Introduction

Piezoelectric materials interconvert electrical and mechanical energies and are widely used for sensors, actuators, transducers, and energy harvesters1,2,3. These piezoelectric devices are crucial in various fields, including biomedical engineering, aerospace, and wireless energy transmission4. Piezoceramics, represented by lead zirconate titanate (Pb(Zr, Ti)O3, PZT) family, dominate the piezoelectric market due to excellent electrical properties and large-scale industrial production5. Considering the toxicity of lead (Pb) and the need for environmental protection, research on lead-free piezoceramics to replace Pb-based ones is imperative6. Among these, potassium sodium niobate ((K, Na)NbO3, KNN)-based lead-free piezoceramics stand out due to the significant progress in their piezoelectric coefficient (d33), electro-strain, and temperature stability achieved through phase boundary engineering (PBE), texturing, defect engineering, and composite ceramics7.

For high-power applications such as ultrasonic motors, cleaners, and transducers, piezoceramics are expected to have both high d33 and mechanical quality factor (Qm) (also known as hard piezoceramics) as they operate in resonant mode5. High d33 ensures the good electromechanical properties, while high Qm reduces the heat generation caused by dissipated energy8. However, achieving a balance between d33 and Qm is highly challenging because they have different preferences for extrinsic contributions. Specifically, high d33 prefers more ferroelectric domain switching, while high Qm is usually achieved by pinning domain switching5. This imbalance is more pronounced in KNN-based piezoceramics. For example, although the PBE strategy significantly increases d33 to the level of 400–600 pC/N, it is accompanied by an extremely low Qm (<50)9,10. Traditional acceptor doping (i.e., copper Cu and manganese Mn) and the newly-proposed isolated-oxygen-vacancy strategy greatly improve Qm but fail to ensure high d335,11. Additionally, traditional acceptor doping is mainly implemented on pristine KNN ceramics with low d33 values (e.g., <150 pC/N), resulting in even worse d33 after acceptor doping as expected.

It is known that the traditional orthorhombic-tetragonal (O-T) PBE reduces the energy barrier between O and T phases and promotes domain switching, thereby enhancing dielectric and piezoelectric properties12,13. However, the promoted domain switching also intensifies energy dissipation and seriously worsens Qm (Fig. 1a). Here, we propose a new strategy to balance d33 and Qm of KNN-based ceramics, namely, local Cu-acceptor defect dipoles embedded O-T PBE. Based on the O-T PBE, we innovatively embed local Cu-acceptor defect dipoles for two purposes. On the one hand, doping an appropriate Cu2+ has little effect on the phase transition temperature and can retain the room-temperature O-T PBE14,15. On the other hand, the valence state difference between Cu2+ and Nb5+ provides a dual effect. Specifically, Cu-acceptor doping forms two defect configurations, namely, dimeric \({({{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{\bullet \bullet })}^{{\prime} }\) and trimeric \({({V}_{O}^{\bullet \bullet }-{{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{\bullet \bullet })}^{\bullet }\)16,17,18. The former primarily forms an electric dipole moment (i.e., defect dipoles, PD) to pin domain switching in addition to an slight elastic distortion, whereas the latter mainly has an elastic distortion with no electric dipole. The local elastic distortion can cause the increased local structural heterogeneity, consequently further reducing the domain size and lowering the polarization energy barrier19,20. That is to say, our strategy can construct a “flat and deep” potential well. The “flat” region is contributed by the retained O-T PBE and the introduced local structural heterogeneity, ensuring high d33, while the “deep” wall comes from the dimeric PD that prevents long-range polarization rotation and pins adjacent domain wall motion (Fig. 1b). Besides, CuO, as the sintering aid, can enhance the sinterability and consequently improve density and grain size of KNN-based ceramics21. The appropriate increase in grain size can reduce the lattice stress caused by the domain switching, favoring piezoelectric properties22. Thus, both high d33 and Qm are expected in such a system. We carefully select 0.95(K0.48Na0.52)NbO3-0.05(Bi0.5Na0.5)HfO3-0.2%molFe2O3-x%molCuO (abbreviated as KNN-BNH-xCu) ceramics to implement our strategy because the KNN-BNH ceramic matrix exhibits the O-T PBE and high piezoelectricity23,24. Encouragingly, KNN-BNH-xCu ceramics successfully validate our strategy and exhibit balanced d33 (~340 pC/N) and Qm (~256) at x = 1.0. Moreover, the related physical mechanisms are analyzed from the multi-scale perspective (i.e., local-mesoscopic-macroscopic) by using combined theoretical and experimental techniques.

