Piezoelectric and pyroelectric properties of Mn-doped 0.36Pb(In1/2Nb1/2)O3–0.36Pb(Mg1/3Nb2/3)O3–0.28PbTiO3 ceramics

0.5 mol% Mn-doped 0.36Pb(In1/2Nb1/2)O3–0.36Pb(Mg1/3Nb2/3)O3–0.28PbTiO3 (Mn-PIMNT) ceramics were fabricated by a two-step precursor method. The phase structure, morphologies, and temperature-dependent dielectric, ferroelectric, piezoelectric, and pyroelectric properties were studied. The results indicated that the Mn-PIMNT ceramics had pure perovskite phase and uniform grain distribution. Meanwhile, it exhibited ultrahigh piezoelectric coefficient d33 of 235 pC/N, high-power figure of merit (FOM) of 60,160 pC/N, large remnant polarization Pr of 34.57 µC cm−2 and coercive field EC of 12.97 kV cm−1, which were much better than that of binary Mn-PMNT ceramics. Moreover, by Mn ions-doping, giant pyroelectric coefficient p up to 4.8 × 10–4 C m−2 K−1 was obtained and the figure of merits for the detectivity Fd reached 2.317 × 10–5 Pa−1/2, higher than PIMNT ceramics. Combined with the outstanding piezoelectric and pyroelectric properties as well as high ferroelectric rhombohedral to tetragonal phase transition temperature Trt (up to 146 ℃) and ferroelectric tetragonal to cubic phase transition TC of 188 ℃, it is suggested that the Mn-PIMNT ceramics are an excellent candidate for piezoelectric and pyroelectric devices.

To expand the operating temperature range, the second-generation relaxor ferroelectrics, represented by yPIN-(1-x−y)PMNxPT (PIMNT) having higher T C than PMNT and PZNT systems, were developed. Nevertheless, the global piezoelectric and pyroelectric performances were inferior [7]. In comparison, the following third-generation relaxor ferroelectrics represented by Mn-doped PIMNT were further reported which simultaneously possess much better temperature stability over a wide temperature range and higher performance [11,12]. The Mn ions doping in PIMNT are Mn 2+ or Mn 3+ , which can be called a "hard dopant". They enter the crystal lattice and replace the high-valent cations of B-site (for example Ti 4+ , Mg 2+ , Nb 5+ , and In 3+ ions) [13]. As a result, oxygen vacancies are created to keep the system electrically neutral. The defect dipoles formed by the oxygen vacancies can decrease the mobility of charge carriers on the domain wall and disrupt the stability of the ferroelectric domain [14,15]. This led to a substantially decreased dielectric and mechanical loss, and meanwhile enhanced piezoelectric and pyroelectric properties [12,13,[15][16][17][18].
In previous works, it can be noted that the pyroelectric investigations on the third-generation relaxor ferroelectrics were mainly focused on the single-crystal form, few works reported on the pyroelectric properties of Mn-PIMNT ceramic systems which are much easier to be commercialized with low cost and better mechanical processability [12]. Under these considerations, the purpose of this article is to explore the third-generation Mn-doped PIMNT ceramics having good piezoelectric and pyroelectric performance as well as high T rt and T C for piezoelectric and pyroelectric devices. We focus on studying the phase and domain structures, dielectric relaxation behavior, ferroelectric, piezoelectric, and pyroelectric properties together with strain behavior. The results indicated that the obtained Mn-doped PIMNT ceramics had ultrahigh piezoelectric and pyroelectric performance together with good thermal stability.

