Phase–composition and temperature dependence of electrocaloric effect in lead–free Bi0.5Na0.5TiO3–BaTiO3–(Sr0.7Bi0.2□0.1)TiO3 ceramics

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Phase–composition and temperature dependence of electrocaloric effect in lead–free Bi0.5 Na0.5 TiO3 –BaTiO3 –(Sr0.7 Bi0.2 䊐0.1 )TiO3 ceramics Feng Li a,b,c , Guorui Chen d , Xing Liu a , Jiwei Zhai a,∗ , Bo Shen a , Huarong Zeng b , Shandong Li d,∗ , Peng Li a , Ke Yang a , Haixue Yan e,∗ a

Functional Materials Research Laboratory, School of Materials Science & Engineering, Tongji University, 4800 Caoan Road, Shanghai 201804, China Key Laboratory of Inorganic Functional Materials and Devices, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China c University of Chinese Academy of Sciences, Beijing 100049, China d College of Physics, Qingdao University, 308 Ningxia Road, Qingdao 266071, China e School of Engineering and Materials Science, Queen Mary University of London, Mile End Road, London, E1 4NS, UK b

a r t i c l e

i n f o

Article history: Received 1 May 2017 Received in revised form 19 June 2017 Accepted 20 June 2017 Available online xxx Keywords: BNT–based ceramic Lead–free Relaxor ferroelectric Electrocaloric

a b s t r a c t The (0.94–x)Bi0.5 Na0.5 TiO3 –0.06BaTiO3 –x(Sr0.7 Bi0.2 䊐0.1 )TiO3 (BNT–BT–xSBT, 0 ≤ x ≤ 0.24) solid solution ceramics were synthesized via a conventional solid–state reaction method and the correlation of phase structure, piezoelectric, ferroelectric properties and electrocaloric effect (ECE) was investigated in detail. The ECE in lead–free BNT–BT–xSBT ceramics was measured directly using a home–made adiabatic calorimeter with maximum adiabatic temperature change T = 0.4 K with x = 0.08 under the electric field E = 6 kV/mm at room temperature. The position of maximum ECE was found in the vicinity of nonergodic and ergodic phase boundary, where the maximum change in entropy occurs as a result of the field–induced phase transformation between the ergodic and long–range ferroelectric phase. Besides, the mechanism for the shift of ECE peak is discussed in detail. Finally, the temperature dependence of ECE for BNT–BT–xSBT (x = 0, 0.04 and 0.08) was also investigated. This work may present a guideline for designing BNT–based ferroelectric relaxor ceramics for EC cooling technologies. © 2017 Elsevier Ltd. All rights reserved.

1. Introduction The ECE is based on the entropy or temperature change occurring in polar crystals when an electric field is applied or removed under adiabatic conditions [1,2]. As the ECE can be accomplished by the simple operating force of an external electric field compared with the magnetocaloric/elastocaloric effect induced by large magnetism/stress [3,4], the research of ferroelectric refrigeration based on ECE have gained a resurgence since a giant ECE with adiabatic temperature change T ≈ 12 K was obtained in Pb(Zr0.95 Ti0.05 )O3 ceramic thin film and P(VDF–TrFE) copolymer thin film [5,6]. Afterwards, the ECE have been extensively studied in a variety of ferroelectric and antiferroelectric ceramics and great progress on ECE has been made [7–17]. However, volatility and toxicity of lead limit their extensive applications. Recently, eco–friendly lead–free (Bi0.5 Na0.5 )TiO3 (BNT)–based ferroelectric relaxor ceramics have attracted great interest due to

∗ Corresponding authors. E-mail addresses: [email protected] (J. Zhai), [email protected] (S. Li), [email protected] (H. Yan).

