Concurrently enhanced piezoelectric performance and curie temperature in stressed lead-free Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃ ceramics

Abstract

Eco-friendly, lead-free BaTiO₃-based piezoelectrics are critical for sustainable electronics, but improving their piezoelectric properties often compromises Curie temperature (T_C). To address this trade-off, we implemented an innovative stress engineering approach by introducing a secondary phase BaAl₂O₄ in Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃ (BCTZ) ceramics. The thermal expansion mismatch between BCTZ and BaAl₂O₄ induces internal stress within the BCTZ matrix, causing significant lattice distortion and phase fraction modulation, which improves both T_C and the piezoelectric coefficient (d₃₃). Additionally, the local electric field and Al³⁺ doping in ABO₃ lattice further enhance d₃₃. Optimized BCTZ ceramics achieve d₃₃ of 650 ± 16 pC N⁻¹, d₃₃^* of 1070 pm V⁻¹, and T_C of 96.5 ± 1.0 °C, placing them at the forefront of lead-free BaTiO₃-based piezoelectrics. This study underscores the effectiveness of bulk stress engineering via a secondary phase for enhancing lead-free piezoelectric ceramics, paving the way for developing high-performance piezoelectric ceramics suitable for broad temperature applications.

Introduction

Over the past few decades, the piezoelectric materials market has predominately relied on lead-based piezoelectric ceramics. However, escalating environmental concerns have spurred extensive endeavors to explore lead-free alternatives. Among these, Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃ (BCTZ) ceramics have emerged as promising candidates due to their high piezoelectric coefficient (d₃₃ > 500 pC N⁻¹)^1,2,3, garnering widespread attention. Many efforts, such as optimizing sintering processes⁴, employing doping strategies^5,6,7, and incorporating second-phase materials⁸, have been carried out to further enhance the performance of BCTZ-based ceramics. Nevertheless, several challenges persist that hinder their practical application³. Specially, it is widely recognized that trade-offs among performance parameters exist in most piezoelectric ceramics, making it difficult to achieve simultaneous improvement across multiple parameters, particularly the concurrent achievement of high Curie temperature (T_C) and high d₃₃. A primary concern is the relatively low T_C (~90 °C) of BCTZ piezoelectric ceramics, which limits their operational temperature range. Notably, enhancing piezoelectricity often comes at the expense of T_C (Supplementary Fig. 1)^{7,8,9,10,11,12,13}, highlighting the necessity to improve both piezoelectric properties and T_C concurrently to broaden the applications of lead-free BCTZ-based piezoceramics.

Strain engineering of perovskite oxide thin films has emerged as an exceptionally powerful technique for manipulating ferroelectric behavior^14,15,16,17. Schlom et al. were the first to demonstrate that introducing epitaxial strain and stress in SrTiO₃ films grown on perovskite-type substrate can improve the T_C by inducing lattice distortion, thereby stabilizing the ferroelectric phase¹⁸. Subsequently, various methods have been developed to modulate the in-plane strains of ferroelectric films, allowing for precise control over various properties such as crystal symmetry, domain structure, defect concentration, and phase ratio^17,19,20. Despite its effectiveness in tailoring the physical properties of ferroelectric films, the application of strain engineering in the field of piezoelectric ceramics remains largely unexplored.

Inspired by strain engineering in ferroelectric thin films, we present a novel stress engineering strategy aimed at simultaneously enhancing the T_C and d₃₃ of BCTZ-based ceramics. Specifically, lead-free ceramics composed of Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃ and a small amount of BaAl₂O₄ were fabricated through the doping of AlN particles. The underlying causes of the secondary phase with varying doping amounts and particle sizes on the piezoelectric performance and T_C are systematically investigated. Encouragingly, this approach yielded an ultrahigh piezoelectric property (d₃₃ = 650 ± 16 pC N⁻¹, d₃₃^* = 1070 pm V⁻¹) along with an excellent T_C (96.5 ± 1.0 °C). This study provides an example of simultaneously elevating both the T_C and d₃₃ of piezoelectric ceramics, offering a promising avenue for future research on high-performance lead-free piezoelectric ceramics.

Results

Piezo- and ferroelectric properties

Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃–x mol% AlN (50 nm) (BCTZ–xAlN–50, x = 0.25, 1.00, 1.50, and 1.75), Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃–x mol% AlN (300 nm) (BCTZ–xAlN–300, x = 0.25, 0.50, 0.75, 1.00, 1.25, 1.50, 1.75, and 2.00), and Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃–x mol% AlN (3000 nm) (BCTZ–xAlN–3000, x = 0.25, 1.00, 1.50, and 1.75) ceramics were prepared using AlN with sizes of 50 nm, 300 nm, and 3000 nm (Supplementary Fig. 2) by a conventional solid-state reaction method, respectively. The comparison of d₃₃ and T_C of BCTZ‒xAlN‒50, BCTZ‒xAlN‒300, and BCTZ‒xAlN‒3000 ceramics is depicted in Fig. 1a. The T_C was determined by temperature dependence of dielectric constant (ε_r‒T) and dielectric loss (tanδ‒T) curves illustrated in Supplementary Figs. 3‒5. For BCTZ‒xAlN‒50, a significant reduction in T_C is observed, while d₃₃ shows a slight improvement. For BCTZ‒xAlN‒300 ceramics, the d₃₃ initially increases as the AlN doping concentration increases up to 1.75 and then decreases beyond 1.75. Notably, the BCTZ‒1.75AlN‒300 ceramic achieves an ultrahigh d₃₃ of 650 ± 16 pC N⁻¹, which is ~30% higher than that of pure BCTZ ceramic (503 ± 15 pC N⁻¹). Meanwhile, the T_C of BCTZ‒xAlN‒300 ceramics initially decreases with the AlN doping concentration, declining from 90.2 ± 0.8 °C in pure BCTZ to 88.2 ± 0.9 °C in BCTZ‒0.25AlN‒300. Subsequently, as the doping concentration increases, T_C exhibits a rapid rise, reaching a maximum of 96.5 ± 1.0 °C at x = 1.75. Remarkably, doping BCTZ ceramics with AlN can simultaneously improve both d₃₃ and T_C, exceeding those of previous reported BaTiO₃-based lead-free piezoceramics, including BCTZ, BCTSn, and BCTHf-based compositions (summarized in Supplementary Fig. 6 and Supplementary Table 1; Supplementary Note 1). Our findings demonstrate significant progress in piezoelectric properties for BaTiO₃-based ceramics, alleviating the trade-off relationship between d₃₃ and T_C, making them highly competitive to replace traditional lead-based materials (Fig. 1b)^{3,7,21,22,23,24,25,26,27}. Meanwhile, BCTZ‒xAlN‒3000 ceramics exhibit a significant increase in T_C accompanied by a slight suppression in d₃₃. These results imply that the incorporation of AlN, with varying dopant sizes and concentrations, exhibits significant influence on ceramic performance, and subsequent analysis will delve deeper into this impact.

