Electroacoustic alignment of robust and highly piezoelectric nylon-11 films

Abstract

Nylon-11—a common and widely used material, is a promising non-fluorinated piezoelectric polymer given its mechanical strength, chemical stability and elasticity. Nevertheless, it typically possesses low piezoelectric performance, which severely limits its use for energy generation applications. Here, we demonstrate the synthesis of highly aligned, piezoelectric nylon-11 films through an energy-efficient (≈1–5 Wh), single-step platform that facilitates the electroacoustic coupling associated with MHz-order nanoscale vibrations during the crystallisation process. Uniquely, such coupling allows the simultaneous induction of (i) nylon’s piezoelectric ${\delta }^{\prime }$-phase, (ii) long-range crystalline ordering, (iii) an ordered hydrogen-bonded network, and (iv) dipole alignment, which we directly probe using time-resolved operando synchrotron grazing-incidence wide-angle X-ray scattering and high-resolution infrared spectroscopy. We show that the material produces a piezoelectric voltage coefficient (g₃₃ = 427 × 10⁻³ Vm N⁻¹) that surpasses the performance of all piezoelectric polymers reported to date. The film’s exceptional mechanical resilience is evident from its stable performance over 20,000 compression cycles at 50 N and its ability to withstand vehicular loads.

Nylon-11 is a common and durable polymer but possess low piezoelectric properties. Here, the authors use mechanical accelerations and strong electric fields to induce crystallization, hydrogen-bonding and dipole alignment in Nylon-11 films, achieving high piezoelectricity.

Introduction

Piezoelectric polymers, given their ability to convert mechanical pressure into electrical energy and vice versa without limitations imposed by conventional rigid piezoelectric ceramics, are a key focus of research in the development of more efficient and versatile platforms for autonomous sensing networks, wearable and implantable devices, and low-power functional sources^1–5. Amongst these, poly(vinylidene difluoride) (PVDF) has been considered the state-of-the-art for piezoelectric polymers, given its exceptional piezoelectric and mechanical properties^6,7. PVDF is nevertheless a polyfluorinated alkyl substance and therefore constitutes a significant environmental hazard^8,9. On the other hand, nylon—a non-fluorinated, durable, lightweight yet high-strength synthetic polymer introduced in the 1930s—is an extremely versatile material whose mechanical resilience and heat resistance have enabled its widespread use in applications as diverse as aircraft carrier arrestor cables, military gear and even space suits. In addition to its exceptional high-force tolerance and impact resistance, nylon (and, in particular, nylon-11—its most studied ferroelectric variant) is, in principle, piezoelectric. This is due to its non-centrosymmetric structure, high dipole moment (3.7 D), and efficient molecular packing enabled by an extensive hydrogen-bonded network. Despite these, however, nylon-11’s potential as a practical piezoelectric material has yet to be realised.

This, in particular, stems from difficulties in controlling nylon-11’s crystallisation, which could result in five distinct phases—α, ${\alpha }^{\prime }$, γ, δ, and ${\delta }^{\prime }$^10–12. Among these, the ${\delta }^{\prime }$-phase possesses the most promising piezoelectric properties but remains challenging to obtain. This is because of the intricate processing techniques required for its synthesis, which include uniaxial stretching of melt-quenched films¹³ or solution quenching via spin-coating and vacuum-assisted solvent removal¹⁴. Despite these complex fabrication strategies, none have thus far successfully led to a well-ordered hydrogen-bonded network in its structure—an essential characteristic for optimising polarity, and consequently, piezoelectric performance^15,16. Further, piezoelectricity is also influenced by both long-range bulk crystal ordering^12,17,18 and dipole alignment, the latter requiring the application of high electric fields¹⁴. Recent attempts utilising electrospinning to achieve dipole alignment of the ${\delta }^{\prime }$-phase have enabled high-field poling, but introduced significant morphological heterogeneity; additionally, long-range bulk crystal ordering or a well-ordered hydrogen-bonded network has not been demonstrated^10,19,20. Altogether, a method that optimises the piezoelectric performance of nylon-11 by collectively facilitating (i) crystallisation of the piezoelectric ${\delta }^{\prime }$-phase, (ii) long-range crystalline ordering, (iii) hydrogen-bonded network ordering and (iv) dipole alignment, has yet to be realised.

Herein, we address this gap with a synthesis platform that facilitates the coupling between intense mechanical accelerations (10⁸ m s⁻² order)²¹ with strong electric fields (10⁸ V m⁻¹ order)^22–24 in the form of nanometre-amplitude surface reflected bulk waves (SRBWs—a hybrid counterpart to surface acoustic waves (SAWs) that additionally possesses a bulk wave component)²⁵ to synchronously achieve long-range crystalline order, optimised hydrogen-bond networks, and molecular dipole alignment within the nylon-11 ${\delta }^{\prime }$-phase—in a single-step and under ambient conditions. Validated by operando grazing-incidence wide-angle X-ray scattering (GIWAXS) and infrared (IR) spectroscopy, we show that the method is capable of producing robust nylon-11 films with outstanding piezoelectric performance (g₃₃ = 427 × 10⁻³ Vm N⁻¹), surpassing that reported for all piezoelectric polymers thus far.

Results

The electromechanical synthesis platform, along with the characteristics of the highly-aligned piezoelectric ${\delta }^{\prime }$-phase nylon-11 films it enables us to synthesise, are schematically illustrated in Fig. 1. As described in the ‘Methods’ section, SRBWs (utilised due to its energy efficiency compared to SAWs)²⁵ are generated by applying an oscillating radio-frequency (RF) electrical signal at the resonant frequency (10 MHz) to an interdigital transducer (IDT) patterned on a single-crystal piezoelectric substrate (lithium niobate (LiNbO₃), chosen for its high piezoelectric coupling coefficient and exceptional chemical resistance during solvent processing). The SRBWs are coupled into the nylon-11 precursor solution, which was drop-casted on the substrate, to drive crystallisation via solvent evaporation under ambient conditions. To glean insight on the mechanism that underpins the effect of the acoustoelectric coupling on the crystallisation process, we carried out the synthesis under three distinct conditions: (1) the full electromechanical coupling arising from the SRBW (EM-SRBW), (2) pure mechanical vibration (M-SRBW), obtained by attenuating the evanescent electric field in the liquid phase with a gold screening layer atop the substrate, and, (3) conventional solvent casting at 60 °C as the control.

Fig. 1. SRBW synthesis and characterisation.

