Biomimetic actuator crafted from a relaxor ferroelectric polymer nanocomposite

Domain: chemistry
Published: 2026-03-02
Editorial outcome: major_revision

Abstract
Soft actuators have garnered widespread attention due to their distributed actuation and higher rortolerance. However, fabricating biomimetic actuators with small-size, simple structure, low energy consumption, and multi-functionality remains challenging. Here, we present a monolayer nanocomposite achieved by incorporating polymer dots into a polyvinylidene fluoride terpolymer. Through interfacial hydrogen bonding and temperature modulation, a gradient distribution of highly polarized regions is achieved within the nanocomposite. This yields ultra high electromechanical performance, with an actuation strain of 14.4% and an output mechanical energy density of 1.92 J cm−³ at 100 MV m−¹. Multifunctional soft actuators are formed from this nanocomposite, weighing only 50 mg, and can skillfully imitate both caterpillar crawling and butterfly flight, unlocking the potential for multimodal locomotion. This insect-sized bionic actuator consumes a mere 3-8 mW. The combination of simple architecture and low energy consumption may pave the way for future development of truly bionic soft robots.

Introduction
Inspired by the movements and habits of small creatures in nature, manufactured as centimeter-scale crawling or flight robots. However, significant efforts have been dedicated in recent years to developing current insect-sized robots typically have complex transmission miniature robots1. In particular, centimeter-scale robots resembling structures, which significantly limit their further miniaturization, insects can perform actions such as crawling2–4, jumping5–8 and mechanical robustness, and reduction of energy consumption. flying9,10. Compared with relatively big traditional robots, these insect- Material selection and design are fundamental to actuation sized soft actuators are not only small in size and lightweight, but also performance19. To ensure that insect-size actuators can handle have low energy consumption and low cost. The application areas for complex terrains and meet diverse operational demands, the ideal insect-size soft actuators encompass medical, communications, agriculture, and military. For instance, they can be precisely maneuvered actuator material should be capable of performing multiple loco- and deployed in hazardous or hard-to-reach areas, such as scanning for motion modes under electrical stimulation. In addition, current soft signs of life during search and rescue missions. In forests, swamps, or robots typically rely on the synergy of multiple material components underground caverns, they can collect environmental data, including to ensure smooth operation. Factors such as the flexibility of the temperature, humidity, and gas composition. Previous studies passive layers and rigid supporting structures, and the bonding have described actuators, which can be actuated chemically11, between the distinctive materials, all affect the overall robotic per- pneumatically12, hydraulically, thermally13,14, or electrically15,16. Among formance. The transmission between the different layers and these different soft actuators, electrostrictive actuators have shown structures can lead to excessive energy consumption. Bionic soft muscle-like energy density, lightweight, flexibility with non-geometric robots using multifunctional materials are expected to emerge that constraints, and fast response17,18, promising to be designed and can integrate actuation and transmission while supporting multiple motion modes.

Polyvinylidene fluoride (PVDF) and its various derivatives, film, we fabricated an insect-sized soft motor capable of multiple including copolymers, terpolymers, and tetrapolymers, have shown motion modes. The bionic caterpillar actuator can crawl at a speed of promise for electrostrictive actuation20,21. However, these relaxor fer- 4.5 cm s−1 with an ultra-low input power of only 3.62 mW, an order of roelectric polymers still have limitations, such as much smaller magnitude lower than reported crawling devices of comparable size. mechanical energy densities compared to piezoelectric ceramics due The biomimetic butterfly actuator flaps its wings at a frequency of to inefficient outcoupling of large in-plane expansion in the soft 8 Hz, achieving a lift-off height of 2–4 mm, with an input power of material to generate useful strain and force outputs22,23.

Herein, we report an all-organic electrostrictive nanocomposite 7.98 mW. Both the crawler and butterfly were made solely from this comprised of a relaxor ferroelectric terpolymer P(VDF-TrFE-CFE) and nanocomposite film, without any additional mechanisms to assist in gradiently distributed polymer dots (PDs). The incorporation of CFE crawling or flying, and the total weight of the actuator is only ~50 mg.

