Low-entropy amorphous dielectric polymers for high-temperature capacitive energy storage

Abstract

Electrostatic capacitors based on polymer dielectrics are essential components in advanced electronic and electrical power systems. An urgent challenge, however, is how to improve their capacitive performance at high temperatures to meet the rising demand for electricity in a harsh-environment present in the emergent applications such as electric vehicles, renewable energy, and aerospace systems. Here, we report a low-entropy amorphous polymer with locally extended chain conformation comprising high-T_g poly(2,6-dimethyl-1,4-phenylene oxide) (PPO) blended with polystyrene (PS) that exhibits an energy density as high as 5.5 J cm⁻³ with an efficiency of >90% at an electric field of 600 MV m⁻¹ at 150 °C, outperforming the existing dielectric polymers. Our results reveal that regulating the conformation entropy of polymer chains introduces a favorable locally extended polymer chain conformation, resulting in dense chain packing with short-range ordered but long-range disordered microstructures, and inhibits the transport of electrons in dielectric polymers, consequently, leading to the substantial improvements of capacitive performance at elevated temperatures. This low-entropy approach is scalable, general, ultra-low-cost and simple, paving the way for mass fabrication of high-performance and high-quality polymer films required for high-temperature film capacitors.

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Broader context

Electrostatic energy storage capacitors are essential components in modern electrical and electronic systems. Compared to ceramic dielectrics, polymer dielectrics offer higher breakdown strengths, greater reliability, scalability, and lightweight properties, making them ideal for high-voltage applications such as power electronics, power conditioning, and pulsed power. However, current dielectric polymers are limited by relatively low working temperatures, which fail to meet the rising demands of electrical energy storage under the extreme conditions found in electric vehicles, aerospace systems, and power grids. Although high-T_g polymers filled with inorganic nanofillers have shown promising elevated-temperature capacitive performance, the scalability of these nanocomposites remains a significant challenge. They are typically prepared by tedious, complex, and time-consuming methods that are incompatible with the current melt-based processes used for capacitor films. Herein we present a highly scalable dielectric polymer with low-entropy molecular conformations that exhibits superior capacitive performance at elevated temperatures and high applied electric fields. This low-entropy approach is scalable, general, ultra-low-cost, and simple, facilitating the mass production of high-performance and high-quality polymer films required for high-temperature electrostatic energy storage capacitors.

Introduction

Electrostatic polymer film capacitors are critically important elements in modern electronic devices and power systems for energy storage and regulation because of their distinctive features of ultrahigh power densities, high-voltage endurance, good reliability and facile processability.^1–5 However, the current polymer dielectrics are limited to relatively low working temperatures, and thus lead to the energy storage capacitors failing to meet the demand for stable operation under the harsh-temperature conditions present in the emergent applications such as electric vehicles, aerospace power electronics and underground oil/gas exploration.^6–11 For example, while the near-engine-temperature in electric vehicles can reach above 120 °C,¹² the operating temperature of biaxially oriented polypropylene (BOPP) film capacitors, which are currently used in power inverters of electric vehicles, is below 105 °C.^13,14 Thus, cooling systems have to be employed to decrease the working temperature from above 120 °C to about 70 °C, which brings extra weight, volume and energy consumption.

