions, as shown in Figure 4(a). These spectra are analogous to other Mndoped BNT-based ceramics and demonstrate that the Mn element has been incorporated into the octahedral environment at the Ti positions [30]. The EPR signal of 0 Mn ceramic is absent because of its diamagnetic nature. The 0.5-2 Mn ceramics exhibit a hyperfine structure with typical six-line splitting but with different EPR intensities. This signal originates from the Mn 2+ as a type of free Mn Ti . The EPR signal of Mn 4+ lies below the detect limit for the EPR spectra in this study. In order to confirm the oxidation state of the Mn element, XPS of Mn 2p spectra for 1 Mn ceramic is presented in the inset of Figure 4(a). The overlapped spectra are also fitted by the Gaussian function to determine the valent state of the Mn dopant. It can be seen that two peaks with different binding energies are observed, indicating the existence of two Mn valent states. The main peak around 641.3 and 653.4 eV can be assigned to Mn2p 3/2 and Mn2p 1/2 , respectively. Specifically, the peak at the lower binding energy of 640.2 eV is assigned to Mn 2+ , the peak at the higher binding energy of 641.9 eV is probably assigned to Mn 4+ , which is in a good agreement with the literature [31]. Therefore, the mix Mn 2+ /Mn 4+ valent states are obtained in Mn-doped BNT-BT-SBT ceramics, further suggesting that the Mn element has diffused into the perovskite B-site in the ceramics. It is obvious that the inhomogeneous EPR signals can be observed, which indicates the presence of the nearest-neighbor oxygen vacancies. With increasing Mn content, the intensity of the EPR signal is enhanced, which is related with the increasing concentration of Mn 2+ and corresponding oxygen vacancies. Moreover, the induced oxygen vacancies are expected to be easily trapped by the free Mn 2+ ion since free Mn 2+ has a strong force to attract V ·· 0 , which is evidenced by the slight asymmetry of the EPR signal that indicates that oxygen vacancies have been trapped by the Mn Ti defect center and form the Mn Ti − V ·· 0 complex [29]. Under this condition, incorporation of Mn will result in the presence of defect complexes Mn Ti − V ·· 0 in the ceramics, which is schematically illustrated in Figure 4(b). The conjugation between Mn 2+ and oxygen will create a P D , which is in parallel with spontaneous polarization (P s ). When the external electric field is applied to the ceramics, P s switches instantaneously with the cycling of external electric field yet the P D exhibits a much slower switching speed compared to P s . So the P D will act as an internal field to switch the domain back to its original state with the release of electric field [11].
The RT P-E loops of the all samples are shown in Figure S2. Only slim P-E loops are observed, which signal the pure ergodic phase for all ceramics [32]. Moreover, the addition of 0.5 Mn and 1 Mn is beneficial to reduce the energy loss for the ceramic, which is related to the appearance of Mn Ti − V ·· 0 , as shown in the inset of Figure S2. The 1 Mn ceramic is selected to study its energy storage behavior because of its enhanced theoretical BDS according to Weibull distributions. As shown in Figure 5(a,b), bipolar P-E hysteresis loops with current density J-E loops of 1 Mn ceramic are measured under different electric fields at RT and 100°C. Quasi-hysteresisfree loops are obtained and the P max gradually increases with the increase in electric field. Generally, the recoverable energy storage density is calculated according to the formula W R = P max P r EdP, the energy storage efficiency η = W R /(W R + W L ) and W S = W R + W L [33], which is illustrated in Figure 5(g). It is obvious that both the stored and recoverable energy storage density W S and W R increase with augmenting electric field. As shown in the inset of Figure 5(a,b), the back-forward and forward E B and E F appear under high electric field strengths, indicating that the ferroelectric order is induced with the coherence of PNRs and then contributes to the large polarization. It is found that the W R for 1 Mn ceramic could reach as high as 1.07 J/cm 3 under moderate electric field of 9 kV/mm at RT and also maintains 0.86 J/cm 3 under 8 kV/mm at 100°C. Interestingly, within various electric field strengths at RT and 100°C, the η is no less than 90%, indicating the ultra-low energy loss during domain switching. This may be the contribution of the dipole complex Mn Ti − V ·· 0 , as discussed above. It should be noted that the η becomes low under an electric field of 8 kV/mm at 100°C, probably due to the enhanced conductivity at high temperatures and it may be also the mechanism for the decreased BDS. A normalized parameter ξ = W R /E is introduced to characterize the energy storage ability. The ξ could reach 119 J/(kV m 2 ) together with a large η = 92% under moderate electric field (9 kV/mm), which is much higher than other BNT-based ceramics, as shown in Figure 5(d,e).
According to literature survey, the relationship of W R ∼ η and W R ∼ ξ is plotted in Figure 5(a,b) among the reported BNT-based ceramics. Generally, there are three approaches to enhance W R : (i) Modulating the phase portion of ferroelectric phase and ergodic phase to maximize the P ( P = P max − P r ) and minimize the W L at moderate electric field. Usually, both of W R and η exhibit common performances.
(ii) A large P max accompanied by a large W L under high electric field. Both of the W R and W L is quite large, however, the η is very low. (iii) Enhancing the measuring electric field for the ceramics within the ergodic phase. The W R can be enhanced compared to I. However, η remains low still.
As a matter of fact, both the W R and η are the two important parameters to measure the energy storage performance in the application of power electronics. What is more, the ceramics should not be repeatedly operated under high electric field, especially for BNT-based material with a large strain. Therefore, the parameter ξ instead of W R is expected to be a benchmark for the energy storage ability in different BNT-based ceramics. The simultaneously high ξ and η is as a result of the quasihysteresis-free loops with high polarizations, as indicated in Figure 5(i).

Conclusions
Mn doping is beneficial to enhance BDS and W R = 1.07 J/cm 3 , ξ = 119 J/(kV m 2 ) and η > 90% are simultaneously obtained in 1 Mn ceramics. The high performance of the ceramic originates from the quasihysteresis-free loops with high polarization as result of the appearance of the defect dipoles, which is confirmed by synergistic measurements, including the XPS, complex impedance and EPR analysis. This work may provide another avenue to design simultaneous high-energy storage density, responsivity and efficiency in BNT-based ceramics.

Disclosure statement
No potential conflict of interest was reported by the authors.