High-efficiency synthesis of high-performance K0.5Na0.5NbO3 ceramics

Powder Technology 346 (2019) 248–255

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High-efficiency synthesis of high-performance K0.5Na0.5NbO3 ceramics Bi Chen a,b, Pengfei Liang c, Di Wu a, Xumei Zhao a, Xiaoshuang Qiao a, Zhanhui Peng a, Lingling Wei d, Xiaolian Chao a,⁎, Zupei Yang a,⁎ a Key Laboratory for Macromolecular Science of Shaanxi Province, Shaanxi Key Laboratory for Advanced Energy Devices, School of Materials Science and Engineering, Shaanxi Normal University, Xi'an 710062, Shaanxi, PR China b School of Chemistry and Chemical Engineering, Yulin University, Yulin 719000, Shaanxi, PR China c College of Physics and Information Technology, Shaanxi Normal University, Xi'an, 710062, Shaanxi, PR China d School of Chemistry and Chemical Engineering, Shaanxi Normal University, Xi'an, 710062, Shaanxi, PR China

a r t i c l e

i n f o

Article history: Received 25 October 2018 Received in revised form 18 December 2018 Accepted 17 January 2019 Available online 20 January 2019 Keywords: (K0.5Na0.5)NbO3 ceramics Microstructure Dielectric properties Ferroelectricity Mechanochemical activation

a b s t r a c t Pure (K0.5Na0.5)NbO3 (KNN) ceramics with high density, fine and uniform-size grains were prepared by mechanochemical activation-assisted process. The time of synthesis is only 100 min, which is 72% – 93% shorter than the 6–24 h of the conventional solid-state method. Compared to samples prepared by conventional solid-state method, both the microstructure evolvement and electric properties were explored in detail. Results show the electric properties was significantly improved. Moreover, the dielectric and ferroelectric properties of obtained KNN ceramics exhibit strong dependence on the crystal size of the initial powders. The optimized ceramics HKNN100 showed a quite high energy storage performance, i.e., large electric energy storage density (Wtol = 1.612 J/cm3) and recoverable energy storage density (Wrec = 0.431 J/cm3), which can be mainly ascribed to the large dielectric breakdown strength (DBS = 110 kV/cm). Our works demonstrated that mechanochemical activation-assisted method possesses advantages for high-efficiency preparation of KNN or KNN-based ceramics. © 2019 Elsevier B.V. All rights reserved.

1. Introduction With growing concern over world ecological requirements and increasing legislative restriction on using lead and lead-based materials, lead-free (K, Na)NbO3 (KNN) became one of the most studied ferroelectric oxides with a perovskite-type crystalline structure [1,2,3] due to its outstanding electrical properties. In last decades, researches of KNNbased ceramics were focused mostly on how to enhance the electric properties by optimizing sintering condition [4–8] or constructing phase boundaries [9] using different additives, such as equipollent metals substitution at the A/B cites [10–14], and aliovalent metals doping at the Nb site [15–16]. Nevertheless, little attention has been paid to how to improve the time-costing preparation process itself. Solid-state reaction method, which employs high-temperature calcination, was usually used to prepared KNN ceramics [1,17]; however, it has obvious shortcomings such as low diffusion rate through a product layer, loose contacts between the particles of components, non-uniform particle size distribution, long ball milling time [1–2] and high volatility of the constituent alkali elements [11,18]. Therefore, it is urgent to explore an improved preparation method. Some new methods were tried for above-mentioned purposes, i.e., self-propagating high⁎ Corresponding authors at: School of Materials Science and Engineering, Shaanxi Normal University, Xi'an 710062, Shaanxi, PR China. E-mail addresses: [email protected] (X. Chao), [email protected] (Z. Yang).

https://doi.org/10.1016/j.powtec.2019.01.039 0032-5910/© 2019 Elsevier B.V. All rights reserved.