Experimental procedures
ceramics were prepared by a two-step precursor method. All the reagent grade raw materials of In 2 O 3 (99%), Nb 2 O 5 (99.99%), PbO (99%), TiO 2 (98%), and MnO 2 (97.5%) were commercially supplied from Sinopharm Chemical Reagent Co., Ltd (Shanghai, China). MgO was obtained by heating 4MgCO 3 ·Mg(OH) 2 ·5H 2 O at 1000 °C for 2 h. All the raw materials were weighed according to the formula. During the fabrication process, In 2 O 3 and Nb 2 O 5 powders with a molar ratio of 1:1 were first ball-milled with alcohol for 6 h and then calcined at 1100℃ for 6 h to prepare InNbO 4 . MgNb 2 O 6 was synthesized by MgO and Nb 2 O 5 powders with a molar ratio of 1:1 at 1000 ℃ for 4 h. Then the powders of PbO, MgNb 2 O 6 , InNbO 4 , TiO 2 , and MnO 2 were ball-milled again with alcohol for 6 h, then dried and calcined at 850 ℃ for 2 h to form pure perovskite phase. The powders were then sintered in the temperature range of 1225 °C to 1265 °C for 2 h. Combined with the density, shrinkage, and electrical properties of sintered Mn-PIMNT ceramics, the optimal sintering temperature was determined to be 1245 °C. The ceramics were polished and thermally etched at 900 ℃ for 2 h to release the surface stresses. Then the samples were painted with silver electrode and poled in silicone oil at room temperature for 15 min with 40 kV cm −1 .
The crystal structures of Mn-PIMNT ceramics were characterized by an X-ray diffractometer (XRD, D8-Advanced, CuKα radiation). The microstructure was observed by field emission scanning electron microscopy (FESEM, S4800, Hitachi). The domain structures were observed by a piezoresponse force microscope (PFM, MFP-3D, Asylum Research, America). The dielectric relaxation behavior along with impedance and admittance spectrum were determined using a precision impedance analyzer (HP4294A, Agilent, America). The P-E hysteresis loops and S-E behavior were examined by ferroelectric analyzer (TF2000, aixACCT, Germany). Piezoelectric coefficients (d 33 ) were directly measured by the quasi-static d 33 meter (ZJ-4AN, China). The pyroelectric coefficient (p) of the sample was obtained by the Byer-Roundy method with a Keithley 6517A electrometer [12]. Figure 1 shows the XRD patterns of Mn-PIMNT ceramics. Pure perovskite phase and sharp diffraction peaks can be observed. The radii of the B-site cations In 3+ , Mg 2+ , Nb 5+ , and Ti 4+ ions in PIMNT were 0.80 Å, 0.72 Å, 0.78 Å and 0.60 Å [19], respectively, and the radii of Mn 2+ and Mn 3+ were 0.67 Å and 0.64 Å, respectively [20,21]. The Mn ions entered the crystal lattice and replaced the high-valent cations of the B-site due to the similar radii. The inset of Fig. 1 shows the SEM image of Mn-PIMNT ceramics. From the SEM image, dense microstructure was formed and the ceramics had well-defined grains. Homogeneous grain size distribution could be observed and the average grain size was about 3.15 µm, which was calculated from the software of Nano Measurer. Figure 2 shows the PFM images of the Mn-PIMNT ceramics at room temperature with a scanning size of 8 × 8 µm 2 . Two different colors in the vertical PFM phase images represented the upward and downward polarization directions, respectively. At the grain boundaries, strip domains were observed to nucleate, which may reduce activation energy due to defects (e.g., grain boundaries) and can be observed in the inset of Fig. 1 of the SEM surface image [22]. The domain across the grain boundary may be affected by the internal residual stress and electric field  [23,24]. Furthermore, nano-sized domain structures can greatly increase the domain wall density, which is beneficial to polarization rotation and domain conversion, thereby enhancing piezoelectric response [25]. Figure 3a shows the dielectric constant ε r and dielectric loss tanδ of poled Mn-PIMNT ceramics from 25 ℃ to 250 ℃ at different frequencies. At room temperature, the values of ε r and tanδ were 790 and 0.01 at 100 Hz and after the Mn-doping, the ε r was lower than the undoped PIMNT ceramics [7,13]. With the temperature increasing, the ε r first increased slowly from room temperature to 125 ℃ and followed a sharp increase to 146 ℃. During this temperature range, the frequency dispersion was not obvious because the thermal disturbance of the dipole moment caused by the temperature did not exceed the effect of the poling field [26]. Two dielectric anomalies corresponding to the T rt = 146 ℃ and T C = 188 ℃ can be observed. The T C corresponding to ferroelectric phase to paraelectric phase transition was 63 °C higher than that of the binary Mn-PMNT single crystals [27]. The dielectric anomaly in tanδ coming from the conductivity and dielectric relaxation could also be detected [28]. The good temperature stability of the Mn-PIMNT ceramics were favorable for high-temperature ferroelectric devices over a broad temperature usage range.

