The ever-increasing demand for energy catalyzes a critical need for proficient and robust energy storage systems [1, 2]. Dielectric film capacitors, distinguished by ultrahigh power density, excellent rate capability and great reliability, are essential components in advanced electronics and electrical power systems, such as pulse power supplies, filters, and converters [3−5]. In particular, next-generation electronic devices toward miniaturization and high integration highlights the necessity for capacitors with high energy storage density and high operating temperatures. For instance, capacitors, a crucial element of power inverters in electric vehicles, accounts for up to 35% of the inverter volume and 25% of its weight [6] and are required to operate at a temperature of 140 °C or above [7−9]. Compared with the traditional dielectric ceramics for high-temperature capacitors, polymer dielectrics are more preferred due to their inherent advantages including high breakdown strength, light weight and facile processability [10−15]. Biaxially oriented polypropylene (BOPP), the mainstream commercially available dielectric polymer to date, however, suffers from relatively low discharged energy density (~ 2 J/cm³) and maximum working temperature below 140 °C [3, 7]. It is thus required an increase in capacitance volume and the integration of auxiliary cooling systems to ensure efficient operation, falling far short of the escalating demand on dielectric materials in compact high-power electronics [16, 17].
To address these urgent needs, inorganic fillers have been incorporated into a variety of polymers with high glass transition temperatures (T g), including polycarbonate (PC), polyetheretherketone (PEEK), polyimide (PI), polyetherimide (PEI), and crosslinked divinyltetramethyldisiloxane-bis(benzocyclobutene) (c-BCB), to exploit high-temperature dielectric polymer composites [18−22]. Particularly, PI, recognized as one of the most prevalent high-temperature engineering plastics, is distinguished by its remarkably high T g, exceptional resistance to both thermal and chemicals, great mechanical strength, and outstanding insulating properties [23−25], making it a promising candidate for high-temperature film capacitor applications. It is known that energy density is highly dependent on the dielectric constant (k) and breakdown strength. While high-k nanofillers such as barium titanate (BT) [26], strontium titanate (ST) [27] and titanium dioxide (TiO 2) [28] are widely utilized to enhance the dielectric constant of the composites, the substantial mismatch in k between nanofillers and polymer matrices distorts the electric fields. The distortion often leads to a pronounced increase in leakage current and a decrease in breakdown strength, especially at elevated temperatures, impeding the improvement in energy density. More recently, wide band-gap nano-fillers, represented by boron nitride nanosheets (BNNS), alumina (Al 2 O 3), magnesium oxide (MgO) and hafnium oxide (HfO 2), have been demonstrated to effectively improve the breakdown strength of composites, thereby facilitating superior high-temperature energy storage capabilities [29−32]. For example, the PI/Al 2 O 3 and PI/HfO 2 composites with large bandgap exhibit higher breakdown strength of 422 and 397 MV/m, respectively, with respect to high-k TiO 2 based PI composite (340 MV/m) [29]. Furthermore, it is suggested that manipulation of fillers geometry, including spherical, fibrous or linear, and planar or lamellar shapes, as well as their spatial arrangement within the matrix, such as oriented, arrayed and multiphasic distributions, could boost the dielectric capacitive performance of polymer nanocomposites [27, 28, 33−35]. BT nanofibers with in-plane distribution [36], in comparison with their nanoparticle counterparts [26], demonstrate marked advantages in inhibiting in-plane charge carries shifting and reducing out-plane leakage current, thereby presenting an enhanced breakdown strength and energy storage performance.