temperature synthesis [19–20], shock waves [21], mechanochemical activation [22–24], etc. Among them, mechanochemical activation has drawn intensive attentions since it shows obvious advantages in the aspects of shorter synthesized time, less volatilization of alkaline species, better chemical homogeneity [22,25], more refined crystals of synthesized powder [26–30]. As the key procedure of mechanochemical activation, the ball milling process via repeated welding and fracturing technique [31–32] cause the generation of defects in solids which accelerate the migration of defects in the bulk, increase the number of contacts between particles, and help refining the grains [33,34]. Meanwhile, the high energy ball milling process also induces cumulative kinetic energy to overcome the energy barrier of activating the reaction [24,35], decreased the calcination temperature to restrain the volatilization of alkaline species. The cumulative kinetic energy is relate to ball-impact energy, ΔEb∗, the ball impact frequency, νt, the milling time, t, and the powder weight, mp, as shown in Eq. (1) [36]: Ecum ¼

ΔEb ν t t mp

ð1Þ

It has been revealed the Ecum is proportional to the milling time t when the other parameters are invariant. As reported [37], the milling time and speed in a high energy ball-milling process have a great influence on the obtained particle size, which is eventually reflected on the microstructure and electric properties of the ceramics. Moreover, It

B. Chen et al. / Powder Technology 346 (2019) 248–255

was reported that the refined powders after ball milling processes need lower calcined temperature [38], and exhibited fewer volatility of the constituent alkali elements [22,39,40,41]. In this paper, we reported that (K0.5Na0.5)NbO3 (KNN) can be fabricated via a mechanochemicalassisted activation route with much shorter milling time than literatures. We further explored the microstructure evolution of initial powder, calcined powder and KNN ceramics with different milling time. At last, the dielectric and ferroelectric properties of the samples at different milling time was investigated systematically, calculated and compared energy storage density W of KNN ceramics which prepared by mechanochemical activation-assisted method or conventional solid-state reaction.

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uncalcined powders and calcined KNN powders (850 °C for C24h; 700 °C and 850 °C for H20min, H100min, H360min, respectively.) were investigated using a X-ray diffractometer (XRD, Philips, Eindhoven, the Netherlands) at a scanning rate of 10°/min in the range of 20–60°, using Cu-Kα radiation (λ = 1.5406 Å) at 40 kV and 50 mA. Raman spectrophotometer (Raman; in Via Reflex, Renishaw, UK) was performed in the range 0–1000 nm. The Surface and Section morphologies of the KNN samples were observed using Field Scanning Electron Microscopy (FE-SEM; SU-8020, Hitachi, Japan). Dielectric measurements of the samples were carried out at a temperature range of 25–500 °C using an LCR meter (TH2818; Tonghui, Changzhou, China) at 1 kHz–200 kHz. Ferroelectric properties were measured by a ferroelectric analyzer (TF-2000; Aix ACCT, Aachen, Germany).

2. Experiments 3. Results and discussions (K0.5Na0.5)NbO3 powders were synthesized using mechanochemical activation-assisted route. Powders of K2CO3, Na2CO3 and Nb2O5 (with a purity is 99.99%, Sinopharm Chemical Reagent co., Ltd) were used as raw materials and weighed according to the stoichiometry. The powders were then blended in hyperpure water in air with a Zirconia vial/ ball set, of the ball-to-powder weight ratio 8:1. The mixing process was conducted in a Fritsch Vario-Planetary Mill (pulverisette P7™), in which the sun wheel and the grinding jar rotate in the opposite directions with speed ratio 1:2. The high-energy ball milling was set that the rotational directions of the sun wheel and jar reverse every 2 min with a rest interval of 8 min to avoid excessive heating. The rotation speed was 300/600 rpm for 20–360 min. The slurry was dried in a freeze dryer and calcined at 700 °C or 850 °C for 9 h to obtain the precursor of KNN powders. Afterwards, the powders were then ball-milled again, followed by drying and cooling. The obtained products were pressed as a binder followed by cold isostatic pressing into pellets with a diameter of 10 mm at 200 MPa. The pellets were sintered at 1120–1140 °C for 2 h in a sealed alumina crucible. Differentiated by milling time, the obtained ceramics with milling time of 100 min and 360 min were labeled as HKNN100, HKNN360, respectively. In comparison, the samples prepared by conventional solid-state reaction process with milling 24 h were labeled as CKNN. Particle size distribution was analyzed using a laser particle size analyzer (Model BI-90Plus, Brookhaven, USA). And the morphologies of the particles were exhibited using Field Transmission Electron Microscopy (FTEM; Tecnai G2 F20, FEI, Netherlands). Phase structure of the