Fig. 1: Schematic illustrations of conventional O-T PBE and our local Cu-acceptor defect dipoles embedded O-T PBE strategy.
Fig. 1: Schematic illustrations of conventional O-T PBE and our local Cu-acceptor defect dipoles embedded O-T PBE strategy.
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a For the conventional O-T PBE strategy, only d33 can be enhanced due to the easy polarization rotation and domain switching, leaving an unbalanced development of d33 and Qm. b For our local Cu-acceptor defect dipoles embedded O-T PBE strategy, both d33 and Qm can be simultaneously improved because of the coupling effect of local Cu-acceptor defect dipoles and O-T PBE. Note that PBE is the abbreviation of phase boundary engineering.

Results and discussion

Phase structure

X-ray diffraction (XRD) patterns show that all samples have a typical perovskite structure. At high doping content, a minor CuO phase is observed (Supplementary Fig. 1), indicating that Cu2+ ions have entered into the lattice of KNN-BNH ceramic matrix. Enlarged XRD patterns in the 2θ range of 45–46° show the changes in position and relative intensity of {200}pc diffraction peaks with different doping levels (Fig. 2a). When x = 0-1, the diffraction peaks first shift a lower angle and then to a higher angle at x = 1.5-3. This can be explained by the replacement of Nb⁵⁺ ions (with a radius of 0.64 Å when coordination number is 6) by Cu²⁺ ions (with a radius of 0.73 Å when coordination number is 6) according to Bragg’s Law (2dsinθ = λ)25. As x ranges from 1.5 to 3, the diffraction peaks shift to a higher angle, which may be due to the existence of the secondary CuO phase affecting the lattice structure. The relative intensity of the {200}pc diffraction peaks also changes with x. For the undoped sample (x = 0), the intensity ratio of (200)pc to (002)pc peak is approximately 2:1, indicating the existence of the T phase10. As x increases, the I200/I002 ratio first increases in the range of x = 0-1 but then changes irregularly in the range of x = 1.5-3. Thus, the sample with x = 1 seems be to a critical point.

Fig. 2: Phase structure study.
Fig. 2: Phase structure study.
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a Enlarged view of XRD pattern with 2θ = 45-46°. b ε′−T, ε″−T, and tan δ−T curves of unpoled samples, measured at 500 kHz. Frequency-dependent ε′-T, ε″-T, and tan δ-T curves of KNN-BNH-xCu with (c) x = 0 and (d) x = 1. Enlarged view of Rietveld refinements for powder XRD patterns of KNN-BNH-xCu with (e) x = 0 and (f) x = 1. (g) Variations of υ1 mode position with x. The insets of (e, f) show the phase fraction of O and T phases, and the inset of (g) is the schematic of υ1 mode of NbO6 octahedron.

To further study the phase structure, temperature-dependent dielectric properties (including εʹ, εʺ, and tan δ) are analyzed. All samples show a dielectric anomaly near 65 °C in their εʹ-T curves due to the O-T phase transition (i.e., TO-T) (Fig. 2b), coinciding with their T-phase featured {200}pc diffraction peaks at room temperature. Notably, Cu-doped KNN-BNH ceramics show another weak dielectric anomaly below the TO-T in εʹ-T curves, which is more significant in the corresponding εʺ-T curves. Specifically, the εʺ-T curve of the undoped KNN-BNH ceramic shows an anomaly near TO-T, while the εʺ-T curves of all Cu-doped KNN-BNH ceramics exhibit an anomaly near −50 °C (much lower than their TO-T). Obviously, these low-temperature anomalies of Cu-doped samples are not related to any phase transitions.

To further analyze the low-temperature anomalies, εʹ-T and εʺ-T curves are measured over a frequency range of 0.1–500 kHz (Fig. 2c, d and S2). The undoped KNN-BNH ceramic shows dielectric anomalies near TO-T in both εʹ-T and εʺ-T curves, similar to canonical ferroelectric ceramics. The abnormally high εʺ and tan δ values near TO-T under low frequency can be attributed to space charges, as both values reduce sharply but do not shift with increasing frequency14,26. In contrast, doping CuO gradually broadens the dielectric anomaly of TO-T, especially at x = 2 and 3 (Supplementary Fig. 2), indicating the enhanced degree of dielectric relaxation. All Cu-doped KNN-BNH ceramics exhibit an obvious dielectric anomaly near −50 °C in their εʹ-T curves (especially under low frequencies). Importantly, these low-temperature anomalies are highly frequency-dependent in their εʺ-T curves, similar to the dielectric relaxation behavior of relaxor ferroelectrics27,28. For example, at x = 1, the temperature difference (∆T) of the local maximum εʺ is about 18 °C (Fig. 2d). Similar phenomena are also observed when comparing Cu-doped and Cu-undoped KNN ceramics (Supplementary Fig. 3). Meanwhile, Cu-doped KNN-BNH ceramics show higher tan δ values at T < −50 °C (far below TO-T) (Fig. 2b), indicative of more local structural heterogeneities19. Therefore, these low-temperature anomalies are not related to phase transitions but inherently attributed to the addition of CuO, which disrupts the long-range ferroelectric order and induces more local structural heterogeneities.