A Schematic depiction of the SRBW device on which highly-aligned ${\delta }^{\prime }$-phase nylon-11 films are synthesised. B The long-range ordering of the material’s crystal structure (resulting in net dipole polarisation in the ‘3’ direction) through the electromechanical coupling associated with the SRBW (EM-SRBW) is confirmed through C operando synchrotron grazing-incidence wide-angle X-ray scattering (GIWAXS) patterns (left), obtained at an incidence of 0.12° to the top surface of the film; also shown is an example of the azimuthal angle (right), which characterises the anisotropy of the Debye-Scherrer rings in the scattering pattern. D Schematic of the film’s aligned dipoles, and molecularly-ordered hydrogen-bonded network, as verified through E the shift in the EM-SRBW nylon-11's amide I and amide II bands toward lower and higher wavenumbers, respectively, in the operando synchrotron high-resolution IR spectra.

Crystalline lattice configuration (phase)

Nylon-11 potentially crystallises into five distinct phases: α, ${\alpha }^{\prime }$, γ, δ, and ${\delta }^{\prime }$^10–12. The α- and ${\alpha }^{\prime }$-phases possess tightly ordered hydrogen bonds that restrict dipole rotation, thus rendering them non-piezoelectric^26,27. The γ-phase, on the other hand, has limited piezoelectricity given its moderate polarity, which enables weak dipole switching²⁸. The δ- and ${\delta }^{\prime }$-phases share pseudo-hexagonal chain structures, with the δ-phase featuring laterally-ordered hydrogen bonds that form distinct polymer sheets, thereby maximising its dipole moment per unit cell. The δ-phase, however, exists only at elevated temperatures (>140 °C) before reverting to the α-phase upon cooling^16,29. On the other hand, the ${\delta }^{\prime }$-phase exhibits a disordered and smectic-like arrangement, enabling dipole reorientation under an electric field due to its weak mobile hydrogen bonds^30,31. As alluded to earlier, however, obtaining the ${\delta }^{\prime }$-phase is inherently difficult, given that nylon-11 is thermodynamically stable as either γ- or α-phase when crystallised using traditional solvents (i.e. m-cresol, formic acid, trifluoracetic acid (TFA))¹⁹.

Crystalline phases of the solvent-cast control, M-SRBW, and EM-SRBW nylon-11 films, synthesised from a 3 wt% nylon-11 solution in TFA/acetone (60:40 mol%), were analysed using synchrotron wide-angle X-ray scattering (WAXS). A comparison of the (100)/(010) reflections in Fig. S1, which correspond to the distance between inter-chain and inter-sheet hydrogen bonds³², respectively, reveals a distinct shift toward higher q values for the M-SRBW and EM-SRBW films. The control film exhibits a peak at 15.22 nm⁻¹ with a corresponding d-spacing of 0.413 nm. This peak can nevertheless be seen to shift to 15.47 nm⁻¹ with a d-spacing of 0.406 nm for both the M-SRBW and EM-SRBW films. Since it is not possible to distinguish between the γ-, $\gamma \prime$-, ${\alpha }^{\prime }$-, or ${\delta }^{\prime }$-phases from the (100)/(010) planes¹⁴, we additionally examine the low-angle (001) reflection, assigned to the arrangement of the amide groups along the polymer chain. These reveal d-spacings of 1.42 nm (q = 4.42 nm⁻¹) for the control film and 1.36 nm (q = 4.61 nm⁻¹) for the M-SRBW and EM-SRBW films (Fig. S1). These q values align with those reported in literature, indicating the prevalence of the γ-phase in the solvent-cast control and the ${\delta }^{\prime }$-phase in both the SRBW-synthesised films¹⁴. Additionally, they suggest that the crystallisation is influenced primarily by the mechanical component of the SRBW excitation responsible for giving rise to the compressive strain gradients that drive ${\delta }^{\prime }$-phase crystallisation via rapid solvent evaporation, with no noticeable contribution from the electric field.

To assess the influence of solvent composition on phase formation, synchrotron WAXS was also conducted on films prepared with the solvent-cast and EM-SRBW methods using three different solvents (50:50 mol% TFA/acetone, m-cresol and formic acid). Examination of the (001) reflections in the EM-SRBW film prepared with the 50:50 mol% TFA/acetone solvent mixture revealed a d-spacing of 1.39 nm (q = 4.51 nm⁻¹), which suggests a mixed ${\delta }^{\prime }$/γ-phase configuration in contrast to the γ-phase observed in the solvent-cast control (d-spacing = 1.42 nm; q = 4.42 nm⁻¹) (Fig. S2A); that mixed phases were obtained could be due to the mixture’s higher vapour pressure compared to the lowest value obtained with the 60:40 mol% TFA/acetone ratio¹⁴ that was likely to be most effective in promoting ${\delta }^{\prime }$-phase crystallisation. m-cresol solvent-cast control films, on the other hand, exhibited an $\alpha /{\alpha }^{\prime }$-phase (d-spacing = 1.15 nm; q = 5.44 nm⁻¹) while those synthesised with the EM-SRBW possessed a mixture of $\alpha /{\alpha }^{\prime }$ (d-spacing = 1.17 nm; q = 5.37 nm⁻¹) and γ-phases (d-spacing = 1.41 nm; q = 4.46 nm⁻¹) (Fig. S2B). The formic acid solvent-cast film displayed the $\alpha /{\alpha }^{\prime }$-phase (d-spacing = 1.16 nm; q = 5.41 nm⁻¹) in contrast to the EM-SRBW film that was dominated by the $\gamma /\gamma \prime$-phase (d-spacing = 1.42 nm; q = 4.42 nm⁻¹) with some $\alpha /{\alpha }^{\prime }$ contributions (d-spacing = 1.17 nm; q = 5.37 nm⁻¹) (Fig. S2C). We therefore omitted further assessment of films synthesised with both m-cresol and formic acid, given that they contained the γ-phase, which is weakly piezoelectric, and the $\alpha /{\alpha }^{\prime }$-phase, which is prone to dielectric breakdown prior to dipole polarisation^14,19,28.

Long-range crystal macrostructure

The synchrotron operando GIWAXS data (experimental configuration shown in Fig. S3A) provide insight into the real-time dynamics associated with the long-range crystalline ordering, as schematically shown in Figs. 1B and 2A. Typically, as the anisotropy within the diffraction arc becomes more prevalent, the degree of crystal orientation increases, conveyed visually by the appearance of partial Debye-Scherrer rings in the GIWAXS pattern³³. The anisotropy is characterised by the azimuthal angle (χ), which describes the angular distribution of scattered intensity along the diffraction ring. In the case of the solvent-cast control, a uniform intensity distribution is observed, forming a complete Debye-Scherrer ring (Fig. 2B), indicative of an isotropic arrangement rather than anisotropy. This suggests randomly oriented crystalline planes, as further corroborated by the constant χ intensity profile of the (100)/(010) plane. In contrast, as the EM-SRBW film transitions towards a crystalline state, a partial Debye-Scherrer ring is observed (Figs. 1C and 2C), with an out-of-plane (OOP) orientation for the (100)/(010) plane evidenced by the increased scattering intensity at 90° as a function of χ. Additionally, qualitative examination of the (001) reflection reveals similar anisotropy within its diffraction arc, suggesting in-plane (IP) ordering along the polymer chain. This further underscores the enhanced long-range crystalline order facilitated by the SRBW acoustoelectric coupling.