Noteworthily, charge–discharge efficiency (η, η = U e/(U e + U l) × 100%, U e: discharged energy density, U l : energy loss) is a key criteria for evaluating the dielectric capacitive performance. High η, that is, low U l, is of great significance to the operating efficiency and reliability of capacitors. Although it is anticipated that the engineering polymer mentioned before would retain dielectric stability at elevated temperatures owing to their high T g and thermal stability, their η declines significantly due to exponential increase of conduction current with increasing temperature and electric field [37−39]. For instance, at 100 MV/m, the leakage current density of PEI/HfO 2 increase from 2.5 × 10−11 A/cm 2 at 25 °C to 1.3 × 10−9 A/cm 2 at 150 °C, leading to a sharp decrease in η from 97% at 25 °C to 72% at 150 °C at 400 MV/m [40]. Moreover, sustainable Joule heat inside the dielectrics generated by large conduction current may cause thermal runaway and deteriorate capacitors lifetime if not being dissipated efficiently [7, 38, 41]. It has been elucidated that a mere 2 °C elevation in the temperature of an electronic component can diminish its performance stability by 10% [42]. However, polymer dielectrics are typically poor heat conductors with relatively low thermal conductivity (0.2–0.5 W/(m·K)) [7, 38, 43−45]. Improving thermal management ability is therefore also desirable for polymer dielectrics to achieve high-temperature energy storage performance and reliable operation except for hindering conduction current. Until now, the incorporation of BNNS, known for their high thermal conductivity, into polymers to control electrical conduction and facilitate heat dissipation, represents the most viable approach for optimizing high-temperature energy storage capabilities [7, 46−48]. For example, integrating BNNS significantly boosts the thermal conductivity of the c-BCB polymer, facilitating a more homogeneous temperature distribution within the polymer capacitor. This composite achieves a η of 97% at 150 °C and 200 MV/m, comparable to that of BOPP at 70 °C [7]. Nevertheless, the k of BNNS is relatively low (i.e., ~ 4), and the production process for BNNS is intricate, primarily involving liquid-phase exfoliation, a method that is both inefficient and challenging to scale.
Herein, a scalable polyimide nanocomposite containing two-dimensional (2D) alumina nanoplates (Al 2 O 3-NPLs) has been developed via in situ polymerization for high-temperature capacitive energy storage. Compared with BNNS, Al 2 O 3 possess better dielectric insulation properties (e.g., large band gap of 8.6 vs. 5.97 eV, high k of 9.5 vs. 4) and acceptable thermal conductivity of 30 W/(m·K), which emerges as a more cost-efficient and readily processable alternative. The superiority of Al 2 O 3 nanoplates in improving field-dependent electrical conduction and thermal management ability, and thus enhancing the elevated-temperature energy storage properties of the composites has been uncovered experimentally and theoretically, along with the comparison with nanoparticles and nanowires filled composites. Benefiting from the simultaneously suppressed conduction loss and increased thermal conductivity, the resulting PI/Al 2 O 3-NPLs composite exhibits a high-temperature capability up to 200 °C and remarkable η over a wide temperature range. Furthermore, the composite demonstrates exceptional thermal and cyclic stability, showcasing its potential utility under extreme environments.
The chemicals 4,4'-diaminodiphenyl ether (ODA, Analytical Reagent), pyromellitic dianhydride (PMDA, 99%), and N,N-dimethylacetamide (DMAc, anhydrous, 99.5%) were purchased from Sigma-Aldrich. Al 2 O 3 nanoparticles with an average diameter of approximately 30 nm were obtained from US-Nano Materials Inc. Al 2 O 3 nanoplates and nanowires were synthesized using a well-established hydrothermal method reported before [19]. The nanostructured Al 2 O 3 nanofillers were used without additional surface modification. The ODA and PMDA monomers were desiccated under vacuum at 80 °C for 12 h prior to film casting.
PI nanocomposites are fabricated via in situ polymerization followed by solution casting. Initially, a certain volume content of Al 2 O 3 nanofillers were dispersed in entirely soluble ODA-DMAc mixture to form a homogeneous suspension using a needle-tip ultrasonic processor (175 W) for 1 h. Subsequently, a predefined quantity of PMDA, maintaining a molar mass ratio of ODA to PMDA at 1:1.05, was introduced stepwise to the mixture of ODA and Al 2 O 3 suspension, followed by further stirred vigorously for 12 h under a nitrogen atmosphere to produce a homogeneous poly(amic acid) (PAA)-Al 2 O 3 precursor solution. This precursor was then evenly cast onto clean glass slides and dried at 70 °C for 6 h under vacuum. After complete solvent removal, the as-prepared films undergone a sequential thermal imidization at 100, 200, and 300 °C for 1 h each. Finally, the films are detached from the glass slides and vacuum dried at 105 °C for 12 h to eliminate any residual moisture or solvent, resulting in nanocomposite films with an average thickness of 12 μm.