superior electrostrain properties [18–20], however, the research on EC properties is at an early stage. Particularly, due to the extra contribution of polar nanoregions (PNRs) accompanied by high dielectric constant, large ECE in BNT–based lead–free relaxor ferroelectrics could be expected [21]. Therefore, exploring ECE in BNT–based ferroelectric relaxor ceramics is of great significance. One main obstacle for understanding ECE in BNT–based ceramics could be the inconsistent measurement methods. Both of indirect and direct characterization on ECE have been reported in BNT–based ceramics and significant differences in the results were observed [10,22,23]. The indirect method refers to the ECE calculated based on Maxwell relation: 1 T = − C

E2  T E1

∂P ∂T

 dE

(1)

E

where P is the polarization, T is the temperature, E1 and E2 are the initial and final electric field strength, C is the heat capacity and  is the density of the material. As a matter of fact, the indirect thermodynamic method generates some concerns when applied to non–equilibrium systems such as nonergodic relaxors. There-

http://dx.doi.org/10.1016/j.jeurceramsoc.2017.06.033 0955-2219/© 2017 Elsevier Ltd. All rights reserved.

Please cite this article in press as: F. Li, et al., Phase–composition and temperature dependence of electrocaloric effect in lead–free Bi0.5 Na0.5 TiO3 –BaTiO3 –(Sr0.7 Bi0.2 䊐0.1 )TiO3 ceramics, J Eur Ceram Soc (2017), http://dx.doi.org/10.1016/j.jeurceramsoc.2017.06.033

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fore, direct method for determining ECE in BNT–based ceramics is preferred [24]. The ECE of BNT–based binary systems, such as Bi0.5 Na0.5 TiO3 –Bi0.5 K0.5 TiO3 , Bi0.5 Na0.5 TiO3 –KNbO3 has been measured directly using a modified differentiate scanning calorimeter (DSC) by Goupil, et al. [25,26]. However, the emphasis is played much on the temperature–dependent ECE for a single composition, and the ECE peak is always located at the high temperature that beyond the useful range for practical applications. In our previous study, ECE of Bi0.5 Na0.5 TiO3 –0.06BaTiO3 system was investigated and its ECE peak was in the vicinity of ferroelectric–relaxor transition temperature (TFR ≈ 100 ◦ C) [27]. Through the substitution of the third component such as SrTiO3 , (Sr, Bi)TiO3 , the increasing of relaxor degree and local random field can modulate the TFR /Tf toward room temperature (Tf –freezing temperature) [28–30]. One question haunts us that whether the maximum ECE occurs when TFR /Tf is tailored near room temperature or just in that case in a finite conditions? If it is true, then we can predict the position of largest ECE in BNT–based ceramics and provide a guideline to design BNT–based compositions for solid state refrigeration. In this study, the ECE of BNT–BT–xSBT pseudo–ternary system is measured directly using a precise home–made adiabatic calorimeter and the evolution of phase structure, piezoelectric, ferroelectric and ECE as a function of SBT content are investigated in detail. Particularly, the relation between TFR /Tf and ECE is constructed in order to clarify the question presented above. The temperature stability of ECE in nonergodic and ergodic relaxor ferroelectric ceramics (x = 0, 0.04 and 0.08) is also investigated.

etched by thermal treatment at 1060 ◦ C for 30 min. The dielectric properties were measured using an LCR meter (Agilent E4980A, Santa Clara, CA) from 25 ◦ C to 400 ◦ C. A ferroelectric testing system (Precision LC, Radiant Technologies, Inc. Albuquerque, NM) was used to measure the polarization-electric field (P–E) hysteresis loops. The piezoelectric constant d33 was measured using a quasi-static d33 meter (ZJ-6A, Institute of Acoustics, Beijing, China). Silver paste was painted on major sides of the discs and fired at 600 ◦ C for 30 min as the electrodes. For the ECE characterizations, a home–made adiabatic calorimeter was used that allows high resolution measurements to determine temperature change, the schematic sample configuration for the T measurement is displayed in Fig. 1.