To address these urgent needs, a variety of high thermal stability engineering polymers^15,16 with high glass transition temperature (T_g) such as polyimides (PIs), polycarbonates (PCs) and fluorene polyesters (FPEs), have been exploited as high-temperature dielectric polymers. Unfortunately, although the high-T_g polymers show stable polarization over a broad temperature, all the polymers show poor charge–discharge efficiencies at the applied high electric and thermal fields, which resulted from sharply increased conduction loss attributable to the thermally and electrically assisted charge injection, excitation and transport.^4,7 Doping appropriate amounts of wide-bandgap inorganic fillers such as boron nitride nanosheets (BNNS)⁷ and aluminium oxide (Al₂O₃)^17,18 in high-T_g polymers to form polymer nanocomposites have proved effective in promoting the high-temperature capacitive performance. It is found that the introduced nanofillers are effective in hindering the charge transport to decrease the conduction loss under elevated temperatures and high applied electric fields. It is further observed that the nanofillers can induce the surrounding polymers to form much tighter and more orderly aggregate structures, which restrains the movement of molecular chains and reduces the entropy of local molecular conformations, resulting in the formation of deep-level localized states at the interfacial zones between nanofillers and the polymer matrix and hindering the charge transport in polymer nanocomposites.^10,19 These results reveal that the entropy of molecular conformations is a key factor influencing charge transport, which affects the conductivity and energy loss of polymer dielectrics. However, the nanocomposite approaches developed to tune the entropy of polymer molecular conformations at the interface regions introduce complications. High loading of nanofillers in the composites significantly compromises film flexibility, leading to voids and cracks in films. This makes the manufacture of large-scale, low-cost, and uniform polymer films challenging.^1,2

Here we depart from the earlier approaches and describe a scalable amorphous dielectric polymer with a low-entropy of molecular conformations, which was prepared by using a general polymer blending strategy. The low-entropy amorphous polymer comprising high-T_g poly(2,6-dimethyl-1,4-phenylene oxide) (PPO) blended with polystyrene (PS) exhibits superior capacitive performance at elevated temperature and high applied electric field. The rationale of our material design is that by using strong van der Waals interactions between the phenyl ring of the PS and the phenylene ring of PPO to “stretch” the main chain at the molecular level, we achieve a significant reduction of molecular conformation entropy in an amorphous polymer with randomly oriented yet superiorly packed extended polymer chains, which can inhibit charge transport via deep-level localized states. This is different from the reported high-temperature polymer nanocomposites which only tune conformation and arrangement of the polymer chain in the particle/matrix interface. Consequently, such a low-entropy amorphous dielectric polymer exhibits a discharged energy density of 5.5 J cm⁻³ with a charge–discharge efficiency of >90% at 150 °C, outperforming the existing dielectric polymers and showing a more than two-fold increase in discharged energy density compared to polyetherimide (PEI), one of the best high-temperature dielectric polymers.⁶ Moreover, compared to reported approaches, the film processing, i.e., polymer blending, is very simple, straightforward and low cost; thus the approach presented paves the way for mass fabrication of high-performance and high-quality polymer films required for high-temperature film capacitors.

Results and discussion

Nano-structures and chain-packing morphology of low-entropy films

To test our idea, we fabricated a series of dielectric films with different ratios of PPO/PS (wt%/wt%) and systematically studied the nanostructure and chain-packing morphology of the blends. The PPO/PS blend miscibility was first studied by differential scanning calorimetry (DSC), atomic force microscopy (AFM) and scanning electron microscopy (SEM). As shown in Fig. S1 (ESI†), the PPO/PS blends at all compositions show a single glass transition temperature (T_g), and the T_g values for the blends are in good agreement with the theoretical values predicted via the Flory–Fox equation²⁰ (Fig. 1A), indicating that the blends are miscible at the molecular level.²¹ Tapping mode AFM data (Fig. 1B, C and Fig. S2, ESI†) and the microstructure picture of SEM (Fig. S3, ESI†) show no sign of phase separation of the two polymers in the blends.

Fig. 1 Structures and structural characterization of low-entropy films. (A) Glass transition temperature (T_g) of PPO/PS blends versus PS content. The experimental T_g of the samples was deduced from DSC curves as shown in Fig. S1 (ESI†). The dashed line in (A) represents the fitting of T_g using the Flory–Fox equation , where T_g1 and T_g2 are the T_gs of polymer 1 and polymer 2, respectively, and w₁ and w₂ are their corresponding weight percentages in the blends. Taping-mode AFM data for the PPO/PS 75%/25% blend, phase (B) and topography (C) images with a scanning area of 10 × 10 μm. (D) FT-IR spectra in the region 1220–1140 cm⁻¹ of PPO/PS blends with different blending ratios. (E) Excess interchain spacing of PPO/PS blends calculated from the peak positions of the X-ray profiles in Fig. S5 (ESI†) versus PS content. (F) Density of PPO/PS blends versus PS content. (G) Change in specific heat capacity during glass transition (ΔC_p) for PPO/PS blends with different ratios. (H) elastic moduli of neat PPO, PS and the 75%/25% PPO/PS blend.