Fig. 1 shown the distributions of particle sizes versus milling time were investigated. Form Fig.1, it can be seen the mean particle size of the raw materials was 1364 nm. After milling for 20 min, the mean particle size was promptly decreased to 143 nm. As the milling time increased, the particle size distribution became narrower, while the mean particle size stayed almost unchanged, with average diameters of 113, 94, 118, 125 and 122 nm, respectively. When extending the milling time to 240 min and 360 min, the mean particle size slightly went up to 200 nm and 171 nm. These results revealed that high-energy ball mill method can reduce the particle size to submicron meter effectively, and that long milling time can make particle agglomeration again. Fig. 2 revealed that TEM images of the powders obtained by according milling times. From Fig. 2(a–e) we noticed that the irregular particles were first sliced into several regular ones then reunite again together with prolonged milling time. Furthermore, compared with particle by conventional ball milling 24 h (marked as C24, Fig. 2(f)) , we can get the similar particle morphology of high-energy ball milling for only 20 min (marked as H20, Fig. 2(a)). Fig. 3 shows XRD patterns of the powders after different milling time. Fig. 3(a1, b1, c1 and d1) shows XRD patterns of the uncalcined powders of H20min, H100min and H360min. It is shown that the characterized peaks of Nb2O5, Na2CO3 and other synthetic products K3Nb7O19 and K3NbO4 (Eq. (2) and (3)) by standard PDF card(#27-1003, #28-0317, #37-0451, #38-1499, #261325, #30-0964), that implied the mechanochemical reaction has been occurred at the high-energy ball milling

Fig. 1. Lognormal distribution of particle sizes by variation of high-energy ball milling times(a-h), and its corresponding of mean diameters (i).

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Fig. 2. TEM images of the powders obtained by different milling times, (a-e) milling 20 min, 60 min, 100 min, 240 min, 360 min by high-energy ball mill process, respectively (marked as H20, H60, H100, H240, H360, respectively); (f) milling 24 h by conventional method (marked as C24).

process. But, it isn't detected the peak of K0.5Na0.5NbO3 that means no KNN phase formed, although it has been proved mechanical energy capable of providing the thermal energy needed to start a reaction and synthesize products directly [35]. In addition, we can seen prolonging the milling time to 360 min, it is observed amorphization character of the reactants, which is consistent with literature [18]. Generally, the synthesis temperature of KNN is at 700–950 °C40, among this range, 850 °C is the favorable temperature based on the literature results [11,40,42,43,44], which avoiding to formed the liquid Krich phase, increasing the density of KNN ceramics via a solid state sintering process, and then, no occurring abnormal grain growth [42]. As a contrast, all the powders were calcined at 700 and 850 °C, respectively. After 700 °C calcined, the powder of H100min obtain the single KNN phase. Other materials display a typical orthorhombic symmetry at room temperature [27], that means KNN becoming predominant

phase. And after 850 °C calcined, the powder of C24 obtain the single KNN phase which is consistent with the literature results [11,40,42,43,44]. However, H20min produced a distinct second phase at 700 or 850 °C, which may be due to a further increased temperature resulting in by-products generated [Eq. (3) and (4)]. H100min and H360min also traced to a slight impure phase at 850 °C[Eq. (3)]. All the possible reactions are shown in the following chemical Eqs. (2), (3) and (4). yields

