Introduction

Dielectric film capacitors with ultrahigh power density are widely applied in the fields of power transmission, pulsed power source, and electric equipment1,2,3,4,5,6,7. Biaxially oriented polypropylene (BOPP) film is one of the most successfully commercialized capacitor materials, but its mechanical strength and thermostability are insufficient, which will easily lose mechanical stability when operation temperature exceeds 100 °C, finally restricting its applications in numerous high-temperature scenes, including electric vehicles, oil/gas explorations and aerospace systems and so on8,9,10,11. Therefore, the high-temperature resistant polymer films (glass-transition temperature (Tg) generally ≥200 °C) have been one of the key research directions in this field12,13,14,15. Polyimide (PI) dielectric film is the most widely used high-temperature resistant polymer material. Through in-depth studies in recent years, PI has achieved relatively high energy-storage level at high-temperature environments, but it is still challenging to meet currently higher requirements in practical applications16,17,18,19,20; Particularly, the relatively low mechanical strength of PI film gradually shows its disadvantage when facing more complicated and hasher application scenarios.

It should be noted that the existed problems of PI above are mainly from its intrinsic backbone structure: First, the imide ring of PI chain generally only owns the relatively weak interchain interactions, thus leading to low mechanical strength. However, the superimposed bending, electrical, and thermal stresses on capacitors always require high mechanical strength of films to prevent breakdown failure10,21,22,23,24. Second, PI chain generally contains a lot of aromatic nucleus structures with strong conjugation effect and small energy gap, making it challenging to restrict electron excitation and carrier formation based on Poole–Frenkel effect, and the subsequent transport process of carrier will produce high conduction loss25,26,27,28. In addition to conduction loss, the polarization loss will also generate owing to the friction interactions among macromolecular chains29. The two overmuch losses cause that the energy-storage level of PI film is always limited up to now. Third, existed studies have conveyed a consensus that reducing interchain conjugation effect and friction interaction can indeed obviously enhance the energy-storage level, but the further reduced interchain interactions also cause obviously inferior mechanical property, including widely used polyetherimide (PEI) film, which has evolved into an intractable and contradictory problem in dielectric energy storage field.

To date, little valuable all-organic design strategy can solve the key problem of polymer dielectrics above. Therefore, ignoring the polymer bulk provisionally, adding high-quality nanofiller into polymer may synergistically strengthen mechanical and energy-storage performances of polymer dielectrics, but the prerequisite is that the excellent properties of the nanofillers are required to fully transfer into the nanocomposite material, which presents huge challenge for nanocomposite materials design and preparation so far30,31,32,33,34. Meanwhile, the nanocomposite material design route also has its difficulty to match current production process toward large-scale fabrication of high-quality polymer film capacitors. Therefore, despite the greater difficulties, it is still urgent to develop an all-organic route to better synergistically strengthen the mechanical and energy-storage performances of polymer dielectrics.

To address these key issues above, this work will purposefully remodel the intrinsic backbone structure (dianhydride and diamine) of PI. Targeted simulation calculations are firstly adopted to predict better alternative units to enhance the mechanical strength and energy-storage level. The dianhydride of PI is transformed into diacyl chloride to form amido bond, which not only introduces abundant and strong hydrogen bonds to enhance mechanical strength, but also slightly improves the energy-storage level of PI film, because introduced hydrogen bonds separately increase interchain interactions and do not reduce the energy gap, which is different from other π–π and charge-transfer (CT) interactions. Then, the traditional diamine is further transformed into fluorine-containing diamine with large energy gap and strong self-lubricating ability that can effectively restrict carrier formation and reduce losses, thereby dramatically strengthening the energy-storage level of PI film.

For the predicted all-organic polymer structure (PTFMB) above through stepwise molecular engineering design, the special complexing behavior of Lewis acid on PTFMB chain as a “transfer station” is further applied to ameliorate chain dispersibility in solution, and then Lewis acid is subsequently removed by uncomplexing effect to construct well-organized and compact chain stacking during postprocessing. Moreover, the Lewis acid content regulation and high-temperature-assisted chain arrangement are also adjusted to sharply enhance film crystallinity, which further strengthens the energy-storage level of PTFMB film. Finally, compared with traditional PI (or PEI) films, the fabricated PTFMB film exhibits a synergistic enhancement in mechanical strength and high-temperature energy-storage performances. This work demonstrates that the synthesis of a targeted high-temperature resistant polymer dielectric can synergistically strengthen the anticorrelated mechanical strength and energy-storage level at harsh environments, which also offers an all-organic simulation prediction and experiment design strategy for large-scale production of high-quality polymer dielectrics.

Results and discussion

Simulated predictions for stepwise molecular engineering design

As shown in Fig. 1a, Supplementary Figs. 1 and 2, and Supplementary Note 1, it is apparent that the inherently two core units of dianhydride and diamine determine the related mechanical and dielectric properties of polyimide (PI or PEI) films. Here, simulated calculations are firstly adopted to predict better alternative structure units to substitute the two traditional core units above for synergistically strengthening both performances of PI or PEI film.

Fig. 1: Stepwise molecular engineering design and computational simulations.
figure 1

a Schematic diagram of interchain interactions between different segments within different polymer molecules. b DFT calculations on HOMO and LUMO energy levels and predicted bandgap energy (Eg) from single polymer chain to two adjacent chains of PI, PABI, and PTFMB77. c Molecular planarity coloring map, three-dimensional structures in amorphous cell and interchain hydrogen bonds within different polymer structures. d Number of hydrogen bonds (H-bonds) and strong H-bonds. e Linear fitting of MSD-t curves. f The cohesive energy density (CED) and diffusion coefficient (D) of PI, PABI. and PTFMB.

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For mechanical strength, the diamine containing benzimidazole structure (PABZ) capable of forming robust interchain hydrogen bonds has been introduced into PI chain in our previous work35, which can enhance the mechanical strength and thermostability compared to the widely used PEI, almost without the ability of forming interchain hydrogen bonds (Fig. 1a, Supplementary Fig. 3, and Supplementary Note 2). The formation of interchain hydrogen bonds in PI model with PABZ structure was also observed, accompanied by a total of 225 hydrogen bonds (Fig. 1d and Supplementary Table 1). The newly introduced interchain hydrogen bonds can improve the mechanical strength of PI film, but the currently presented value is still relatively low (generally <250 MPa, Supplementary Fig. 3). This can be attributed to the fact that the imide rings formed in PI chain still lead to relatively weak interchain forces, though PABZ is present.