Fig. 2. Long-range crystalline order.

A Schematic illustration of the synchrotron operando GIWAXS experiment representing the isotropy of the solvent-cast control and the anisotropy of the EM-SRBW nylon-11 macrostructure. 2D GIWAXS patterns showing B random ordering in the solvent-cast control and long-range crystal order in C the EM-SRBW films, evidenced by complete and partial Debye-Scherrer rings (top), respectively. The bottom of B, C represents the intensity along the (100)/(010) plane. Long-range ordering was evaluated through D Herman’s orientation function (f), which provides a quantitative progressive indication of the long-range crystal ordering over time for both the solvent-cast control and EM-SRBW films (ambient conditions).

Furthermore, Herman’s orientation function (f) is utilised to measure the degree of orientation or long-range order in the films, derived from scattering intensities along χ (e.g. the (100)/(010) plane)³⁴. For the EM-SRBW film, it can be seen that f asymptotes at a value of approximately 0.32, in contrast with the lack of any preferential long-range order in the control film (Fig. 2D). As observed from Fig. S3B–E, anisotropy is present even at lower EM-SRBW (combined mechanical and electric) field strengths, with long-range order demonstrably improving (f rises from 0.23 to 0.32) as the field strength is increased from 6.9 to 12.9 dBm. These results are consistent with previous studies on piezoelectric polymers, such as that of glycine–polycaprolactone nanofiber films (f = 0.38)¹⁸, which highlighted the role of controlled structural ordering in enhancing piezoelectric properties. Altogether, the time-resolved scattering patterns in both the OOP and IP directions, along with their analysis in Fig. 2B, C, offer comparable insights into the structural organisation of the EM-SRBW film. In particular, they indicate the formation of a preferential ‘edge-on’ crystallite orientation in the film, which can be attributed to the orthogonal alignment of dipoles relative to the surface of the LiNbO₃ substrate³⁵.

Hydrogen bond ordering

Hydrogen bond ordering of the nylon-11 films was evaluated through synchrotron operando high-resolution infrared (IR) spectroscopy (Fig. 1D, E), and corroborated through complementary techniques such as solid-state nuclear magnetic resonance (NMR), differential scanning calorimetry (DSC) and WAXS. The real-time IR spectra for completely crystallised nylon-11 films in Fig. S4 showed the amide I peak, which can be attributed to the C=O stretching mode, and the amide II peak associated with the in-plane N–H bending and –N–CO– stretching modes of the central amide. For the solvent-cast control film, the amide I and II peaks appeared at 1637 cm⁻¹ and 1549 cm⁻¹, respectively (Δ ≈ 88 cm⁻¹). In contrast, these peaks can be seen to shift to 1635 cm⁻¹ and 1553 cm⁻¹ (Δ ≈ 82 cm⁻¹) for the M-SRBW films, and to 1634 cm⁻¹ and 1555 cm⁻¹ (Δ ≈ 79 cm⁻¹) for the EM-SRBW films. This trend is representative of increased hydrogen bond ordering, which is seen when amide I and II bands typically move to lower and higher wavenumbers, respectively^36,37. In comparison, for ordered solution-quenched nylon-11 films reported in the literature¹⁴, the amide I and II peaks appear at 1634 cm⁻¹ and 1547 cm⁻¹, respectively (Δ ≈ 87 cm⁻¹), indicating lower hydrogen bond ordering compared to the EM-SRBW films. To delineate possible thermal effects introduced by the EM-SRBW platform, we also conducted operando mid-IR measurements under purely thermal conditions (noting a maximum temperature rise to 40 °C during EM-SRBW synthesis at 14.6 dBm; see Fig. S5C inset), which showed that increasing temperature decreases hydrogen bond order, as evident from the shift in amide peaks further apart from Δ ≈ 84 cm⁻¹ (Fig. S5A) to Δ ≈ 90 cm⁻¹ (Fig. S5B) at elevated temperatures (80 °C). In contrast, operando EM-SRBW experiments showed that higher field strength (along with increasing substrate heating) increased the order of the hydrogen-bonded network instead, as can be seen from the decreasing peak separation (Fig. S6), indicating that the dominant mechanism is electroacoustically-driven as opposed to being thermally-induced.

The solid-state NMR spectra of the nylon-11 aliphatic chains in Fig. S7 predominantly show CH₂ segments in the gauche conformation (29.8 ppm) for the solvent-cast control, and a higher proportion of those in the trans-conformation (33.0 ppm) for the M-SRBW and EM-SRBW films. Given both the similar crystallinity across all films (Fig. S8) and the tendency of the CH₂ chains in the amorphous regions to adopt gauche bonding due to conformational disorder in the methylene segments²⁸, it can be inferred that the CH₂ chains adopt a crankshaft structure between amides in the crystalline region of the solvent-cast nylon-11 control. In contrast, the M-SRBW and EM-SRBW nylon-11 films assume a more ordered all-trans zigzag conformation—a distinct structural difference that aligns with the aforementioned IR spectral findings.

Parenthetically, DSC measurements show an increase in melting temperature from 172 °C to 183 °C and 189 °C (Fig. S8) for the solvent-cast control, M-SRBW and EM-SRBW films, respectively. The result suggests an enhancement in the thermal stability of the crystal lattice³⁸ and further confirms the formation of a robust intermolecular hydrogen-bonded network³⁹ with the acoustoelectric coupling. This is corroborated by a slight shift in the (100)/(010) peak positions toward higher q values and lower d-spacing for the M-SRBW and EM-SRBW films (compared to the solvent-cast control), as observed in the WAXS analysis in Fig. S1. Such a shift arises from enhanced intramolecular ordering, where hydrogen-bonded amide groups assemble into two-dimensional sheets⁴⁰.

In summary, the observed reduction in inter-sheet spacing (WAXS), shift in the amide peaks (mid-IR), trans-conformation (NMR), and higher melting temperature (DSC) collectively indicate that the EM-SRBW-synthesised nylon-11 possesses a tightly packed, ordered hydrogen-bonded network. This enhanced ordering strengthens hydrogen bond interactions, which are known to improve the mechanical properties of the film^41,42. Altogether, it can be seen that the M-SRBW and EM-SRBW nylon-11 films both possess some of the key characteristics (i.e., piezoelectric phase, long-range crystalline order and hydrogen-bonded network order) required to render the material piezoelectric, with further enhancement afforded by the improvement in the hydrogen bonding order and thermal stability with the EM-SRBW films.