3K 2 CO3 þ 7Nb2 O5 → 2K 3 Nb7 O19 þ 3CO2 ↑

ð2Þ

yields

ð3Þ

yields

ð4Þ

3K 2 CO3 þ Nb2 O5 → 2K 3 NbO4 þ 3CO2 ↑ K 2 CO3 þ 3Nb2 O5 → 2KNb3 O8 þ CO2 ↑

Fig. 3. XRD patterns of the powders obtained by different milling times (Group a for milling 24 h with conventional method; Group b-d for milling 20 min, 100 min and 360 min with high energy ball milling, respectively; Group 1 for un-calcining; Group 2, 3 for calcining at 700 and 850 °C, respectively).

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Fig. 4. XRD patterns of CKNN, HKNN100, HKNN360 ceramics (a) and its enlargement at 44–47°(b).

The summary reaction of alkali carbonates and niobium oxide is described by Eq. (5): yields

K 2 CO3 þ Na2 CO3 þ 2Nb2 O5 → 4K 0:5 Na0:5 NbO3 þ 2CO2 ↑

ð5Þ

Moreover, the milling time has a important effect on single KNN phase synthesis. From Fig.3(b2 and c2), a trace of second phase are observed in the powders of H20min at the 2θ range 20° to 60°, and tiny traces of impurities are showed in the powders of H360min (Fig. 3(d2 and d3)). This indicates that the short milling time cause insufficient contacting of the raw material, and the prolong milling time cause impurities. Therefore, 100 min is the best time for milling. In summary, the powder of H100min calcined at 700 °C can obtained a single phase KNN, which means high-energy ball milling process not only enhanced the efficiency of milling process, but also decreased the calcined temperature from 850 °C to 700 °C. Fig. 4(a) shows the XRD patterns of CKNN, HKNN100 and HKNN360 ceramics (The powder of H20min will not be further discussed since it didn't make pure KNN phase). Different from the precursor powder, it can be seen from split (200) peak about 2θ = 45° that all samples are (220)o and (002)o [42]. The ratio of relative intensities of (220)o and (002)o for all samples are near 2:1, as shown in Fig. 4(b). It is implied a typical orthorhombic structure for all samples. Moreover, the (220)o and (002)o diffraction peaks are found to move toward higher angles with increasing milling time, suggesting that there is a slight lattice

251

shrinkage [45]. This is probably because small amount of undesirable impurities were introduced in the high-energy milling process, which is confirmed in Fig. 3(d2, c3 and d3). These impurities also resulted in a low symmetry orthorhombic (200) phase for HKNN100 and HKNN360. Some research results confirmed the low symmetric orthorhombic phase is favorable for electrical properties [16]. Meanwhile, the diffraction peaks of HKNN100 and HKNN360 were greatly broadened and reduced, indicating great refinement in grain size. To explore the effect of high energy ball milling time on the phase transition of the powders, the Raman spectrum was performed for the uncalcined powders and 700 °C calcined powders which obtained by conventional methods (milling for 24 h) or high energy ball milling process (milling for 20 min, 60 min, 100 min, 240 min, 360 min, respectively), respectively, as shown in Fig. 5(a) and Fig. 5(b). Commonly, the vibrational modes of isolated cations and coordination polyhedrons [46–47] were used to described the Raman spectra. It is generally believed that the translational modes of K+/Na+ cations and the internal modes of NbO6 octahedrons contribute to the vibrations [46]. In that case, the octahedron which in Oh symmetry was consist of 6 internal modes such as A1g(ν1) + Eg(ν2) + 2F1u(ν3,ν4) + F2g(ν5) + F2u(ν6) modes. Among these modes, the A1g(ν1), Eg(ν2), and F1u(ν3) modes are stretching modes and the F1u(ν4), F2g(ν5), F2u(ν6) are bending modes [47]. For the uncalcined powders (Fig. 5(a)), the Raman spectrum doesn't show typical vibrations corresponding to a perovskite phase. The peaks were identified as internal vibrational ν5 (about 228 cm−1), ν4 (about 312 cm−1), ν2 (about 645 cm−1), ν1 (about 687 cm−1), ν3 (about 814 cm−1) and ν1 + ν5 (about 899 cm−1) modes of pseudo NbO6 octahedra. It is noted that the peaks intensity of ν1 and ν5 become weaker, but the coupling peaks intensity of ν1 + ν5 become stronger with the ball milling times increased from 20 min to 360 min. Generally, only A1g(ν1), Eg(ν2), and F2g(ν5) modes were Raman activated among the internal modes of NbO6 octahedron. But we can see from the Fig. 5(a), the uncalcined powders, whether it is conventional ball milling or high-energy ball milling, it doesn't form the symmetry of NbO6 octahedron. At the same time, the corresponding change was the full width at half maximum (FWHM) for ν1 gradually decreased from H20min to H360min, although they all higher than the FWHM (ν1) of peak C24h. Fig. 5(b) shows the milling times dependent Raman spectra were collected for the 700 °C calcined powders. The Raman spectrum shows