Therefore, diacyl chloride was selected to substitute the dianhydride component of PI. The transformation from imide rings to amido bonds in the modified polymer structure (PABI) introduces interchain hydrogen bonds, and Fig. 1d shows the existence of 1453 hydrogen bonds in PABI model. Moreover, the strong/weak hydrogen bond in polymer structure is further analyzed and defined based on the statistical bond lengths of different hydrogen bonds (Fig. 1d and Supplementary Note 3). The result indicates that PABI owns the higher percentage of strong hydrogen bond (36.7%) than that of PI only with 20.4%. In addition, different from π–π and CT interactions, the abundant and strong hydrogen bond interactions above seem not to increase the energy gap, and simulation calculations indicate that the energy gap of PABI (3.85 eV) even has a little increase relative to that of PI (3.21 eV) (Fig. 1b). Therefore, the introduced abundant and strong hydrogen bonds hold significant potential for synergistically enhancing the mechanical and dielectric properties of polymer film.

For energy-storage level, through analyzing the backbone structure, it is found that traditional high-temperature resistant polymers, originally engineered for high thermal stability, consist of numerous aromatic nucleus structures, thereby forming conjugated planar segments that will produce strong interchain π–π stacking and CT interactions with diminished energy gaps. For PABI above, the dihedral angle (an angle between two benzene rings) of backbone structure is about 7.51°, exhibiting strong conjugation, so the energy gap is only 3.85 eV (Fig. 1b and Supplementary Fig. 4). Notably, lowering of energy gap inevitably causes compromised high-field performance and poor energy density, especially at high-temperature environments.

To weaken interchain conjugation effects, a diamine with biphenyl structure—featuring two benzene rings in a mutually perpendicular configuration—is chosen to replace the traditional diamine above. Supplementary Fig. 4 shows the diamine with biphenyl structure contains electron-donating –CH3 group, and the calculated dihedral angle of formed structure is 87.85°, which nearly maximizes the distortion degree and avoids the conjugation of backbone structure, thereby effectively increasing energy gap to 4.60 eV (Supplementary Fig. 5). Then, a strong electron-absorbing –CF3 group is further introduced to substitute –CH3 group above. The calculated dihedral angle of 77.73° in formed polymer structure (PTFMB) has a subtle decrease (Supplementary Fig. 4), which should be attributed to that the F atom with strong electronegativity can form hydrogen bond with the surrounding H atom to affect corresponding dihedral angle (Supplementary Fig. 6). However, the F atom simultaneously owns a strong ability to limit electron excitation and carrier formation, which can thus increase the energy gap of PTFMB up to a higher value of 4.68 eV (Fig. 1b)36,37.

Notably, current theoretical studies on bandgap simulations predominantly employ simplified small-molecule systems38,39,40,41. However, the electron delocalization extent in simplified systems is significantly lower than that in practical polymer systems, consequently yielding overestimated bandgap values. Therefore, we further conducted a systematic examination on bandgap evolution in dimeric systems, which indeed revealed reduced bandgap values of 2.74 eV (PI), 3.56 eV (PABI), and 3.91 eV (PTFMB) (Fig. 1b and Supplementary Fig. 7, 8). The introduced interchain interactions (e.g., hydrogen bonding, π–π stacking, and charge transfer) substantially decreased the bandgaps across all systems, but PTFMB system consistently exhibited higher bandgap characteristic than that of PI and PABI counterpart.

Moreover, the molecular planarity determined by molecular planarity parameter (MPP) and span of deviation from plane (SDP) affects carrier transport behavior (Supplementary Note 4)42. Figure 1c demonstrates that PABI exhibits an exceptional planarity with an MPP of 0.295 Å, indicating tight atomic distribution along the ideal plane, and an SDP of 1.244 Å, reflecting minimal surface undulations and uniform two-dimensional extension, owing to its highly symmetric conjugated aromatic framework. In contrast, PTFMB shows markedly increased non-planarity (MPP = 0.971 Å, SDP = 3.661 Å), owing to its twisted biphenyl structure disrupting π-conjugation continuity and substantial steric hindrance from -CF₃ group, inducing out-of-plane distortions. The dimer system also indicates that PI exhibits molecular stacking characteristic (Supplementary Fig. 14b), where the high-density conjugation overlap promotes significant delocalization of electron clouds between molecules.

In addition, our previous work has proposed an experiential formula about frictional work (W) between polymer chains under electric field based on numerous experiment and simulation data (Fig. 1e, f, Supplementary Figs. 9 and 10, and Supplementary Notes 5–7)43,44:

$$W\propto A\times {CED}\times D$$
(1)

where A is the constant, CED is the cohesive energy density, D is the diffusion coefficient. D describes the amount of diffusing atoms passing through a unit area per unit time, reflecting the diffusing rate of diffusing atoms.

$${{{{\rm{D}}}}}_{\infty }=\frac{1}{6{{{{\rm{N}}}}}_{\infty }}{{{\mathrm{lim}}}}_{{{{\rm{i}}}}\to \infty }\frac{{{{\rm{d}}}}}{{{{\rm{dt}}}}}{\sum}_{{{{\rm{i}}}}=1}^{{{{{\rm{N}}}}}_{{{{\rm{a}}}}}} < {{|{{{\rm{r}}}}}_{{{{\rm{i}}}}}\left({{{\rm{t}}}}\right)-{{{{\rm{r}}}}}_{{{{\rm{i}}}}}\left(0\right)|}^{2} > $$
(2)

Equation (2) can be approximated by Eq. (3):

$$\,D=\frac{a}{6}$$
(3)

Where a is the derivative of MSD with time (Supplementary Fig. 9). It is found that PTFMB possesses the smallest CED and D, thus contributing to the smallest W compared with that of PI and PABI. The study demonstrates that PTFMB effectively reduces the interchain friction energy expenditure when the polar groups undergo orientation under electric field, thereby reducing the polarization loss45,46,47,48. Meanwhile, PTFMB also owns a larger interchain free volume because of large dihedral angle as demonstrated by simulation calculation and experimental results, which further reduces the polarization loss (Supplementary Figs. 11 and 12 and Supplementary Note 8).