Fig. 1: Piezo- and ferroelectric properties of BCTZ–xAlN–50 ceramics.

a Comparison of the d₃₃ and T_C of BCTZ–xAlN–50 ceramics, BCTZ–xAlN–300 ceramics, and BCTZ–xAlN–3000 ceramics. Error bars show standard error of mean. b Comparison of d₃₃ and T_C with the other piezoceramics^{3,7,21,22,23,24,25,26,27}. c The electric field dependent d₃₃^* for BCTZ–xAlN–300 ceramics. d Comparison of d₃₃^* and d₃₃ with the other piezoceramics^{28,29,30,31,32}. e PFM amplitude and phase hysteresis versus voltage loops of BCTZ ceramic and BCZT−1.75AlN−300 ceramic. f1 Temperature-dependent d₃₃^* of BCTZ and BCTZ–1.75AlN–300 ceramics. f2 Comparison of the normalized d₃₃^* between BCTZ and BCTZ–1.75AlN–300 ceramics under different temperatures. Normalized d₃₃^* is defined as the ratio of the d₃₃^* value measured at various temperatures to the d₃₃ value measured at the initial room temperature. g1 The in-situ d₃₃ as a function of temperature for BCTZ and BCTZ–1.75AlN–300 ceramics. g2 Comparison of the normalized d₃₃ between BCTZ and BCTZ–1.75AlN–300 ceramics at different temperatures. The inset is the enlarged view of the normalized d₃₃ between 25 and 40 °C.

The unipolar strain-electric field (S‒E) curves of the BCTZ‒xAlN‒300 ceramics are presented in Supplementary Fig. 7. The converse piezoelectric coefficient (d₃₃^* = S_max/E_max, where S_max is the maximum strain and E_max is the maximum electric field) derived from the unipolar S‒E curves at varied electric fields is detailed in Fig. 1c. The doping of AlN markedly improves the strain in the ceramics, with the most pronounced effect observed in the BCTZ‒1.75AlN‒300 ceramic. This particular ceramic achieves the highest d₃₃^* = 1070 pm V⁻¹ at an electric field of 10 kV cm⁻¹, which is 0.6 times greater than that of the pure BCTZ ceramic (680 pm V⁻¹). Figure 1d compares the d₃₃ and d₃₃^* values of lead-based and lead-free piezoelectric ceramics, indicating that the high-performance BCTZ−1.75AlN−300 ceramics in this work have greater potential for electro-mechanical conversion compared to other lead-based and lead-free piezoelectric ceramic systems^{28,29,30,31,32}. Specifically, when compared to traditional lead-based materials, such as PZT-5H, our BCTZ−1.75AlN−300 ceramics offer a balanced combination of high piezoelectric performance and moderate thermal stability. These properties showcase them promising potential as a viable alternative to lead-based ceramics. Notably, the local polarization-electric field hysteresis loops of BCZT and BCTZ−1.75AlN−300 ceramics were recorded by SS-PFM measurements, reflecting the local piezoelectric response (Fig. 1e). The phase hysteresis loops exhibit a characteristic 180° contrast, indicating a fully reversible polarization dynamic. The amplitude of both samples shows butterfly loops describing the local displacement as a function of voltage, which is typical of the piezoelectric response in ferroelectric materials. The BCTZ−1.75AlN−300 sample demonstrates an increased amplitude compared to the undoped BCZT ceramics, offering direct microscopic evidence of the significantly enhanced piezoelectric response in BCTZ−1.75AlN−300 ceramics. The polarization-electric field (P‒E) loops of the BCTZ‒xAlN‒300 ceramics reveal that BCTZ−1.75AlN−300 ceramic exhibits both the largest maximum polarization (P_max) of 18.9 μC cm⁻² and largest remnant polarization (P_r) of 9.54 μC cm⁻² among all the ceramics (Supplementary Figs. 8‒9, Supplementary Note 2). Additionally, the P‒E and S‒E curves for the BCTZ‒xAlN‒50 and BCTZ‒xAlN‒3000 ceramics are depicted in Supplementary Figs. 10‒11, respectively. When compared to BCTZ‒xAlN‒300 samples with equivalent doping levels, both BCTZ‒xAlN‒50 and BCTZ‒xAlN‒3000 ceramics exhibit lower strain and polarization values. The d₃₃^* behavior of the BCTZ‒xAlN‒50 and BCTZ‒xAlN‒3000 ceramics corresponds to the changes in d₃₃ observed in their respective ceramics. The optimal d₃₃^* values for BCTZ‒xAlN‒50 and BCTZ–xAlN–3000 ceramics are achieved at doping levels of x = 1.75 and x = 2.00, respectively, under an electric field of 10 kV cm⁻¹, with values of 860 and 902 pm V⁻¹, accordingly.

