Introduction

Ferroelectrics, characterized by reversible spontaneous polarization, possess unique properties that make them well-suited for advanced wearable electronics, such as health-monitoring devices, electronic skin, and bio-integrated systems1,2,3,4,5,6. As wearable electronics evolve, the demand for materials that are elastic enough to conform to body movements while minimizing discomfort is growing7,8,9,10. However, traditional inorganic ferroelectric ceramics and common polymer-based ferroelectrics often face challenges due to their limited elastic recovery and high stiffness, which hinder their application in flexible devices11,12,13,14,15. This has sparked a surge of interest in developing ferroelectrics with improved elasticity. The primary challenge lies in the inherent trade-off between the disordered molecular structures needed for flexibility and the ordered crystalline domains essential for ferroelectric properties. Integrating both characteristics into a single material has been proven to be difficult. As a result, creating ferroelectric elastomers that can overcome this conflict is critical yet remains a significant challenge.

Elastification of ferroelectric materials is typically achieved through physical blending16,17, structural engineering18,19,20,21, or intrinsic elastification22,23,24,25. Among these, intrinsic elastification has emerged as the most promising approach, offering the potential for ferroelectric elastomers with both high strain endurance and excellent elasticity. Intrinsic elastification can be achieved by designing polymers with minimal crosslinking and covalent bonding in the amorphous regions, which effectively balances mechanical flexibility with ferroelectric performance. Recent breakthroughs, such as Gao et al.’s work on crosslinked poly(vinylidene fluoride-trifluoroethylene) (P(VDF-TrFE)), have led to the creation of intrinsically ferroelectric elastomers with strong ferroelectric responses and impressive elastic recovery26. However, the high TrFE content (45 mol%) and the introduction of crosslinking agents significantly reduce the Curie transition temperature and remanent polarization, limiting the material’s applicability, particularly in harsh environments. To overcome these limitations, polyvinylidene fluoride (PVDF) copolymers with lower TrFE content (20 mol%) have been employed, aiming to improve the Curie transition temperature to 85 °C27, yet this approach still faces challenges. Additionally, PVDF terpolymers based ferroelectric elastomers with even lower Curie temperatures have been explored, but a major breakthrough in improving the Curie transition temperature of ferroelectric elastomers remains elusive28,29,30.

Though copolymer components are favorable for easily forming the ferroelectric phase31, they often lower the Curie temperature and compromise thermal stability. Developing intrinsic ferroelectric elastomers that offer both thermal stability and performance in extreme environments is still a significant hurdle. Moreover, improving polarization switching dynamics and optimizing ferroelectric hysteresis loops are essential for advancing these materials toward practical devices. Achieving rapid polarization switching is particularly critical for enhancing the sensitivity of ferroelectric responses to external stimuli, which is crucial for wearable electronics and bio-integrated systems32,33,34,35,36.

In this context, PVDF homopolymer stands out for its ability to exhibit polar phases with a high Curie temperature, even exceeding its melting point. However, PVDF is known to exist in at least five distinct crystalline forms, with the nonpolar α phase being the most stable and easily formed under standard processing conditions, polar β and γ phases are more challenging to obtain and requiring specific conditions37,38. Traditional methods like stretching and blending, commonly used to promote polar phases, are not suitable for fabricating intrinsic ferroelectric elastomers39,40,41,42. Our research group has developed a approach using a smart memory effect to induce the formation of a polar phase. By controlling the heating and cooling process, we can manipulate the α crystal state, enabling a transition of it to the γ phase without the need for additives or external fields43,44,45. This breakthrough leads to the successful fabrication of an intrinsic ferroelectric elastomer that combines a high Curie temperature with elastic recovery and rapid polarization switching.

Herein, we fabricated an intrinsic ferroelectric elastomer with high Curie transition temperature and fast polarization switching based on PVDF homopolymer by incorporating a smart memory effect with the slight crosslinking of the amorphous regions. Polyethylene glycol (PEG) diamine is chosen as the soft crosslinker. The ferroelectric elastomer maintains stable ferroelectric responses in a broad temperature range, up to 150 °C—1.6 times higher than that of copolymer-based ferroelectric elastomers—significantly expanding the usage temperature window of ferroelectric devices. Under 30% strain, the ferroelectric elastomer exhibits an elastic recovery of 85%, making it well-suited for flexible and wearable electronics. Moreover, crosslinked PVDF films maintain a stable ferroelectric performance even under strains up to 200%. Notably, the stretched elastic crosslinked films with oriented structure exhibit an ultralow coercive field, which is only half that of PVDF, significantly reducing the operating voltage and energy consumption of ferroelectric devices. Moreover, stretching shortens the ferroelectric switching time by two orders of magnitude, allowing faster domain switching at lower electric fields. These properties make the material an ideal candidate for intelligent, high-performance wearable flexible electronics and memory devices.