2. Experimental procedure

3. Results

The BNT–BT–xSBT (x = 0–0.24) ceramics were synthesized by a conventional solid–state reaction method using high–purity chemicals: Bi2 O3 (99.0%), Na2 CO3 (99.8%), TiO2 (99.0%), SrCO3 (99.0%), and BaCO3 (99.0%) as raw materials, and the detailed experimental procedures can be found elsewhere [29]. The phase structures were analyzed by an X–ray diffractometer (XRD, D/Mzx–rB; Rigaku, Tokyo, Japan) with CuK␣1 radiation. Raman spectra were measured by a Horiba Lab-Ram iHR550 spectrometer. The microstructure was observed using a scanning electron microscope (SEM, JEOL JSM6490LV, Tokyo, Japan) and the samples surface were polished and

3.1. Phase and microstructure

Fig. 1. The schematic sample configuration for the T measurement.

Fig. 2(a) shows the XRD patterns of unpoled BNT–BT–xSBT (x = 0–0.24) ceramics at room temperature. A pure perovskite structure is obtained for all compositions and no appreciable second phase appears. In order to give an insight into the phase structure evolution of BNT–BT–xSBT system, a locally magnified (111) and (200) peak for unpoled and poled ceramics are presented in Fig. 2(b) and (c). It is well known that the composition x = 0 (BNT–0.06BT) locates in the vicinity of rhombohedral–tetragonal

Fig. 2. (a) Room–temperature XRD patterns of unpoled BNT–BT–xSBT ceramics (x = 0–0.24); (b) and (c) the locally magnified (111) and (200) diffraction peaks for unpoled and poled ceramics, respectively.

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Fig. 3. The grain morphology of the polished and thermally etched samples for BNT–BT–xSBT (x = 0, 0.08, 0.16 and 0.24) sintered at their optimum temperature.

morphotropic phase boundary (MPB) [31], however, the coexistence of a cubic and polar phase in unpoled BNT–0.06BT leads to a tiny of non–cubic distortion and these structural characteristics are the same to the other compositions in unpoled state, as shown in Fig. 2(b). Upon the application of an electric field, PNRs correlate with each other, finally inducing a ferroelectric phase. The splitting of (111) and (200) peaks for x = 0–0.04, shown by the arrows in Fig. 2(c), indicates the coexistence of tetragonal and rhombohedral phase. When the SBT content increases to x = 0.08–0.24, single (111) and (200) peaks can be observed, however, presence of a small hump on the lower angle side of (200) peak was found, as shown by diamond symbol in Fig. 2(c). The hump appearing on the lower angle side of the (200) peak is probably not related to tetragonal phase, but to a tetragonal distortion of the ergodic phase induced by the external electric field. With increasing SBT content, the intensity of hump gradually weaken. Similar phenomenon have been also witnessed in Bi0.5 Na0.5 TiO3 –BaTiO3 –SrTiO3 and Bi0.5 Na0.5 TiO3 –BaTiO3 – Bi(Zn0.5 Ti0.5 )TiO3 systems [32,33]. A thorough inspection of the position of (200) peak marked by the dashed line in Fig. 2(b), indicates a slight shift toward lower angles with increasing SBT content, which is an indication of lattice expansion. This is due to the relatively small ionic radii of Na+ compared with Sr2+ in the A site (CN = 12, RNa + = 1.39 Å < RSr 2+ = 1.44 Å) [34]. Fig. 3 shows the grain morphology of the polished and thermal etched samples with x = 0, 0.08, 0.16 and 0.24 sintered at their optimum temperature, respectively. The high density and well–developed grains with pore–free microstructure are obtained via the SEM examination. The additions of SBT content have not significantly affected the grain growth. Using a linear intercept method, the grain size is in the vicinity of 1.2–1.4 ␮m. 3.2. Raman spectra, dielectric properties and P–E hysteresis loops Fig. 4(a) and (b) shows the room–temperature Raman spectra of unpoled and poled BNT–BT–xSBT (0 ≤ x ≤ 0.24) ceramics. The deconvolution of six Lorentzian peaks is presented to better determine the locations and variations of all the Raman modes. Three Raman vibration modes can be assigned to A–site (∼135 cm−1 ),