In addition, visualization analysis by the independent gradient model based on Hirshfeld partition (IGMH) method further shows that a continuous interaction region spanning the entire conjugated plane is formed in PI and PABI49 (Fig. 2a). In contrast, PTFMB presents the molecular plane deviating from coplanar alignment, and the related IGMH analysis reveals the fragmented interaction regions and distinct interaction interruption bands at the biphenyl axis twist.

Fig. 2: Further computational simulations on intermolecular interactions and electrical states of different structures.
figure 2

a Visualization of interchain interactions of PI, PABI, and PTFMB structures though IGMH method. b Hole&ele distributions of three polymer structures. c Chole&Cele distributions of three polymer structures. d Excitation energy of the first excited state (S1) for three polymer structures. e Charge transfer matrix. f Total transfer amounts of electrons for the first five excited states for three polymer structures.

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Moreover, the electron transitions from the ground state to the first five excited states were calculated, respectively50. As illustrated in Supplementary Fig. 13 and Supplementary Note 9, PI exhibited the lowest excitation energy, while PTFMB had the highest thresholds, which indicates that PI undergoes electron excitation with greater ease compared to that of PTFMB. For the first excited state (S1), the hole distribution representing the spatial origin of excited electrons reflects significant excitation activity in specific regions (Supplementary Fig. 14). The excitation in PI and PABI predominantly localizes on diamine moiety, whereas PTFMB exhibits additional excitation on the benzene rings of acyl chloride group (Supplementary Fig. 15). Conversely, the electron distribution reveals distinct spatial patterns (Supplementary Fig. 14d). PI concentrates electrons on dianhydride segment, and PABI displays delocalized electron redistribution, and PTFMB shows electron accumulation on acyl chloride and benzene rings adjacent to amide bonds. The integration of hole and electron distributions generates hole-electron overlap map, while spatial smoothing of these distributions can yield Chole and Cele profiles. As shown in Fig. 2b, c, PI shows minimal spatial overlap between holes and electrons, indicating pronounced CT characteristic. In contrast, PABI and PTFMB demonstrate substantial overlap regions, reflecting hybrid excitation behaviors combining localized and CT mechanisms. To quantitatively characterize the overlap between holes and electrons, the S-, D-, H-, and T-index parameters were defined (Supplementary Note 10). As shown in Supplementary Table 2, the S-index of PI was smaller than that of PABI and PTFMB, demonstrating enhanced charge separation. PTFMB showed the smallest D-, H-, and T-indices, indicating minimal centroid separation, limited spatial delocalization, and negligible charge separation.

To elucidate the direction of electron transfer, we numbered all non-hydrogen atoms, as shown in Supplementary Fig. 16 and Supplementary Note 11. In PI, the holes localize on the diamine segment while electrons concentrate on the imide rings, with minimal spatial overlap, whereas PABI shows broader electron distribution, and the electron transitions of PTFMB are restricted to nitrogen atoms and the adjacent benzene rings, excluding –CF₃ groups and other diamine carbons. We further employed charge transfer matrix (CTM)-based heatmaps (i.e., color-coded matrix plots) for visualization to analyze how molecular sites and inter-site coupling influence electron excitation characteristics (Supplementary Note 12). As shown in Fig. 2e and Supplementary Figs. 17–19, for PI and PABI, the high-intensity regions in the first excited state predominantly occupy Quadrant II, signifying dominant electron transfer from diamine to dianhydride segments. Notably, PABI shows extended delocalization patterns, as evidenced by the broader distribution of pale blue regions compared to PI. PTFMB shows the fundamentally altered behavior: (1) Significant contraction of delocalized regions (reduced pale blue areas) correlates with diminished inter-segment electron transfer. (2) Enhanced localized self-excitation within individual segments manifests through intensified Quadrant I/III signals. These observations collectively indicate that fluorinated substituents in PTFMB restrict the long-range charge transfer while promoting intra-segment excitonic confinement. The interfragment charge transfer method was further employed to quantify the electron transfer between diamine and dianhydride (diacyl chloride) fragments during electron excitations. As shown in Supplementary Fig. 20, Fig. 2f, and Supplementary Note 13, PI exhibits the strongest CT characteristic (90.57%) and highest electron transfer magnitude (2.779 eV), while replacing dianhydride with diacyl chloride in PABI and PTFMB significantly reduces CT contributions (54.19% and 51.09%) and charge transfer (0.824 eV and 0.554 eV), where PTFMB shows the weakest charge redistribution.

In a word, such PTFMB by stepwise molecular engineering design above owns special molecular compositions, interchain interactions, steric configurations, and electrical characteristics, which thus presents great potentials to synergistically strengthen the mechanical and energy-storage performances of polymer dielectrics.

Matched synthesis route design

On the strength of above-mentioned simulation results, it has been preliminarily predicted that the alternative structure units of diacyl chloride and fluorine-containing diamine can be applied to substitute traditional dianhydride and diamine of PI for synergistically strengthening the mechanical strength and energy-storage level. For such PTFMB structure, the matched synthesis route will be explored to fabricate the high-quality polymer dielectric film. However, different from traditional PI polymerization system generally only in single solvent, the PTFMB polymerization system produces the gel phenomenon in single solvent during polymerization (Fig. 3aI), which severely prevents subsequent polymerization process. The simulation results above have consistently demonstrated that a significant number of robust hydrogen bonds were incorporated into PTFMB, so the formed PTFMB chains in early polymerization easily generate agglomeration because of strong interchain interactions, thereby causing a rapid gelation phenomenon. It is further found that the addition of Lewis acid (like LiCl) can efficiently prevent the generation of gel phenomenon, which thus guarantees the completion of polymerization process51,52,53,54,55,56.