The conformations of the polymer chains and the packing behavior in the blends were evaluated experimentally. FT-IR spectroscopy was employed to characterize chain conformation in PPO/PS blends, and the spectral data are presented in Fig. 1D and Fig. S4 (ESI†). We noticed that the band widths associated with these vibrational transitions are narrower in the polymer blends than in the homopolymers (Fig. 1D). It suggested that the conformational freedom or at least the distribution of conformations available to PPO in blends is limited compared to the neat PPO, and the degree of conformational freedom available to neat PPO is high.^22,23 The conformation of the PPO chain is relevant to the conjugation of benzene π electrons that is, in any aromatic ether the conjugation of the benzene π electrons with the lone-pair electrons on the oxygen depends upon the angle the plane of the benzene ring makes with the plane of the C–O–C bond.²² The greatest conjugation occurs at 0° while the lowest occurs at 90°, which means that the locally extended polymer chain conformation is beneficial for the conjugation of benzene π electrons and makes the system more stable. X-ray diffraction (XRD) data revealed that the 75%/25% PPO/PS blend exhibits a broad amorphous scattering peak at the 2θ angle of 13.1° (Fig. S5, ESI†). The corresponding interchain spacing is smaller than the value predicted by the linear relationship to the PPO and PS blending ratio. A 75%/25% PPO/PS blend exhibited the largest excess distance (i.e., the deviation of the actual interchain spacing from the theoretical interchain spacing based on linear relationship) among all the blends (Fig. 1E).

The density of the blend films was also measured, and the 75%/25% PPO/PS blend exhibited the highest density, which was 1.08 g cm⁻³, higher than 0.96 g cm⁻³ for neat PPO films and 0.91 g cm⁻³ for the neat PS films used in this study (Fig. 1F). Fig. 1G revealed that the change in specific heat capacity (ΔC_p) during the glass transition was the lowest as the content of PS approaches 0.25. In general, a change of polymer chain conformation from a highly self-entangled state to an extended backbone state is the primary contributor to the ΔC_p during a polymer's glass transition. Therefore, the minimum ΔC_p of the blend with PPO/PS 75%/25% ratio (Fig. 1G) implied that PPO backbones are more extended for this blend compared with pristine polymers and other blends before the glass transition.²⁴ Our results showed that the blend film of 75%/25% PPO/PS ratio has the most extended polymer chain conformation among all the blend compositions of PPO and PS studied, which leads to the lowest ΔC_p and the largest excess interchain distance as observed in XRD. Clearly, the deformed extended chain-packing morphology improves the packing density. Increased elastic modulus and thermal conductivity of the blend, compared with PPO and PS, were also observed (Fig. 1H and Fig. S7, ESI†), and were consistent with the more extended chain-packing morphology and the increase in density. These results illustrate that the mechanical and thermal performance of dielectric polymers can also be tuned and improved by the low-entropy polymer chains with locally extended conformation.