Results
Fabrication and interaction mechanism
Fig. 1a shows a schematic of the fabrication process for the relaxor ferroelectric polymer nanocomposite with a gradient distribution of all-trans conformations, which consists of relaxor ferroelectric P(VDF-TrFE-CFE) terpolymers and polymer dots (PDs) (See Method section for details). Briefly, PDs were initially prepared by hydrothermal cross-linking with citric acid and ethylenediamine following the classical bottom-up approach24. And PDs were homogeneously mixed with pre-dissolved PVDF terpolymer solution by tip ultrasonic. After that, the solution was drop-casted onto a clean glass and followed by a progressive heating process (60–120 °C) to evaporate the solvent. When the temperature exceeds 60 °C, there is a significant difference in solvent evaporation rates between the upper (in contact with air) and the lower (in contact with a glass substrate) surfaces of the composite film, which results in a higher proportion of all-trans conformations in the lower layer compared to the upper layer25. Note that this optimized composite film is defined as terpolymer/PDs-G, and the pure terpolymer film prepared using a progressive heating process is defined as terpolymer-G. After peeling off the film from the glass substrate, carbon nanotubes (CNTs) dispersion was spray-coated on two surfaces of the polymer film to achieve the elliptically shaped active area as an electrode. By utilizing the abundant functional groups on the PD surface to form hydrogen bonds with the fluorine atoms in the polymer chains, the proportion of all-trans conformations in the composite film is increased (Fig. 1b), resulting in the construction of films with a gradient distribution of all-trans conformations (Fig. 1c). The nanocomposite terpolymer/PDs-G film prepared using this strategy exhibits a giant strain in the presence of an electrical field (Fig. 1d). Under an applied electric field of 42 MV m−1, the terpolymer/PDs-G film folds together directly, while the terpolymer-G film undergoes only minor deformation (Supplementary Movie 1).

Fig. 1e and Supplementary Fig. S1a illustrate the thickness strain (S 3) of nanocomposites with different concentrations of PDs as a function of electric field. As PD content increases, the strain initially rises and then decreases, and the composite with 0.5 wt% PDs content has the highest strain. This is due to the best hydrogen bonding interactions between PDs and the terpolymer in the 0.5 wt% PDs doped composites. To assess the robustness and standardization potential of our gradual-heating process, we rigorously evaluated its batch-to-batch reproducibility. Multiple independent batches of the terpolymer/PDs-G (c = 0.5 wt%) composite were fabricated under identical conditions. The electrostrictive strain values of these new batches were highly consistent with the data from our original batch reported above (Supplementary Fig. S2). The exceptionally low standard deviations relative to the mean values confirm that our fabrication process is stable, straightforward to standardize, and yields highly reproducible, high-performance materials. To analyze the reason for the superior performance of the gradient-distributed structure composite, we investigated the effect of PDs introduction and the impact of the gradient structure on the performance of the nanocomposites using the control variable approach, respectively. First, to study the influence of PDs on the composite from a microscopic perspective, we prepared uniformly distributed composite films (named as terpolymer/PDs) and terpolymer films by the slow solvent evaporation method at a constant temperature of 60 °C. The Fourier transform infrared (FT-IR) spectrum of PDs (Fig. 1f) confirms the presence of amide functional groups, as evidenced by their vibrational fingerprints centered at 1562 cm−1 (N-H in-plane bending), 1704 cm−1 (C=O stretching of the amide bond), and 3300 cm−1 (N-H stretching). As shown in Fig. 1g, after incorporating PDs into the terpolymer, the bending absorption peak of the N-H (in-plane bending) shifted to 1550 cm−1, and a very obvious absorption peak appeared at 1636 cm−1 which increased with the increase of PDs content. This phenomenon is attributed to the formation of hydrogen bonds between the N-H of the amide bond in the PDs and the fluorine atom in the terpolymers26, and hydrogen bonds can induce the formation of all-trans conformation in PVDF-based polymers27. The gradient distribution of the all-trans conformation, rather than the spatial distribution of PDs, is the primary outcome of our gradual-heating process. While the PDs are uniformly dispersed due to good compatibility with the terpolymer matrix, the thermal gradient during film formation creates a corresponding solvent evaporation rate gradient. The bottom surface of the film, experiencing a slower evaporation rate, facilitates a more orderly arrangement of polymer chains into the all-trans conformation. The PDs, rich in hydrogen-bonding sites, locally amplify this effect by stabilizing the all-trans conformation around them. Thus, the driving force is the evaporation rate gradient templated by the thermal gradient, which directly programs the gradient in polymer chain conformation.