To evaluate the thermal stability of the ceramics, we measured the unipolar strain and in-situ d₃₃ as a function of temperature. Temperature-dependent unipolar S–E curves of BCTZ and BCTZ–1.75AlN–300 ceramics are examined (Supplementary Fig. 12), and the corresponding d₃₃^* are summarized in Fig. 1f1. As the temperature increases, the strain in BCTZ–1.75AlN–300 ceramics decreases slowly, with an overall variation of less than 14% over the temperature range from 25 to 60 °C. The data obtained from multiple measurements of BCTZ–1.75AlN–300 ceramics demonstrates good repeatability (Supplementary Fig. 13). In contrast, the strain of BCTZ ceramics decreases sharply, exhibiting a larger variation of 26%. Additionally, the in-situ thermal stability of d₃₃ from 25 to 100 °C and the corresponding normalized d₃₃ retention are shown in Fig. 1g. As the temperature increases, the d₃₃ of BCTZ ceramic decreases monotonically, whereas the d₃₃ of BCTZ–1.75AlN–300 ceramic initially increases and then decreases. At 60 °C, the BCTZ–1.75AlN–300 ceramic exhibits a d₃₃ of 550 pC N⁻¹ and a d₃₃ retention of 85%, which is higher than that of BCTZ ceramics and other AlN-doped BCTZ ceramics (Supplementary Fig. 14), demonstrating its strong potential for broad-temperature applications.

Additionally, the ε_r–T and tanδ–T curves of BCTZ–xAlN–300 ceramics are measured to accurately determine the phase transformation temperatures (Supplementary Fig. 4). All ceramics exhibit three anomalies corresponding to the phase transitions from ferroelectric rhombohedral to orthorhombic (T_R-O), orthorhombic to tetragonal (T_O-T), and tetragonal to paraelectric cubic phases (T_C), which are summarized in Supplementary Fig. 15. The MPB at room temperature mitigates polarization anisotropy by facilitating polarization rotation between the tetragonal (001)_T and the rhombohedral (111)_R phase states, contributing to the outstanding piezoelectric and dielectric properties (Supplementary Fig. 16)^33,34.

Phase structure and micromorphology

The as-synthesized ceramics were characterized using high-resolution Synchrotron X-ray diffraction technique (SXRD), as illustrated in Supplementary Figs. 17–18. Taking BCTZ–xAlN–300 ceramics as examples, all samples exhibit a typical perovskite structure. A weak diffraction peak observed at 2θ angle of 10.74° corresponds to the (202) planes of BaAl₂O₄ with hexagonal symmetry and P6322 space group (JCPDS No.17-0306). The same phenomenon can also be observed in both BCTZ–xAlN–50 and BCTZ–xAlN–3000 ceramics. Despite the introduction AlN in the precursor of BCTZ, its inherent instability at high temperatures (>800 °C) leads to a reaction between AlN and BCTZ during high-temperature sintering, resulting in the formation of BaAl₂O₄ (Supplementary Fig. 19 and Supplementary Note 3). This indicates that the actual composition of the secondary phase in BCTZ–xAlN ceramics is BaAl₂O₄, and AlN in the BCTZ–xAlN ceramics is solely utilized for descriptive annotation.

The impact of secondary phase BaAl₂O₄ on the phase transformation of BCTZ–xAlN ceramics was investigated. Initially, Fig. 2a depicts the contour plots of the (110) and (111) diffraction peaks of BCTZ–xAlN–300 ceramic. It is evident that the (110) and (111) diffraction peaks of BCTZ–xAlN ceramics shift to higher degrees as the dopant AlN particles increase, indicating a gradual enhancement in the shrinkage of BCTZ lattice. Subsequently, Rietveld refinement was applied to the high-resolution SXRD patterns to accurately determine the crystal structure. BaTiO₃, with its tetragonal (T) phase characterized by P4mm space group (JCPDS No.05-0626), the orthorhombic (O) phase by the Amm2 space group (JCPDS No.81-2200), and the rhombohedral (R) phase by the R3m space group (JCPDS No.85-0368), was used as the initial structural model for these refinements (Supplementary Fig. 20). The refinement results confirmed that all ceramics exhibit a coexistence of R, O, and T phases. The phase fractions for all ceramics are summarized in Fig. 2b. As the AlN doping concentration increases, the fractions of T phase and O phase of BCTZ–xAlN–300 ceramics decrease initially, accompanied by a continuous increase in the R phase. Notably, BCTZ–1.75AlN–300 ceramics exhibit the highest fractions of R phase and the largest α angle, suggesting enhanced crystal asymmetry (Fig. 2c). Furthermore, this ceramic has the largest c/a ratio, indicating significant lattice distortion along the c-axis and substantial local displacement of B-site ions^5,35. Correspondingly, the full-width at half-maxima (FWHM) of (110) and (111) diffraction peaks in BCTZ–1.75AlN–300 ceramic, confirming maximal crystal asymmetry and lattice distortion⁷. These structural features contribute to an increased dipole moment, thereby enhancing piezoelectric performance (such as d₃₃ and d₃₃^*).

Fig. 2: Phase and microstructure characterization of BCTZ–xAlN ceramics.

a Evolution of (110) and (111) diffraction peaks as a function of the AlN doping content. b Phase fractions, (c) α angle of R phase, lattice constant (c/a) of T phase, and FWHM for the (110) and (111) planes of BCTZ–xAlN–300 ceramics obtained from the Rietveld refinement. d Contour plot of in-situ variable temperature SXRD for the (111) and (200) diffraction peaks of BCTZ–xAlN–300 ceramics. e Average grain size of the BCTZ–xAlN ceramics. f TEM image of BCTZ–1.75AlN–300 ceramic (Inset is the average domain size of the BCTZ–xAlN ceramics). g SEM image and EDS mapping image of BCTZ–1.75AlN–300 ceramic (the yellow-circled area represents BaAl₂O₄). HRTEM images and corresponding SAED patterns of (h) BCTZ and (i) BaAl₂O₄ in BCTZ–1.75AlN–300 ceramics.