Dielectric properties and charge transport

To characterize the properties of the low-entropy amorphous polymer blends, we first measured the dielectric properties of the films over a frequency range of 100 Hz to 1 MHz and temperature range of 30 to 200 °C. As PS content increases, the dielectric permittivity of the PPO/PS films increases firstly and then decreases (Fig. 2A and B), in agreement with the trend of density of the films. Moreover, the largest dielectric permittivity occurs at the composition of PPO/PS 75%/25% (Fig. 2B), due to the most densely packed low-entropy polymer chains with locally extended conformation that results in the largest number of dipoles per unit volume of the film in this composition. We observed a high stability of permittivity over the frequency range measured, that is, with the variation of 1.1% for low-entropy films (Fig. 2A and C). The permittivity of PPO/PS 75%/25% also exhibits a much higher stability over the temperature range from room temperature to 200 °C, e.g., the variation of 2.4% compared with 10.1% of the neat PPO films (Fig. 2D). In addition, we found that the loss tangent of the low-entropy films is suppressed over the measured temperature and frequency ranges (Fig. 2A, C and D), for example, ∼0.15% at room temperature and 1 MHz, compared to that (∼0.24%) of the neat PPO films. The low dielectric loss is favorable for developing high-efficiency energy storage dielectrics. We ascribe the evolution of the dielectric performances to the densely packed low-entropy polymer chains with locally extended conformation, which results in higher permittivity and low loss and good frequency/temperature stability.

Fig. 2 Dielectric measurements of PPO/PS blend films. (A) Dielectric properties of PPO, PS and their blend with ratio of 75%/25% as a function of frequency at room temperature. (B) The dielectric constant of PPO/PS blends at 1 kHz and room temperature vs. PS content. (C) Dielectric properties of PPO, PS, and the 75%/25% blend as a function of frequency at 150 °C. (D) Temperature-dependent dielectric constant and dielectric loss at 1 kHz of PPO, PS, and the 75%/25% blend.

We then measured the leakage current and breakdown properties of the low-entropy films and compared them with the neat PPO films at elevated temperatures. The leakage current density of the low-entropy film of PPO/PS 75%/25% is suppressed by more than two orders of magnitude compared with that of the neat PPO film, for example, 3.09 × 10⁻⁸ A cm⁻² for the PPO/PS 75%/25% compared with and 1.78 × 10⁻⁶ A cm⁻² of the neat PPO at an electric field of 300 MV m⁻¹ and the temperature of 150 °C (Fig. 3A). Fig. 3B shows the temperature-dependent electrical conductivity, indicating that PPO/PS (75%/25%) exhibits the highest activation energy (0.99 eV) compared to pristine PPO (0.92 eV) and PS (0.84 eV). This suggests the greatest charge carrier trap depth in the low-entropy PPO/PS (75%/25%) blend. We derived the characteristic breakdown strength E_b of the low-entropy film from the Weibull distribution analysis which is presented in Fig. 3C. An E_b of 662 MV m⁻¹ was achieved in the low-entropy films of PPO/PS 75%/25% at a temperature of 150 °C, which is substantially enhanced by 65.5% compared to the neat PPO of E_b of 400 MV m⁻¹ and is higher than those of most of currently reported high-temperature polymers, for example, 377 MV m⁻¹ of PEI,¹¹ 400 MV m⁻¹ of PEEU¹⁷ and 641 MV m⁻¹ of CS-ODA²⁵ (Fig. 3D). The large enhancement of breakdown strength of the low-entropy film compared to the neat PPO is consistent with the increased Young's modulus originating from the dense packing of low-entropy polymer chains and significantly decreased leakage current, as described by the widely accepted dielectric breakdown mechanisms.²⁶

Fig. 3 Electric conduction behaviors and dielectric breakdown strength. (A) The current density of PPO and PPO/PS 75%/25% with increasing electric field at 150 °C. The dotted curves fit to hopping conduction model. (B) Temperature-dependent electrical conductivity at 200 MV m⁻¹. (C) Weibull statistic of the breakdown strength of PPO and PPO/PS 75%/25% at 150 °C. (D) Comparison of the breakdown strength between PPO/PS 75%/25% and the current high-temperature neat dielectric polymers at 150 °C. (E) A schematic illustration of the low-entropy amorphous polymer shows locally extended polymer chain conformations, resulting in dense chain packing with short-range ordered but long-range disordered microstructures. This configuration shortens the electron hopping distance, thereby inhibiting electron transport in dielectric polymers.