Results and discussion

Structure of elastic PVDF via slight crosslinking

The crosslinking reaction between diamine and vinylidene fluoride (VDF) was selected to construct the elastic network. Initially, PVDF polymer was dissolved in N, N-dimethylformamide (DMF) along with varying ratios of Polyoxyethylene diamine (PEG diamine) crosslinker and the resulting solution was then casting as thick films. After removal of the solvent under vacuum, the crosslinked PVDF was obtained through thermal crosslinking. Figure 1a illustrates the crosslinking mechanism of elastic polymer. Upon heating, VDF units undergo dehydrofluorination via bimolecular elimination to generate double bonds. Subsequently, these double bonds engage in Michael addition with amino groups, forming C−N single bonds, followed by further dehydrofluorination to establish imine bonds (C=N), thereby completing the crosslinking reaction.

Fig. 1: The preparation of ferroelectric elastomers.
Fig. 1: The preparation of ferroelectric elastomers.The alternative text for this image may have been generated using AI.
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a The mechanism of preparing elastic ferroelectrics. b DSC curve of the PVDF/PEG diamine blend indicating in-situ crosslinking and the crystallinity of crosslinked films prepared at different crosslinking temperatures (inset of Fig. 1b). c FTIR spectra of pristine PVDF, blended and crosslinked PVDF/PEG diamine films. The inset in (c) is a magnified view of the cyan shaded area showing the characteristic C=N absorption peak. d XPS C1s spectra of crosslinked film and PEG diamine. Intensity is given in arbitrary units (arb. units), representing relative signal strength.

As shown by the differential scanning calorimetry (DSC) curve of the PVDF/PEG diamine blended film in Fig. 1b, the appearance of a broad exothermic peak confirms the occurrence of the crosslinking reaction, which starts at 190 °C and ends at 290 °C. Following the crosslinking process at different crosslinking temperatures, the films underwent a non-isothermal crystallization treatment, and their crystallinity was subsequently assessed through DSC analysis (Fig. S1). By comparing the crystallinity of films prepared at different crosslinking temperatures (inset of Fig. 1b), it is found that 240 °C is optimum crosslinking temperature for yielding the highest crystallinity of 42%, and is thus selected as the crosslinking condition to finalize the transition from plastic to elastic materials.

The elastic crosslinking networks were characterized using Fourier transform infrared (FTIR) and X-ray photoelectron spectroscopy (XPS) techniques. The presence of a broad peak corresponding to the stretching of the imine bonds at 1645 cm−1 in the crosslinked PVDF films (Fig. 1c) confirms the formation of the imine crosslinking bond. Additionally, XPS analysis (Figs. 1d and S2) also reveals the complete conversion of amino groups from the crosslinker (PEG diamine) into imine groups, evidenced by the characteristic shift from amino peak to imine peak. Moreover, the crosslinked membrane exhibits resistance to a serious of common organic solvents such as cyclohexanone, acetone, DMF, and N,N-dimethylacetamide (DMAc). After immersion in these solvents, minimal swelling was observed (Table S1), elucidating the crosslinked films displayed superior stability.

Further, the crystalline structure of the pristine and crosslinked PVDF which were non-isothermally crystallized from 240 °C (crosslinking temperature), was identified (Fig. S3). As for the pristine sample, the characteristic absorption bands of α crystals at 763 and 796 cm−1 are evidenced in FTIR spectra, suggesting the generation of mainly α crystals in the non-isothermal crystallization process. This is because that α phase is the most kinetically stable phase and thus most easily forms. In contrast, the crosslinked films display not only α phase reflections but also distinct characteristic peaks of the γ phase locating at 833 cm−1, with their intensities steadily increasing as the crosslinking ratio rises. At a crosslinking ratio of 10%, the γ phase content reaches approximately 68%. In addition, the DSC heating curves of the crosslinked films reveal the presence of an α phase melting peak alongside a higher-temperature γ phase melting peak, indicating that crosslinking effectively promotes the formation of the γ phase (Fig. S4). Moreover, the crystallization temperature and melting point decrease with the increasing crosslinking ratio. This behavior is attributed to the introduction of crosslinking points and the formation of intermolecular interactions, which restrict the mobility of PVDF chain segments26,46,47. Additionally, the presence of structural imperfections introduced by crosslinking reduces the overall crystallinity and results in a lower melting point26. Furthermore, to quantitatively support these observations, we performed non-isothermal crystallization kinetics analyses based on Jeziorny method48,49. As shown in Fig. S5, the crystallization rate of the 10% crosslinked film is lower than that of pristine PVDF at comparable cooling rates46,47.

During the above-mentioned conventional non-isothermal crystallization treatment, although the insertion of PEG favors some formation of γ form crystals in the crosslinked film, a certain amount of α crystals still persist. To further facilitate the formation of polar phase, the smart memory effect, induced by a partial melting of initial crystals, was employed to manipulate the polymorphs of PVDF during the recrystallization process. This approach firstly requires partial melting of the α crystals, and then at proper partially melting condition inducing the heterogeneous amorphous states with T3GT3G’ conformation, which consequently act as nuclei of γ-crystals during cooling process (the thermal protocol can be seen in Fig. S6). The intrinsic mechanism driving the α-γ phase transition is attributed to the matching between the thermal dynamics of molten α phase and the thermal dynamics of the conformational transition within amorphous chains43,44,45. By controlling the annealing temperature, the ability of the α-γ phase transformation is regulated by the partial melting state of the samples.