Moreover, in-situ variable-temperature high-resolution SXRD has been utilized to investigate the crystallographic phase evolution of BCTZ–1.75AlN–300 ceramics (Supplementary Fig. 21). Contour maps of expanded SXRD patterns for the (111) and (200) diffraction peaks illustrate the temperature-induced phase transformations in BCTZ–1.75AlN–300 ceramics during heating process (Fig. 2d). Remarkably, the ceramics initially undergo a sequential disappearance of the R, O, and T phases upon heating, ultimately transforming into a paraelectric cubic (C) phase. This observation is consistent with the ε_r–T and tanδ–T results. This phase transformation leads to a decrease in d₃₃^* and d₃₃ as the temperature rises, as observed in Fig. 1f, g.

To further examine the local structure, the surface morphology and average grain size of thermally etched BCTZ–xAlN ceramics were analyzed using SEM (Supplementary Figs. 22–26). All ceramics exhibit a highly dense microstructure, with a high relative density (above 95%) and low porosity (summarized in Supplementary Table 2 and Supplementary Note 4), which contributes to their electrical properties^27,36. The pristine BCTZ ceramic exhibits irregularly shaped grains with an average size of ~16.4 μm. The average grain size of BCTZ–xAlN–50, BCTZ–xAlN–300, and BCTZ–xAlN–3000 ceramics decreases markedly as AlN doping concentration increases (Fig. 2e). This continuous reduction in size can be attributed to the hindering effect of BaAl₂O₄ particles on the grain boundaries during the sintering process. The reduction in grain size leads to a refined domain configuration (Fig. 2f, Supplementary Fig. 27 and Supplementary Note 5). Specifically, the domain size decreases from 210 nm in the undoped BCTZ ceramic to 180 nm in the BCTZ–1.75AlN–300 sample. Considering that the domain size is proportional to the square root of the domain wall energy, this refinement facilitates the polarization orientation under an external electric field and the orderly arrangement of ferroelectric domains, thereby enhancing the extrinsic piezoelectric response of ceramics^22,37,38.

Energy dispersive X-ray spectroscopy (EDS) mapping further reveals the distinct distribution of the elements along grain boundaries in BCTZ–xAlN ceramics (Fig. 2g, Supplementary Figs. 28–30, and Supplementary Note 6). BCTZ–xAlN–50 ceramics exhibit nano-sized, plate-like BaAl₂O₄ particles that aggregate at the grain boundaries of the BCTZ matrix. In contrast, both BCTZ–xAlN–300 and BCTZ–xAlN–3000 ceramics feature BaAl₂O₄ particles with a micro-sized, strip-like shape. The average grain size distribution of these particles is summarized in the Supplementary Fig. 31. As the doping concentration increases or the average size of the AlN dopant particles increases, the average sizes of the BaAl₂O₄ secondary phase particles gradually increase. For example, at a doping concentration of x = 1.75, the grain size of BaAl₂O₄ in BCTZ−1.75AlN−3000 ceramic is 17 μm, which is larger than the 9 μm observed in BCTZ–1.75AlN−300 ceramic.

Scanning transmission electron microscopy (STEM) and the corresponding EDS elemental mapping of BCTZ–1.75AlN–300 ceramic were performed, with the results shown in Supplementary Fig. 32. The distribution of O and Ba elements appears uniform throughout the mapping area, whereas Al elements are predominantly concentrated in the central region. This concentration suggests the presence of two distinct co-existing phases. Furthermore, the high-resolution TEM (HRTEM) image of the BCTZ–1.75AlN–300 ceramics reveals two different sets of lattice fringes (Fig. 2h–i). The interplanar spacings of 0.28 nm correspond to both the (110) and (101) crystal planes of BCTZ, while those of 0.51 nm and 0.39 nm are associated with the (001) and (200) crystallographic planes of BaAl₂O₄, respectively. The corresponding selected area electron diffraction (SAED) patterns, recorded along the [$\bar{1}11$] and [$01\bar{1}$] directions, respectively, provide further confirmation of the high crystallinity of both BCTZ and BaAl₂O₄ phases.

Stress analysis

During the cooling stage of ceramic sintering, the disparity in thermal expansion coefficients between BCTZ (~11.3 × 10⁻⁶ K⁻¹) and BaAl₂O₄ (~5.8 × 10⁻⁶ K⁻¹) induces stress within the grains of BCTZ (Supplementary Fig. 33, Supplementary Note 7)^14,39. This stress plays a crucial role in determining the properties of ceramics. The stress distribution within BCTZ and BCTZ–xAlN ceramics was simulated by solving solid mechanics equilibrium conditions based on a finite element method using COMSOL Multiphysics® Software, as depicted in Fig. 3a, b. The irregular particles represent BCTZ, while the strip-like particles correspond to BaAl₂O₄. In the BCTZ ceramics, no significant concentration of internal stress is observed across the grains. In contrast, the BCTZ–xAlN ceramics exhibit a pronounced heterogeneous stress distribution concentrated within the BCTZ grains, whereas the stress within the BaAl₂O₄ particles remains relatively low.

Fig. 3: Stress in the BCTZ–xAlN ceramic.

Finite-element simulation of internal stress distribution within (a) BCTZ and (b) BCTZ–xAlN ceramics. c1 HRTEM image of BCTZ and corresponding strain maps obtained by GPA along the (c2) xx direction and (c3) yy direction. d1 HRTEM image of BCTZ–1.75AlN–300 and corresponding strain maps obtained by GPA along the (d2) xx direction and (d3) yy direction. Raman mapping images of ceramic integrating with the bands of 523 cm⁻¹: (e1) BCTZ, (f1) BCTZ–1.75AlN–300, (g1) BCTZ–1.75AlN–3000, and (h1) BCTZ–1.75AlN–50 ceramics. Raman mapping images of ceramic integrating with the bands of 390 cm⁻¹: (f2) BCTZ–1.75AlN–300, (g2) BCTZ–1.75AlN–3000, and (h2) BCTZ–1.75AlN–50 ceramics. Corresponding contour plots of specific Raman spectra from points A to B in the bands range of 100–1000 cm⁻¹ and 450–650 cm⁻¹, respectively: (e2–e3) BCTZ, (f3–f4) BCTZ–1.75AlN–300, (g3–g4) BCTZ–1.75AlN–3000, and (h3–h4) BCTZ–1.75AlN–50 ceramics. i Schematic of unit cell of pristine and strained BaTiO₃ bulk with (1 × 1) periodicities. j The total energies of the R, O, T, and C phases of BaTiO₃ as a function of strain. k The difference in the total energies of O, T, and C phases of BaTiO₃ and its R phase as a function of strain.