Fig. 3: Characterization of polymer chain dispersion and film orientation.
figure 3

a Diagram of “complexing/uncomplexing” process between LiCl and PTFMB chain. b Dynamic viscosity of PTFMB in DMAc/LiCl solutions with the different LiCl contents. c Schematic of the relative energies of Li+ with one or two PTFMB model compounds. d Schematic of polarizing infrared test. e FTIR spectra of PTFMB-4%LiCl film with different polarization angles of 0 ~ 180°. f Polar diagrams of absorption as a function of incident polarized light angles.

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In order to clarify the interactions between solvent/LiCl and PTFMB chain, the PTFMB solutions with different LiCl contents were prepared. It is found that the viscosity of PTFMB solution gradually decreases with increased LiCl content, and the formed gel effect has vanished when the LiCl content reaches 3% (Fig. 3aII, b and Supplementary Fig. 21). The simulated calculations provide compelling evidence that LiCl can produce complexing interactions with the amido bond of PTFMB chain, destroying the strong interchain hydrogen bond interactions, which thus prevents the generation of gel phenomenon to accomplish polymerization process (Fig. 3c and Supplementary Note 14). In addition to promoting polymerization process, the complexing effect also guarantees the good dispersibility of single PTFMB chain in solvent, facilitating the formation of well-organized and compact chain packing structures during the subsequent processes of solvent removal and film fabrication (Fig. 3aI–aIII). It should be noted that, deprived of complexing-assisted dispersibility, PTFMB chains will produce aggregation, thereby leading to unordered and incompact chain packing, which is disadvantageous to the mechanical and dielectric properties of PTFMB film (Supplementary Fig. 22). The polarized FTIR test was applied to assess the orientation degree of chain packing based on calculated Dichroic ratio (R) (Fig. 3d–f, Supplementary Fig. 23, and Supplementary Note 15):

$$R=\frac{{I}_{0}+{I}_{180}}{{I}_{90}}$$
(4)

The presented result indicates that the “complexing/uncomplexing” behavior of LiCl endows PTFMB film with a higher orientation degree (R = 6.74) than that of PTFMB film without LiCl regulation (R = 4.88). Finally, the LiCl complexing effect assists to fabricate initial PTFMB film with well-organized and compact chain packing, and the subsequent uncomplexing of LiCl guarantees that strong interchain hydrogen bonds regain to fabricate high-quality PTFMB film (Fig. 3aIII and Supplementary Fig. 24).

It is known that the enhanced crystallinity of polymer dielectric film is advantageous for minimizing dielectric loss and elevating breakdown strength. We further found that LiCl content also influences the crystallinity of PTFMB film, which can be attributed to the enhanced orientation degree above. PTFMB-0% LiCl shows a broad diffuse peak around 23.2° and a narrower but low absorption peak near 18°, representing a low crystallinity. As the content of LiCl increases, the absorption peak near 18° gradually becomes narrower and higher, indicating an increase in crystallinity. (Fig. 4a, b and Supplementary Fig. 25). Supplementary Fig. 26 shows that the crystallinity of PTFMB film gets up to the maximum value (30.2%) at an appropriate LiCl content of about 4% compared with that of PTFMB without LiCl regulation (14.8%). It is further found that the crystallinity of PTFMB film can be altered through changing annealing temperatures, attributed to the high-temperature-assisted chain rearrangement. The crystallinity of PTFMB film almost linearly increases with increased annealing temperature, further increasing the crystallinity up to 40.52% (Fig. 4c, Supplementary Fig. 27, and Supplementary Note 16). In addition, the crystallinity size (D) is calculated by the integral breadth of corrected peak-profile according to Scherrer’s equation:

$$D=\frac{K\lambda }{\cos \theta \delta (2\theta )}$$
(5)

where K is a constant ≈ 1, δ(2θ) is the integral breadth expressed in radian, λ is the wavelength of radiation, and θ is the Bragg angle. Figure 4c shows that the crystallinity size of PTFMB film gradually increases with increased annealing temperature. These results prove that PTFMB films with different crystallinities and sizes can be regulated through controlling LiCl content and annealing temperature (Fig. 4d).

Fig. 4: Characterization of film crystallinity.
figure 4

a, b XRD patterns of PTFMB-0%LiCl and PTFMB-4%LiCl films. c Crystallinity and size of crystallites of PTFMB films with various different annealing temperatures. d Diagram of the crystallization process of PTFMB. e SAXS patterns and one-dimensional integral curves for SAXS patters of PTFMB-0%LiCl and PTFMB-4%LiCl films. f WAXS patterns and one-dimensional integral curves for WAXS patters of PTFMB-0%LiCl and PTFMB-4%LiCl films.

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Synchrotron radiation small angle X-ray scattering (SAXS) patterns test is further applied to study the orientation and crystallinity of PTFMB film. It is evident that PTFMB-4%LiCl exhibits enhanced diffraction streaks along the equatorial axis (Fig. 4e and Supplementary Fig. 28), whose half-width at half-maximum is lower that of PTFMB-0%LiCl, so the orientation degree of PTFMB-4%LiCl is higher than that of PTFMB-0%LiCl57. In addition, PTFMB-4%LiCl also exhibits a new crystallinity peak around 19° in Synchrotron radiation wide angle X-ray scattering (WAXS) patterns, aligning with the distinct crystallinity peak situated near 18° in the aforementioned XRD spectrum of PTFMB film, which further demonstrates the high crystallinity in PTFMB-4%LiCl compared to that of PTFMB-0%LiCl (Fig. 4f).