X-ray diffraction (XRD) pattern shows the (110/200) reflection of composites shifting towards higher angles (Fig. 2a). This indicates that the introduction of PDs reduced the interchain spacing of the terpolymer. By comparing FT-IR spectra, we observed an increase in the all-trans conformation within the terpolymer following the introduction of PDs (Fig. 2b). To visualize the polar network in the nanocomposite more directly, we performed atomic force microscopy-infrared spectroscopy (AFM-IR) on both the terpolymer and the nanocomposite28. In the chemical image of the terpolymer/PDs, the polar network structure was clearly visible (Fig. 2c), whereas the terpolymer lacked a polar network due to the absence of the all-trans conformation (Fig. 2d). Local infrared spectroscopy confirmed that PDs introduction indeed induces the formation of the all-trans conformation in the terpolymer (Fig. 2e and 2f). In contrast, due to all-trans conformation defects, homogeneous chemical mapping existed for the terpolymer (Fig. 2d). The thermal properties of the composites with varying PDs contents were characterized by differential scanning calorimetry (DSC). The second heating cycle revealed that both the melting temperature and melting enthalpy of the composite films exhibited a gradual decrease with increasing PDs content (Supplementary Fig. S3). This trend indicates that the incorporation of PDs, through hydrogen bonding, significantly influences the microstructure of the polymer matrix, likely by altering the chain packing and organization associated with the all-trans conformation. We measured the S 3 of these uniformly distributed all-trans conformation films (Supplementary Fig. S4 and S5), which were 5.6% (10 Hz) and 6.7% (1 Hz) for terpolymer/PDs and 2.5% (10 Hz) and 3.1% (1 Hz) for the pure terpolymer at 100 MV m−1. The electrostrictive performance of composites was significantly enhanced by introducing PDs. The electrostrictive strain was characterized at 1 Hz and 10 Hz to ensure direct comparability with established literature and to assess the frequency-dependent response20,21. The marginally higher strain at 1 Hz is typical of relaxor ferroelectrics, reflecting the enhanced polarization alignment at lower frequencies. The maintained performance at 10 Hz confirms the material’s capability for faster actuation, highlighting its versatility across different operational regimes. It is noteworthy that the absolute electromechanical performance of P(VDF-TrFE-CFE) can be influenced by several intrinsic and extrinsic factors, including the specific monomer ratio and film-processing history. Therefore, the most critical and rigorous assessment of any enhancement strategy is a direct comparison against a control sample prepared and characterized under identical conditions, as performed in this work. All strain data reported herein were derived from multiple parallel measurements, with error bars confirming the high reproducibility of our results. We obtained the polarization response for the terpolymers and nanocomposites under unipolar and triangular electric fields (Supplementary Fig. S6 and S7), and the slender hysteresis loops showed the relaxor ferroelectric properties of the terpolymer/PDs. The composites exhibited higher polarization than that of pure terpolymers under electric fields, and the terpolymer/PDs displayed higher polarization of 0.064 C m−2 than the terpolymer of 0.057 C m−2 under an electric field of 100 MV m−1 at 1 Hz.