The heterogeneous stress within the BCTZ grains can lead to lattice strains. The geometric phase analysis (GPA) method was conducted to determine the internal stress in the selected areas of HRTEM images of BCTZ and BCTZ–1.75AlN–300 (Fig. 3c, d). Both the unstressed BCTZ and the stressed BCTZ–1.75AlN–300 ceramics could be clearly observed in in-plane (ε_xx) and out-of-plane (ε_yy) strain maps. The BCTZ ceramic exhibited no obvious lattice distortion and strain within the grain. In contrast, GPA analysis at the interface of BCTZ–1.75AlN–300 ceramics reveals a substantial distribution of internal strain within the BCTZ region. These images provide direct evidence of the presence of secondary-phase-induced stress in BCTZ–1.75AlN–300 ceramics.

Moreover, Raman mapping spectra were collected at various spatial positions across the surface of BCTZ and BCTZ–xAlN ceramics to visually characterize the effect of internal stress on these ceramics. For both BCTZ and BCTZ–xAlN ceramics, the modes at 150 and 210 cm⁻¹, attributed to A₁(TO), are the characteristics of the R phase of BaTiO₃⁴⁰. The modes at 290 cm⁻¹, assigned to E(TO + LO), along with those at 450 cm⁻¹, serves as indicators of the O phase⁴¹. The modes at 350, 523, and 720 cm⁻¹ correspond to B₁(TO)/E(TO + LO), E(TO)/A₁(TO), and A₁(LO)/E(LO), respectively, representing the distinctive characteristics of the T phase^6,41. These Raman spectroscopy results indicate that both BCTZ and BCTZ–xAlN ceramics exhibit multi-phase coexistence, which is consistent with the observations from the SXRD, ε_r–T and tanδ–T. Figure 3e1 presents a color-coded Raman map that depicts the intensity around 523 cm⁻¹ within a 20 × 20 μm² region of the BCTZ ceramic. Specific Raman spectra and the corresponding contour map from points A to B are detailed in the Supplementary Fig. 34 and Fig. 3e2–e3. Although there is a slight variation in peak intensity for BCTZ at all selected points, the positions of the peaks remain largely unchanged. This indicates a uniform distribution of stress within BCTZ ceramic, without any significant stress concentrations.

For the BCTZ–1.75AlN–300 ceramics, an additional Raman map depicting intensity around 430 cm⁻¹ was collected (Fig. 3f and Supplementary Fig. 35). The peaks at 244 cm⁻¹, 430 cm⁻¹, 580 cm⁻¹, 670 cm⁻¹, and 780 cm⁻¹ correspond to the F_2g(1), E_g, F_2g(2), F_2g(3), and A_1g modes, respectively, indicative of BaAl₂O₄⁴². By utilizing the characteristic modes of both BCTZ and BaAl₂O₄, we identified a single BaAl₂O₄ grain and two BCTZ grains along trajectory from point A to B. Changes induced by stress changes in the lattice of ceramic can affect the vibrational frequencies of phonon modes, leading to shifts in vibrational modes^43,44. The E(TO)/A₁(TO) mode at ~523 cm⁻¹ is used to characterize the stress due to its relatively high intensity in all Raman spectra. Notably, for BCTZ–1.75AlN–300 ceramic, there is a noticeable shift in the peak positions of the mode at ~523 cm⁻¹ (indicated by red arrow), indicating the presence of stress within the BCTZ particles. Similarly, noticeable peak shift in the BCTZ–1.75AlN–3000 sample (marked by the red arrow) also indicates stress within the BCTZ particles due to micro-sized BaAl₂O₄ second phase (Fig. 3g and Supplementary Fig. 36). This stress is particularly prominent in samples with larger AlN particle sizes, where the formation of the secondary phase is more pronounced, contributing to substantial stress within the BCTZ matrix. This stress distorts the perovskite structure and significantly alters the lattice distortion in the ROT phases, thereby enhancing the T_C of the ceramic. It also modifies the phase fractions of the ROT phases, improving the piezoelectric response. Consequently, both T_C and the piezoelectric response are simultaneously enhanced in the ceramics, consistent with observations of T_C enhancement in stressed ferroelectric thin films^18,20. Conversely, due to nano-sized BaAl₂O₄ second phase, BCTZ–1.75AlN–50 ceramic exhibits relatively uniform unstressed BCTZ particles with no significant shift in the peak positions (Fig. 3h and Supplementary Fig. 37).

To further elucidate the effects of strain on BCTZ ceramics, we employed first-principles calculations within DFT to assess its influence on phase structural stability. Using BaTiO₃ as a representative model, different strain (0% to 10%) were applied to the unit cell (Fig. 3i). We calculated the total energies of the R, O, T, and paraelectric C phases of BaTiO₃, with higher total energy indicating lower stability (Fig. 3j). The results show that with 0−10% strain, the R phase consistently has the lowest total energy and highest stability, followed by the T phase, then the O phase, with the C phase being the least stable. The energy differences (ΔE) between the O, T, C phases and the R phase increase with strain, highlighting the higher stability of the R phase under stress compared to the other phases (Fig. 3k). This finding aligns well with experimental observations in BCTZ−xAlN−300 ceramics with x < 1.75. An increase in stress correlates with a gradual increase in the R phase content, a slight decrease in the O phase content, and a reduction in the T phase content, indicating that strain contributes to the enhanced piezoelectric response. Notably, the C phase is the most destabilized under strain, suggesting that strain inhibits the C phase’s increase due to temperature rise, thereby improving the ceramic’s temperature stability.