Results and discussion
The modified Curie-Weiss law with the following equation was used to analysis the diffuse of the phase transition, where the ε, ε m , T m , γ, and C * are the relative permittivity, the maximum relative permittivity, the degree of diffuse, and the Curie-Weiss constant, respectively. To study the relaxation characteristics of Mn-PIMNT ceramics, ln(1/ε − 1/ε m ) and ln(T − T m ) are shown in Fig. 3b. The fitted values indicated that γ at 0.1, 1, and 10 kHz were 1.927, 1.939, and 1.997, respectively. It is known that if the fitted value of γ is 1, it stands for a normal ferroelectric system. While the fitted value of γ is 2, it represents an ideal relaxor ferroelectric system [14,16]. So, the results suggested that they had strong relaxor characteristics. The figure that γ increased with increasing frequency, which indicated the diffuse behavior of the medium depended on the frequency. The reason was attributed to the dipole deflecting could not keep up with the electric field. Moreover, the dipole will lag further with the increase of frequency. Figure 4 shows the P-E hysteresis loops of Mn-PIMNT ceramics measured at 10 Hz. At room temperature, the values of remnant polarization P r and coercive field E C were 34.57 µC cm −2 and 12.97 kV cm −1 , respectively. Compared to PMNT ceramics [18], the E C was substantially improved, making Mn-PIMNT suitable in high-power applications. With the increasing electric field amplitude (E 0 ), the P r , E C , and hysteresis area < A > all increased gradually, indicating that domain motions and polarization switching became completely. Due to "hard" Mn-doping, the ceramics exhibited significantly enhanced P r and E C , compared to the pure PIMNT ceramics (P r = 27.3 µC cm −2 and E C = 8.2 kV cm −1 ) [7]. The values of P r and E C changed with the electric field and comparatively saturated at 50 kV cm −1 . In addition, we further studied temperature-dependent the P-E hysteresis loops at 40 kV cm −1 . The polarization axis is slightly asymmetric and the internal bias field of E i was calculated by the equation, where the E i was generated by the acceptor oxygen vacancy charge-defective dipole, the E C+ and E C-represent the positive and negative coercive field of P-E hysteresis loops, respectively [13]. The E i was calculated to vary between 0.15 and 0.39 kV cm −1 . With increasing temperature, the P r , E C , E i , and < A > of P-E loops decreased, indicating the shape of a dielectric material and the polarization switching became easier. Furthermore, as the temperature rising to 120 ℃, P r and E C decreased approximately linearly. The slope of P r changed near T rt and T C was attributed to the structural phase transition or the reduction of the long-term polar order. It was worth noting that when the temperature was much higher than T C and reached 200 ℃, a slight hysteresis characteristic could still be detected from the P-E hysteresis loop, which indicated that microdomains still existed where the temperature was much higher than T C . Similar phenomenon also existed in the Mn-PIMNT single crystals [29]. Figure 5 shows the S-E behavior measured at 10 Hz. The unipolar strain behavior showed an S-shape in Fig. 5a [ [30][31][32]. The bidirectional field-induced strain curves exhibited a typical butterfly shape in the inset of Fig. 5a. The normalized piezoelectric strain coefficient d * 33 , which is an important parameter for actuator applications, is calculated by the equation, The electric field dependence of strain and piezoelectric strain d * 33 for Mn-PIMNT ceramics are presented in Fig. 5b. The unipolar strain S max and d * 33 were 0.22% and 355 pm V −1 at an electric field of 60 kV cm −1 , respectively. The d *