With the annealing temperature rises, as shown in Fig. S7, characteristic peaks of α phase at 796 and 763 cm−1 initially reduce and then increase. Crosslinked PVDF exhibits much sharper reduction of α phase at around 166 °C. Along with the weakening of peak intensity of α phase, the γ phase characteristic peaks at 833 cm−1 increase significantly. When the temperature rises above 170 °C, the α phase reappears43,44,45. The γ phase fraction (\({F}_{\gamma }\)) was quantitatively calculated using the Lambert-Beer law50. As shown in Fig. 2a and Table S2, \({F}_{\gamma }\) exhibits the maximum value for all samples when annealing temperature is set at the proper temperature. By contrast with pristine PVDF, the crosslinked PVDF exhibits a much higher γ phase content over a broader range of annealing temperatures. Nearly 100% γ phase is easily obtained for the crosslinked films in a wider temperature, whereas pristine PVDF films still retain a certain fraction of the α phase.

Fig. 2: Crystalline properties.
Fig. 2: Crystalline properties.The alternative text for this image may have been generated using AI.
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a The γ phase fraction (\({F}_{\gamma }\)) of pristine PVDF (0% crosslink) and 5% crosslinked films after thermal-treated at different temperatures and recrystallized during cooling to the room temperature. Data are presented as mean values and error bars represent the standard deviation (SD) of three independent measurements (n = 3, where n is the number of samples that are used to generate statistical data). (b) XRD, c FTIR, and d crystallinity of pristine PVDF and varying crosslinked films thermal-treated at 170 °C. e CF stretching and rocking region recorded in the FTIR spectra of pristine and 10% crosslinked PVDF. f Energy barrier for the α-γ phase transition of pristine and crosslinked PVDF. The two structural formulas shown in the figure illustrate the characteristic conformations of the α (TGTG′) and γ (TTTG) phases. g Schematics illustrating the intrinsic mechanism of crosslinked PVDF form polar phases.

Furthermore, the variation of the crystalline properties with crosslinking ratios was investigated at the optimum annealing temperature of 170 °C. The wide-angle X-ray diffraction profiles (Fig. 2b) indicate that with enhancing crosslinking ratio, several α/δ phase diffraction peaks evolve to two distinct diffraction peaks that can be attributed to γ/δ phases. Given the similarity between the diffraction peaks of (110)α, (110)γ, (110)δ, (020)δ, and (020)α, FTIR spectra were recorded to further confirm the crystalline structure. As shown in Fig. 2c, the crosslinked PVDF exhibits only characteristic IR peaks of γ phase located at 833 and 1234 cm−1, indicating that nearly pure γ phase is obtained after partial melting and recrystallization. Since the crystallinity of γ phase plays a critical role in ferroelectric and piezoelectric properties, DSC curves of PVDF samples were obtained. As shown in Fig. S8, the heating curve of crosslinked films exhibits a single melting peak of the γ phase, signifying the formation of nearly pure polar phase. By contrast, pristine PVDF reveals two melting peaks, indicating the mixture of α and γ crystals. The crystallinity of the films at varying crosslinking ratios was quantified by integrating the melting peaks. As depicted in Fig. 2d, crystallinity decreases with increasing crosslinking ratio owing to the insertion of PEG segment. Notably, although the overall crystallinity of crosslinked PVDF decreases, the fraction of γ phases in crosslinked samples is still higher than that of pristine PVDF as the crystallinity of pristine PVDF is derived from the contributions of both the α and γ crystals. Additionally, the atomic force microscope (AFM) observation was conducted to reveal the morphology evolution at the nanoscale. As shown in Fig. S9, the pristine PVDF and varying crosslinked PVDF films exhibited a flake-like morphology. Moreover, the size of the flake-like morphology decreased to a certain degree with increasing crosslinking ratio (Fig. S10a). Additionally, the crystallite size along the (110)γ plane, calculated using the Scherrer equation, decreases from 8.9 nm in the pristine film to 4.3 nm at a 10% crosslinking ratio (Fig. S10b). Here, we selected a maximum crosslinking ratio of 10%, as increasing the ratio to 15% results in phase separation (Fig. S11), reduced crystallinity, and a diminished capacity to induce the polar phase, ultimately compromising the mechanical, ferroelectric, and piezoelectric properties.