The micro-morphologies of nanostructured Al 2 O 3 and PI nanocomposites were characterized using transmission electron microscopy (TEM, JEOL JEM-2100Plus) and scanning electron microscope (SEM, FEI Scios 2, USA). X-ray diffraction (XRD) was carried out on a PANalytical Xpert Pro MPD θ–θ diffractometer using Cu K α radiation with a wavelength of 1.54 Å. Fourier transform infrared (FT-IR) spectra are acquired with a FTS-8010 Fourier transform infrared spectrometer from Varian Digilab, performed at room temperature in attenuated total reflection (ATR) mode, featuring a resolution of 4 cm−1 and a wavelength range of 4000 to 400 cm−1. Thermal gravimetric analysis (TGA) and differential scanning calorimetry (DSC) tests were conducted using a METTLER TOLEDO TGA/DSC3+ analyzer and a TA instrument model DISCOVER DSC250, respectively, at a heating rate of 10 °C/min under a nitrogen atmosphere. Young's moduli were determined from the stress–strain curves measured using a TA RSA G2 Dynamic Thermo-mechanical Analyzer (DMA) with a linear stretching rate of 0.02%/s. The films were cut into rectangular strips measuring 40 mm in length and 5 mm in width. For electrical measurements, both sides of the nanocomposite films were sputtered with gold electrodes of a diameter of 6 mm for the dielectric spectra and a diameter of 3 mm for the rest tests. Dielectric spectra over a frequency range of 100 Hz to 1 MHz and a temperature range from 25 to 200 °C were obtained using a Hewlett Packard E4980A precision LCR meter. Dielectric breakdown strength was assessed using a Trek 610C instrument under a DC voltage ramp of 500 V/s, with analysis based on a two-parameter Weibull distribution model from at least 15 samples for each test. Electrical displacement–electric field (D–E) loops were measured by a modified Sawyer-Tower Circuit with a high-voltage amplifier system (PolyK Technologies) using a triangular unipolar wave at 10 Hz. Both breakdown strength and D–E loop tests were conducted in a Dow Corning HT insulating oil bath. Conduction current was collected through a Hewlett Packard 4140B pA meter. Thermal diffusivity (α) was measured by a NETZSCH LFA467 laser flash analyzer, and the in-plane thermal conductivity (TC) was calculated according to the formula TC = α·C p·ρ, where C p represents the specific heat capacity acquired from DSC tests, and ρ is the density. The samples were cut into discs with a diameter of 25 mm, pre-treated with a thin layer of fine graphite powder on both surfaces. The electric field distribution, current density distribution and steady-state temperature distribution were simulated by the finite element calculations as detailed in the Electronic Supplementary Material (ESM). The simulation models were shown in Fig. S1 in the ESM.
The two-dimensional hexagonal Al 2 O 3 nanoplates with an average thickness of 30 nm and a width of 800–1000 nm were synthesised successfully by hydrothermal method, as verified by SEM and XRD (Figs. 1(a)–1(c)). The observed diffraction peaks in XRD pattern at 2 θ of 67°, 45° and 37° correspond to the typical face-centered cubic structure for γ-Al 2 O 3 (JCPDS No. 50-0741) [19]. No impurity peaks are observed, indicating the good crystallinity and high purity of Al 2 O 3-NPLs. The as-prepared Al 2 O 3-NPLs were then utilized to prepare the polyimide nanocomposites by in situ polymerization. The chemical structures of the nanocomposites are identified by FT-IR spectra shown in Fig. 1(d). The disappearance of characteristic PAA peaks at 1624 cm−1 (C=O stretching) and 3252 cm−1 (N–H stretching), and the formation of new peaks for imide groups at 1755, 1720, 1366 and 723 cm−1 arising from C=O asymmetric stretching, C=O symmetric stretching, C–N asymmetric stretching and C=O bending, respectively, indicate the completely thermal imidization from PAA to PI. The broad absorption peaks at around 500–900 cm−1, which are assigned to the AlO 6 octahedra vibration, confirm the incorporation of Al 2 O 3 fillers into PI matrix. Besides, the presence and distribution of Al 2 O 3-NPLs in the PI nanocomposites has also been affirmed by cross-sectional SEM images in Fig. S2 in the ESM. As clearly shown in Fig. 1(e), Al 2 O 3 nanoplates are uniformly dispersed in the polymer matrix without obvious agglomeration, yielding high-quality nanocomposites with a thickness of ~ 12 μm. Notably, the solution-casting composite films show a mostly parallel alignment of Al 2 O 3 nanoplates to the PI matrix, similar to the previously reported results [28, 49, 50], which is desirable for hindering the development of breakdown paths and thus achieving excellent dielectric performance.