Mechanical and energy-storage properties analysis

In-situ heating FTIR test is applied to study the variations of hydrogen bonds, especially at high-temperature environments. Noda’s rules effectively clarify the fluctuations in peak intensity that occur within both synchronous (left column) and asynchronous (right column) correlation spectra of N–H groups58 (Supplementary Note 17). For PI, where Φ(3371, 3354) > 0 and Ψ(3371, 3354) < 0, it suggests that the peak gradually shifts from 3354 cm−1 to 3371 cm−1 during heating (Fig. 5a–c). For PTFMB, where Φ(3343, 3333) > 0 and Ψ(3343, 3333) < 0, it indicates that the peak gradually shifts from 3333 cm−1 to 3343 cm−1 (Fig. 5d–f). Both the wavenumbers of –NH groups forming hydrogen bonds in PTFMB (3333 cm−1 and 3343 cm−1) are lower that of PI (3354 cm−1 and 3371 cm−1) at 40 °C and 360 °C, reflecting the stronger hydrogen bonds in PTFMB than that of PI. Furthermore, as the temperature increases from 40 °C to 360 °C, the wavenumber shift of –NH groups in PTFMB is only 10 cm−1, which is smaller than that in PI (17 cm−1) and PABI (24 cm−1) (Fig. 5g, h and Supplementary Fig. 29). The result demonstrates that PTFMB possesses a stronger ability to maintain robust hydrogen bonds even at elevated temperatures. This exceptional ability should be ascribed to the high degree of orientation and crystallinity in PTFMB film, which effectively curtails potential deterioration of related performances.

Fig. 5: Characterization of in-situ heating FTIR and mechanical strength.
figure 5

a The dynamic infrared spectrum and peak shift of H-bonded N-H groups during heating from 40 °C to 360 °C of PI film. b, c The synchronous and asynchronous infrared spectra of PI film, with red and blue regions representing positive and negative correlation intensities, respectively. d The dynamic infrared spectrum and peak shift of H-bonded N-H groups during heating from 40 °C to 360 °C of PTFMB film. e, f The synchronous and asynchronous infrared spectra of PTFMB film. g H-bonded N-H groups stretching vibration peak positions from 40 °C and 360 °C for different samples. h Stress-strain curve. i Mechanical strength of PI, PABI and PTFMB films. Error bars represent the standard deviations, and the number of replicate experiments for the error bars is 5. j Comparison of mechanical strength among reported dielectric polymers. k Diagram of metallized film capacitor.

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It should be noted that the high mechanical strength of polymer film is crucial, because the superimposed bending, electrical, and thermal stresses are constantly generated during operation of film capacitors, posing significant challenges. The existence of strong hydrogen bonds in PTFMB film will increase the mechanical strength, and the high orientation and crystallinity degree in PTFMB film also have a synergetic strengthening effect (Supplementary Fig. 30). Therefore, the fabricated PTFMB film presents a higher mechanical strength of 378.3 MPa than that of PI film (208.7 MPa) and PABI film (285.4 MPa) (Fig. 5i and Supplementary Fig. 31). The presented mechanical strength of PTFMB film has also exceeded currently reported polymer films in dielectric energy storage field (Fig. 5j, Supplementary Table 3, and Supplementary Note 18)16,59,60,61,62,63. Notably, high temperature will obviously aggravate the stress-induced damage effect (Fig. 5k), but as demonstrated above, PTFMB film has a greater capacity to maintain strong hydrogen bonds at high-temperature environments, which continuously endows it with high mechanical strength, making it suitable for high-temperature dielectric energy storage applications.

Furthermore, UV–vis tests show PABI film has a little higher energy gap (2.75 eV) that of PI film (2.63 eV) (Fig. 6d), which is consistent with simulation trend above, so the breakdown strength (455 MV/m) and discharged energy density (1.36 J/cm3) of PABI film also have a slight increase compared with that of PI film (362 MV/m and 0.26 J/cm3) at 150 °C (Fig. 6f, Supplementary Figs. 34 and 35 and Supplementary Notes 19 and 20). The presented results suggest that the introduced abundant and robust hydrogen bonds not only significantly enhance the mechanical strength of polymer film but also do not compromise its breakdown strength and energy storage performances, even showing a slight improvement, which is different from other interchain interactions like π–π and CT interactions.

Fig. 6: Energy-storage properties tests of different films.
figure 6

a–c Photographs of PI, PABI, and PTFMB films and corresponding solutions. d Calculated optical energy gaps of PI, PABI, and PTFMB films based on UV–vis results. e Weibull distribution of breakdown strength of PTFMB films at different annealing temperatures. f Weibull distribution of breakdown strength of PI, PABI, and PTFMB films at 150 °C. g Dielectric constant and loss (tanδ) at 150 °C of PI, PABI and PTFMB films. h Dielectric loss at 150 °C and 100 Hz of PI, PABI, and PTFMB films. i Discharged energy density (Ud) and efficiency (η) at 150 °C of PI, PABI, and PTFMB films. j D-E curves at 350 MV/m and 150 °C. k The maximum discharged energy density with efficiency above 90% (Uη90) and the maximum discharged energy density (Umax) at 150 °C of PI, PABI and PTFMB films. l Comparison of discharged energy density at 150 °C with efficiency above 90% and Young’s Modulus among this work and reported high-temperature dielectric polymer films2,59,60,61,63,64,65,66,67.

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On the foundation of PABI structure, UV–vis/energy gap test further demonstrations that PTFMB film has a higher energy gap (3.30 eV) than that of PABI film by avoiding conjugated polymer backbones (Fig. 6d), which also is consistent with the simulation trend above. Meanwhile, the yellow PI and PABI films have also transformed into transparent PTFMB film because of reduced photoelectron excitation, which also reflects a significant increase in energy gap (Fig. 6a–c). Supplementary Fig. 36 presents AFM characterizations for PI, PABI, and PTFMB films, revealing high surface smoothness across all films. The larger energy gap of PTFMB has a strong ability to effectively restrict electron excitation for enhancing breakdown strength. Notably, as demonstrated above, the crystallinity of PTFMB film can be regulated, which can further regulate the related breakdown strength. The small crystallinity of PTFMB film at low annealing temperature of 80 °C only owns low breakdown strength (633 MV/m), and the higher crystallinity at increased annealing temperature gradually increases breakdown strength of PTFMB film up to 862 MV/m (Fig. 6e and Supplementary Note 21). However, when the annealing temperature exceeds a certain threshold (up to 400 °C), the high temperature will inevitably damage polymer structure. Consequently, although the crystallinity of PTFMB film is further enhanced, its breakdown strength exhibits an inverse decline (778 MV/m). In addition, Fig. 6g, h also shows that PTFMB film owns the lower dielectric loss at energy-storage test condition (150 °C, 100 Hz), which also reduces the polarization loss in PTFMB film as well.