Local electric field and defect configuration

To further investigate the impact of the secondary phase BaAl₂O₄ on BCTZ–xAlN ceramics, a spatially resolved analysis of surface potential was conducted using KPFM to assess the local contact potential difference within the BCTZ–1.75AlN–300 ceramic. The contact potential difference (V_CPD) quantifies the variation in work functions between the AFM tip (W_tip) and the sample surface (W_sample), which can be expressed by the following equation:

$${V}_{{CPD}}=\frac{{W}_{{tip}}-{W}_{{sample}}}{e}$$

(1)

Where e is the elementary charge. Examination of the topography of the BCTZ–1.75AlN–300 ceramic, as depicted in Fig. 4a, reveals a single BaAl₂O₄ grain situated at the boundary of BCTZ grain. The corresponding surface potential map, illustrated in Fig. 4b, highlights the potential contrast between the BaAl₂O₄ and BCTZ grains, with the BaAl₂O₄ exhibiting a higher surface potential characterized by a smaller work function. The potential profile illustrates a local potential difference of ~65 mV between the BCTZ matrix and the BaAl₂O₄, as evident from the potential versus distance curve (Fig. 4c). This observed potential disparity induces a local electric field at the interface of BCTZ and BaAl₂O₄. However, this phenomenon is not observed in BCTZ ceramic (Supplementary Fig. 38, Supplementary Note 8). Notably, in the P‒E loops of the AlN-doped BCTZ ceramics, the enhanced asymmetric characterization of the coercive electric field also confirms the presence of an internal electric field (Supplementary Fig. 39 and Supplementary Note 9). This internal electric field potentially leads to subtle structural distortion in BCTZ, further influencing the phase structure of ceramic^45,46,47 and stabilizing the local ordering of ferroelectric domains⁴⁸. These effects enhance the piezoelectric response of BCTZ–xAlN ceramics.

Fig. 4: Local electric field and defect configuration of ceramics.

a Topographies, and (b) surface potential mapping of BCTZ–1.75AlN–300 ceramics measured using KPFM. c Surface potential profiles along the long edge of the region defined by the blue dashed box, as derived from the surface potential mappings. Inset is the schematic diagram of the KPFM testing. d XPS spectra of O1s in BCTZ and BCTZ–1.75AlN–300 ceramics. e The concentration of oxygen vacancies (O2/O1 ratios). f Impedance spectra of BCTZ, BCTZ–xAlN–300 ceramics at 500 °C. g Fitted E_a values of grains and grain boundaries in BCTZ–xAlN–300 ceramics.

To investigate the chemical states of the BCTZ–xAlN ceramics, the O elements in all BCTZ–xAlN–300 ceramics were examined using XPS spectra (Fig. 4d and Supplementary Fig. 40). The spectra can be fitted with four distinct peaks at binding energy of 529.0, 530.5, 531.5 and 532.5 eV, assigned to lattice oxygen (O1), highly oxidative oxygen species closely related to the surface oxygen vacancies (O2), the hydroxyl groups or the surface adsorbed oxygen (O3), and adsorbed molecular water (O4), respectively^49,50,51,52. The relative concentration of surface oxygen vacancies was estimated from the ratio (O2/O1) of oxygen species related to surface oxygen vacancies to lattice oxygen species⁵¹. It is observed that the oxygen vacancies concentration in BCTZ–xAlN–300 ceramics initially decreases and then increases with increasing x, reaching a minimum at x = 1.75 (Fig. 4e). These oxygen vacancies concentration change may be related to the incorporation of Al³⁺ into the lattice of BCTZ.

Initially, the formation of oxygen vacancies in all ceramics can be attributed to oxygen evaporation during the high-temperature sintering process, which can be described by the following reaction.

$${{{\rm{O}}}}_{{{\rm{O}}}}^{\times }\to \frac{1}{2}{{{\rm{O}}}}_{{{\rm{2}}}}+{{{\rm{V}}}}_{{{\rm{O}}}}^{\cdot \cdot }+2{{{\rm{e}}}}^{{\prime} }$$

(2)

In ceramics, two types of charge carriers can be generated: oxygen vacancies (${{\mbox{V}}}_{{\mbox{O}}}^{\cdot \cdot }$) and free electrons (e′). For BCTZ–xAlN–300 ceramics with x ranging from 0.25 to 2.00, AlN doping leads to incorporation of Al³⁺ into the lattice, occupying A-sites in the perovskite structure and forming the extrinsic defect ${{\mbox{Al}}}_{{\mbox{Ba}},{\mbox{Ca}}}^{\cdot }$. This incorporation alters the phase structure of the ceramics, enhancing the piezoelectric properties but also causing a decrease in the T_C of the ceramics. Simultaneously, the oxygen vacancies concentration in the ceramics is significantly reduced. The corresponding defect chemical equation is as follows.

$$3{{{\rm{A1N}}}}+{{{\rm{B}}}}{{{\rm{a}}}}_{{{\rm{Ba}}}}^{\times }+\frac{5}{2}{{{\rm{O}}}}_{{{\rm{2}}}}+{{{\rm{V}}}}_{{{\rm{O}}}}^{\cdot \cdot }+{{{\rm{e}}}}^{{\prime} }\to {{{\rm{A}}}}{1}_{{{\rm{Ba}}}}^{\cdot }+{{{\rm{BaA}}}}{1}_{2}{{{\rm{O}}}}_{4}+\frac{3}{2}{{{\rm{N}}}}_{2}+{{{\rm{O}}}}_{{{\rm{O}}}}^{\times }$$

(3)

Hence, the BCTZ–1.75AlN–300 ceramic exhibits a significantly reducing oxygen vacancies concentration, which also contributes to the enhanced resistivity of ceramics. This is verified by the impedance spectra (Fig. 4f and Supplementary Fig. 41). Furthermore, these impedance spectra are fitted to evaluate the electrical microstructure of ceramics (Fig. 4g, Supplementary Figs. 42‒43, and Supplementary Note 10). For all ceramics, the activation energy (E_a) at the grain boundaries is higher than within the grains, indicating that oxygen vacancies migration is more difficult at the grain boundaries. As AlN doping concentration increases, the E_a for both grains and grain boundaries in BCTZ–xAlN–300 ceramics increases significantly, indicating oxygen vacancy migration suppression. Notably, the BCTZ–1.75AlN–300 ceramic exhibits the highest E_a for both grains and grain boundaries, contributing to its exceptionally high piezoelectric performance.