Effect of polymer films with a gradient distribution of all-trans conformation on actuation strain
After clarifying the interaction between PDs and terpolymers, we analyzed the mesoscopic structure of the films with a gradient of all-trans conformation. The FT-IR spectra of the upper and lower surfaces of the terpolymer films before and after the introduction of PDs confirmed that all-trans conformations were concentrated on one side (Fig. 3a). The absorption peak intensity of the lower surface of the composite at 1286 cm−1 was significantly higher than that on the upper surface of the composite, whereas the absorption peak intensity of the lower surface of the terpolymer at 1286 cm−1 was slightly increased compared to that on the upper surface. Fig. 3b shows the grazing incidence XRD patterns of the upper and lower surfaces of terpolymers and terpolymer/PDs composites. The reflection of the lower surface of the polymer film before and after adding PDs was shifted towards a higher angle compared to the upper surface, indicating a decrease in the interchain spacing between the lower surfaces of the terpolymer and composites. The introduction of PDs increases the number of all-trans conformations, while gradual heating promotes the uneven distribution of these increased all-trans conformations. The spatial chemical analysis of the cross-sections of terpolymer-G and terpolymer/PDs-G based on AFM-IR was shown in Fig. 3c-f (laser irradiation frequency of 1286 cm−1, corresponding to all-trans conformation). Five points were selected from bottom to top on the cross-section of the materials, and local infrared spectroscopy confirmed that the cross-section of the terpolymer/PDs-G gradually decreased in all-trans conformation from bottom to top (Fig. 3e), with a slow change process rather than a sudden change. The local infrared spectrum of the cross-section of the terpolymer-G was shown in Fig. 3f, and the all-trans conformation also changed slowly. The interconnection network of polar phases can be clearly seen in the chemical image of the terpolymer/PDs-G composite (Fig. 3c). On the contrary, owing to the deficiency of all-trans conformation, homogeneous chemical mapping was present in the terpolymer (Fig. 3d). The AFM-IR of the upper and lower surfaces of the terpolymer-G films before and after the introduction of PDs more intuitively confirmed the phenomenon of concentrated distribution of all-trans conformation on one side (Supplementary Fig. S8). The polar phase was most pronounced on the lower surface of the composite film by laser irradiation at 1286 cm−1. We performed high-resolution transmission electron microscopy (HRTEM) on the upper and lower layers of the cross-section of the terpolymer/PDs-G composite (Fig. 3g, 3j). The fast Fourier transform indicated that the terpolymer at the interface between the lower layer of the terpolymer/PDs-G nanocomposite (Fig. 3h) exhibited more all-trans conformation than the upper layer (Fig. 3i). We compare the actuation strain of all-trans conformational changes originating from enhanced hydrogen bonding interactions in a gradient-distributed structure. Fig. 3k, S1b, and S9 show the S 3 of terpolymer/PDs-G under different electric field strengths. At 100 MV m−1, the S 3 of the terpolymer/PDs-G film reaches 13.1% at 10 Hz and 14.4% at 1 Hz, whereas that of the terpolymer-G film is only 3.0% at 10 Hz and 3.2% at 1 Hz, which is an improvement of approximately 330-350%. As shown in Fig. 3l, the terpolymer/PDs-G film exhibits the largest actuation performance when compared to inorganic, organic, and composite materials reported in the literature20,22,29–32. In addition, as the frequency increases, the peak of the dielectric constant (ε r) of the terpolymer/PDs-G film shifts to higher temperatures, indicating its relaxor ferroelectric properties (Supplementary Fig. S10). 0.5 wt% of terpolymer/PDs-G shows a higher ε r than that of terpolymer-G at 1000 Hz and room temperature, which contributes to the enhancement of its electrostrictive performance (Supplementary Fig. S11). Meanwhile, the P-E loops of the terpolymer-G and terpolymer/PDs-G nanocomposites under unipolar and bipolar electric fields (Supplementary Figs. S12 and S13) display slim hysteresis curves, further confirming their relaxor ferroelectric behavior. Moreover, under an electric field of 100 MV m−1 and a frequency of 1 Hz, the polarization of the composite material reaches 0.070 C m−2, which is higher than 0.059 C m−2 for terpolymer-G.

Electrostriction coefficient (Q 33) is a key parameter for evaluating the electrostriction properties of materials. Based on the electrostrictive strains (Fig. 3k and Supplementary Fig. S1b) and polarizations (Supplementary Fig. S12), we derived the Q 33 of nanocomposites and pristine terpolymer under different electric fields (Supplementary Fig. S14) from the equation: S 3 = Q 33 P2. Q 33 of terpolymer/PDs-G is more than twice as much as that of the terpolymer-G. At 100 MV m−1 and 1 Hz, the |Q 33 | of terpolymer/PDs-G is 44.3 m4 C−2 while that of terpolymer-G is only 10.6 m4 C−2. The electromechanical coupling factor (k 33) can be calculated from output mechanical energy with k 33 = output mechanical energy / input electrical energy33. At 100 MV m−1 and 1 Hz, the k 33 of terpolymer/PDs-G is 0.370 while that of terpolymer-G is 0.0794 (as shown in Supplementary Fig. S15). For effective operation of electric actuators, the driving material must have excellent mechanical energy density (E m) to generate sufficient force to move the load. Mechanical energy density is defined as E m = 0.5 YS² (Y, elastic modulus; S, actuation strain). The large strain and modulus (0.186 GPa) in the terpolymer/PDs-G produced 1.92 J cm−3 (Supplementary Fig. S16), which is nearly twice that of ferroelectric polymers and polymer composites, and even comparable to ferroelectric ceramics.