Moreover, when exposed to a high-temperature environment at 150 °C, PTFMB film still maintains a high breakdown strength of 682 MV/m, which has an increase of 88% and 50% compared to that of PI film (362 MV/m) and PABI film (455 MV/m), respectively (Fig. 6f). The discharged energy density of PTFMB film (6.06 J/cm3) at a charge-discharge efficiency of 90% also presents a large increase of 346% and 2231% compared with that of PABI film (1.36 J/cm3) and PI film (0.26 J/cm3) at 150 °C (Fig. 6i, k), and the largest discharged energy density of PTFMB film (10.29 J/cm3) is also higher than that of PI film (3.39 J/cm3) and PABI film (2.19 J/cm3).

Tracing origins, first, the large energy gap and high crystallinity in PTFMB film effectively restrict electron excitation and transport under high temperature and strong electric field, and Supplementary Fig. 35 also shows that PTFMB film owns the lower leakage current density, which reflects reduced conduction loss of PTFMB film. Second, the high crystallinity and introduced –CF3 groups can reduce the frictional work of polarization when polar groups deflect orientation under electric field as demonstrated above. Third, the high mechanical strength also guarantees that PTFMB film can maintain high structure stability even under the coexistence of large electric and thermal stresses. Finally, the high mechanical strength, coupled with reduction in conduction and polarization losses, significantly contributes to the high energy-storage level of PTFMB film.

It is worth noting that the currently reported polymer dielectrics exist an intractable and contradictory problem of synergistically enhancing the mechanical strength and energy-storage level. Especially for all-organic structure designs, increasing energy-storage level often comes at the expense of mechanical strength. This is because previous design approaches have primarily focused on either widening energy gap or reducing interchain interactions within polymer structure, without addressing both aspects simultaneously. Our work integrates robust hydrogen bonding interactions, which not only maintain but also marginally enhance energy-storage level. In addition, we optimize the aggregation structure to achieve high film crystallinity, thereby balancing and further elevating mechanical and energy storage performances. Finally, the high mechanical strength (378.3 MPa), breakdown strength (682 MV/m), energy density/efficiency (6.06 J/cm3, η = 90%) at 150 °C in PTFMB film have a synergistical enhancement (81%, 88%, 2231%) than that of PI film, which also has the similar enhancement (373%, 49%, 418%) than that of reported PEI film. Although the energy density is not the highest, the comprehensive performances of PTFMB film present its preponderance among existed polymer dielectric films, especially for mechanical strength in this field to our knowledge, as shown in Fig. 6l, Supplementary Table 4, and Supplementary Note 222,59,60,61,63,64,65,66,67.

Assessment for devices and application

The self-healing characteristic is significant during operation of film capacitor. EDS and SEM tests show that one-time breakdown generates large-area gasification of Au electrode, while carbonized conductive region is smaller. As a result, a wide insulating isolation strip forms between Au electrode and carbonized conductive region. The self-healing characteristic guarantees continuous operation of PTFMB film without obvious performance reduction (Fig. 7a–c and Supplementary Fig. 37).

Fig. 7: Performances tests in multiple application scenarios and devices fabrications.
figure 7

a Schematic illustration of self-healing process. b SEM image of PTFMB film after electrical breakdown. c EDS analysis of PTFMB film after electrical breakdown. d The breakdown strength of PTFMB film before and after drying. e Electric displacement-electric field loops of PTFMB film before and after corona discharge. f Ud and η of PTFMB film with cycle number at 300 MV/m and 150 °C. g Preparation of a multilayer polymer capacitor. h The capacitance of PTFMB film capacitors at 25 °C. i The capacitance of different capacitors as a function of frequency at 25 °C. j, k A Film capacitor operating at AC-DC stacking condition in electric vehicle. l Ud and η of the PTFMB film with cycle number at DC:50 MV/m +AC:200 MV/m at 150 °C.

Full size image

In addition, the introduced –CF3 groups also improve the hydrophobicity of PTFMB film (Supplementary Fig. 38), and the atmospheric moisture exposure has a negligible impact on breakdown strength of PTFMB film (Fig. 7d), but that weakens the breakdown strength of PABI film. The good hydrophobicity of PTFMB film has its advantage in practical applications especially in water-sensitive electronic and electrical fields. Meanwhile, the tests on different regions of PTFMB film show little difference in energy storage performance, indicating good film uniformity (Supplementary Fig. 39). Moreover, PTFMB film also demonstrates its relatively stable energy storage performance by corona discharge treatment (IEC60343 standard), highlighting its potential for complicated electrical environments (Fig. 7e).

Furthermore, the electric energy can rapidly charge into PTFMB film within 2.5 μs, and the stored energy can also rapidly release only at 11.1 μs with a high power density of 519 MW/kg during charging-discharging process, which is the advantage of film capacitor among other energy storage devices (Supplementary Fig. 40). Figure 7f shows that the energy storage density/efficiency of PTFMB film have little change through repetitive charging-discharging tests up to 10,000 times, reflecting its long service life. We further fabricated multilayer polymer film capacitor, alternating layer-by-layer stacking by metal electrode and film (Fig. 7g–i and Supplementary Fig. 41). PTFMB film capacitor delivers a relatively stable capacitance, which surpasses that of fabricated BOPP and PET film capacitors.

It should be noted that most existed studies focus on alternating current (AC) tests for polymer dielectrics, but many applications like electric vehicles and power systems usually subject film capacitors to AC-DC stacking conditions (Fig. 7j, k). Figure 7l and Supplementary Fig. 42 show that PTFMB film can still maintain stable operation and relatively high energy-storage level at AC-DC stacking condition compared to individual AC field condition, indicating its high application potentials in these fields as well.

In summary, through directly focusing on the shortcomings of two core structure units of PI, we develop a targeted all-organic synthesis route to fabricate crystalline polyamide dielectric with the coexistence of strong hydrogen bonds and high energy gap. Subsequently, a matched polymerization and molding process are developed, where the “complexing/uncomplexing” regulation of Lewis acid is appplied to guarantee well-organized and compact chain packing, and the film crystallinity is further improved through regulating Lewis acid content and annealing temperature. The resulting polymer film demonstrates a synergistic improvement in mechanical strength and energy-storage level compared to conventional PI (or PEI) films. This work provides an all-organic simulation and experimental method for creating advanced polymer dielectrics toward high-temperature dielectric energy storage applications, encompassing new-energy vehicles, photovoltaic grid integration systems, and oil-gas explorations68,69,70,71,72,73,74,75,76.