To understand the mechanisms of the greatly decreased leakage current density and improved breakdown strength in the low-entropy PPO/PS 75%/25% film, we studied the charge transport behavior through analysis of the current density using different conduction models.²⁷ As shown in the fitting curves in Fig. 3A, the J–E data of the neat PPO and the low-entropy PPO/PS 75%/25% were fitted well by the hopping conduction model with a coefficient of determination (R²) both larger than 0.99, implying that hopping conduction is the predominant conduction mechanism at high temperatures and high electric fields for the tested samples. The derived hopping distance (λ) decreases from 0.76 nm in the neat PPO to 0.52 nm in the 75%/25% PPO/PS blend. The energy gained by mobile charges between two traps is directly proportional to the hopping distance λ. The significantly reduced hopping distance λ reduces the energy gained by mobile charges, thereby improving the breakdown strength and lowering conduction losses.^28,29 These experimental results indicate that the low-entropy amorphous microstructure of the polymer blend, which is short-range ordered but long-range disordered, effectively scatters charge carriers and inhibits the transport of mobile electrons (Fig. 3E). Consequently, this reduces the leakage current density and enhances the dielectric breakdown strength.

Energy storage properties of low-entropy films

We next studied the high-field capacitive energy storage properties at high temperatures based on the measurement of the unipolar electric polarization–electric field (P–E) loops of the dielectric films, as presented in Fig. S8–S12, ESI.† The slimmer P–E loops of the low-entropy polymer of PPS/PS 75%/25% are indicative of lower conduction loss than that of the commercial high-T_g polymer dielectrics, such as PEI and PC. As summarized in Fig. 4A and B, PPO/PS 75%/25% outperforms all the commercial high-T_g polymer dielectrics at elevated temperatures in terms of the discharged energy density (U_e) and the charge–discharge efficiency (η). For example, U_e of PPO/PS 75%/25% exceeds 5.5 J cm⁻³ under 600 MV m⁻¹ with a η of larger than 90% at 150 °C, which is more than twice that of neat PPO of 2.3 J cm⁻³, accompanied by an electric field value more than 1.5 times higher than that of neat PPO at 90% η. Furthermore, the U_e of PPO/PS 75%/25% at 90% η and 150 °C is more than twice that of PEI (2.6 J cm⁻³), the best high-temperature dielectric polymer. The high efficiency is linked to the low dielectric loss and suppressed leakage current in the low-entropy amorphous polymer. Fig. 4C compared the energy storage performance at 90% η of our low-entropy amorphous polymer with those of the state-of-the-art all-organic thermoplastic dielectric polymers,^10,30–37 for example, o-POFNB³⁰ with an U_e of 4.8 J cm⁻³, PEI/NTCDA³¹ with an U_e of 4.2 J cm⁻³, and ht-PEKNA³⁷ with an U_e of 3.1 J cm⁻³. Obviously, the low-entropy amorphous polymer film shows higher U_e than those of commercially available capacitor films and comparable U_e with the latest reported state-of-the-art all-organic thermoplastic dielectric polymers, and stands out for its scalable, low-cost, and simple polymer blending strategy, thus paving the way for the mass fabrication of high-performance and high-quality polymer films required for high-temperature film capacitors.

Fig. 4 Electrostatic energy storage performance. (A) Discharged energy density (U_d) and (B) charge–discharge efficiency (η) of PPO/PS 75%/25% and commercial dielectric polymers measured at 150 °C. (C) Comparison of discharged energy density at 150 °C and above 90% efficiency between this work and that of other reported high-temperature all-organic thermoplastic dielectric polymers. (D) Cyclic performance of neat PPO and PPO/PS 75%/25% film capacitors at 150 °C and 400 MV m⁻¹. (E) SEM image and EDS analysis of PPO/PS 75%/25% film after electrical breakdown. (F) P–E loops of PPO/PS 75%/25% film before and after the self-healing process at 350 MV m⁻¹.