Due to the highly close TO-T values, Rietveld refinements are carried out for powder XRD patterns of ceramics with x = 0 and 1 (Fig. 2e, f and S5, Supplementary Table 1). A combination of O and T phases can well describe the measured data, confirming the coexistence of O-T phases. Although the TO-T values remain almost unchanged, the Rietveld refinements show a slight increase in the T phase content from 40.4% (at x = 0) to 46.8% (at x = 1), which is consistent with the variations in bulk XRD patterns. To further explore the influence of CuO on the lattice structure, structure-sensitive Raman spectra are studied (Supplementary Fig. 6)29. The Lorentz function is used to fit the Raman peaks within 500–700 cm−1 to extract the υ1 vibration mode that represents the double-degenerate symmetric O-Nb-O stretching vibration30. As x increases, υ1 first shifts to a higher wavenumber at x = 0–1 and then to a lower wavenumber at x = 1.5–3 (Fig. 2g). A shift of υ1 to higher wavenumbers indicates the enhanced B-O bond energy31. Therefore, doping CuO enhances the Nb-O bond energy except for at high content (i.e., x = 3). The detailed explanation will be discussed later. Moreover, the Raman spectra also prove that x = 1 is a critical point.

Microscopic-mesoscopic-local structures

Multi-scale structural studies are then conducted to reveal the physical mechanisms. Both average grain size and relative density increase with x, well aligning with our expectation (Supplementary Fig. 7 and Supplementary Table 2). The excess CuO accumulates at the grain boundaries and causes the relative density exceeding 100% at x = 3 due to its much larger density (6.3 ~ 6.9 g/cm3) than KNN-BNH ceramic matrix (4.6 g/cm3). Both SEM-EDS and TEM-EDS observations confirm uniform distribution of all elements in the KNN-BNH-1Cu sample, further proving that Cu atoms have completely entered the lattice of KNN-BNH matrix (Supplementary Fig. 8–11 and Supplementary Table 3). Subsequently, ferroelectric domains are studied by performing out-of-plane PFM (OP-PFM) measurements. The domain configurations highly depend on the content of CuO. The Cu-undoped sample (i.e., x = 0) exhibit complex domain configurations containing both small nano-sized domains (~170 nm) and large micron-sized domains (~1 μm) (Fig. 3a), akin to those of the reported ceramics with multi-phase coexistence32,33. In contrast, the KNN-BNH-1Cu sample shows overwhelmingly high-density nano-sized domains with a size of ~100 nm (Fig. 3b). These nano-sized domains have been reported to rapidly respond to external electric fields and thus enhance piezoelectricity34. However, the KNN-BNH-3Cu sample displays the fragmented and isolated domains and much low-density domain walls (Fig. 3c), consistent with its increased degree of dielectric relaxation. The domain configurations are further confirmed in randomly selected different areas (Supplementary Fig. 12–14). Switching spectroscopy PFM (SS-FPM) measurements are performed on several random points to study local domain switching (Fig. 3 and Supplementary Fig. 15). All samples show amplitude loops with a typically butterfly shape and phase loops with a maximum phase-contrast of 180°, indicating sufficient local domain switching. It is noted that adding CuO makes the amplitude loops become asymmetrical, indicating the existence of local defect dipoles. Interestingly, as x increases, the local coercive voltage (i.e., ∆VC = V + C-V−C) decreases, while the maximum amplitude value increases. Thus, doping with CuO not only introduces the defect dipoles but also facilitates local domain switching, which is expected to maintain high d33 and improve Qm simultaneously.