33
increased first and then decreased with the further increase of the electric field, meanwhile, the strain for Mn-PIMNT increased [31]. The increase in d * 33 at low field was ascribed to the contribution of more non-180° domain wall movement [30]. However, at high electric field levels, the electrical domains were partially clamped by the electrical field and the lattice extension may saturate, resulting in a reduction in the unipolar S max and the d * 33 [30,31]. Figure 5c, d show unipolar strain curves of Mn-PIMNT under different temperatures measured at 10 Hz and 40 kV cm −1 . The unipolar S max and d * 33 first increased to a peak value and then followed a decrease instead. Meanwhile, unipolar S max and d *

33
showed a similar tendency and reached up to ~ 0.274% and 697 pm V −1 at 170 °C. In addition, according to the IEEE standard, the electromechanical coupling coefficient of Mn-PIMNT was calculated from the following formula [33]: where k p is planar electromechanical coupling, f r represents the resonant frequency and f a is anti-resonant frequency in radial vibration mode. The parameter k p was calculated to be 43.1% by the impedance spectrum in Fig. 6a. The mechanical quality factor Q m was calculated by the equation, According to the admittance spectrum as shown in Fig. 6b, f − 1 2 and f + 1 2 represent frequencies at -3 dB down the peak admittance. The parameter Q m was calculated to be 256. In addition, the figure of merit (FOM) of the output characteristic parameters for high-power piezoelectric applications [18] was also calculated by the following formula: The direct piezoelectric effect of d 33 was tested to be 235 pC/N and the FOM was calculated to be 60,160 pC/N. Table 1 summarized the fundamental performance parameter  [7,18]. The reason can be attributed to Mn-doping, which introduced oxygen vacancies and formed the defect dipoles. They disrupted the stability of the ferroelectric domain as well as reduced the mobility of charge carriers on the domain wall, subsequently inducing piezoelectric hardening characteristics [14,15]. Figure 7 shows the variation of the pyroelectric coefficient (p) with the temperature for the Mn-PIMNT ceramics. The figures of merit (FOMs) for the current responsiv- ( 0 r tan ) are calculated to assess the performance of pyroelectric materials. The volume-specific heat C v and the permittivity of free space ε 0 are 2.5 × 10 6 J m −3 K −1 [27] and 8.85 × 10 -12 F m −1 , respectively [12]. The value of p was 4.8 × 10 -4 C m -2 K -1 at room temperature, which was twice as large as that of LiTaO 3 single crystals [34]. With temperature increasing to 130 ℃, the p increased to 2.3 × 10 -3 C m -2 K -1 . Furthermore, the figures of merits of F i , F v , and F d were calculated to be 1.84 × 10 -10 m V -1 , 0.028 m 2 C -1 , and 2.317 × 10 -5 Pa -1/2 , respectively. A comparison of the pyroelectric parameters of the Mn-PIMNT and other typical ferroelectric ceramics at 1 kHz is summarized in Table 2. Compared to undoped PIMNT ceramics, the Mn-PIMNT ceramics exhibited enhanced F v and F d than other ferroelectric ceramics [17,35,36], which are desirable for uncooled infrared detection applications.

Conclusions
In summary, the third-generation relaxor ferroelectric Mn-PIMNT ceramics were fabricated and the temperature dependence of dielectric, ferroelectric, piezoelectric, and pyroelectric properties of Mn-PIMNT ceramics were investigated. Excellent piezoelectric coefficient d 33 of 235 pC/N, high-power figure of merit (FOM) of 60,160 pC/N, large remnant polarization P r of 34.57 µC cm −2 and coercive field E C of 12.97 kV cm −1 were obtained along with giant pyroelectric coefficient p up to 4.8 × 10 -4 C m −2 K −1 and detectivity F d as high as 2.317 × 10 -5 Pa −1/2 . The strain nonlinearity and domain wall motion under a large driving field were revealed and the unipolar strain up to ~ 0.274% and d * 33 as high as 697 pm/V were obtained at 170 °C. High ferroelectric rhombohedral to tetragonal phase transition temperature T rt (up to 146 ℃) and ferroelectric tetragonal to cubic phase transition T C of 188 ℃ make the Mn-PIMNT ceramics an excellent candidate for piezoelectric and pyroelectric device applications.