Ti O bond (200–400 cm−1 , Band A and B) and TiO6 octahedra (500–600 cm−1 , Band C and D) in the range of 100–700 cm−1 , respectively, as clearly shown in Fig. 4(c) (poled sample) [35–37]. Broad and diffused Raman modes are observed for all compositions as a result of the disorder induced by random occupancy of cations Na/Bi/Sr/Ba in the A site [38,39]. On the one hand, modes around 230–320 cm−1 of Ti O bond and 540–620 cm−1 of TiO6 octahedra vibrations become narrower for poled ceramics compared to the unpoled samples with x = 0, 0.02 and 0.04, which indicates the rising polarity of the unit cells after poling. However, when the SBT content is increased to x = 0.08, no obvious change occurs for poled and unpoled ceramics. On the other hand, as shown by dashed lines in Fig. 4(b), modes of Ti O bond and TiO6 octahedra vibrations gradually separate from each other with increasing SBT content. In order to quantify the modes variation, the peak position of the A, B, C and D bands, as well as their corresponding width between band A and B (WB-A ), band C and D (WD-C ) as a function of SBT content are plotted in Fig. 4(d) and (e). Band B generally shifts towards high wavenumbers, while band A shifts to low wavenumbers, enlarging the WB-A which indicates an increasing octahedral distortion. As a matter of fact, the Ti O vibration is related to the dynamics of PNRs and the steadily widened WB-A illustrates the increasing relaxor characteristics [40]. The variation of the modes near 530 cm−1 and 620 cm−1 associated with the TiO6 vibrations is analogous to the discussion above, the increasing divergence between Band C and D indicates the decrease in unit–cell anisotropy. Additionally, the decreased intensity of the Raman spectrum and the sudden increase in WD-C when x is increased to 0.08 indicates the ferroelectric-to-relaxor phase transformation for poled BNT–BT–xSBT ceramics. The temperature dependence of dielectric constant and loss tangent for poled BNT–BT–xSBT (0 ≤ x ≤ 0.24) ceramics at 1 kHz is summarized in Fig. 5(a) (for simplicity, the loss tangent curves for x = 0.08–0.24 are not included since they just present just a shoulder instead of peak). The dielectric constant of all samples exhibits a shoulder at Ts and a broad dielectric maximum at Tm . The dielectric shoulder at Ts is ascribed to the thermal evolution of

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Fig. 4. (a) and (b) Raman spectra for unpoled and poled BNT–BT–xSBT ceramics (0 ≤ x ≤ 0.24); (c) the deconvolution of Raman spectrum fitted according to Lorentzian peak functions; (d) and (e) peak position of A, B, C and D bands and the evolution of width between band A and B (WB-A ), band C and D (WD-C ) as a function of SBT content. Table 1 The variation of Tm , TFR (Tf ) as a function of SBT content. x ◦

Tm ( C) TFR /(Tf ) (◦ C)

0

0.02

0.04

0.08

0.12

0.16

0.20

0.24

282 101

283 70

288 46

283 (31)

283 /

278 /

277 /

272 /

PNRs with R3c and P4bm symmetry, which is correlated with the relaxation of PNRs in the rhombohedral phase, and Tm is associated with the tetragonal PNRs emerging from the rhombohedral PNRs [41]. Additionally, another dielectric anomaly at TFR that is independent of frequency is detected for x = 0–0.04, which also corresponds to the sharp loss tangent curves, as pointed out by the dashed line in Fig. 5(a). Besides, contrarily to nearly invariant Tm values, the field–induced dielectric anomaly TFR steadily decreases from 101, 70 and 46 ◦ C for x = 0, 0.02 and 0.04, respectively, and can not be detected above the room temperature for x ≥ 0.08, as shown in Fig. 5(a). Another important parameter Tf –defined as the onset temperature of the emergence of nonergodic PNRs from the ergodic matrix during cooling, can   be obtained using the Vogel–Fulcher (VF) EA s −Tf )

