Optimizing Sintering Temperature for Enhanced Piezoelectric Performance in PMT-PNT-PZT Ceramics

1. Introduction

Piezoelectric materials enable efficient interconversion between electrical and mechanical energy, making them essential functional components in transducers, actuators, sensors, and energy harvesters [1,2,3,4,5,6]. Among them, lead-based ceramics such as PZT have been widely adopted owing to their excellent piezoelectric and electrical properties [7,8,9,10,11,12,13,14]. To further improve piezoelectric performance, lead-based relaxor ceramics have been developed, leveraging their distinctive microstructures to achieve enhanced functional behavior.

PZT-based solid-solution piezoelectric ceramics are often modified or doped for specific applications. The introduction of rare earth elements Sm³⁺ and Eu³⁺ in PMN-PT ceramics has proven effective in improving dielectric properties [15,16,17]. Similarly, the incorporation of Mn and Fe has been reported to effectively enhance the mechanical quality factor [18,19,20,21]. Another strategy involves forming relaxor solid solutions, such as Pb(Ni_1/3Nb_2/3)O₃-PbZrO₃-PbTiO₃ [22], Pb(Sn_1/3Nb_2/3)O₃-Pb(Zn_1/3Nb_2/3)O₃-Pb(Zr,TiO₃) [23], Pb(Sc_1/2Nb_1/2)O₃-PbTiO₃ [24] to enhance piezoelectric performance. While many studies on PZT-based ceramics focus on increasing the piezoelectric constant d₃₃ or the planar electromechanical coupling factor k_p , only a limited number aim to achieve a balanced improvement in both d₃₃ and Q_m. This balance is crucial for high-power applications, as d₃₃ influences the sensitivity of ultrasonic transducers, while Q_m is associated with heat generation under high drive conditions [25,26]. Therefore, the coordination of d₃₃ and Q_m is of great significance for high-power application materials [27]. In this work, 0.006Pb(Mn_1/3Ta_2/3)O₃-0.114Pb(Ni_1/3Ta_2/3)O₃-0.43PbZrO₃-0.45PbTiO₃ (PMT-PNT-PZ-PT) is designed to simultaneously enhance d₃₃ and Q_m, showing promising potential for high-power device applications.

Sintering temperature plays a critical role in determining the properties of piezoelectric ceramics. MnO₂-doped PZT-PZN ceramics were meticulously prepared at a high temperature of 900℃, exhibiting an impressive d₃₃ of 330 pCN^-1 and a remarkable Q_m of 1000 [28]. The Pb(Mn_1/3Nb_2/3)O₃-Pb(Ni_1/3Nb_2/3)O₃-Pb(Zr_0.50Ti_0.50)O₃ (PMN-PNN-PZT) ceramics were sintered at 900 °C, exhibiting a d₃₃ of 346 pCN^-1 and a Q_m of 1130 [29]. The PMN-PZT-Li₂CO₃ ceramics were sintered at a temperature of 940 °C, exhibiting a high quality factor Q_m (2264), a high Curie temperature T_c (317 °C) and an impressive dielectric constant (1216) [30]. The PSNT-Mn with LiBiO₂ ceramics were sintered at 950 °C, resulting in d₃₃ = 340 pCN^-1, Q_m = 800, T_c = 263 °C [31]. PNN-PZN-PMN-PZ-PT ceramics were sintered at 950 °C, exhibit excellent piezoelectric properties d₃₃* = 503 pmV^-1, Q_m = 471 [32]. By exploring the influence of the sintering temperature (T_s) on the structure of PMT-PNT-PZ-PT ceramic, the best T_s was found to obtain good electrical properties.

In this work, 0.4Pb(Ni_1/3Ta_2/3)O₃ -0.6PbTiO₃ exhibits a high dielectric property, and it induces lead vacancies by incorporating Ni²⁺ and Ta⁵⁺ into the B site of PZT, facilitating domain switching and effectively enhancing the piezoelectric performance of the ceramic, 0.006Pb(Mn_1/3Ta_2/3)O₃-0.114Pb(Ni_1/3Ta_2/3)O₃-0.43PbZrO₃-0.45PbTiO₃ (PMT-PNT -PZ-PT) is designed to simultaneously enhance d₃₃ and Q_m, showing promising potential for high-power device applications.