Fig. 3: Mesoscopic ferroelectric domain structure.
Fig. 3: Mesoscopic ferroelectric domain structure.
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Amplitude and phase images of OP-FPM measurements and SS-PFM measurements of (a) KNN-BNH-0Cu, (b) KNN-BNH-1Cu, and (c) KNN-BNH-3Cu ceramics. The scanning area of is 5 μm × 5 μm, and the white bars represent 1 μm. For a better view, the amplitude images are shown with colorful and gray, respectively.

We further conduct TEM/STEM imaging for the KNN-BNH-1Cu sample to study the local structure. TEM images reveal multifarious ferroelectric domains, including irregular nano-domains of ~50 nm (Fig. 4a), hierarchical nano-domains of ~20–50 nm (Fig. 4b and Supplementary Fig. 16), and lamellar domains ranging from 200–400 nm (Supplementary Fig. 16). These nanodomains were widely reported in lead-free and lead-based piezoceramics with high piezoelectricity35, indicating superior piezoelectric properties at x = 1. Then the local atomic structure were studied using the STEM-HAADF measurements and simulating polarization vectors (Fig. 4c, d). The calculated polarization vectors show the coexisting R/O and T phases along the [100] direction, with an average polarization displacement of 27.0 pm (Fig. 4c–e). The [110]-orientated STEM HAADF image also shows the nanoscale coexistence of R, O, and T phases (Supplementary Fig. 18). The phase gradually transforms on the nanoscale, leading to abundant nano-domains and easy local domain switching. Interestingly, the local T phase mainly locates at the areas with more A-site atom intensity (Fig. 4f). This phenomenon may be rationalized by the large ionic radius of K+ ions and remarkable off-center displacement of Bi3+ ions, building larger cube-octahedral cages that allow Nb5+ to displace along the [001] direction36,37. Additionally, the c/a ratio and lattice parameter maps reveal distinctly local structural heterogeneities associated with the addition of these chemical dopants like BNH and CuO (Fig. 4g, h).

Fig. 4: Mesoscopic and atomic scale study of KNN-BNH-1Cu ceramics.
Fig. 4: Mesoscopic and atomic scale study of KNN-BNH-1Cu ceramics.
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a, b TEM images of KNN-BNH−1Cu ceramics. c STEM-HAADF image along the [100] direction, covered by simulated polarization vectors. d Schematic diagrams of the ABO3 unit cell along the [100] zone axis and the color wheel of simulated polarization vectors in (c). e Statistical diagrams of polarization displacement and angle of simulated polarization vectors in (c). (f-h) A-site atom intensity map, c/a ratio map, and lattice parameters maps corresponding to (c). The statistics of domain size in (a, b) is along the black line segments.

Electromechanical properties

As x increases in KNN-BNH-xCu ceramics, d33 gradually decreases while Qm first increases and then decreases, with optimal properties at x = 1 (i.e., d33 = 340 pC/N and Qm = 256) (Fig. 5a). For KNN-based piezoceramics, a d33 higher than 300 pC/N is usually accompanied by a low Qm value (e.g., <150). Thus, the balanced d33 and Qm values obtained here are superior to those of previously-reported KNN-based ceramics (Fig. 5b)11,38,39,40,41,42,43. The hardening of piezoceramics increases the phase angle, which is also observed in these samples. The KNN-BNH-1Cu sample shows a higher phase angle of 86.7° compared to the KNN-BNH sample (Fig. 5c, d), indicating a more saturated poling process44. This increased phase angle is also observed for other Cu-doped KNN-BNH samples even at the high doping content (Supplementary Fig. 20). A radar map is used to compare the electrical properties between KNN-BNH-0CuO and KNN-BNH-1CuO samples (Fig. 5e). The KNN-BNH-1CuO sample has a d33 value only decreasing by 1/5 but a Qm increasing by a factor of 4, proving much better overall performance.

Fig. 5: Electromechanical properties.
Fig. 5: Electromechanical properties.
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a Variations of d33 and Qm values with x. The error bar represents the standard deviation. b Comparison of d33 and Qm values between this work and other reported KNN-based ceramics. Frequency-dependent phase angle and impedance spectra of poled (c) KNN-BNH-0Cu and (d) KNN-BNH-1Cu ceramics. e Comparison of comprehensive properties of KNN-BNH-0Cu and KNN-BNH-1Cu ceramics. f P-E loops and (g) polarization values of unpoled KNN-BNH-1Cu under the first two electric cycles. h Schematic diagrams showing domains and defect dipoles corresponding to different points in (f). fa and fr in (c) are the anti-resonance and resonance frequencies, respectively. PS and PD in (h) represent the spontaneous polarization and the defect dipole polarization, respectively.