Besides, we also investigated the isothermal crystallization behavior of the crosslinked films. Compared to pristine PVDF, the crosslinked films maintain excellent ability to induce polar phase. As depicted in Figs. S12 and 13, the γ phase content increases progressively with increasing crosslinking ratio after thermal treatment at 155 °C for 24 h. When the crosslinking ratio reaches 10%, a pure γ phase is achieved, as confirmed by both FTIR and POM analyses. In contrast, the pristine PVDF film predominantly exhibits α phase, with only a minor presence of γ phase. Even at an isothermal crystallization temperature as low as 150 °C, the 10% crosslinked films predominantly generate a substantial amount of γ phase, with only negligible α-crystals (Figs. S14 and 15). Furthermore, we further investigated the isothermal crystallization kinetics of the pristine PVDF and crosslinked films. As shown in Fig. S16, the crosslinked films do not exhibit an induction period during isothermal crystallization at 155 °C. However, pristine films require an induction period of 86 min before crystallization begins. This indicates a marked increase in the crystallization rate of crosslinked PVDF compared to the pristine films. The Avrami equation were employed to further quantify the crystallization rate, revealing that the rate constant of the 10% crosslinked films is 10−3.99, which represents an enhancement of an order of magnitude compared to that of pristine PVDF.

The enhanced capability of inducing polar phase in crosslinked films stems from the formation of intermolecular forces. The absorption peak positions of CF bonds in the FTIR spectra confirm the generation of intermolecular forces. As shown in Fig. 2e, the CF vibration peak at 1178 cm−1 shift to lower wavenumbers, indicating reduced electron cloud density around CF bonds due to the presence of intermolecular forces. Furthermore, molecular dynamics simulations were conducted to calculate the energy barrier for the α-γ phase transition of pristine and crosslinked PVDF. As depicted in Fig. 2f, crosslinking markedly lowers the energy barrier for the α-γ phase transition. Consequently, the strong electronegative CF bonds in PVDF molecular chains interact with the ether linkages(COC) segments in PEG chains via dipole-dipole forces, thereby benefiting dipole orientation and facilitating the formation of polar phases (Fig. 2g).

Properties of the crosslinked PVDF

The ferroelectric properties of the crosslinked films after treated with the smart memory effect were characterized by polarization-electric field (P−E) tests. As depicted in Fig. 3, compared to pristine PVDF, the maximum polarization (\({P}_{\max }\)) and remanent polarization (\({P}_{r}\)) of the crosslinked films initially increase significantly with crosslinking ratio when it is small and then decrease with further increasing crosslinking ratio (Fig. S17 and Fig. 3b). Notably, for the optimum ferroelectric film with a crosslinking ratio of 5%, the values of \({P}_{r}\) and \({P}_{\max }\) are 7.3 μC/cm2 and 12.3 μC/cm2, respectively—more than twice those of the pristine PVDF film (3.0 μC/cm2 and 5.7 μC/cm2). Moreover, the \({P}_{r}\) value of the crosslinked films represents the highest value reported among ferroelectric elastomers to date. Significantly, in extensively studied copolymer-based crosslinked systems, the incorporation of crosslinking segments typically sacrifices dipole density, leading to reduced polarization intensity. In contrast, our slightly cross-linked homopolymer PVDF facilitates the formation of polar phases, thereby enhancing ferroelectric performance. Even at a crosslinking ratio of 10%, the \({P}_{r}\) and \({P}_{\max }\) values remain higher than those of pristine PVDF. Besides, the crosslinking lowers down the coercive field (\({E}_{c}\)) within a certain crosslinking ratio. The coercive field of the 5% crosslinked films reaches approximately 120 MV/m, closely aligning with the P(VDF-TrFE) based ferroelectric elastomer. The reduced coercive field can be attributed to the decreased crystallite size. Under an alternating electric field, the domain switching in smaller crystals involves weaker coupling forces, thus facilitating easier domain switching and lower coercive field51,52. The introduction of a crosslinked network similarly leads to a reduction in crystallite size. As shown in Fig. S10, the crystallite size decreases from 8.9 nm in the pristine PVDF film to 4.3 nm at a crosslinking ratio of 10%. This crystallite size reduction contributes to the observed decrease in coercive field. In summary, compared to the pristine PVDF film, slightly crosslinked films exhibit more defined rectangular P-E loops and enhanced polarization properties including higher \({P}_{r}\) and lower coercive field.

Fig. 3: Ferroelectric properties.
Fig. 3: Ferroelectric properties.The alternative text for this image may have been generated using AI.
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a P–E loops of pristine PVDF (0%) and crosslinked PVDF films (5% & 10%). b The polarization values of samples with different crosslinking ratios. c Variation of dielectric constant and crystallinity as a function of crosslinking ratio. The blue shaded region represents the crosslinking ratio corresponding to the optimal ferroelectric performance. d Temperature-dependent FTIR spectra of 10% crosslinked PVDF films recorded during heating from 30 °C to 200 °C represented by different colors. The ferroelectric phase disappears at approximately 160 °C. e Temperature-dependent P–E loops and f temperature-dependent dielectric constant of 10% crosslinked PVDF films. g Polarization value and γ phase infrared absorption intensities as a function of temperature for crosslinked PVDF and P(VDF-TrFE). Blue and yellow shaded areas correspond to the stable temperature regimes of the ferroelectric and paraelectric phases, respectively. The inset schematic illustrates the evolution of the crystalline structure during the heating process. h Comparison of thermal stability and \({P}_{r}\) values between previously reported elastic ferroelectric materials based on P(VDF-TrFE), Poly(vinylidene fluoride-trifluoroethylene-chlorofluoroethylene) (P(VDF-TrFE-CFE)), Poly(vinylidene fluoride-co-chlorotrifluoroethylene) (P(VDF-CTFE)) and PVDF achieved in this work.