law: f = f0 exp − k(T

[42]. Ts instead of Tm should be used in

VF laws since Ts is related to the polar rhombohedral PNRs [41]. Besides, Ts is determined by a close–up “shoulder” measured at various frequencies from the loss tangent curves due to the broad and diffused dielectric curves. Fig. 5(b) shows the temperature dependence of dielectric constant and loss tangent upon cooling for x = 0.08 at various frequencies, the VF fitting of the Ts as a function of various frequencies is shown in Fig. 5(c). The Tf of x = 0.08 is fitted to be 31 ◦ C, which is in the vicinity of room temperature. The variation of Tm , TFR (Tf ) as a function of SBT content is listed in Table 1.

Fig. 6(a)–(h) depicts the P–E hysteresis loops together with corresponding J–E curves for x = 0–0.24 at room temperature. Saturated P–E loops accompanied by a single sharp J–E peaks J1 can be observed for x = 0 and 0.02, indicating that nonergodic phase is dominant in these samples. By comparison, when the SBT content reaches to x = 0.04, two J–E peaks denoted as J1 and J2 are found, which can be attributed to the coexistence of nonergodic and ergodic phase. Recently, based on in–situ TEM observation, Guo et al. verified that peak J2 corresponded to the ferroelectric–relaxor transition with the disruption of ferroelectric domains into nanodomains, while peak J1 marked the relaxor–ferroelectric transition with the coalescence and growth of nanodomains into large lamellar domains [43]. The appearance of the J1 and J2 is the sign of the coexistence of ergodic and nonergodic relaxor phase [44–46]. With further addition of SBT content to x = 0.08, the increase in fraction of ergodic phase accompanied by the enhanced local random field make the J2 peaks appear in the unloading branch of the J–E loops alongside a sudden drop of remanent polarization Pr . Finally, J–E curves without the presence of clear peaks are observed for x = 0.20 and 0.24 composition. Fig. 6(i) shows the electric field values E1 and E2 corresponding polarization current peaks J1 and J2 as a function of SBT content. Particularly, the value of E2 is approximately zero for x = 0.08, as shown by the arrow, suggesting that a specific proportion of non-

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Fig. 5. (a) Temperature dependence of dielectric constant and loss tangent for poled BNT–BT–xSBT (0 ≤ x ≤ 0.24) ceramics at 1 kHz upon heating; the inset shows the locally magnified dielectric curves for x = 0.04; (b) temperature dependence of dielectric constant for x = 0.08 at 1–100 kHz upon cooling; the inset shows the corresponding loss tangent curves, and Ts is determined by a close–up “shoulder” measured at various frequencies. (c) Vogel–Fulcher fitting of the Ts as a function of various frequencies for x = 0.08.

Fig. 6. (a–h) P–E hysteresis loops together with J–E curves measured for BNT–BT–xSBT (x = 0–0.24) ceramics; (i) the electric field E1 and E2 for the corresponding current peaks J1 and J2 as a function of SBT content.

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Fig. 7. (a) The temperature change signal versus time measured directly for poled BNT–BT–xSBT (x = 0–0.24) under the electric field of E = 6 kV/mm; (b) direct electrocaloric temperature change T under various electric field strength E = 1–6 kV/mm as a function of SBT content.