Dipole alignment

Analysis of the solvent-cast nylon-11 control films with dual AC resonance tracking (DART) piezoresponse force microscopy (PFM) revealed a wide phase distribution of the deflection signal of the cantilever probe across the piezoelectric domains of the crystal (Fig. 3A(i)), which is indicative of randomly oriented dipoles with little to no alignment (Fig. 3A(ii, iii))⁴³. In the M-SRBW film in which crystallisation of the nylon-11 film occurs under pure mechanical vibration, we observe partial alignment in the dipole distribution, as seen from a 150° (out-of-plane) phase shift in dipole orientation (Fig. 3B(i–iii)). In contrast, the PFM phase distribution can be seen to narrow significantly for the EM-SRBW configuration, in which the full acoustoelectric coupling arises from the mechanical excitation and evanescent electric field. This is accompanied by a 20° (out-of-plane) phase shift (Fig. 3C(i–iii)), indicating highly-oriented dipole alignment⁴³ driven by the combined coupling between the mechanical and electrical fields, the latter (10⁸ V m⁻¹ order) notably exceeding nylon-11’s coercive field but remaining below its dielectric breakdown threshold^14,26.

Fig. 3. Dipole characterisation.

(i) Surface dipole characterisation through dual AC resonance tracking (DART) piezoresponse force microscopy (PFM), as measured through the (ii) dipole phase, and, subsequent (iii) line profile (top) and PFM phase distributions (bottom) of the A solvent-cast control, B M-SRBW and C EM-SRBW films, revealing the randomly aligned γ-phase, partially aligned ${\delta }^{\prime }$-phase, and highly aligned ${\delta }^{\prime }$-phase dipoles, respectively.

The film’s ability to generate electric charge in response to applied mechanical stress was additionally quantified by the effective piezoelectric coefficient (d_33,eff), measured using DART-PFM. In the solvent-cast nylon-11 control, we obtained d_33,eff = 1.20 pm V⁻¹. This value increased by 106% to 2.47 pm V⁻¹ for the M-SRBW configuration, and by 170% to 3.24 pm V⁻¹ for the EM-SRBW film (Fig. S9), alluding to the superior macroscopic performance and hence practical utility of the SRBW-synthesised films. This enhancement surpasses the piezoelectric response of nylon-11 nanofibers (d_33,eff = 2.1 pm V⁻¹)⁴⁴ and closely matches that of ${\delta }^{\prime }$-phase nanowires (d_33,eff = 3.22 pm V⁻¹)⁴², indicating the intrinsic piezoelectricity of the ${\delta }^{\prime }$-phase EM-SRBW nylon-11 films.

Macroscopic piezoelectric properties

The electrical output of nylon-11 piezoelectric energy generation devices fabricated using each type of film (i.e., the electrode-coated solvent-cast control, M-SRBW and EM-SRBW films following their detachment from the LiNbO₃ substrate) was measured under cyclic compression to assess their macroscale piezoelectric properties (Fig. 4A). Importantly, we adhered to an in-contact compressive testing regime for all measurements on the piezoelectric devices⁴⁵. By preventing separation and hence ensuring continuous contact throughout the measurements, we eliminate any electrostatic forces that can lead to inflated measured piezoelectric values^46,47. The solvent-cast control can be seen from Fig. 4B to exhibit minimal charge generation (0.6 pC) under a Δ10 N load (≈200 kPa). Whereas the M-SRBW film only provided a moderate enhancement (3.7 pC; Fig. 4C), the EM-SRBW film was instead observed to deliver a substantial performance improvement, generating 110 pC (Fig. 4D) and hence underscoring its superior piezoelectric response.

Fig. 4. Macroscale device piezoelectric properties.

A Schematic illustration of the in-contact compressive testing regime (Δ10 N load—eliminating contact separation to prevent unintended electrostatic forces) used to characterise the macroscale piezoelectric performance of the nylon-11 films. Current and charge output for B the solvent-cast control (the insets showing a magnified view of the current and charge), and the C M-SRBW and D EM-SRBW devices. E Calculated piezoelectric coefficient, d₃₃, from the measured charge output for the EM-SRBW device. Measured voltage and calculated power density for F the solvent-cast control, and the G M-SRBW and H EM-SRBW devices.

By extracting the slope of the measured charge output as a function of the applied variable force conditions, it is also possible to determine the macroscale piezoelectric coefficient (d₃₃). Given its miniscule current output, we obtained a negligible d₃₃ value for the solvent-cast control films (Fig. S10). The M-SRBW samples, on the other hand, recorded a modest d₃₃ value of approximately 0.41 pC N⁻¹ (Fig. S11), which can be attributed to partial dipole alignment of the piezoelectric ${\delta }^{\prime }$-phase. In contrast, given its significantly higher current and charge output, the EM-SRBW samples yielded a d₃₃ value of 11.26 ± 0.3 pC N⁻¹ (Figs. 4E and S12), which is comparable to state-of-the-art fluoropolymer materials that have been reported to date in the literature (Table S1)^48,49.

To show the advantages of the combined mechanical and electrical fields simultaneously inherent in the EM-SRBW processing, we performed comparative studies on solvent-cast (γ-phase) and M-SRBW (δ’-phase) films subject to sequential DC electrical poling following their crystallisation. Solvent-cast films exposed to a 200 MV m⁻¹ field exhibited no measurable piezoelectric response (Fig. S13) and likely suffered dielectric breakdown, as confirmed by the sudden increase in current measurements from the DC poling unit. Even at a reduced field of 100 MV m⁻¹, the response remained minimal, with a d₃₃ value of only 0.04 pC N⁻¹ (Fig. S14), consistent with the intrinsic weakly-piezoelectric nature of the γ-phase^19,28 and/or pinholes observed in the film (Fig. S15)¹⁴. On the other hand, the ${\delta }^{\prime }$-phase associated with the M-SRBW films under the same 200 MV m⁻¹ field showed a reduced piezoelectric response (d₃₃ = 0.36 pC N⁻¹; Fig. S16), underscoring the greater effectiveness of the simultaneous phase induction, dipole alignment and electrical poling processing (components of the EM-SRBW) over sequential crystallisation followed by DC poling.