2. Materials and Methods

0.006Pb(Mn_1/3Ta_2/3)O₃-0.114Pb(Ni_1/3Ta_2/3)O₃-0.43PbZrO₃-0.45PbTiO₃ (abbreviated as PMT-PNT-PZ-PT) ceramics were fabricated via the conventional solid-state method. PbO (99.9%), ZrO₂ (99.99%), NiO(99%), Ta₂O₅ (99.99%) and MnO₂ (99%) were used as the starting materials. All the starting materials were weighed based on the stoichiometric ratio and then milled in the alcohol with zirconia balls for 12 h. The resulting mixture was dried and pre-burned at 750 °C for 2.5 h, and the calcined powder was ball-milled for 24 h again. The powder was mixed with 7 wt.% polyvinyl alcohol (PVA) as a binder, which provides cohesion and green strength for shaping. The mixture was then uniaxially pressed into gray disks with a diameter of 13 mm. Subsequently, the prepared disks underwent a debinding process at 550 °C to completely remove the PVA. Finally, the disks were embedded in a sacrificial powder with the same composition as the ceramic matrix to prevent reaction and deformation, followed by sintering for 2.5 hours at different temperatures (T_s = 1180, 1200, 1250, and 1270 °C). To facilitate electrical properties characterization, the samples were polished and coated with silver to serve as electrodes.

The crystalline phases of the ceramics were evidenced by XRD (D/max-rB 12kW X-ray diffractometer). The fractured surface micromorphology of the sintered samples was determined by the scanning electron microscopy (SEM, SU5000). The temperature dependence of dielectric constant and loss were measured using a LCR test instrument (Agilent, E4980A, Santa Clara, CA, USA) from the room temperature to 450 ℃ at 0.1 kHz, 1 kHz, 10 kHz, 100 kHz, 1 MHz. For piezoelectric properties, the ceramics were poled in a silicone oil bath at 150 °C under a DC field of 3 kV/mm for 15 minutes. Poling at this elevated temperature lowers the coercive field, thereby facilitating domain alignment under the applied electric field. The piezoelectric coefficient d₃₃ was measured by a quasi-static meter (ZJ-4A). The ferroelectric hysteresis loops of the ceramic samples were measured by the ferroelectric test system. (premier II, Radiant Tech, Albuquerque, USA). The mechanical quality factor (Q_m) and electromechanical coupling factor (k_p) were calculated based on the IEEE standards using the Agilent 4294A Precision Impedance Analyzer, as follows:

$$Q_m={\displaystyle \frac{1}{4\pi (f_a-f_r)R_1(C_0+C_1)}}$$

(1)

$${\displaystyle \frac{1}{k_p^2}}=0.395\times {\displaystyle \frac{f_r}{f_a-f_r}}+0.574$$

(2)

3. Results and Discussion

Figure 1 presents the XRD patterns of PMT-PNT-PZ-PT ceramics sintered at different temperatures (1180 °C, 1200 °C, 1250 °C, and 1270 °C). It can be seen that the positions of the most intense diffraction peaks remain consistent across all sintering temperatures, with all samples exhibiting characteristic peaks of a perovskite structure. This indicates that the doped ions have been successfully incorporated into the perovskite lattice, forming a stable solid solution. The crystal structure is identified as tetragonal for the PMT-PNT-PZ-PT ceramics, as evidenced by a (200)/(002) peak intensity ratio of approximately 1:2. Furthermore, the intensity of the (110) diffraction peak varies with sintering temperature, suggesting that the sintering process influences the crystallinity of the ceramics.

Figure 2 (a–d) displays the SEM micrographs of PMT-PNT-PZ-PT ceramics sintered at different temperatures. The samples sintered at 1180 °C and 1200 °C exhibit a dense microstructure, well-developed grains, and no visible pores or abnormal grain growth, suggesting excellent ceramic quality and optimal electrical properties at these temperatures. The grain size as a function of sintering temperature is summarized in Figure 2 (e and f). As the sintering temperature increases, the average grain size shows a clear upward trend, indicating that higher temperatures promote grain growth. However, the emergence of pores at elevated temperatures likely reduces the densification of the ceramics, thereby degrading their electrical performance. In the sample sintered at 1270 °C, the grains are closely bonded and exhibit a tendency to coalesce, which can be attributed to the partial melting of over-sintered grains.

Figure 3(a-d) illustrates the temperature dependence of the dielectric constant and loss for PMT-PNT-PZ-PT ceramics measured from room temperature to 450 °C at frequencies of 0.1 kHz, 1 kHz, 10 kHz, 100 kHz, 1MHz. A single permittivity peak is observed near the Curie temperature (approximately 300 °C) for all compositions, indicating a ferroelectric–paraelectric phase transition. At room temperature, the ceramics exhibit a purely tetragonal phase, which is consistent with the XRD analysis. As frequency increases, the loss (tanδ) values gradually decrease due to polaron relaxation processes associated with localized oxygen vacancies present in the crystal lattice [33]. At the low frequency, tanδ start to rapidly increase beyond T_c due to thermally-activated space charge conduction behavior [34].