Fig. 5. Raman patterns of the powders obtained by different milling times; a, b for uncalcined and 700 °C calcined for 5 h at milling 20 min, 100 min and 360 min, respectively; c for corresponding to the FWHM (ν1) of a, b; d for Raman patterns of CKNN, HKNN100, HKNN360 ceramics; e for fitting curves of c on 450–750 cm−1; f for FWHM (ν1) of e.

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Fig. 6. FE-SEM micrographs of thermally etched surface of CKNN(a1), HKNN100(b1), HKNN360(c1) ceramics; cross-sectional microstructures of CKNN(a2), HKNN100(b2), HKNN360(c2) ceramics.

typical vibrations corresponding to a perovskite phase. The peak in the region lower than 200 cm−1 can be assigned to the translational modes of Na+/K+ cations and rotational mode of the NbO6 octahedra [46,48]. The ν6 mode of NbO6 octahedra become Raman activated and also appear at this region, but compared with peak of the powder C24h, the peaks intensity of the powder by different high-energy ball milling times (H20min-H360min) was weaker with increased the ball milling times from 20 min to 360 min. Compared with the uncalcined powders, there are three main changes occurred at ν5, ν1 and ν4, ν3, corresponding to the peak of ν5 shift from 225 to 258 cm−1, ν1 shift from 689 to 611 cm−1, and ν4, ν3 move and merge with ν5, ν1. This may be caused by increasing in binding strength that due to the shortening of the distance between Nb5+ and its coordinated oxygens and distortion

of O-Nb-O angles [47,49]. It is also demonstrated the formation of the O phase or O-T phase after calcined at 700 °C. The observation seem to agree well with the evidence of XRD. Furthermore, the FWHM (ν1) also demonstrates a influences in different milling duration and whether to calcine, as shown in Fig. 5(c). The results shown that FWHM (ν1) of 700 °C calcined powders far lower than the uncalcined powders which illustrated the strength of the chemical bonds are more consistence and the material are more uniformity. Similar to 700 °C calcined powders, the Raman spectrum of the ceramics of CKNN, HKNN100, HKNN360 also shows typical vibrations corresponding to a perovskite phase (Fig. 5(d)). In order to discuss the differences of the ceramics, the spectrum was fitted into the 2 peaks at 553 cm−1 (ν2) and 612 cm−1 (ν1), respectively (Fig. 5(e)). The A1g

Fig. 7. Temperature dependent dielectric properties of CKNN, HKNN100 and HKNN360 ceramics at various frequencies (a, b, c); The dependences of dielectric constant of the CKNN, HKNN100, HKNN360 at frequencies 100 kHz(d).