The hardening effect stems from the defect dipoles studied by P-E loops of unpoled samples under the first two electric cycles. The KNN-BNH-0Cu ceramics show highly similar P-E loops under two electric cycles (Supplementary Fig. 19). By contrast, the Cu-doped KNN-BNH ceramics with x = 0.5–1.5 show the decreased maximum polarization (Pmax) during the second electric cycle (Fig. 5f and Supplementary Fig. 19), which can be explained by the dynamic activities of defect dipoles. Taking the KNN-BNH-1Cu ceramic as an example (Fig. 5f–h):

  • At point a: ferroelectric domains and defect dipoles randomly distribute in the unpoled state, and defect dipoles align with the direction of domains according to the symmetry-conforming principle45;

  • At point b: applying a positive electric field of +30 kV/cm drives the most domains upwards, except for a few domains close to defect dipoles because orientations of these defect dipoles differ from the electric field;

  • At point d: applying a negative electric field of −30 kV/cm makes some domains downwards, but defect dipoles pin the switching of their adjacent domains, resulting in the smaller Pmax than that measured at point b;

  • At point f: applying a positive electric field of +30 kV/cm again makes less domains upwards due to the pinning of defect dipoles, resulting in much smaller Pmax than that measured at point b.

  • At point h: Since the pinning effect of defect dipoles gets saturated, the unpinned domains can rapidly switch under external electric fields, resulting in the almost equal Pmax values at points f and h.

For Cu-doped KNN-BNH ceramics with x = 2-3, the first two P-E loops highly overlap due to the enhanced dielectric relaxation behavior offsetting the pinning effect of doping CuO.

Formation mechanisms of defect dipoles

X-ray absorption fine structure (XAFS) spectra measurements and the first-principle calculations are further used to reveal the formation mechanism of Cu-related defect dipoles. Cu K-edge and Nb K-edge XAFS measurements (Fig. 6a–c and Supplementary Fig. 21) show that the rising edge position for the Cu K-edge X-ray absorption near edge spectra (XANES) of Cu-doped KNN-BNH samples is closer to that of CuO reference (Fig. 6a), indicating the same valence state (i.e., +2). Cu-doped KNN-BNH samples also show a shoulder peak at 8986 eV (1 s to 4pz transition), suggesting a similar square planar coordination environment as CuO (Fig. 6a)46,47. Fig. 6b, c presents the k2-weighted Cu K-edge Fourier transform EXAFS (FT-EXAFS) spectra and their EXAFS fitting results in R space. For Cu-doped KNN-BNH samples, there is a sharp peak at R ~ 1.5 Å, corresponding to the Cu-O coordination shell. Lack of distinctly sharp peak at R > 2 Å indicates that Cu atoms in Cu-doped KNN-BNH samples have no long-range order in the crystal lattice, unlike the situation of Nb atoms (Supplementary Fig. 21b). Furthermore, the EXAFS fitting results suggest that the Cu atoms in Cu-doped KNN-BNH samples have a coordination number (CN) of 3.8 ~ 4.1, and the average Cu-O bond length is around 1.93 Å (Fig. 6c and Supplementary Table 4). These results confirm that the Cu atoms in Cu-doped KNN-BNH samples possess the CuO4 square planar coordination structure, which unambiguously proves the existence of the trimeric \({\left({V}_{O}^{\bullet \bullet }-{{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{\bullet \bullet }\right)}^{\bullet }\) defect configuration16.

Fig. 6: Formation mechanisms of defect dipoles.
Fig. 6: Formation mechanisms of defect dipoles.
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a Normalized Cu K-edge XANES spectra and b k2-weighted Cu K-edge FT-EXAFS spectra of KNN-BNH-Cu and standard samples (Cu foil, CuO, and Cu2O). c Experimental and fitting FT-EXAFS spectra for KNN-BNH-Cu samples and CuO standard sample. d Lattice structure and e charge density profiles of relaxed KNN-Cu supercells. f Comparisons of Nb-O bond lengths of relaxed KNN and KNN-Cu supercells. g Experimental and simulated Cu K-edge XANES spectra for the KNN-BNH-1Cu sample. h X-band EPR spectra of KNN-BNH-xCu samples. i Synchrotron based XRD patterns of KNN-BNH-1Cu from −100 to 25 °C.