Considering the practical applications, we evaluated the performance reliability and self-healing capability of our low-entropy amorphous polymer films. In the cycling reliability test, the films were charged and discharged repeatedly over 30 000 cycles at an electric field of 400 MV m⁻¹ (much higher than that under the operating condition, i.e., 200 MV m⁻¹, of capacitors in common power systems such as electric vehicles) and a temperature of 150 °C. The energy storage properties are presented in Fig. 4D. The low-entropy amorphous polymer film with the PPO/PS composition of 75%/25% survive after 30 000 cycles with a slight performance degradation, that is, ≤0.6%% for U_e with η maintained at >95% whereas the neat PPO showed low cycling reliability and breakdown in less than 64 cycles at 150 °C and 400 MV m⁻¹. We attribute the excellent cycling reliability to the low leakage current density and high breakdown strength in the low-entropy films, which will promote the broad applicability of these dielectric capacitors.

The PPO/PS 75%/25% film capacitor also exhibits excellent self-healing capability. As evidenced by scanning electron microscopy (SEM) images and energy-dispersive spectroscopy (EDS) mapping (Fig. 4E), after a breakdown occurred in a metalized PPO/PS 75%25% film, the electrodes were mostly evaporated and the carbonization area was isolated at the breakdown hole. Therefore, no conductive path was formed between the carbonized breakdown hole and the electrodes in the low-entropy amorphous polymer. Subsequently, we showed that, following a dielectric breakdown, PPO/PS 75%/25% worked normally in the next charge–discharge cycle, with tiny decreases in U_e (from 2.3 to 2.1 J cm⁻³) and η (from 96.3% to 93.2%) at 150 °C (Fig. 4F). It is empirically understood that, for a polymer with the general formula C_αH_βO_γN_δS_θ, a low ratio of (α + δ + θ) to (β + γ) typically enables the polymer molecules to decompose into gas with less residue of deposited carbon during electric breakdown, and yields a good self-healing ability.³⁸ In this sense, the self-healing of PPO/PS 75%/25% capacitor attributed to the low ratio of (α + δ + θ) to (β + γ), this is 0.89 and 1.0 for PPO and PS, respectively, much lower than those of the aromatic polymer dielectrics, including PEI (1.3), polyimides (1.6) and PEEK (1.33).⁶ The excellent breakdown self-healing capability of PPO/PS 75%/25% offers high reliability to its metallized film capacitors for the practical applicability.

Conclusions

In summary, we demonstrated a low-entropy strategy to enhance the high-temperate capacitive performance of amorphous dielectric polymers. The dense packing behavior with locally extended polymer chain conformation resulted in short-range ordered but long-range disordered microstructures, inhibiting the transport of electrons in dielectric polymers. Thus, a high energy density of 5.5 J cm⁻³ and efficiency of ≥90% were achieved in the low-entropy PPO/PS 75%/25% film at 150 °C, outperforming the existing dielectric polymers and showing a more than two-fold increase in discharged energy density compared to neat PPO. These results pave the way for regulating the conformation entropy of polymer chains to reduce leakage current density and achieve high energy density and high charge–discharge efficiency of high-temperature dielectric polymers. Since the raw materials are readily commercial accessible, and the film processing, that is general polymer blending, is simple, straightforward, and low-cost, this low-entropy strategy shows promise to address the challenge in scalable fabrication of high-performance and high-quality polymer films required for high-temperature capacitive energy storage.

Experimental

Materials

Poly(phenylene oxide) (PPO) used in the experiments was commercially available PX-100L (Mitsubishi Engineering-Plastic Corporation, Japan). Polystyrene (PS) and xylene were purchased from Aladdin Chemistry Co. Ltd, China. All commercially available reagents were used directly without further treatment unless otherwise noted.

Preparation of low-entropy PPO/PS films

To prepare the low-entropy PPO/PS films, a suitable amount of PPO powder and PS pellets with a designed weight ratio, was dissolved in xylene and stirred for over 12 hours at 60 °C to yield a clear and homogeneous solution. Then, the solution was drop-cast on a clean glass slide and kept in a drying oven at 60 °C for 6 hours to remove the solvent. Next, the films were further heated at 120 °C for 12 hours under vacuum to further remove any residual solvent. The cast film was peeled off from the glass substrate in deionized water and then dried at 120 °C in a vacuum oven for another 12 h to obtain the free-standing films.