Considering that the high electric field applied during ferroelectric testing may induce phase transitions, the crystalline structures of the films after polarization under different electric field strengths were further investigated. For the pristine PVDF film, the FTIR spectra (Fig. S18a) don’t show noticeable change in peak positions or intensities with increasing electric field, indicating the dominance of the γ phase with a minor presence of the α phase. To further investigate whether the α phase undergoes an electric-field-induced phase transformation, additional XRD and DSC analyses were conducted. As shown in Fig. S18b, the XRD pattern of the pristine PVDF film polarized at 300 MV/m reveals a marked decrease in the intensity of the (100)α peak and a near disappearance of the (120)α peak. Moreover, a new endothermic peak appears at approximately 153 °C in the DSC heating curve, which corresponds to the phase transition of the δ phase53,54,55. These results confirm that the α phase in the pristine PVDF undergoes electric-field-induced transformation to the polar δ phase. Therefore, the ferroelectric property of pristine PVDF is attributed to the γ and the δ phase56,57. In contrast, the crosslinked films don’t exhibit characteristic α phase peaks in either FTIR or XRD patterns (Fig. S19), confirming the absence of the α phase. Only γ phase absorption bands are observed, even under higher applied electric fields. However, the XRD diffraction patterns of the γ and δ phases are extremely similar. Further analysis of the FTIR spectra reveals a significant attenuation of the 763 cm⁻1 band in crosslinked samples, leaving only γ phase signatures. Given the lack of direct evidence linking the ferroelectric response to γ phase, we cannot exclude the possibility that the δ phase or other new phase contribute to the ferroelectric behavior in crosslinked samples. More in-depth and systematic studies will be done in future work.

The macroscopic polarization can be expressed in terms of the total dipole moment and dielectric constant, as described in Eq. (1):

$$P={d}_{{ij}}\sigma+{\varepsilon }_{0}\chi E$$
(1)

Here, \({d}_{{ij}}\) is the piezoelectric constant, \(\sigma\) is the Young’s modulus, \({\varepsilon }_{0}\) represents the dielectric constant in the absence of dipolar rearrangement, \(\chi\) is the electric susceptibility and \(E\) represents the external electric field. Notably, the first term, \({d}_{{ij}}\sigma\), corresponds to the intrinsic polarization, which is influenced by the dipole moments of molecular chains and the crystallinity of the polar phase. The second term, \({\varepsilon }_{0}\chi E\), represents the field-induced polarization component58,59. Based on the equation, the macroscopic polarization is proportional to the dipole moment, dielectric constant, and Young’s modulus.

Dielectric properties of crosslinked films with varying crosslinking densities are displayed in Fig. S20, which reveals that the dielectric constant gradually increases with increasing crosslinking ratio. Remarkably, at 100 Hz, the dielectric constant of the 10% crosslinked film (30.5) is almost twice that of the pristine PVDF (18.4). This enhancement can be attributed to dipolar polarization introduced by the crosslinker, which improves the dielectric properties and, consequently, enhances the macroscopic polarization of the crosslinked films60,61. Although the dielectric constant is enhanced with higher crosslinking ratios, the dielectric loss also increases accordingly, particularly in the low-frequency region. The polarization intensity of crosslinked films demonstrates the nonlinear dependence on crosslinking ratio, and it exhibits the optimal value at the 5% crosslinking ratio. It is most probably caused by the balance between the increase of dielectric constant and decrease of crystallinity with crosslinking ratio. With further increase in the crosslinking ratio over the optimum value, as shown in Fig. 3c, the insertion of crosslinked segments disrupts the long-range ordered structure of PVDF molecular chains, leading to a notable reduction in crystallinity and modulus, which ultimately hinders further improvements in ferroelectric performance.