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Fig. 8. P-E hysteresis loops at their critical breakdown strength for the ceramics (a) CKNN, (b) HKNN100 and (c) HKNN360 at various electric field; (d) Ps, Pr and Ec of the CKNN, HKNN100, HKNN360 ceramics at electric filed 50 kV/mm.

(ν1) stretching mode of HKNN360 is stronger than HKNN100 and CKNN, and all FWHMs (ν1) of peak less than 38 cm−1, even FWHM (ν1) of CKNN is as low as 28.5 cm−1 (Fig. 5(f)). To check the microstructural dependence of obtained KNN ceramics on milling time, the SEM micrographs of the surface and cross-sectional microstructure for the sample CKNN, HKNN100, HKNN360 were captured, as shown in Fig. 6. It can be observed that the grains in all samples exhibit cubic-like shape, which is a typical character of KNN ceramics [50,22]. The grains of CKNN ceramics are inhomogeneous and loosely arranged, and pores can be found. In comparison, HKNN100 and HKNN360 exhibit obviously smaller grains, indicating the efficacy of high energy-ball milling process in reducing the grain size. Statistics of grain size using nano measurer software and measurements of bulk density by Archimedes method with distilled water [51] were illustrated in Fig. 6(a2, b2 and c2). As well known, smaller particle size of starting powder is beneficial for high densified ceramics with fine grained structure [52], which explained the higher bulk density of HKNN100 and HKNN360 than that of CKNN. The reduction of density from 4.326 g/cm3 of HKNN100 to 4.308 g/cm3 of HKNN360 could be ascribed to the poor sinterability of these agglomerated powders of the latter, which is also indicated by the pores found in thermally etched surfaces and cross-sectional of HKNN360 ceramics. In addition, the presence of impurity phases in HKNN360 induced by the ball milling should also be taken into account [53]. Since porosity could provide electrical conduction path, it shall be then responsible for the degraded electrical breakdown strength [54], as shown in Fig. 8(c). Fig. 7(a,b and c) shows temperature dependent dielectric properties of CKNN, HKNN100 and HKNN360 ceramics at frequency ranging from 1 to 200 kHz. All samples undergo two phase transitions, from orthorhombic to tetragonal (TO-T) and from ferroelectric tetragonal to paraelectric cubic (TC). The phase transition temperatures (197 °C for TO-T and 419 °C for TC) observed in our CKNN ceramic agree well with the previously reported values for KNN ceramics [6,12,55–57]. In comparison, the two dielectric abnormal peaks are located around 195 °C and 371 °C for HKNN100 and HKNN360, implying that highenergy ball milling process can reduce the phase transition temperatures TO-T and TC. This phenomenon in turn indicates that the particle size of the calcined powders has a crucial influence on the dielectric

behaviors of the ceramics, as also reported in other ceramic systems [22,27]. Comparing the dielectric loss of the CKNN with HKNN100 and HKNN360, we can see that HKNN100 possesses the best thermal stability and lowest dielectric loss, which correspond to its highest mass density and least porosity. Room temperature P-E hysteresis loops of CKNN, HKNN100, HKNN360 ceramics measured at 1 kHz are shown in Fig. 8. All the samples were measured at their critical breakdown strength until broken down. From Fig. 9(a), it can be seen that CKNN ceramics exhibit square-shaped P-E loops with a DBS 50 kV/cm, which is a characteristic feature of normal ferroelectric materials [8,58]. Fig. 8(b) showed that the P-E hysteresis loops of HKNN100 ceramics transform gradually from square shape to slim shape as electrical field went to DBS at 110 kV/cm. For HKNN360 (Fig. 8(c)), the P-E hysteresis loops became fatshaped with a DBS 55 kV/cm. In general, the DBS depends on both internal factors such as porosity, grain size, and the generated second phase, and external factors such as sample thickness, sample area, and electrode configuration [59]. Among these factors, the most dominant ones are the grain size and porosity [60]. Despite of the small grain size, HKNN360 didn't exhibit decent DBS since numerous porosity were produced due to particle agglomeration in precursor powders

Fig. 9. Calculated energy storage density W of HKNN100 ceramics.