To investigate whether the deformation effect of protrusion or depression in composites is related to the elastic modulus of the upper and lower surfaces of the film, we measured the elastic modulus of the relatively uniform distribution of all-trans conformation before and after adding PDs (Supplementary Fig. S17). The elastic moduli of the composites were generally higher than those of the terpolymers, with the elastic moduli of the upper and lower surfaces of the terpolymer being 153 MPa and 156 MPa, respectively, while those of the terpolymer/PDs were 170 MPa and 175 MPa for the upper and lower surfaces of the terpolymer/PDs, respectively. The elastic modulus of the materials with a uniform distribution of the all-trans conformations was slightly higher for their lower surfaces than for their upper surfaces. Supplementary Fig. S18 showed the elastic modulus of the upper and lower surfaces of films with all-trans conformation gradient distribution before and after the addition of PDs. The elastic moduli of the upper and lower surfaces of the terpolymer-G were 163 MPa and 167 MPa, respectively, while the elastic moduli of the upper and lower surfaces of the terpolymer/PDs-G composites were 183 MPa and 189 MPa, respectively. The differences in elastic modulus between the upper and lower surfaces of the two were similar to those of terpolymers and composites with a uniform distribution of all-trans conformation. The composites with a gradient distribution of all-trans conformation can realize unidirectional bending actuation. Therefore, the deformation effect of composites that can achieve protrusion or depression was mainly related to the gradient distribution of all-trans conformation, but not to the elastic modulus of the upper and lower surfaces. This asymmetry between the upper and lower surfaces caused the polymer composite to form a structure similar to that of a classic unimorph device, allowing the single-layer polymer film to bend under an electric field. Leveraging this unique characteristic, we have developed a complete caterpillars-to-butterflies biomimetic robotics system.

The manufacture and performance of the caterpillars-to-butterflies bionic actuation system
Nanocomposites with a gradient structure have excellent actuation properties because they form a natural single deformation. The actuator made of this nanocomposite has the potential to create insect-size robots. The breakdown strength of terpolymers and nanocomposites is shown in Supplementary Fig. S20. The introduction of PDs increases the breakdown strength by about 27 MV m−1 to 249 MV m−1. To gain deeper insight into the mechanism behind the enhanced dielectric breakdown strength, we employed thermally-stimulated depolarization current (TSDC) and leakage current measurements (Supplementary Fig. S21 and S22). The TSDC spectrum of the terpolymer/PDs-G composite revealed a main relaxation peak shifted to a higher temperature (24.9 °C) compared to that of the pure terpolymer-G (23.8 °C), indicating the presence of deepened energy traps for charge carriers, which we attribute to the strong hydrogen-bonded interfaces. Furthermore, the composite exhibited a lower leakage current density than the pure polymer under high electric fields (Supplementary Fig. S22). These findings collectively demonstrate that the PDs, via hydrogen bonding, create effective charge traps that suppress charge injection and transport. This suppressed conduction is a key factor in delaying the onset of dielectric breakdown, thereby contributing to the breakdown strength and reliability of the composite. When fabricating dielectric devices, breakdown strength is a critical material property metric, requiring that the applied electric field remain below 20% of the breakdown strength20. For actuators fabricated from our materials, actuation performance at an electric field below 50 MV m−¹ is particularly crucial.

The upper surface of the terpolymer/PDs-G was placed upward, as shown in Fig. 4a. The composite was made into a shape similar to that of a caterpillar. The head and tail of the caterpillar are sharper and narrower compared to the abdomen. A slightly wider abdomen provides sufficient power while crawling, and the head and tail are designed with symmetrical circular arcs to reduce the resistance of the caterpillar to extend and stretch its body during the crawling process. The electrode coating is applied in the middle of the material, and the elliptical electrode was designed to ensure the stable crawling of the caterpillar during the crawling process. The crawling route is designed as a bumpy ladder structure, which is more in line with the various complex road conditions encountered by caterpillars during their survival in nature. Supplementary Fig. S19a shows the complete morphology of the crawling device. In the presence of an electric field, the caterpillars can steadily crawl at a peristaltic frequency of 2 Hz. The caterpillar crawled a distance of 9 cm in 2 s, with a crawling speed of up to 4.5 cm s−1. The operating voltage and current are shown in Supplementary Fig. S23, and the single-cycle energy consumption can be calculated (W = UIt, U, single-cycle voltage; I, single-cycle current; t, single-cycle elapsed time). The single-cycle energy consumption is only 9.04 × 10−4 J. We recorded the crawling process using a high-speed camera (Fig. 4b, Supplementary Movie 2).