Considering the proven Cu-O square planar structure, the first-principles calculations based on density functional theory (DFT) are used to calculate the local crystal structure of Cu-doped KNN (Fig. 6d–f and Supplementary Fig. 22)48,49. The relaxed KNN-Cu supercells show that the oxygen vacancies form at the two top positions of the oxygen octahedron along the c axis, resulting in the Cu-O square planar structure within the a-b plane. This structure drives the Nb5+ closer to O2- along the c (or -c) axis (as marked by the black arrows in Fig. 6d), enhancing Nb-O bond hybridization (Fig. 6e) and showing an elongation along the c axis (Supplementary Fig. 22). Compared to the pristine KNN supercells, doping Cu atom significantly affects the Nb-O bond length of nearest and next-nearest NbO6 octahedron (Fig. 6f and Supplementary Fig. 22), accounting for the observed local structural heterogeneities. Thus, the DFT calculations well explain the increased T phase content and the Raman υ1 vibration shifting to higher wavenumber after adding CuO. Furthermore, the simulated theoretical Cu K-edge XANES spectra of relaxed KNN-Cu supercells are compared with the experimental spectra of KNN-BNH-1Cu sample (Fig. 6g). The simulated spectrum matches well with the experimental spectra, further proving the reliability of the DFT calculations.

The X-band electron paramagnetic resonance (EPR) spectra also show a stronger signal at g = 2.003 for the KNN-BNH-1Cu sample, confirming more oxygen vacancies (Fig. 6h). Similar evidence is observed in the X-ray photoelectron spectra (XPS) (Supplementary Fig. 23). Temperature-dependent synchrotron XRD for the KNN-BNH-1Cu powder sample from −100 to 25 °C shows that as the temperature decreases, {110} diffraction peaks gradually transform from the T-phase feature to the O-phase one, and both {200} and {211} diffraction peaks also show the increased O phase content (as evidenced by the increased intensity of the left peak), as shown in Fig. 6i. Therefore, the low-temperature dielectric relaxation behavior in εʺ-T curves is due to the local structural heterogeneity rather than global phase transition, akin to the situations of relaxor ferroelectrics27. Specifically, the KNN-BNH matrix is dominated by the O phase at low temperatures, while the defect dipoles constrain the adjacent lattice cells with the T-phase feature.

Since the foregoing results and analysis unambiguously prove the existence of trimeric \({\left({V}_{O}^{\bullet \bullet }-{{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{\bullet \bullet }\right)}^{\bullet }\) defect complex, the equal dimeric \({({{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{\bullet \bullet })}^{{\prime} }\) defect should also be observed to keep the electroneutrality condition according to the Kröger-Vink notation50:

$$2C{uO}{\to }^{{Nb}_{2}{O}_{5}}2{{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }+2{O}_{O}^{\times }+3{V}_{O}^{{{\bullet }}{{\bullet }}}$$
(1)
$$2{{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }+3{V}_{O}^{{{\bullet }}{{\bullet }}}\to {\left({V}_{O}^{{{\bullet }}{{\bullet }}}-{{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{{{\bullet }}{{\bullet }}}\right)}^{{{\bullet }}}+{\left({{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{{{\bullet }}{{\bullet }}}\right)}^{{\prime} }$$
(2)

The dimeric ones establish PD and cause an elastic distortion, while the trimeric ones only cause an elastic distortion16. Both dimeric and trimeric defects can hinder the domain switching, especially the dimeric ones due to the generated PD. According to the principle of symmetry-conforming property of point defects, PD aligns with the spontaneous polarization (PS) and restrains adjacent domain switching, achieving overall hardening effect (called as bulk effect)51. Meanwhile, the propagation of new domain walls during polarization process will be impeded by PD, causing the hardening effect at domain walls (called as domain wall effect)52,53. Both effects contribute to the hardening effect of KNN-BNH-Cu ceramics. The elastic distortion caused by both defects (especially the trimeric ones) induces the local structure heterogeneities that disturb the long-range ordering and decrease domain sizes, consequently flattening the energy barrier of the retained O-T PBE. Finally, as we designed, both d33 and Qm can be enhanced via the local Cu-acceptor defect dipoles embedded O-T PBE strategy.