Furthermore, the requirement for performance under extreme conditions underscores the critical importance of thermal stability in ferroelectric phases. In contrast to PVDF copolymers, the homo polar PVDF utilized in this study exhibits a Curie transition temperature exceeding its melting point, thereby substantially broadens the operational temperature window for integrated devices. Temperature-dependent FTIR was employed to assess the stability of the polar phases. Specifically, for P(VDF-TrFE) based ferroelectric elastomer, as the temperature increases from ambient conditions to 200 °C, the intensity of the polar phase characteristic peak located at 833 cm−1 dissipates entirely at 110 °C (Fig. S21), illustrating the failure of ferroelectric phase at high temperature. In contrast, for crosslinked PVDF-based ferroelectric elastomer, the polar phase persists up to 170 °C before fully melting (Fig. 3d). Furthermore, the stability of ferroelectric property was characterized by temperature-dependent hysteresis loop presented in Fig. 3e. As the temperature increases, the enhanced mobility of the dipoles contributes to a gradual increase in both the \({P}_{r}\) and \({P}_{\max }\) values. Moreover, the temperature-dependent dielectric performance further demonstrates excellent thermal stability from room temperature up to the melting point (160 °C) (Fig. 3f). The thermal stability schematics and the variations of γ crystals amount and the \({P}_{r}\) value with temperature were extracted and shown in Fig. 3g to accurately determine the stable temperature range of ferroelectric properties in crosslinked homo PVDF and copolymer P(VDF-TrFE) based ferroelectric elastomer. The ferroelectric properties of the crosslinked films synthesized in this study remain stable within room temperature to 150 °C. The failure temperature of ferroelectric elastomer developed in this work is much higher than that of P(VDF-TrFE) based ferroelectric elastomer (100 °C), representing the highest reported value for intrinsic elastic ferroelectric materials. Clearly, the intrinsic ferroelectric elastomers synthesized in this study exhibit thermal stability across a wider temperature range and achieve a high polarization value as can be seen from Fig. 3h62. This breakthrough addresses critical limitations of conventional ferroelectric elastomers, offering a robust solution for high-performance applications in challenging thermal and mechanical conditions, and paving the way for advanced technologies in wearable electronics and beyond.

Considering that the complex strains are typically encountered during material processing and practical applications, the stability of ferroelectric performance under mechanical stress is particularly critical. Consequently, the ferroelectric responses of elastic ferroelectrics under varying strains were investigated. The 10% crosslinked PVDF films were subjected to tensile strains of 50%, 100%, 150%, and 200% using a universal testing machine and were subsequently fixed for further analysis.

Initially, the crystallization properties of the films before and after stretching were assessed. As shown in Fig. S22a, b, in-situ FTIR and XRD measurements were employed on films stretched up to 200% to confirm the crystalline phase during uniaxial stretching, and the calculated relative contents of different crystalline phases are shown in Fig. S22c. As depicted in Fig. S22a, the initial film exhibits γ-crystals. When the strain reaches 100%, a new absorption band at 1279 cm−1 corresponding to β phase appears. With further increase in the stretching ratio, the β phase content continues to grow. Additionally, the emergence of the (200)β diffraction peak at 2θ = 20.7° in the XRD pattern further confirms the formation of the β phase during the stretching process. These results indicate that stretching induces a γ-β conformational transition. At a stretching ratio of 200%, the Fβ value reached 46.7%. Moreover, the ferroelectric properties of films prepared at highest stretching ratios were characterized in Fig. 4a–c. For the 200% stretched crosslinked films, both the \({P}_{r}\) and \({P}_{\max }\) exhibit an enhancement than untreated crosslinked film, indicating that the ferroelectric responses of these elastic ferroelectrics remain robust under applied strain. Furthermore, the coercive field decreases from 140 to 62 MV/m, enabling the ferroelectric switching of stretched film at much lower electric fields. The formation of higher polarity β phase and strain induced orientation of crystals favors the alignment of ferroelectric domains and consequently enhances the remanent polarization and lowers down the coercive field.

Fig. 4: Ferroelectric properties of 10% crosslinked films before and after stretching.
Fig. 4: Ferroelectric properties of 10% crosslinked films before and after stretching.The alternative text for this image may have been generated using AI.
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a P-E loops. b Polarization trend as a function of electric field. Pink and blue shaded regions correspond to the polarization evolution of the crosslinked films in the unstretched state (0%) and at 200% stretching ratio, respectively. c A comparison of \({P}_{r}\), \({P}_{\max }\), and \({E}_{c}\) values for 10% crosslinked films before and after stretching (under 200% strains). d Schematic diagram of polarization switched under the alternating electric field and the applied electric field program used to track polarization switching. e Polarization switching dynamics at 100 MV/m (blue shaded area indicates the ferroelectric domain switching process), and f the characteristic switching time \({\tau }_{s}\) for crosslinked PVDF films before and after stretching (under 200% strains).

The dynamics of polarization switching are crucial factor that influence the performance and practical applications of ferroelectric materials. Therefore, the switching rates of ferroelectric domains in 10% crosslinked films before and after stretching were characterized. The domain switching process was monitored in situ by applying a series of periodic electric field pulses. As shown in Fig. 4d, firstly, a positive saturation electric field was applied to fully polarize the ferroelectric domains (P1). Next, another positive electric field of the same magnitude was employed to measure the total polarization response (P2). Subsequently, negative electric fields with varying pulse widths were used to achieve domain writing (P3). Finally, a positive electric field of the same magnitude was applied again to measure the corresponding polarization value (P4). By tracking the difference between P2 and P4 (ΔP) as a function of pulse time, the domain switching dynamics can be in situ monitored. As illustrated in Fig. 4e, under a lower electric field of 100 MV/m, the stretched film completes polarization reversal within a narrow time range of 2 ms, achieving a higher polarization value. In contrast, the untreated film fails to reach polarization saturation within 100 ms. Subsequently, the KAI model, which assumes homogeneous nucleation and unrestricted domain growth, was used to describe polarization switching behavior, stating that