Comparisons of the peak-to-peak voltage—measured in-contact, as before, using variable resistors connected in parallel to the circuit—between the films further corroborate the enhanced piezoelectric efficiency of the EM-SRBW films. At 4 GΩ, the solvent-cast control, M-SRBW and EM-SRBW nylon-11 films produced voltages of approximately 0.005 V (Fig. S17), 0.03 V (Fig. S18), and 0.65 V (Fig. S19), respectively. Subsequently, the power density was seen to increase from 0.002 μW cm⁻³ for the solvent-cast control (Fig. 4F) to 0.03 μW cm⁻³ for the M-SRBW samples (Fig. 4G). Notably, EM-SRBW films yielded a power density of 12.5 μW cm⁻³—a remarkable 400-fold enhancement over the M-SRBW film (Fig. 4H), alluding to the dominant role of the electric field intrinsic to the EM-SRBW in the alignment of the dipoles throughout the bulk of the film. To demonstrate the influence of electrostatic forces on the films, the measurements were also carried out at 100 mm s⁻¹ with a 1 mm separation between the sample and the compression head (Δ10 N load). The voltage output increased to 1.7, 2.8, and 3.0 V (Fig. S20) for the solvent-cast control, M-SRBW and EM-SRBW films, respectively, highlighting the necessity of in-contact testing to isolate the piezoelectric contributions inherent to the material. Parenthetically, we also note that the EM-SRBW films (10 μm thick) comprised a void-free, continuous morphology, ensuring a congruent and uniform structure, in contrast to the pinholes and voids found in the solvent-cast control films (Fig. S15) and the thinner (5 and 3 μm thick) EM-SRBW films (Fig. S21), which were likely caused by vapour-induced phase separation^14,50—the latter likely stemming from the intensified acoustic streaming during crystallisation of the thinner films.

Durability and piezoelectric performance

Having established the enhanced piezoelectric performance of EM-SRBW nylon-11 films via controlled laboratory testing, we now systematically compare their mechanical properties and piezoelectric durability under real-world extreme loading. For the former, nanoindentation measurements revealed a significantly higher compressive modulus Y₃ for the EM-SRBW films (Y₃ = 1.99 GPa) compared to the solvent-cast control (0.91 GPa), closely matching that of bulk nylon-11 (Fig. 5A)⁵¹. This enhancement was further supported by bimodal amplitude–frequency-modulation (AM–FM) measurements (Fig. 5B), which revealed an approximate two-fold increase in the effective elastic modulus (E^*), consistent with that expected for a film possessing a more ordered hydrogen-bonded network^38,52. In contrast, the solvent-cast film possessed distinct crystalline domains and amorphous regions, with E^* rising significantly from the amorphous regions to the crystalline domains⁵³, as can be seen from the histogram profile for the solvent-cast control film in Fig. 5B.

Fig. 5. Mechanical properties, long-term durability and extreme loading performance.

A Compressive modulus (Y₃) and B effective elastic modulus (E^*) maps (top) alongside histograms representing the distribution of the data (bottom) for both the solvent-cast control and the EM-SRBW nylon-11 films; the elastic modulus is superimposed onto the morphology acquired from atomic force microscopy (AFM). C Piezoelectric output, measured with finger tapping before (top) and after (bottom) vehicular compression (middle). D Long-term electromechanical durability (over 20,000 compression cycles) under an in-contact 50 N force, with a magnified view of the cyclic voltage response. E Measured piezoelectric voltage coefficient (g₃₃) for EM-SRBW nylon-11 in relation to values reported in the literature for the best-performing piezoelectric polymers.

To evaluate mechanical robustness under extreme stress, the EM-SRBW films were first subjected to localised mechanical perturbation via a simple finger tap (Movie S1), followed by large-scale deformation under high-magnitude loading from a moving vehicle (≈14,000 N) (Movie S2). Following vehicular compression, finger taps were applied to assess retention of the film’s function (Fig. 5C). Separately, long-term operational stability was examined through controlled compression testing, wherein the films can be observed to maintain a consistent piezoelectric response over 20,000 loading cycles under an in-contact force of ∆F = 48 N (Fig. 5D and Movie S3). Additionally, the durability of EM-SRBW films under varying temperature and humidity conditions was evaluated (Fig. S22) using an in-contact dynamic load of Δ20 N (2–22 N), in which the films can be seen to retain functional piezoelectric performance despite environmental stress (see details in the ‘Methods’). Specifically, we observe the d₃₃ value of 11.4 pC N⁻¹ at 25 °C and 35% relative humidity (RH) to decrease to 3.55 pC N⁻¹ at 60 °C and 1.04 pC N⁻¹ at 80 °C (Fig. S22A, B), whereas elevated humidity (91% RH) at 25 °C led to a decrease in d₃₃ to 6.4 pC N⁻¹ (Fig. S22A, C)—the dimunition in performance likely arising due to dipole relaxation^54,55 and device delamination under humid conditions⁵⁶.

Further, we evaluate the piezoelectric voltage coefficient (g₃₃)—a critical parameter for energy harvesting applications—by measuring the film’s relative permittivity (dielectric constant; Fig. S23A)), which was consistent with the range (3–4) reported in the literature at low frequencies (100 Hz) for nylon-11^26,57; the relatively low ${\epsilon }_{33}^{\mathrm{T}}$ (3.09), can be attributed to restricted dipole mobility as a consequence of the strong hydrogen-bonded network⁵⁸, which also gives rise to the remarkable structural integrity of the EM-SRBW films. As such, the EM-SRBW films possessed an exceptional g₃₃ value of 427 × 10⁻³ Vm N⁻¹ (Fig. 5E), surpassing the performance of all piezoelectric polymers reported to date (Table S1). Additionally, we note that the dielectric loss remained low ($\tan \delta <0.03$ at 1 kHz; Fig. S23B), indicating excellent suitability for piezoelectric applications where high dielectric loss (e.g. >0.5 at 1 kHz) can negatively impact piezoelectric film performance^59,60. Finally, we note that the EM-SRBW nylon-11’s figure of merit (FOM = d₃₃ × g₃₃ = 4.81 × 10⁻¹² Pa⁻¹) and electromechanical coupling coefficient (k₃₃ = 0.096) exceed that of conventional nylon-11 films and approach that for PVDF (Table S1). While these values remain below that for ceramics, they nevertheless represent the best performance among flexible non-fluorinated polymers, thereby offering an inexpensive, industrially-accessible option with robust mechanical and piezoelectric properties.

Discussion

We demonstrate the synthesis of highly-aligned nylon-11 films that exhibit exceptional piezoelectric properties via an energy-efficient (≈1–5 Wh), single-step electroacoustic approach. The 10 MHz nanometre-amplitude electromechanical substrate vibrations can be seen to facilitate crystallisation of the polymer that simultaneously enables ${\delta }^{\prime }$-phase formation, long-range crystalline ordering, hydrogen bond network order, and dipole alignment. We show, in particular, that this scalable and cost-effective route—through massive parallelisation enabled by the economies-of-scale associated with mass nanofabrication⁶¹—is capable of producing nylon-11 films with a record-high voltage coefficient of g₃₃ = 427 × 10⁻³ Vm N⁻¹ that surpasses all values that have been reported to date for piezoelectric polymers. Additionally, GIWAXS and IR operando measurements confirm the crystallographic ordering, while separate mechanical testing demonstrates the film’s resilience, whose functionality is maintained even after withstanding the high loads of vehicular compression.