In this work, we propose the local Cu-acceptor defect dipoles embedded O-T PBE strategy to balance the development of d33 and Qm of KNN-based ceramics and validate it with KNN-BNH-xCuO ceramics. This strategy retains the room-temperature O-T PBE in KNN-BNH-xCu ceramics but introduces considerable Cu2+-related defects, namely, dimeric \({({{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{\bullet \bullet })}^{{\prime} }\) and trimeric \({\left({V}_{O}^{\bullet \bullet }-{{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{\bullet \bullet }\right)}^{\bullet }\) defects. The retained O-T PBE is featured with nanoscale multi-phase coexistence and abundant nano-domains due to the introduced local structural heterogeneities, ensuring high d33. The trimeric \({\left({V}_{O}^{\bullet \bullet }-{{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{\bullet \bullet }\right)}^{\bullet }\) defect is unambiguously proved by the XAFS spectra showing the existence of CuO4 square planar coordination structure. Such trimeric defect enhances the Nb-O bond energy and enhances the T phase characteristics, as proved by XRD patterns, Raman spectra, low-temperature synchrotron XRD patterns, and DFT calculations. The corresponding dimeric \({({{Cu}}_{{Nb}}^{{\prime} {\prime} {\prime} }-{V}_{O}^{\bullet \bullet })}^{{\prime} }\) defect forms PD that pins domain switching through the bulk and domain wall effects. Due to the coupling effect of the O-T PBE and the Cu2+-related defect configurations, the KNN-BNH-1Cu sample shows a combination of d33 = 340 pC/N and Qm = 256, surpassing the results of other typical KNN-based piezoceramics. Therefore, our strategy provides a paradigm for simultaneously optimizing d33 and Qm of lead-free piezoceramics that are promising for high-power applications.

Methods

Preparation of ceramics

0.95(K0.48Na0.52)NbO3-0.05(Bi0.5Na0.5)HfO3-0.2%molFe2O3-x%molCuO (KNN-BNH-xCu; x = 0, 0.5, 1, 1.5, 2, and 3), K0.48Na0.52NbO3-0.2%molFe2O3 (KNN), and K0.48Na0.52NbO3-0.2%molFe2O3-1%molCuO (KNN-1Cu) ceramics were synthesized by the conventional solid-state reaction method. High-purity raw materials of K2CO3 (99%), Na2CO3 (99.8%), Bi2O3 (99.999%), HfO2 (99%), Fe2O3 (99.9%), Nb2O5 (99.95%), and CuO (99%) were purchased from the Sinopharm Chemical Reagent Co., Ltd (China). The small amount of Fe2O3 was added to improve the sinterability of ceramics. First, with the purpose of keeping the phase boundary feature, the calcinated KNN-BNH powder was prepared in advance. The raw materials were weighed according to stoichiometric ratio and mixed well with alcohol as medium in zirconia vessel at a speed of 258 r/min on the planetary ball mill for 12 h. The slurries were then dried and calcinated at 850 °C for 6 h. Then, the CuO was added into the calcinated KNN-BNH powder for the secondary ball milling. After secondary ball milling, the powder was bonded by 5 wt.% polyvinyl alcohol (PVA) and then pressed into disks (diameter: 10 mm, thickness: 1 mm) under a pressure of ~300 MPa. The organic PVA was burned at 600 °C for 4 h, and the green disks were sintered at 1070–1090 °C for 3 h. The sintered pellets were spread with silver paste on both sides and then were heated at 600 °C for 30 min to form the electrodes. Finally, the samples were poled in the silicon oil at a direct current electric field of 3 kV/mm for 30 min.

Measurements for structure and performance

The density of ceramics was measured by the Archimedes method. For each composition, five ceramic disks were measured and the average density was calculated. The crystal structure of sintered ceramics was measured by an X-ray diffraction (XRD) equipment (Bruker D8 Advance XRD, Bruker AXS Inc., Madison, WI, Cu Kα). To conduct the Rietveld refinements, the as-sintered ceramics were manually ground into powder and then annealed at 600 °C for 2 h. The powder XRD (PXRD) patterns were collected by using a high-resolution XRD equipment (Empyrean, PANalytical, Netherlands), and the MAUD software was used for conducting the Rietveld refinements. Temperature-dependent synchrotron based XRD patterns were measured by the synchrotron X-ray beam (BL02U2 station of Shanghai Synchrotron Radiation Facility (SSRF)). Raman spectra were measured to study the local vibration modes of NbO6 octahedron (LabRAM HR Evolution micro-Raman spectrometer, HORIBA Jobin-Yvon, France). To observe the ferroelectric domains, the as-sintered ceramics were polished mechanically (DS001, Truer, Shanghai, China). Then, the domain structures of polished ceramics were characterized by a piezoresponse force microscopy (PFM, MFP-3D, Asylum Research, Goleta, CA, USA) with a conductive Pt-Ir-coated cantilever (PPP-NCHPt, Nanosensors, Switzerland) in the contact mode. The microtopography of samples is collected by scanning electron microscope (SEM, GeminiSEM, ZEISS300, Germany). Transmission electron microscopy (FEI Titan Themis G2 300, Thermofisher Scientific) with a probe corrector was used to observe the microstructures at 300 kV.