ergodic and ergodic phases can adjust the Tf (31 ◦ C) close to the room temperature, as discussed above. 3.3. Electrocaloric effect Fig. 7(a) shows the temperature change signal versus time measured directly for poled BNT–BT–xSBT (x = 0–0.24) under an electric field of E = 6 kV/mm, where the appearance of exothermic peaks is due to the application of the electric field while endothermic peaks refer to the field removal. This phenomenon represents the typical characteristic of positive ECE. Recently, although negative ECE in BNT–based systems has been reported, the direct measurement can be a more powerful tool to unravel the nature of ECE in BNT–based systems [22,23,47,48]. The direct electrocaloric temperature change T under various electric field strength E = 1–6 kV/mm as a function of SBT content is plotted in Fig. 7(b). Interestingly, in the low electric field range E = 1–3 kV/mm, the maximum T is located at x = 0.04 increases from T = 0.04 at 1 kV/mm, to 0.07 at 2 kV/mm and finally to 0.14 K at 3 kV/mm. By further increasing the electric field strength to E = 4–6 kV/mm, the ECE is significantly improved and the position of maximum T moves upward to x = 0.08 composition, as shown by dashed arrow. The T equals to 0.23 (4 kV/mm), 0.33 (5 kV/mm) and 0.40 K (6 kV/mm), respectively. Fig. 8(a)–(c) shows the temperature dependence of T at various electric fields for x = 0, 0.04 and 0.08, respectively. It should be noted that at high electric field, T moves to higher temperature that is few degrees higher than TFR , which can be obviously observed in x = 0 and 0.04 compositions. When increasing SBT content to x = 0.08, the T displays a weaker dependence on temperature and exhibits stable ECE properties. The peak value of T gradually decreases with addition of SBT content, 0.92 K for x = 0, 0.63 K for x = 0.04 and 0.40 K for x = 0.08, respectively, as shown in

max −T Fig. 8(d). The instability ␩ = TT × 100% is defined to evalumax ate the stability of ECE with temperature [49]. It can be seen that the instability is lower than 20% in the whole temperature range for x = 0.08 compared to that of x = 0 and x = 0.04 which show larger instability values at temperatures lower than TFR , as displayed in Fig. 8(e).

4. Discussions The coexistence of tetragonal and rhombohedral phases for BNT–0.06BT was firstly reported by T. Takenaka [31]. Substitution of SBT content will continuously weaken the binding force of lattice together with the increasing the local random field, which is witnessed by the splitting of current density peaks in Fig. 6. At the same time, the decrease of TFR can be detected until x ≤ 0.04 which finally disappear for x ≥ 0.08, which is due to the irreversibility between nonergodic and ferroelectric phase and reversibility between ergodic and ferroelectric phase, respectively. That is to say, the nonergodic relaxor phase in x = 0–0.04 can be transformed into long–range ferroelectric order once the electric field is applied and it can not be reversed back to nonergodic relaxor phase when the electric field is released. In contrast, dominant ergodic relaxor phase in x ≥ 0.08 can be transformed into long–range ferroelectric order when the electric field is applied and it can be reversed back to its original state when the electric field is released [50]. This was confirmed by the Raman spectrum, dielectric properties and P–E loops. The easiness of polarization reorientation will contribute to the d33 value when the TFR approaches to room temperature provided that a ferroelectric order can be induced under the applied electric field. The d33 value increases from 138, 186–190 pC/N for x = 0, 0.02 and 0.04, respectively. In this composition region, the entropy change is maximized for x = 0.04. However, further addition of SBT content will drastically weakens the piezo-

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Fig. 8. (a)–(c) Temperature dependence of T at various electric fields; (d) Summary of T temperature dependence and (e) instability for BNT–BT–xSBT (x = 0, 0.04 and 0.08), respectively. [Fig. 8(a) is adapted from Ref [27]. for comparison].