Discussion

In this study, we investigated the influence of AlN with varying particle sizes and doping concentrations on BCTZ-based ceramics. The addition of AlN led to the formation of BaAl₂O₄ secondary phases and incorporation of Al³⁺ into lattice (Fig. 5a). Specifically, the reaction between AlN and BCTZ resulted in BaAl₂O₄ particles forming at the BCTZ grain boundaries, increasing stress within the BCTZ grain (stress effect, Fig. 5b). This stress caused lattice distortion and altered the phase fractions, resulting in simultaneous increases in both T_C and d₃₃. Additionally, the local electric field at BCTZ and BaAl₂O₄ interface enhanced d₃₃ (local electric field effect, Fig. 5c). A small amount of Al³⁺ ions was incorporated into the A-sites of the perovskite lattice (doping effect, Fig. 5d), further modifying the phase structure and improving d₃₃ but often reducing T_C. Both the AlN doping levels (mole ratio) and particle size (reaction activities) generated different stress, local electric field, and doping effects, and thus significantly influencing piezoelectric properties and T_C of the ceramics.

Fig. 5: Schematic illustration for stress, local electric field, and doping effects.

a Schematic illustration of AlN doped BCTZ–xAlN ceramics. Schematic illustration of the influence of (b) stress, (c) local electric field, and (d) doping effects on the performance of ceramics as functions of AlN doping concentration and particle size. The z-axis indicates performance parameters such as d₃₃ or T_C, where red represents higher values and purple indicates lower values.

In the BCTZ–xAlN–3000 ceramics, as the AlN doping concentration increased, d₃₃ initially decreased, then increased, and finally decreased again, while T_C initially rose and subsequently declined (Fig. 1a). For these ceramics, the secondary phase size reached the micrometer scale and significantly increased with higher doping levels, exceeding the dimensions of secondary phases in the BCTZ–xAlN–50 and BCTZ–xAlN–300 samples at the same doping concentrations. This negatively affected the piezoelectric performance of the ceramics. Additionally, the low reactivity of these large particles limited Al³⁺ doping effect, moderating the improvement in d₃₃. At lower doping levels, a substantial amount of the secondary phase formed with minimal Al³⁺ incorporation, leading to a significant rise in T_C due to the dominated stress effect. However, the excessive formation of large BaAl₂O₄ particles inhibited the increase in d₃₃. With further increases in doping concentration (x = 1.00), a slight enhancement in d₃₃ was observed, albeit accompanied by a minor decline in T_C, attributed to the moderate doping effect. At x = 1.75, the doping effect contributed to an increase in d₃₃ that was insufficient to offset the negative impacts of the excessive BaAl₂O₄ formation, resulting in decreases in both d₃₃ and T_C.

Conversely, in BCTZ–1.75AlN–50 ceramics, the introduction of highly reactive nanoscale AlN promotes greater Al³⁺ doping, leading to a significant enhancement of d₃₃. However, this is accompanied by a reduction in T_C due to insufficient stress generation. Thus, with increasing doping concentrations, the d₃₃ of BCTZ–xAlN–50 ceramics increases while the T_C decreases. These contrasting behaviors reveal a crucial aspect: smaller AlN particles tend to be overly reactive, while larger particles are less reactive, resulting in either excessive or insufficient doping effect, respectively. Similarly, lower doping levels lead to inadequate doping and stress effects, while higher doping levels cause excessive doping and stress.

Encouragingly, the high-performance observed in BCTZ–xAlN–300 ceramics can be attributed to an optimal balance between the size of the secondary phase and the incorporation of Al³⁺ into the lattice. As the doping levels increased, d₃₃ initially rose, reaching a maximum at x = 1.75, before declining (Fig. 1a). Meanwhile, T_C exhibited a slight decrease at lower doping levels, followed by an increase peak at x = 1.75, and a subsequent decrease. At low doping levels, the doping effect predominantly dictated the performance. As the doping concentration further increased, the combined effects of stress, electric field, and doping contributed to the overall enhancement of performance. The use of 300 nm AlN powder at a concentration of 1.75 mol% provided the most favorable conditions for enhancing the properties of the ceramics. Moderate particle size generated sufficient stress to enhance T_C, while still allowing controlled Al³⁺ doping, resulting in a simultaneous improvement in both T_C and d₃₃. However, this performance optimization was not without limitations. For BCTZ–2.00AlN–300 ceramics, excessive doping led to a reduction in lattice distortion, consequently diminishing d₃₃. The stress effect insufficiently counteracts the negative impacts of excessive doping effect, resulting in a significant reduction in T_C.

Through optimizing stress, local electric field, and doping effects, the highest piezoelectric properties (d₃₃ = 650 ± 16 pC N⁻¹, d₃₃^* = 1070 pm V⁻¹), a high T_C (96.5 ± 1.0 °C), and remarkable temperature stability, are achieved in BCTZ–1.75AlN–300. Our findings elucidate the mechanisms behind the exceptional piezoelectric properties and desirable T_C of the BCTZ–xAlN ceramics, underscoring the effectiveness of bulk stress engineering through the introduction of a secondary phase in enhancing lead-free piezoelectric ceramics. This approach paves the way for developing high-performance piezoelectric materials suitable for a wide range of temperature applications.