Methods

Materials

3,3′,4,4′-Biphenyl tetracarboxylic diandhydride (BPDA, >99.5%) were obtained from Sunlight Medical Co., Ltd. (Changzhou, China). Dimethylacetamide (DMAc; >99%), lithium chloride (LiCl; >98.5%), and terephthaloyl chloride (TPC; >99%) were obtained from Adamasbeta Co., Ltd. (Shanghai, China). 2,2′-Bis (trifluoromethyl)benzidine (TFMB) and 5-(6)-amino-2-(4-aminobenzene) benzimidazole (PABZ, >99.5%) were obtained from Sunlight Medical Co., Ltd. (Changzhou, China). All of the above monomers were dried at high temperature under vacuum before use.

Synthesis of PI, PABI, and PTFMB solutions

PI: Certain amount of PABZ was added into a three-neck flash containing NMP. After PABZ was completely dissolved in NMP, BPDA was added into the reaction flask and stirred under nitrogen atmosphere until it was clarified. Then stirred 2 h to obtain a clear and transparent polyamic acid (PAA) solution.

PABI: 2.03 × g of PABZ was added into a three necked flask containing 60 × g of DMAc/LiCl (3.5 wt%) mixture, which was followed by stirring at room temperature to completely dissolve PABZ. Then 2.50 × g of TPC was added to the mixture. Finally, the solution was stirred for 2 h to get PABI solution.

PTFMB: 2.66 × g of TFMB was added into a three-necked flask containing 60 × g of DMAc/LiCl (3.5 wt%) mixture, which was followed by stirring at room temperature to completely dissolve TFMB. Then 1.52 × g of TPC was added to the mixture. Finally, the solution was stirred for 2 h to get PTFMB solution.

Synthesis of PI, PABI, and PTFMB films

PI: The PAA solution was uniformly coated on a clean glass sheet by a film coating machine, and heat-treated in a vacuum oven at 80 °C for 2 h, 120 °C for 1 h, 220 °C for 1 h and 350 °C for 0.5 h. Then the film was removed from the substrate by immersion in hot water and dried at 100 °C for a least 2 h.

PABI: The PABI solution was cast in clean glass plates, followed by immersing into coagulation bath to get films. The films were washed with deionized water at 80 °C for 4 h to remove residual LiCl. Then, it was annealed at 360 °C for 0.5 h in vacuum oven. Then the film was removed from the substrate by immersion in hot water and dried at 100 °C for a least 12 h.

PTFMB: The PTFMB solution was cast in clean glass plates, followed by immersing into coagulation bath to get films. The films were washed with deionized water at 80 °C for 4 h to remove residual LiCl. Then, it was annealed at 80, 320, 340, 360, 380, and 400 °C for 0.5 h in vacuum oven. Then the film was removed from the substrate by immersion in hot water and dried at 100 °C for a least 12 h. They are named PTFMB-80, PTFMB-320, PTFMB-340, PTFMB-360, PTFMB-380, and PTFMB-400, respectively. Moreover, the PTFMB was dissolved in the DMAc/LiCl solvent with different LiCl content. The film was further coated and after the LiCl was removed by coagulation bath and boiling, the PTFMB film was obtained by heat treatment. They are named PTFMB-0%LiCl, PTFMB-3%LiCl, PTFMB-4%LiCl and PTFMB-5%LiCl, respectively.

Characterization

1H NMR spectra for all polymers were recorded on a BrukerAVII-400 MHz spectrometer using deuterated chloroform as the solvent. ATR-FTIR spectra of all films were measured using a Perkin Elmer Frontier FTIR spectrometer with a resolution of 1 cm−1. Polarized FTIR (polarized FTIR) was used to collect the infrared spectra of samples. The polarized light of 0°, 30°, 60°, 90°, 120°, 150°, and 180° were used to irradiate the sample to study the intensity change of the infrared absorption band in different regions, and further study the conformation and orientation of long chain or large molecular chain. Dielectric spectra were acquired over a board frequency range from 10 Hz to 106 Hz by Novocontrol broadband dielectric/impedance spectrometer (Novocontrol Technologies GmbH& Co. KG.) with a liquid nitrogen cooling system. To analyze the inter-chain distance in the polymers, XRD patterns were obtained using a high-resolution X-ray diffractometer (Rigaku Ultima IV) with Cu Kα radiation (λ = 1.5418 Å). The d-spacing was calculated using Bragg’s law:

$$n\lambda=2{dsin}\theta$$
(6)

UV–Vis spectroscopy covering the 200–600 nm range was conducted using a UV-2300 spectrophotometer. The structural analysis included scanning electron microscopy (SEM) images of the cross-sections and internal Au electrodes of the capacitors, along with corresponding. EDS mappings for C and Au elements, obtained via LSM-5900LV. Film thickness was measured using a DUALSCOPE MPO thickness gauge, and optical images of the dielectric films were captured using an Olympus SZX16 microscope. Dynamic mechanical analysis (DMA) was recorded on PerkinElmer DMA8000 with a heating rate of 10 °C/min in nitrogen atmosphere. The breakdown strength tests were conducted using the Beijing Guance DDJ-50 kV instrument. The testing conditions involved applying a DC voltage with a ramp rate of 500 V/s at temperatures of 25 °C and 150 °C. Samples were placed between cylindrical electrodes, and the entire electrode system was immersed in insulating oil. Breakdown voltage was recorded and divided by the film thickness to determine the breakdown strength, with at least 10 data points collected for each sample group. The mechanical properties were tested using an INSTRON5967 instrument. Samples were cut into strips measuring 10 × 100 mm and clamped in the tensile grips, with a gauge length of 20 mm. The tensile testing was performed at a rate of 5 mm/min. SAXS measurements were conducted using a Nano STAR-U instrument (Bruker AXS INC). The HI-STAR detector was used, with testing voltage and current set at 40 kV and 650 μA, respectively. A Cu target was used, with an X-ray wavelength (λ) of 1.54 Å. Tests were performed at room temperature. The sample-to-detector distance was 1074 mm, and the scattering vector (q) range was from 0.044 to 2.0 nm−1. For a system composed of sparsely dispersed, randomly oriented particles with consistent size and shape, and uniform electron density within each particle, the scattering intensity can be expressed as follows65,66:

$$I\left(h\right)=4{I}_{e}{V}^{2}{\rho }_{0}^{2}{\varphi }^{2}\left({hR}\right)$$
(7)
$$h=\frac{4\pi \sin \theta }{\lambda }$$
(8)

h is the scattering vector, \({I}_{e}\) is the scattering intensity of a single electron, \({\rho }_{0}\) is the electron density of the particle, V is the volume of the particle, R is the radius of the particle, \(\varphi\) is the scattering function, \(\lambda\) is the wavelength of the incident X-rays, 2\(\theta\) is the scattering angle.

The Guinier formula describes the scattering intensity of a system with M non-interfering particles as follows:

$$I\left(h\right)={I}_{e}{M}_{n}^{2}{e}^{\left(-\frac{{h}^{2}{R}_{g}^{2}}{3}\right)}$$
(9)
$${R}_{g}={\left[\frac{\sum k{ \, f}_{k}^{2}{r}_{k}^{2}}{\sum k{ \, f}_{k}}\right]}^{\frac{1}{2}}$$
(10)

In this formula, n is the total number of electrons in a particle; Rg is the radius of gyration, which is the root mean square distance of each electron from the center of mass of the particle. where rk is the distance of the scattering unit from the center of mass of the particle, and fk is the scattering factor of the k-th scattering unit.

Energy storage density and discharge efficiency were measured based on Sawyer-Tower and virtual ground circuits. This system recorded the displacement-polarization curves of dielectric films under various temperatures and electric field strengths. Prior to testing, gold electrodes with a diameter of 5 mm were evaporated onto both sides of the films to ensure uniform current flow. The testing frequency was set at 100 Hz. For high-temperature energy storage performance tests, the insulating oil was heated using a heating device, and testing commenced after stabilizing the temperature for 10 min to ensure accuracy. The current leakage measurements were conducted using Keithley 2410 electrometer and a Keithley 2290-10 high voltage source. The conduction mechanisms in dielectric polymer films are typically classified into two categories: electrode-limited and bulk-limited mechanisms. Electrode-limited mechanisms include Richardson–Schottky emission and Fowler–Nordheim tunneling. Bulk-limited mechanisms encompass ohmic conduction, Poole–Frenkel emission, jump conduction, and space-charge-limited conduction. Leakage current resulting from Richardson–Schottky emission, Poole–Frenkel emission, and jump conduction increases exponentially with temperature, making them the predominant sources of leakage current under conditions of elevated temperatures and high electric fields. Richardson–Schottky emission describes the transport mechanism where electrons in a metal electrode acquire sufficient thermal energy to surmount the energy barrier at the metal-dielectric interface. Subsequently, they are injected into the dielectric, thereby becoming charge carriers. The resulting current density is:

$$J=A{T}^{2} * \exp \left[-\left(\frac{\varphi -\sqrt{\frac{{q}^{3}}{4\pi \varepsilon }}\sqrt{E}}{{kT}}\right)\right]$$
(11)

Where J is the current density, E is the electric field, A is the Richardson constant, φ is the height of the barrier at the electrode-dielectric interface, q is the charge of the carrier, T is the temperature, ε is the dielectric constant of the dielectric material, and k is the Boltzmann constant.

Poole-Frenkel (P-F) emission describes how electrons in the trap gain energy and become charge carriers inside the dielectric. The resulting current density is expressed as:

$$J={q\mu }{N}_{{{{\rm{C}}}}}E * \exp \left[-\frac{\left({\varphi }_{{{{\rm{B}}}}}-{\beta }_{{\mbox{PF}}}\sqrt{E}\right)}{{kT}}\right]$$
(12)

Where J is the current density, E is the electric field, A is the Richardson constant, φ is the height of the barrier at the electrode-dielectric interface, q is the charge of the carrier, T is the temperature, ε is the dielectric constant of the dielectric material, and k is the Boltzmann constant.

Jump conduction involves the tunneling of charge carriers between localized states through quantum mechanical tunneling effects. In this mechanism, carriers can traverse from one trapping site to another, even when their energy is lower than the maximum energy of the barrier between the sites, due to the tunneling effect.

$$J=2{nq}\nu \lambda * \exp \left(-\frac{{W}_{{{{\rm{a}}}}}}{{kT}}\right)\sinh \left(\frac{\lambda {qE}}{2{kT}}\right)$$
(13)

Where n is the carrier concentration, ν is the escape frequency, λ is the jump distance, W is the activation energy, and the other symbols are the same as previously defined.

PI, PABI, and PTFMB 3D electrostatic potentials were accurately calculated based on density functional theory (DFT). In order to ensure the efficiency and accuracy of the calculation process, the 6–31 G(d) basis set of B3LYP heterocyclic functional and Pople was used to describe the electron exchange integral. The HOMO and LUMO levels of PI, PABI, and PTFMB model compounds were calculated by using Gaussian16. The specific method is B3LYP, the electron exchange integration uses the 6–31 G(d) basis of the Pople, and the energy was corrected using Grimme’s dispersion and damping.

Material Studio software was used to first build a polymer chain of 40 repeat units. Then the chain is geometrically optimized. Then the NVT was simulated at 300, 800, and 300 K respectively, and the single molecular chain with the lowest energy was obtained. Then, an Amorphous Cell containing 15 polymer chains is established using the Amorphous cell module. Then the Forcite module is used to carry out 5 rounds of annealing cycles on the box, and the box is geometrically optimized after each cycle. Then the molecular dynamics simulation of the annealed box was carried out, and finally the amorphous cell was obtained. Then the hydrogen bond was established in the polymer model, and the hydrogen bond formation sites of the three systems were visualized. The number and length of hydrogen bonds formed in polymers were calculated by using hydrogen bond statistics script.