$$\Delta P=2{P}_{s}\left[1-\exp \left(\right.-{\left(\frac{t}{{\tau }_{s}}\right)}^{n}\right]$$
(2)

can be deformed to

$$\log \left[-{{\mathrm{ln}}}\left(1-\frac{\Delta P}{2{P}_{{s}}}\right)\right]=n\log t-n\log {\tau }_{s}$$
(3)

where \({\tau }_{s}\) and n are the characteristic switching time and the geometric dimension for the domain growth, respectively63,64,65,66. The KAI fitting results and the corresponding parameters \({\tau }_{s}\), are shown in Figs. S23 and 4f. After 200% stretching, the geometric dimension (n) increases to approximately 1, suggesting that domain nucleation and growth tend to be one-dimensional. The characteristic switching time of the stretched crosslinked film decreases to 10−2.64 s, compared to 10−0.44 s for the pristine crosslinked film, indicating the crystal alignment induced by stretching facilitates a faster switching rate with the low voltage operation, highlighting their applicability in fields like flexible electronics and memory devices. What’s more, to determine whether the observed fast ferroelectric switching is caused by the reduced coercive field or the intrinsic structure change, we introduced a normalized parameter, ξ, defined as the ratio between the applied electric field and the coercive field of each sample (\(E/{E}_{c}\)). The switching times of the films before and after 200% stretching were compared across a range of ξ values. The polarization switching dynamics under varying electric fields were monitored by in-situ tracking of the polarization as a function of pulse time (Fig. S24a, b). The corresponding fitting curves based on the KAI model are presented in Fig. S24c, d, and the switching time constants (τs) were calculated using Eqs. (2) and (3). As shown in Fig. S24e, under relatively low applied electric fields (ξ < 1.5), the stretched samples exhibit significantly shorter switching times compared to the unstretched samples, indicating that molecular orientation plays a dominant role in enhancing ferroelectric switching. However, when the applied electric field becomes sufficiently high (ξ > 1.5), the switching times for both types of samples converge, suggesting that the electric field strength becomes the controlling factor at high electric field regimes.

Furthermore, the effect of crosslinking density on stress-strain behavior and elastic recovery of the pristine and crosslinked PVDF films is also explored. As the crosslinking density increased, Young’s modulus and yield strength gradual decreases. At a crosslinking ratio of 10%, the modulus decreases to 696 MPa, which is attributed to the incorporation of PEG soft chains as can be seen in Figs. 5a, b and S25. The elasticity improvement is another critical factor for practical applications. To assess this, cyclic stress-strain measurements were performed within a strain range of 10–40%. As shown in Fig. 5c, d, the crosslinked films exhibit much higher elastic recovery rate than pristine PVDF. Notably, 10% crosslinked films achieve elastic recovery ratios exceeding 85% under strains ranging from 10% to 30%. Moreover, after 2000 cycles of stretching and releasing, the elastic recovery ratio remains stable without significant degradation (see Fig. S26). In contrast, pristine PVDF shows limited elasticity, achieving only 54% elastic recovery even at a very lower strain of 15%.

Fig. 5: Mechanical properties.
Fig. 5: Mechanical properties.The alternative text for this image may have been generated using AI.
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a Stress-strain curves of pristine and various crosslinked PVDF films. (Inset) Enlarged view of the area highlighted in gray. b Modulus of pristine and crosslinked films at various crosslinking ratios. Data are presented as mean values and error bars represent the SD of three independent measurements (n = 3, where n is the number of samples that are used to generate statistical data). (c) Cyclic stress-strain curves of pristine homo-PVDF and 10% crosslinked PVDF films under different strains represented by different colors. d Elastic recovery ratios of pristine and various crosslinked PVDF films.

We have developed a ferroelectric elastomer that exhibits thermal stability and rapid polarization switching capabilities, overcoming critical limitations of existing materials. This innovation is achieved by thermally crosslinking a high Curie temperature PVDF homopolymer with PEG diamine. The ferroelectric phase being effectively induced via a smart melt-memory effect. A key breakthrough of this material is its ability to retain ferroelectric properties at temperatures up to 150 °C, expanding its potential for high-temperature and extreme environmental conditions. Additionally, it demonstrates mechanical flexibility, achieving over 85% elastic recovery under 30% strain—ensuring adaptability of wearable devices. Furthermore, we harness stretch-induced dipole orientation to achieved a dramatic enhancement in the domain switching speed, reducing switching time by two orders of magnitude. This advancement enables a fast, low-energy, and thermally stable polarization response, opening avenues for the fabrication of highly sensitive and rapid-response ferroelectric elastomers.