To further analyze the dielectric behavior, a modified empirical formula was employed to evaluate the dielectric dispersion and diffusivity (γ) of the phase transition. The calculation of γ based on the Curie-Weiss law can be formulated as follows

$${\displaystyle \raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{$\epsilon $}\right.}-{\displaystyle \raisebox{1ex}{$1$}\!\left/ \!\raisebox{-1ex}{${\epsilon }_m$}\right.}={\displaystyle \raisebox{1ex}{${(T-T_m)}^{\gamma }$}\!\left/ \!\raisebox{-1ex}{$C$}\right.}$$

(3)

where ε_m is the maximum dielectric constant at T_m. Figure 3(e) further reveals that all samples exhibit characteristic relaxor ferroelectric behavior with a diffuse phase transition, regardless of sintering temperature. As the sintering temperature increases, the diffuseness parameter γ gradually decreases, suggesting that grain growth tends to suppress the relaxor characteristics. Notably, in the sample sintered at 1180 °C, the dielectric loss (tanδ) remains below 0.02 at 10 kHz. This low loss profile is highly favorable for applications such as ultrasonic transducers.

The P-E loops of PMT-PNT-PZ-PT measured at room temperature are shown in Figure 4. With the increase of sintering temperature, P-E loops become asymmetrical gradually. The presence of acceptor-oxygen vacancy defect dipoles leads to the generation of an internal bias field [35,36,37]. The internal bias field can be calculated by

$$E_i={\displaystyle \frac{\left|E_c^{+}-\left|E_c^{-}\right|\right|}{2}}$$

(4)

Figure 4(b) shows the changing rule of $E_c^{+}$, $\left|E_c^{-}\right|$ and E_i as a function of sintering temperature. The E_i increases proportionally with the sintering temperature, whereby higher temperatures lead to elevated concentrations of oxygen vacancies and defect dipoles, ultimately resulting in an augmented internal bias field.

Figure 5 describes the relationship between d₃₃, Q_m and sintering temperature. The d₃₃ value gradually decreases with the sintering temperature, while the Q_m value initially increases and then subsequently decreases. In order to assess the optimal comprehensive properties of PMT-PNT-PZ-PT ceramics, figure of merit (FOM) is defined as FOM = d₃₃ × Q_m [38,39]. According to Figure 5, the optimum-integrated performances of d₃₃ (400 pC/N), Q_m (509), and FOM (203600 pC/N) for the sintering temperature 1180 °C. The observed trends can be explained by the incorporation of Mn ions and the associated generation of oxygen vacancies for charge compensation. These point defects act as pinning centers for domain walls, restricting their motion [40,41]. This pinning effect is responsible for the gradual decrease in d₃₃ and the initial increase in Q_m with sintering temperature, as domain wall contributions to dielectric and piezoelectric responses are reduced while mechanical losses are minimized. For perovskite ferroelectrics, the piezoelectric coefficient d₃₃ can be expressed as [42]:

$$d_{33}=2P_{r}Q{\epsilon }_0{\epsilon }_r$$

(5)

where P_r is residual polarization, ε_r is the dielectric constant, ε₀ is the dielectric constant in vacuum, Q is the electrostrictive coefficient (for the same material Q is the same). When the sintering temperature is 1180 °C, the ceramics have the largest piezoelectric constant (d₃₃). Combined with Figure 5 (b), the variation of d₃₃ and P_rε_r is consistent, so the large d₃₃ comes from P_rε_r. The analysis shows that the Mn enters the lattice at higher temperatures. To uphold the electrical neutrality of the cell, a greater number of oxygen vacancies are generated within the material. The existence of oxygen vacancies hinders the mobility of the domain wall, resulting in a decrease in d₃₃ and k_p, while Q_m exhibits an increase. Incorporation of Mn ions into the system induces ceramic "strengthening" effects, leading to an enhancement in Q_m.

The properties of the ceramics are compared with those of commercial PZT4 ceramics and other ceramics, as presented in Table 1. We effectively reconcile the trade-off between d₃₃ and Q_m, thereby presenting a novel approach to enhance the sensitivity of ultrasonic transducers while minimizing heat generation. Furthermore, it reveals significant advantages in various other properties compared to the aforementioned piezoelectric ceramics.

4. Conclusions

The effects of sintering temperature on the phase structure and electrical properties of 0.006Pb(Mn_1/3Ta_2/3)O₃-0.114Pb(Ni_1/3Ta_2/3)O₃-0.43PbZrO₃-0.45PbTiO₃ (PMT-PNT-PZ-PT) ceramics are comprehensively investigated. The sintering temperature plays a critical role in tailoring the microstructure and electrical performance of PMT-PNT-PZ-PT ceramics. The ceramic sintered at 1200 °C exhibits the smallest grain size along with the highest mechanical quality factor (Q_m) and relative permittivity (ε_r). The piezoelectric constant d₃₃ (400 pC/N) is achieved at 1180 °C, accompanied by a significant P_rε_r (39115 μCcm^-2). Compared to commercial PZT-4 ceramics, the composition sintered at 1180 °C attained an optimal balance between d₃₃ and Q_m, resulting in a superior comprehensive figure of merit (FOM = 2.04 × 10⁵ pC/N). These results not only provide a viable candidate material for high-power and temperature-stable piezoelectric devices but also offer valuable insights into the processing-property relationships in complex perovskite ceramics.