Structural characterization

Differential scanning calorimetry (DSC) data was measured under a nitrogen atmosphere (50 ml min⁻¹) by using a PerkinElmer DSC 4000 at a heating and cooling rate of 10 °C min⁻¹ to study glass transition of PPO/PS blends with various compositions. Tapping-mode AFM (MFP-3D Infinity, Asylum Research) was used to simultaneously obtain topography and phase shift images using the same tip and measurement settings for all samples. The cryo-fractured cross-section micromorphology of the films was characterized via FIB-SEM (FEI Scios, Thermo Fisher). Fourier-transform infrared (FTIR) analysis was performed on a Nicolet iS10 (Thermo Scientific) FTIR spectrometer by averaging 100 scans at a resolution of 2 cm⁻¹. X-ray diffraction (XRD) analysis was performed using a Rigaku MiniFlex600 diffractometer. The radiation source was a Cu K_α source with a wavelength of 1.54 angstroms (Å). The densities of the PPO/PS blends, PPO, and PS were measured by using the AccuPyc II 1340 gas displacement pycnometer system. Specific heat capacity data of the films were measured under a nitrogen atmosphere (50 ml min⁻¹) by using a discovery DSC 25 at a heating rate of 5 °C min⁻¹. The storage modulus of the PPO/PS films was measured by using a DMA instrument (DMA850, TA instruments) equipped with a tensile clamp.

Dielectric and electrical measurements

The typical thickness of films used for electrical characterizations was within the range of 8–15 μm, measured by Millimar C1200M compact length measuring instrument. Gold electrodes with a diameter of 6 mm and thicknesses of 60 nm were sputtered on both sides of the polymer films for all the electrical measurements. Dielectric spectra were acquired over a broad temperature range from 25 °C to 200 °C and a broad frequency range from 100 Hz to 1 MHz by using a Keysight E4980AL LCR meter in conjunction with a Sun Systems environment chamber equipped with a liquid nitrogen cooling system. The breakdown strengths at various temperatures were measured using a high-voltage amplifier system (Trek 610E, United States). The temperature was controlled by utilizing a digital hot plate equipped with a thermal couple. Breakdown strengths were determined by using two-parameter Weibull statistics in more than 15 samples. The leakage currents were measured under various electric fields by using a Keithley 6517B pA meter and a Trek 610E amplifier as the DC voltage source.

Electric polarization–electric field (P–E) loop and discharge measurements

The polarization–electric field unipolar loops (P–E loops) at different temperatures of the films were measured using a modified Sawyer–Tower circuit under a unipolar triangle voltage with a frequency of 100 Hz with a Trek 610E high-voltage amplifier system. All the P–E loop measurements were performed with the sample immersed in insulating silicone oil, whose temperature is controlled by a digital hot plate equipped with a thermal couple. The cyclic charge–discharge tests were performed by using a PK-CPR1502 test system (PolyK Technologies, United States). Electrical conductivity was measured under an electric field of 200 MV m⁻¹ using a Keithley 6517B pA meter and a Trek 610E amplifier as the DC voltage source.

Author contributions

QYZ directed the research and designed the experiment. DML and YQZ performed experiments. QYZ, YNH, SWH, SXD and QMZ participated in the discussion. QYZ and QMZ wrote the paper.

Data availability

The data supporting this study's findings are available from the corresponding author upon reasonable request.

Conflicts of interest

The authors declare no competing interests.

Acknowledgements

This research was supported by the National Natural Science Foundation of China (grant no. 52203096) and the Shenzhen Science and Technology Innovation Program (RCBS20221008093110036 and ZDSYS20220527171402005). The authors thank the assistance with SEM observation received from the Electron Microscope Center of Shenzhen University.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4ee02455a

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