electric constant d33 and remnant polarization Pr as a result of the increasing local random field. For x = 0.08, the d33 value was near zero with the drastic decease of Pr , which can be attributed to the dominance of ergodic phase (a portion of nonergodic phase still exist). Indeed, the easier polarization reorientation predicts easier changes in entropy that lead to a high ECE in the nonergodic phase region. Aside from the polarization contributions, field–induced phase transformations can provide an extra entropy contribution and result in an enhanced ECE [51]. From Fig. 7, one can note that at high electric field (E = 4–6 kV/mm), T of x = 0.08 is much higher than that of x = 0.04. Based on the above discussions, it is reasonable to speculate that the contribution of phase transformation is larger than the contribution of polarization reorientation that result in the shift of ECE peak. To disclose this interesting phenomenon, according to the thermodynamic and statistic analysis by Pirc [52], the entropy change S =

ln˝ 2 P 30 

(2)

where  is the Curie constant in the asymptotic behaviour of the linear dielectric susceptibility which reflects the dipole correlation strengths,  is the number of the polar states, ␧0 is the vacuum dielectric constant and P is the polarization which can be estimated from the hysteresis loop at saturation regime. In particular, for normal ferroelectric materials,  is 6 for tetragonal phase and  is 8 for rhombohedral phase. In relaxor ferroelectrics, the effective dipolar entities are PNRs hence the  of relaxor ferroelectrics is much larger than that of normal ferroelectrics. When the electric field is in the range of E = 1–3 kV/mm, due to the enhanced local random field for x = 0.08 and thereafter, the PNRs can hardly correlate with each other due to EJ1 = 4 kV/mm, thus leading to small saturation polarization. Under these con-

ditions, the  remains very small compared to that of normal ferroelectrics. Nevertheless, once the electric field is increased to 4 kV/mm, the electric field can overcome the local random field, and then a large number of dipole entities  is activated, providing the basis for a large ECE. Although the ECE is maximized for x = 0.04 in the composition range of x = 0–0.04, the magnitude of T is much lower than that of x = 0.08. This suggests that the phase transformation is a dominant factor in this region. With further increasing SBT content to x ≥ 0.12 (nearly pure ergodic phase), the continuously enhanced local random field will degenerate the saturate polarization and the ECE decreases. Therefore, the coexistence of ergodic and nonergodic phase for BNT–BT–xSBT provides two advantages. Firstly, BNT–based ferroelectric ceramics have a large maximum polarization Pmax (∼40 ␮C/cm2 ), the modulation of relaxor degree will not significantly deteriorate Pmax until Tf is tailored near to room temperature. Secondly, the local random field breaks up the global symmetry of the lattice, resulting in a large number of PNRs together with a large number of . This allows a large ECE to be achieved. It can be concluded that at low electric field, the polarization reorientation is a dominant factor, while at high electric fields, the phase transformation is the major contribution to obtain a large ECE. As mentioned above, T moves to higher temperature that is few degrees higher than TFR at high electric field. The Eq. (2) can be also applied to the variation of temperature–dependent ECE for x = 0 and 0.04. Jo et al. proposed that the depolarization and ferroelectric–to–relaxor transition can be identified as two separate processes consisting of the ferroelectric domains detexturing and their subsequent fragmentation into polar nanoregions (PNRs) [53]. The value of  for detextured ferroelectric domains is lower than that of PNRs. At low electric fields (1–4 kV/mm), the dynamic PNRs cannot correlate with each other to form a ferroelectric