Additionally, nylon-11 offers distinct environmental advantages through its recyclability—while preserving its mechanical integrity⁶², and biodegradability⁶³, thereby collectively supporting its use as a sustainable piezoelectric alternative with viable end-of-life pathways. Such endowment of nylon-11—a widely available and affordable commodity polymer—with superior piezoelectric performance through the MHz-order electroacoustic coupling is anticipated to offer a sustainable alternative to piezoelectric fluoropolymers, thereby positioning nylon-11 as a strong candidate for next-generation energy harvesting applications.

Methods

Nylon-11 precursor solution

Unless otherwise stated, all films were synthesised from nylon-11 precursor solution, which was prepared by dissolving 3 mm pellets of 3 wt.% nylon-11 (M_w = 201.31 g mol⁻¹; Sigma-Aldrich Pty. Ltd., Castle Hill, NSW, Australia) in 60 mol% trifluoroacetic acid (TFA) (≥99.8%, CAS 76-05-1; Supelco, Castle Hill, NSW, Australia) and 40 mol% acetone (≥99.8%, CAS 67-64-1, ACS Reagent; Sigma-Aldrich Pty. Ltd., Castle Hill, NSW, Australia) with constant magnetic bead stirring at room temperature for 4 h. Additionally, we prepared three other precursor solutions: 3 wt.% nylon-11 in TFA/acetone (50:50 mol%), 5 wt.% nylon-11 in m-cresol (≥99%, CAS 108-39-4; Merck Life Science Pty. Ltd., Bayswater, VIC, Australia), and 5 wt.% nylon-11 in formic acid (≥95%, CAS 64-18-6; Merck Life Science Pty. Ltd., Bayswater, VIC, Australia).

SRBW platform

SRBW substrates used for the synthesis of the M-SRBW and EM-SRBW nylon-11 films (Fig. 1A) were fabricated by photolithographically patterning a pair of interdigitated transducer (IDT) electrodes onto a 500 μm thick single-crystal (128° Y-rotated X-propagating) piezoelectric lithium niobate (LiNbO₃) substrate (Roditi International Corp. Ltd., London, UK). The IDT comprised 10 nm titanium (Ti) and 200 nm gold (Au) metal layers, whose finger width and gap of λ/4 specify the SRBW wavelength λ = 378 μm and hence frequency f = λ/c_s = 10 MHz, wherein c_s = 3780 ms⁻¹ is the acoustic wave phase speed in LiNbO₃. The IDT for the EM-SRBW platform comprised 68 finger pairs with an aperture of 13.94 mm. For the M-SRBW platform, 20 finger pairs were replaced by the gold layer (200 nm), acting to effectively attenuate the electric field. To generate the SRBW, an RF signal at the resonant frequency (10 MHz) from a signal generator (Multicomp PRO MP750289; Newark Electronics, Richfield, OH, USA) and amplifier (ZHL-5W-1+; MiniCircuits, Brooklyn, NY, USA) is supplied to the IDT.

Nylon-11 film synthesis

Solvent-cast control films from TFA/acetone were synthesised by drop casting 0.1 ml of the aforementioned nylon-11 precursor solution onto the LiNbO₃ substrate, but in the absence of the SRBW excitation. Instead, the film was heated at 60 °C for 60 min. Films from m-cresol and formic acid precursor solutions were solvent-cast at 150 °C and 80 °C, respectively. For the SRBW synthesis (both M-SRBW and EM-SRBW films), the same volume (0.1 ml) was pipetted onto the LiNbO₃ substrate and the device actuated at a power of 14.6 dBm (unless otherwise stated; e.g., for the EM-SRBW nylon-11 film synthesised with formic acid, the power was increased to 32 dBm due to film inhomogeneity observed at 14.6 dBm) for 2 h to subsequently yield approximately 10 μm thick films which could then be peeled off the substrate. The films were constrained by a rectangular barrier (1.4 × 1.6 cm), glued to the LiNbO₃ substrate, resulting in a film surface area of 1.2 cm², which we used for extensive characterisation; larger films (2.2 cm²) were also synthesised but showed a loss in δ’-phase (Fig. S25), with a d-spacing of the (001) plane decreasing to 1.38 nm (q = 4.54 nm⁻¹) compared to 1.36 nm (q = 4.61 nm⁻¹) for the 1.2 cm² films. During synthesis, we measured the indirect heating effect of the SRBW on the precursor solution temperature with an IR camera (IC080 LV; Trotec GmbH, Heinsberg, Germany). For some M-SRBW films, we also carried out post-synthesis DC poling using a high-voltage DC amplifier power supply (Trek Model 677A; Advanced Energy Industries, Inc., Fort Collins, CO, USA).

Film characterisation

DSC (Discovery DSC 250; TA Instruments, New Castle, DE, USA) was employed to quantitatively evaluate the crystallinity and to qualitatively assess the molecular bonding of the nylon-11 films that were synthesised. Briefly, 3–4 mg of the samples were placed in a metal pan and heated to 210 °C at a ramp rate of 10 °C min⁻¹. The degree of crystallinity (χ_c) was determined from

$${\chi }_c=\frac{\mathrm{\Delta }{H}_m}{\mathrm{\Delta }{H}_m^0}\times 100\%,$$ 1

where ΔH_m represents the melting enthalpy of the various nylon-11 films and $\mathrm{\Delta }{H}_m^0=189.05$ J g⁻¹ the heat of fusion of fully crystalline nylon-11⁶⁴.

NMR: Solid-state ¹³C cross-polarisation magic angle spinning NMR data were collected using a spectrometer (500 MHz DD2 spectrometer; Agilent Technologies, Mulgrave, VIC, Australia) equipped with a 4 mm HXY triple-resonance MAS probe (¹³C Larmor frequency 125.65 MHz). Film samples were loaded into 4 mm ZrO₂ rotors with a poly(tetrafluoroethylene) cap and Kel-F drive tip with a rotation speed of 12 kHz. The tancpx pulse sequence was used with a recycle delay of 2 s, 20 ms acquisition time, 1 ms contact time, and TPPM ¹H decoupling. Each experiment consisted of 12,000–14,000 scans over a total experiment time of approximately 8 h. Samples were externally referenced to the methine resonance of adamantane at 29.2 ppm. Line broadening of 50 Hz was applied, and data were processed and plotted using Origin (OriginPro 2023b, 10.0.5.153; OriginLab Corp., Northampton, MA, USA).

Scanning electron microscopy (SEM) (Verios 460L XHR-SEM; FEI, Hillsboro, OR, USA): Nylon-11 film surface topography images were acquired at 3 kV and 1.6 pA. A 5 nm protective iridium coating was deposited atop the film surface prior to SEM characterisation to prevent surface charging.