Methods

Materials

Barium carbonate (BaCO₃, 99%), calcium carbonate (CaCO₃, 99%), and aluminum nitride with three different kinds of average sizes (AlN, 50 nm/300 nm/3000 nm, 99%) were purchased from Shanghai Macklin Biochemical Co, Ltd Barium zirconium oxide (BaZrO₃, 99%) was purchased from Alfa Aesar Co, Ltd. Titanium dioxide (TiO₂, 99%) was purchased from Tianjin Fuchen Chemical reagents Co, Ltd. All the chemicals were used as received without further purification.

Preparation of ceramics

Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃–x mol% AlN (300 nm) (BCTZ–xAlN–300, x = 0, 0.25, 0.50, 0.75, 1.00, 1.25, 1.50, 1.75, and 2.00) ceramics were prepared by a conventional solid-state reaction method. The raw materials (BaCO₃, CaCO₃, BaZrO₃, and TiO₂) were weighed according to stoichiometric ratios and ball-milled in ethanol for 12 h. The slurry was dried at 75 °C and then calcined at 1150 °C for 3 h, yielding single-phase BCTZ powders. Afterwards, both the BCTZ powders and the commercial AlN nanoparticles with the average size of 300 nm were weighed according to the formula of BCTZ–xAlN–300 (x = 0, 0.25, 0.50, 0.75, 1.00, 1.25, 1.50, 1.75, and 2.00). Each mixture was ball-milled again in ethanol for 12 h and dried at 75 °C. All specimens were pressed into disks with a diameter of 10 mm under 10 MPa and sintered at 1400–1450 °C for 3 h in air. The heating rate was 3 °C min⁻¹ and cooling rate was 1 °C min⁻¹ (sintering temperatures are summarized in Supplementary Table 2). Similarly, Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃–x mol% AlN (50 nm) (BCTZ–xAlN–50, x = 0.25, 1.00, 1.50, and 1.75) and Ba_0.85Ca_0.15Ti_0.9Zr_0.1O₃–x mol% AlN (3000 nm) (BCTZ–xAlN–3000, x = 0.25, 1.00, 1.50, and 1.75) were prepared using AlN nanoparticles with sizes of 50 nm and 3000 nm, respectively, following the same method.

Characterization

High-resolution synchrotron X-ray diffraction (SXRD) and in-situ variable temperature SXRD measurements were performed at the Powder Diffraction beamline of the Australian Synchrotron, using a wavelength of 0.590246 Å and a scan speed of 0.2256° min⁻¹. The crystal structures and lattice parameters of the ceramics were analyzed using GSAS-EXPGUI Rietveld refinement software^53,54,55. Scanning electron microscopy (SEM) images and EDS mapping were obtained using a field emission SEM (Gemini 300, Zeiss, Germany). Transmission electron microscopy (TEM) images, HRTEM images, and EDS mapping were obtained by a JEOL JEM2100F TEM with an EX-24063 JGT EDS. The average relative density and porosity of the ceramics were determined using the Archimedes method. The residual stress and phonon anharmonicity in BCTZ–xAlN ceramics were evaluated by a Raman spectroscope (LabRAM HR Evolution, Horiba, France). X-ray photoelectron spectroscopy (XPS) was performed on a photoelectron spectrometer (ESCALAB, 250Xi, Thermo Fisher, USA). Thermogravimetric analysis (TG) and differential scanning calorimetry (DSC) were carried out for BCTZ and BCTZ–1.75AlN–300 powders at a heating rate of 20 °C min⁻¹ (STA449F5 Jupiter, NETZSCH, Germany). Thermal expansion coefficient of ceramic was measured using a Thermal dilatometer (DIL1412STD, Orton, USA). Switching Spectroscopy Piezoelectric Force Microscope (SS-PFM) and Kelvin probe force microscopy (KPFM) measurements were performed using a Dimension Icon Atomic Force Microscope (Bruker, USA).

Electrical property measurements

For electrical measurements, an Ag electrode (~0.8 cm in diameter, ~17 µm thick) was screen-printed on the ceramic surface and then fired at 550 °C for 30 min. The ceramic disks were polarized in a silicone oil bath at room temperature under a DC field of 4 kV mm⁻¹ for 30 min. Quasi-static piezoelectric coefficient (d₃₃) was subsequently quantified using a quasi-static piezoelectric constant meter (ZJ–3A, Institute of Acoustics, China). Planar electromechanical coupling factor (k_p), relative permittivity (ε_r), dielectric loss (tanδ) at 1 kHz, and impedance spectra were determined using a capacitance meter (Agilent 4294 A, Agilent, Santa Clara, USA). Temperature-dependent dielectric properties were measured from −50 °C to 200 °C (Agilent E4980A, Agilent, Santa Clara, USA). Ferroelectric polarization–electric field hysteresis (P–E) loops and electric strain–electric field (S–E) curves were measured using a ferroelectric analyzer (aixACCT TF Analyzer 3000, Aachen, Germany) with a laser interferometer vibrometer (SP–S120/500, sios Mebtechnik GmbH, Germany) at a frequency of 1 Hz. In-situ d₃₃ of ceramics was measured using the high-temperature in-situ piezoelectric tester (TZDM-D33T, Harbin Julang Technology Co. Ltd, China).

Computational methodology

The total energies of the R, O, T, and C phases of bulk BaTiO₃ under different strain conditions were calculated using first-principles simulations within density functional theory (DFT)^56,57. The projected augmented wave pseudopotentials method was used as implemented in the Vienna Ab initio Simulation Package (VASP)^58,59. The exchange correlation energy was calculated using the generalized gradient approximation (GGA) of the Perdew-Burke-Ernzerhof form ref. ⁶⁰. The plane wave cutoff energy was set to 500 eV. The Monkhorst-Pack k-point mesh was taken as 13 × 13 × 13⁶¹. Utilizing the conjugate gradient method, the plane lattice constant and atomic coordinates were fully relaxed until the energy and force converged to 10⁻⁵eV and 10⁻² eV/Å, respectively.

Reporting summary

Further information on research design is available in Nature Portfolio Reporting Summary linked to this article.