The real part of the dielectric constant (εʹ), the imaginary part of the dielectric constant (εʺ), and dielectric loss (tan δ) were tested by a broadband dielectric spectrometer (Novocontrol Concept 80, Novocontrol Technologies GmbH, Germany) under the temperature from −150 to 200 °C at frequencies ranging from 0.1 to 500 kHz. A ferroelectric analyzer (aixACCT TF Analyzer 2000, Germany) was used to measure the polarization-electric field hysteresis loops (P-E) and strain-electric field curves (S-E) at room temperature. Electron paramagnetic resonance spectra (EPR, Bruker, Emxplus, Germany) and X-ray photoelectron spectra (XPS, XPS-ultracold, Kratos, England) were measured to study the oxygen vacancies. The piezoelectric coefficients (d33) and mechanical quality factor (Qm) were detected by a Berlincourt meter (ZJ-3A, Institute of Acoustics, Chinese Academy of Sciences) and an Impedance analyzer (Model PV70A Co., Ltd., Bandera Electronics, China), respectively. The mechanical quality factor (Qm) was calculated by the resonance-antiresonance method according to the following formula54:

$${Q}_{m}\approx \frac{{f}_{a}^{2}}{2\pi {f}_{r}{Z}_{r}{C}^{T}\left({f}_{a}^{2}-{f}_{r}^{2}\right)}$$
(3)

To reveal the reliability of measured performance, we measured d33 and Qm values of samples from different batches for each composition, and calculated the mean value and standard deviation (see Supplementary Tables 5-9). The small standard deviation manifests the stable performance from batch to batch and reliable measurements.

The Cu K-edge and Hf L3-edge X-ray absorption fine structure (XAFS) spectra, including X-ray absorption near edge spectra (XANES) and extended X-ray absorption fine structure (EXAFS), were collected at XAFCA beamline of Singapore synchrotron light source (SSLS)55. The Nb K-edge XAFS spectra were collected at the XAS beamline of Australian Synchrotron. For energy calibration, the Cu foil and Nb foil were measured simultaneously with the samples. All the data were processed according to the standard procedures in Athena module within the Demeter software packages. The EXAFS fitting was conducted by Artemis software using FEFF6 theoretical scattering path calculation code56. The Cu K-edge XANES spectra for the relaxed structural model of the KNN-BNH-1Cu sample was also simulated by the FDMNES software57.

Atomic-resolution STEM image analysis

The atom positions were determined using a two-dimensional Gaussian peak-fitting method before calculating the atom displacement, the c/a ratio, and the lattice parameters. For the image along with the [100] zone axis, the displacement of the B-site atom in a perovskite structure was defined as the difference between its real position and the reference site based on the nearest four A-site atoms. For the image along with the [110] zone axis, the B-site atom displacement was obtained by calculating the difference between its real position and the averaged position of the two neighboring A-site atoms. The peak intensities of both A- and B-sites atoms were defined as the averaged intensity within an adaptive mask of each column, where the center of the mask is the position of the atom column and the radius is the full width at half maximum.

First principles calculations

The first-principle calculations are conducted based on density functional theory (DFT) as implemented in Vienna Ab initio Simulation Package (VASP)58,59. The Perdew-Burke-Ernzerhof revised for solids (PBEsol)60 exchange-correlation functional with projector-augmented wave (PAW)61 is employed. The Cu 3p64s13d10, K 3s23p64s1, Nb 4p65s24d3, and O 2s22p4 orbitals are treated as valence electrons for calculations. A plane wave expansion with a cut-off energy of 520 eV is used to represent the wave function. Γ-centered k-point meshes with a grid of spacing 0.04 × 2π Å−1 are selected for Brillouin zone sampling. The Kohn-Sham orbitals are iteratively updated in the self-consistency cycle until an energy convergence of 10−4 eV is obtained, and geometry optimizations proceed until the residual atomic forces are below 0.01 eV/Å. We select a 4 × 4 × 4 perovskite supercell containing 320 atoms, with [111]-oriented K/Na ordering and [110] polarization as the initial structure. A Nb atom is then replaced by a Cu atom, and two adjacent oxygen atoms are removed to maintain a Cu-centered parallelogram configuration. After full relaxation, a stable Cu-doped KNN supercell is obtained.