Operando mid-infrared (IR) spectroscopy was carried out on the THz beamline at the Australian Nuclear Science and Technology Organisation (ANSTO) facility at the Australian Synchrotron using an IFS125HR spectrometer (Bruker Corp., Billerica, MA, USA) with a diamond attenuated total reflectance (ATR) crystal (GladiATR; Pike Technologies, Madison, WI, USA) and a liquid-nitrogen cooled mid-band mercury cadmium telluride (MCT) detector. Spectra were recorded at room temperature over a spectral range of 4000–600 cm⁻¹ with 4 cm⁻¹ spectral resolution, and analysed using OPUS 8.0.19 (Bruker Optik GmbH, Ettlingen, Germany). A 100 μl nylon-11 precursor droplet was pipetted onto the ATR crystal, covered by an interfacing SRBW device, and operando measurements were taken every 4 s for 60 min under varying SRBW power levels.

WAXS and oper ando GIWAXS experiments were performed at the SAXS/WAXS beamline at the Australian Synchrotron. The beam energy was fixed at 15 keV for both WAXS/GIWAXS measurements. For GIWAXS, an incident angle of 0.12° was used. The scattered X-rays were recorded using an in-vacuum (<1 × 10⁻⁵ mbar) Pilatus 2M detector (Dectris AG, Baden-Daettwil, Switzerland) with a total exposure time of 1 s, achieved by averaging and merging three separate exposures taken at three different detector positions into a single image, to remove the gaps between sensors. Measurements were calibrated using silver behenate. Background signal from the substrate scatter was minimised by keeping the incident angle of the X-ray beam close to the critical angle. Data collection and analysis were conducted using scatterBrain software (ANSTO, Clayton, VIC, Australia). The d-spacing was calculated from WAXS measurements using the scattering vector (q) representing the momentum transfer in reciprocal space, as described by:

$$q=\frac{2\pi }{d}.$$ 2

Herman’s orientation function

$$f=\frac{3⟨{\cos }^2\chi ⟩-1}{2},$$ 3

wherein

$$⟨{\cos }^2\chi ⟩=\frac{{\sum }_0^{\pi /2}I\left(\chi \right){\cos }^2\chi \sin \chi }{{\sum }_0^{\pi /2}I\left(\chi \right)\sin \chi },$$ 4

was used to evaluate the degree of orientation (long-range crystalline order) by analysing the azimuthal intensity profile [I(χ)] along the nylon-11 (100)/(010) reflection. The function varies from 1 (indicating perfect alignment parallel to the reference direction) to −0.5 for perfect perpendicular alignment, while a value of 0 represents a randomly oriented sample.

Mechanical characterisation

The compression modulus (Y₃) of the solvent-cast and EM-SRBW nylon-11 films was assessed using nanoindentation with a controlled load of 2 mN, utilising a triboIndenter (TI 950; Hysitron, Eden Prairie, MN, USA); all experiments employed a Berkovich-4 shaped diamond indenter and a standard transducer (SNS-0138-330). To improve statistical reliability, 20 indentations were conducted on each sample probe, following a method reported in the literature⁶⁵. Y₃ values were extracted from the slope of the load–displacement curve during the unloading phase of each indentation cycle⁶⁵. Amplitude modulation–frequency modulation (AM–FM) atomic force microscopy (Cypher S AFM; Oxford Instruments, CA, USA) was also performed to measure the effective elastic modulus (E^*) of the nylon-11 films. A TESPA-V2 cantilever tip (Bruker Pty. Ltd., Preston VIC, Australia) with nominal spring constant k ≈ 37 N m⁻¹ was used for all measurements. Calibration was performed using a standard polystyrene-low-density polyethylene (PS-LDPE-12M; Bruker Corp., Billerica, MA, USA) reference sample, which has a known Young’s modulus of 2 GPa for the polystyrene (PS) region and 0.1 GPa for the elastomer region⁶⁶. Initially, the AFM system was thermally tuned to determine the resonance peaks of the cantilever. A force curve was then acquired to establish tip-sample contact mechanics. Following this, the standard PS-LDPE-12M reference sample (Fig. S24) was measured to validate the system calibration before proceeding to the analysis of the nylon-11 films.

Local dipole measurements

PFM (Asylum Research MFP-3D Infinity; Oxford Instruments, Santa Barbara, CA, USA) using Dual AC Resonance Tracking Piezo Force Microscopy (DART-PFM) contact mode with a contact resonance frequency shift of approximately 260 kHz was employed for the local piezoelectric measurements. Data acquisition and analysis were conducted using the open source software IGOR Pro (v6.32A; WaveMetrics Inc., Lake Oswego, OR, USA). A conductive Pt/Ir-coated AFM tip (SCM-PIT-V2, resonance frequency 75 kHz, spring constant 3 N m⁻¹, radius ≈25 nm; Bruker Pty. Ltd., Preston VIC, Australia) was used for the DART-PFM measurements. The nylon-11 sample was adhered using carbon tape to a conductive Au/Cr substrate, which was grounded to the PFM stage prior to the measurement, in which scans were conducted at a frequency of 1 Hz over an area of at least 5 × 5μm, utilising 256 pixels per line. An AC driving amplitude is swept from 1 to 5 V whilst the tip was in contact with the nylon-11 sample to calculate the quantitative vertical piezoelectric coefficient of the nylon-11 films, referred to as the effective piezoelectric constant (d_33,eff). Unlike DART hysteresis PFM, which frequently experiences the influence of electrostatic interactions during polarisation switching, particularly in the ON-field state⁶⁷, our approach involved the use of a DC voltage bias to nullify the surface potential of the nylon-11 and thus reduce the impact of electrostatic effects during DART scanning PFM⁶⁸. To accurately mitigate and correct for DART scanning electrostatic effects, we conducted local Kelvin probe force microscopy scans to find the local surface potential and thus apply an opposing DC voltage prior to the PFM measurements to offset the surface potential⁶⁸. In addition, we also conducted pre-measurements to calculate the effective piezoelectric coefficient (d_33,eff) of periodically-polarised LiNbO₃ (PPLN; Asylum Research; Oxford Instruments, Santa Barbara, CA, USA) to further ensure reliability and accuracy in the measurement. This ensures that the PFM cantilever was correctly calibrated through the measurements of a known standard.