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long–range order due to the local random field, determining the T is maximized at TFR . Once the electric field is sufficiently increased to induce the PNRs to form long–range ferroelectric order, the large number of all PNRs contributes to the T, whose maximum shifts to high temperature. A further temperature increase would continuously strengthen local random field, impede the appearance of ferroelectric order and reduce the T. The above discussion may explain the phenomena for the shift of ECE peak under higher electric field. 5. Conclusion In summary, ECE in lead–free BNT–BT–xSBT (x = 0–0.24) ceramics were measured directly using a home–made adiabatic calorimeter. A maximum adiabatic temperature change T = 0.4 K was found at x = 0.08 under an electric field of 6 kV/mm at room temperature. The coexistence of nonergodic and ergodic relaxor phase leads to a maximum change in entropy for the field–induced phase transformation between the ergodic and long–range ferroelectric phase. The shift of the ECE peak at higher electric field is mainly due to the difference in the number of the polar states  in different states. Besides, the crossover from nonergodic to ergodic relaxor phase, resulted from an increasing SBT content, determined an increased temperature stability of the ECE, with a variation lower than 20% in the temperature range from RT to 100 ◦ C for x = 0.08. This work may present a guideline when designing BNT–based ferroelectric relaxor ceramics for temperature–insensitive EC cooling technologies. Acknowledgements This work was supported by Shanghai Municipal Science and Technology Commission funded international cooperation project under No. 16520721500 and National Key R&D Program of China (2016YFA0201103). The author S. D. Li thanks the National Natural Science Foundation of China No. 11674187. References [1] X. Moya, S. Kar-Narayan, N.D. Mathur, Caloric materials near ferroic phase transitions, Nat. Mater. 13 (2014) 439–450. [2] M. Valant, Electrocaloric materials for future solid-state refrigeration technologies, Prog. Mater. Sci. 57 (2012) 980–1009. [3] C. Bechtold, C. Chluba, R.L. de Miranda, E. Quandt, High cyclic stability of the elastocaloric effect in sputtered TiNiCu shape memory films, Appl. Phys. Lett. 101 (2012) 091903. [4] S.A. Nikitin, G. Myalikgulyev, A.M. Tishin, M.P. Annaorazov, K.A. Asatryan, A.L. Tyurin, The magnetocaloric effect in Fe49Rh51 compound, Phys. Lett. A 148 (1990) 363–366. [5] A.S. Mischenko, Q. Zhang, J.F. Scott, R.W. Whatmore, N.D. Mathur, Giant electrocaloric effect in thin-film PbZr0.95 Ti0.05 O3 , Science 311 (2006) 1270–1271. [6] B. Neese, B.J. Chu, S.G. Lu, Y. Wang, E. Furman, Q.M. Zhang, Large electrocaloric effect in ferroelectric polymers near room temperature, Science 321 (2008) 821–823. [7] J.T. Li, Y. Bai, S.Q. Qin, J. Fu, R.Z. Zuo, L.J. Qiao, Direct and indirect characterization of electrocaloric effect in (Na,K)NbO3 based lead free ceramics, Appl. Phys. Lett. 109 (2016) 162902. [8] Y. Bai, D. Wei, L.J. Qiao, Control multiple electrocaloric effect peak in Pb(Mg1/3 Nb2/3 )O3 -PbTiO3 by phase composition and crystal orientation, Appl. Phys. Lett. 107 (2015) 192904. [9] F.L. Goupil, A. Berenov, A.K. Axelsson, M. Valant, N.M. Alford, Direct and indirect electrocaloric measurements on <001>-PbMg1/3 Nb2/3 O3 -30PbTiO3 single crystals, J. Appl. Phys. 111 (2012) 124109. [10] F.L. Goupil, R. McKinnon, V. Koval, G. Viola, S. Dunn, A. Berenov, H.X. Yan, N.M. Alford, Tuning the electrocaloric enhancement near the morphotropic phase boundary in lead-free ceramics, Sci. Rep. 6 (2016) 28251–28256. [11] Z.P. Xu, Z.M. Fan, X.M. Liu, X.L. Tan, Impact of phase transition sequence on the electrocaloric effect in Pb(Nb,Zr,Sn,Ti)O3 ceramics, Appl. Phys. Lett. 110 (2017) 082901. [12] Y. Zhao, X.H. Hao, Q. Zhang, A giant electrocaloric effect of a Pb0.97 La0.02 (Zr0.75 Sn0.18 Ti0.07 )O3 antiferroelectric thick film at room temperature, J. Mater. Chem. C 3 (2015) 1694–1699.

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Please cite this article in press as: F. Li, et al., Phase–composition and temperature dependence of electrocaloric effect in lead–free Bi0.5 Na0.5 TiO3 –BaTiO3 –(Sr0.7 Bi0.2 䊐0.1 )TiO3 ceramics, J Eur Ceram Soc (2017), http://dx.doi.org/10.1016/j.jeurceramsoc.2017.06.033