Methods

Materials

PVDF (Mw = 5.3 × 105 g mol−1) was purchased from Sigma–Aldrich (Missouri, USA). DMF, DMAc, Acetone, Cyclohexane, and PEG diamine 1000 Da were obtained from Beijing Chemical Factory (Beijing, China). All chemicals used were of analytical grade and had a purity of 99%.

Preparation of crosslinked PVDF thick films

PVDF and PEG diamine (content varying according to different crosslinking densities) were dissolved in DMF to form a homogeneous solution. A thick film was prepared by casting on a glass substrate and placed in a vacuum oven for drying at 70 °C for 24 h. After volatilizing most of the solvent, the film was crosslinked at 240 °C in a vacuum oven for 1 h and then transfer directly to room temperature. The thickness of the film was approximately 30 μm. Thermal annealing treatment (melt memory effect, shown in Fig. S6) was employed to induce polar phase. The procedure is to first heat up to different temperatures (164, 166, 168, 170, 172, and 174 °C) and annealing for 10 min, and then reduce to room temperature. The rising and cooling rates are 5 °C/min.

Characterization

The chemical structure of the crosslinked samples was characterized through FTIR and XPS spectroscopy. FTIR were performed on a Spectrum 100 FTIR spectrometer (Perkin-Elmer) with a resolution of 4 cm−1, while XPS data were acquired using a monochromatic Al-Kα source on a Kratos AXIS SUPRA X-ray photoelectron spectrometer.

The detailed morphologies of pristine and crosslinked PVDF films were characterized by atomic force microscopy (AFM) and Optical Microscopy under crossed polarizers (POM). XRD experiments were performed on 1W1A Diffuse X-ray Scattering Station, Beijing Synchrotron Radiation Facility (BSRF-1W1A). FTIR is also used to further confirm the crystal structures. For samples containing a mixture of α and γ phases, the relative fraction of the PVDF γ phase (\({F}_{\gamma }\)) in different samples was calculated according to Eq. (4)50

$${F}_{\gamma }=\frac{{A}_{\gamma }}{\left(\frac{{K}_{\gamma }}{{K}_{\alpha }}\right){A}_{\alpha }+{A}_{\gamma }}$$
(4)

where Aα and Aγ are the absorbances of the bands at 763 cm−1 and 833 cm−1, Kα and Kγ are the absorption coefficients of related bands with values of 0.365 μm−1 and 0.150 μm−1, respectively46. For samples containing a mixture of α, β, and γ phases, the relative contents of β and γ phases were determined using the following procedure. First, the combined fraction \({F}_{\left(\beta+\gamma \right)}\) was calculated using Eq. (5).

$${F}_{\left(\beta+\gamma \right)}=\frac{{A}_{(\beta+\gamma )}}{\left(\frac{{K}_{\left(\beta+\gamma \right)}}{{K}_{\alpha }}\right){A}_{\alpha }+{A}_{(\beta+\gamma )}}$$
(5)

where the \({A}_{\alpha }\) and \({A}_{(\beta+\gamma )}\) are the absorbances of the band at 763 cm−1 and 840 cm−1, Kα and \({K}_{\left(\beta+\gamma \right)}\) are the absorption coefficients of related bands with values of 6.1 × 104 cm2 mol−1 and 7.7 × 104 cm2 mol−1, respectively. Then, the individual fractions of β and γ were determined separately using Eqs. (6) and (7), respectively.

$${F}_{\beta }={F}_{\left(\beta+\gamma \right)}\times \frac{{A}_{1275}}{{A}_{1275}+{A}_{1234}}$$
(6)
$${F}_{\gamma }={F}_{\left(\beta+\gamma \right)}\times \frac{{A}_{1234}}{{A}_{1275}+{A}_{1234}}$$
(7)

The melting behavior of PVDF and crosslinked films was characterized through a TA Instruments Q2000 differential scanning calorimeter (DSC) under nitrogen environment with 10 °C/min of thermal ramp. The theoretical computations were done using Gaussian16 program to calculate the energy barrier for the α-γ phase transition of pristine and crosslinked PVDF. Geometrical optimization and frequency calculation are evaluated by B3LYP(D3BJ)/def2-SVP, while using M06-2x(D3)/def2-TZVP when calculating energy. The energy changes are corrected with zero-point energy.

The mechanical property measurement was conducted on a universal tensile test machine (Instron 5943), and the cyclic stress-strain curve was obtained by repeatedly stretching the sample with the same instrument. In the ferroelectric testing, P-E loop test was conducted using the Premiere II ferroelectric tester from Radiant Technologies, Inc. (USA). In this test, an alternating electric field with a sine waveform was applied at a frequency of 10 Hz. For dielectric spectroscopy testing, the samples were equipped with Au electrodes (diameter 2 cm) grown by magnetron sputtering. A wide-frequency impedance spectrometer (DMS2000, Balab, China) was employed to measure the dielectric properties in the temperature range of 30 to 200 °C, at a rate of 2 °C/min, over the frequency range of 10−1–107 Hz.