Macroscale piezoelectric device fabrication

Piezoelectric devices were fabricated by first detaching or peeling the films from the LiNbO₃ substrate, then depositing electrodes onto each of the solvent-cast control, M-SRBW and EM-SRBW-synthesised films via electron beam deposition (PRO Line PVD 75; The Kurt J. Lesker Company, Jefferson Hills, PA, USA). This process involved depositing a 10 nm Cr layer and a 100 nm Au layer onto the films to constitute the two electrodes sandwiching the film. A shadow mask with a 0.56 cm² active area was first used to define the electrode placement on both sides of the material. Copper foil tape was then attached as a contact point to securely connect with the Cr/Au coating. The wires were soldered onto the copper tape before making this connection to avoid any heat damage. Finally, insulating polyimide tape (Kapton^®; DuPont Company, Wilmington, DE, USA) was applied to both surfaces to fully enclose the films as a final protective layer.

Macroscale piezoelectricity quantification

Macroscale piezoelectric measurements of each piezoelectric device (Fig. S26A) involved subjecting them to in-contact repetitive compressive force and measuring the output voltages and currents. Briefly, for voltage measurements, a compressive dynamic load with a minimum of approximately 2 N and a maximum of 12 N (ΔF = 10 N; 200 kPa) was applied to the active area (0.5 cm²) of the piezoelectric device (film surface area of 1.2 cm²) at a frequency of 1 Hz (velocity of 10 mm s⁻¹) using a dynamic testing instrument (ElectroForce 2500; Bose Corp., Eden Prairie, MN, USA). For current measurements, all parameters remained constant except for the maximum applied load, which was varied between 12 and 42 N (ΔF = 10–40 N; 200–800 kPa). Under varying temperature and humidity, a dynamic load of 2–22 N (ΔF = 20 N; 400 kPa) at 1 Hz with an in-contact velocity of 100 mm s⁻¹ was applied. Temperature was controlled using a thermoelectric Peltier cooler module (MPADV-127-140170-S; Premier Farnell Limited, Sydney, NSW, Australia).

Volumetric power density (P_D) calculations

The electrical output, in this case, voltage, was determined by connecting a known variable resistor (1 kΩ–4 GΩ) in parallel with the piezoelectric device. Readings were acquired with a precision source meter unit (SMU-B2912B; Keysight Technologies, Mulgrave, VIC, Australia) at 0.1 s intervals (Fig. S26B), which has an input impedance >10 GΩ. Voltage data were captured using QuickIV Measurement software (Keysight Technologies, Colorado Springs, CO, USA) (Fig. S26B). From these measurements, the power output density of the device can then be calculated from

$$P_{\mathrm{D}}=\frac{V_{pp}^2}{A_{e}Rt},$$ 5

wherein V_pp is the peak-to-peak voltage, R the load resistance, t the film thickness, and A_e the effective surface area of the cylindrical impactor in contact with the electrode surface (0.56 cm²), respectively. To demonstrate the piezoelectric device’s performance, we subjected it to high-pressure dynamic vehicular compression (2016 Ford Focus Sport; Ford Motor Co., Detroit, MI, USA) and finger tapping before and after the operation. The electrical output was measured using the source metre unit, demonstrating efficient energy conversion under high-pressure mechanical stress.

Piezoelectric charge coefficient (d₃₃) calculations

Independently, we corroborate the aforementioned piezoelectric charge coefficient with the measured charge output from varying in-contact dynamic compression of an aluminium force head; the resultant slope of the charge–force relationship corresponds to the d₃₃ value of the material. Importantly, in-contact preloading negates contributions arising from electrostatic contact electrification interactions⁴⁷. The devices were placed with conductive tape on an acrylic stage. Since the area of the bottom and top electrodes (deposited on top of the nylon-11 film) is equal in size, it is possible to obtain the piezoelectric coefficient from the charge generated from a given dynamic compressive force:

$$d_{\mathrm{33}}=\frac{Q_{sample}}{F_{dynamic}}.$$ 6

In this case, a dynamic mechanical tester (ElectroForce 2500; Bose Corp., Eden Prairie, MN, USA) was utilised for the compression testing. Signals were recorded using a picoammeter (6485; Keithley Instruments, Cleveland, OH, USA), which has an input impedance of 1 GΩ, extracted using a digital data acquisition toolbox (NI-DAQ 9223; National Instruments Corp., Austin, TX, USA) with input impedance >1 GΩ at a sampling rate of 1 kHz, and analysed using DAQExpress software (National Instruments Corp., Austin, TX, USA).

Piezoelectric voltage coefficients (g₃₃) were calculated from experimentally determined d₃₃ values and the film’s relative permittivity as follows:

$$g_{33}=\frac{d_{33}}{{\epsilon }_0{\epsilon }_{33}^{\mathrm{T}}},$$ 7

where ε₀ is the permittivity of free space and ${\epsilon }_{33}^{\mathrm{T}}$ the relative permittivity (dielectric constant) of the material at a constant stress.

Electromechanical coupling coefficients (k₃₃) were determined through the energy conversion relationship:

$$k_{33}^2=\frac{d_{33}^2{Y}_3}{{\epsilon }_0{\epsilon }_{33}^{\mathrm{T}}},$$ 8

where Y₃ is the compressive modulus measured from the nanoindentation experiment.

Measurement of dielectric properties

The dielectric properties of the nylon-11 films were investigated using an LCR meter (E4980AL; Keysight Technologies, Mulgrave, VIC, Australia). The frequency was swept from 100 Hz to 1 MHz at 0.5 V, with the probes connected directly to the electrodes of the film. Three individual samples were measured for each nylon-11 film to ensure consistency. The relative permittivity (${\epsilon }_{33}^{\mathrm{T}}$) of the nylon-11 film was obtained from the measured capacitance (C) from

$${\epsilon }_{33}^{\mathrm{T}}=\frac{tC}{A_p{\epsilon }_0},$$ 9

for a thickness t = 10 μm and overlapping electrode area A_p = 0.56 cm². The dielectric loss ($\tan \delta$), on the other hand, was calculated from

$$\tan \delta =\frac{1}{\omega RC},$$ 10

with C being the parallel capacitance, R the resistance and ω the angular frequency. All dielectric measurements were performed at constant stress and ambient conditions (≈38% RH).

Author contributions

Conceptualisation: R.K., Y.E., A.R.R. and L.Y.Y.; methodology: R.K., P.C.S., A.R.R. and L.Y.Y.; formal analysis: R.K., Y.E., P.C.S., A.R.R. and L.Y.Y.; investigation: R.K. and A.R.R.; visualisation: R.K., Y.E., P.C.S., A.R.R. and L.Y.Y.; supervision: P.C.S., A.R.R. and L.Y.Y.; writing—original draft: R.K., Y.E. and A.R.R.

Peer review

Peer review information

Nature Communications thanks Chris R. Bowen, who co-reviewed with Zihe Li, and the other, anonymous, reviewer(s) for their contribution to the peer review of this work. A peer review file is available.

Data availability

All data supporting the findings of this study are available from the corresponding author upon request.

Competing interests

The authors declare no competing interests.

Supplementary information

The online version contains supplementary material available at 10.1038/s41467-025-66389-1.