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Table 1 Comparison of various parameter for reported KNN ceramics synthesized by different techniques. Samples

Hot pressing

Mechanochemical activation

L. Egerton [66]

Conventional synthesis Hansu Birol [67]

Kepi Chen [42]

◆This work

R. E. Jaeger [68]

Rajan Singh [27]

Chongtham Jiten [22]

★ This work

Milling time (h)

/

24

6

24

/

32

20

Calcine temperature (°C) Grain size (nm) Bulk density (g/cm3) εr (1 kHz) Tc (°C) d33 (pC/N) DBS (kV/cm) Wtol (J/cm3) Wrec(J/cm3)

850 / 4.25 290 420 80 / / /

825 ~6000 4.30 ~480 400 110 90 / /

800 ~650 / / / 128.3 / 0.344 0.039

850 560 4.275 560 419 80 50 / /

/ / 4.46 ~340 420 160 / / /

550 ~105 / 680 365 95 45 / /

900 ~148 / 792 ~399 114 30 / /

~1.67 (100 min) 700 280 4.326 731 371 78 110 1.612 0.431

[23], which resulted in low bulk density and high leakage conductance at high electric fields. Remanent polarization Pr, saturated polarization Ps and coercive field Ec of the CKNN, HKNN100, HKNN360 ceramics were acquired from their P-E loops at electric field of 50 kV/cm, and shown in Fig. 8(d). The CKNN ceramics showed a Pr value ~9.14 μC/cm2 and a Ec value ~13.8 kV/cm2, respectively. Generally, high remanent polarization is related to the high internal polarizability, strain, electromechanical coupling, and electro-optic activity [54], while the coercive field Ec is related to the grain size of the ceramics, i.e., higher Ec usually represent smaller grain size and vice versa. HKNN100 and HKNN360 samples have improved Pr and Ec values, which is consistent with our aforementioned analysis of Fig. 6. In particular, we calculated the energy storage performance of HKNN100 ceramics, according to the P-E hysteresis loops, the results as shown in Fig. 9. The electric energy storage density Wtol, recoverable energy density Wrec, and loss energy density Wloss were calculated using the following equations [58,61–65]: Z W tol ¼

EdP 0≤E≤Emax 0

Z W rec ¼

P max

Ps

EdP Pr

W loss ¼ W−W rec

ð6Þ

ð7Þ ð8Þ

The obtained Wtol, Wrec and Wloss are 1.612 J/cm3, Wrec 0.431 J/cm3 and Wloss 1.181 J/cm3, individually. Furthermore, to comprehensively evaluate the efficiency of the preparation process and the performances of obtained K0.5Na0.5NbO3 ceramics in this work, we made a systematic comparison of both synthesis parameter and sample performance between literatures and our mechanochemical activation synthesis, as shown in Table 1. One can clearly see that the milling time of this work (HKNN100) was greatly shortened (100 min), no matter compared with conventional synthesis (6–24 h) or other high energy ball milling processes (20 −32h). Furthermore, the electrical performance of our HKNN100 sample is also better than what was reported in the references [22,27,66,17]. 4. Conclusions Single-phase and fine-crystalline KNN ceramics were successfully synthesized by the mechanochemical activation-assisted method. As compared to the conventional solid-state method and reported mechanochemical activation-assisted method, this work not only significantly reduced the synthesis time but also largely improved the electrical performance. The time of synthesis is only 100 min, which is 0.3% – 27.8% of the milling time of the conventional method. Eventually, high density (4.326 g/cm3) and submicron sized grains (0.28 μm) were

simultaneously realized in our HKNN100 ceramics, which exhibits large DBS (110 kV/cm), Wtol (1.612 J/cm3) and Wrec (0.431 J/cm3) values. Our results suggest mechanochemical activation-assisted synthesis method is a promising strategy for high-efficiency preparation of KNN-based ceramics.

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