In contrast to the above process, by changing the shape of the film and control conditions, going through such a cocooning process, we also produced a bionic butterfly flight actuator from this material. As shown in Fig. 4a, the composite was made to resemble the shape of a butterfly. The head and tail of the butterfly were designed in a streamlined shape. The reason for the small, curved, streamlined shape of the head is that it carries the head of the butterfly when the wings swing upward. Then, when the wings were swung downward, the large, curved, streamlined shape of the tail was designed to allow the butterfly to take off from the ground. We coated the electrodes in the middle area of the material and designed the electrode area in an elliptical shape because the elliptical active area was more conducive to driving the wings on both sides. In the presence of an electric field, the wings achieved reversible flapping behavior. The operating voltage and current are shown in Supplementary Fig. S24, and single-cycle energy consumption was only 4.99 × 10−4 J. Supplementary Fig. S25a shows the complete morphology of the flight device. We recorded the actuation process using a high-speed camera (Fig. 4c, Supplementary Movies 3 and 4). Without any additional auxiliary equipment, the active actuation area in the center of the film alone can carry up to 20 times its own weight, allowing the flight device to reach a height of 2-4 mm above the ground. To precisely evaluate its actuation performance, we systematically characterized the variation of the angular displacement of the butterfly wing with the applied electric field. We systematically characterized the angular displacement of the butterfly wing as a function of the applied electric field (Supplementary Fig. S26). As shown in Supplementary Movie 5, the actuator demonstrates rapid, reversible, and highly repeatable flapping behavior under a 1 Hz electric field. The flipping angle of the wing steadily increases with the electric field strength. Quantitative data reveal that the angular displacement grows from a slight movement at low fields to a substantial angle of approximately 90° at an electric field of 50 MV m−1. This significant angular displacement not only visually confirms the success of our gradient design in amplifying local high strain into effective macroscopic motion but also signifies that this all-organic actuator has reached a state-of-the-art level in its output capability.

The two insect-sized actuators are lightweight. The bionic caterpillar actuator weighs 55.4 mg, and the bionic butterfly actuator weighs 44.8 mg. As shown in the Table S1, both actuators are significantly lighter than those reported in existing studies3,17,34–40, demonstrating a substantial weight advantage for their class. To quantitatively evaluate the energy efficiency of our actuators, we calculated the input power per cycle from the recorded voltage and current transients during operation, defined as P = W/t (W, single-cycle energy consumption; t, single-cycle elapsed time). The most rigorous validation of our material strategy comes from a direct comparison under identical actuator designs. We fabricated bio-inspired caterpillars and butterflies using both the terpolymer/PDs-G composite and the control terpolymer-G. Remarkably, to achieve a comparable actuation strain, the terpolymer/PDs-G actuator requires a driving field of only 30 MV m−1, whereas the terpolymer-G actuator necessitates 100 MV m−1 (Fig. 3k). Their operating voltage and current are shown in Supplementary Figs. S27 and S28, and their motion behavior is depicted in Supplementary Movies 6 and 7. This drastic reduction in the required operating field directly translates to a substantial decrease in input power. Specifically, the input power for the terpolymer/PDs-G caterpillar actuator is 2.532 mW, which is an order of magnitude lower than the 26.32 mW required by the terpolymer-G caterpillar at their respective driving fields. A similar power reduction was confirmed for the butterfly actuators (5.248 mW vs. 57.94 mW). This internal comparison provides definitive evidence that our composite material enables ultra-low power consumption. Furthermore, compared with other reported electric soft actuators (Fig. 4d, Table S1)3,17,34–38,41,42, our terpolymer/PDs-G actuators (consuming 3.62 mW and 7.98 mW at higher performance regimes) also demonstrate highly competitive, low input power, underscoring their overall efficiency and practical potential. The achievement of two locomotion modes using a single material with ultra-low energy consumption allows for a targeted approach to meet different needs, which is unprecedented. The composite has enormous application value in the field of soft robot actuation and provides ideas for the profound development of sensors.

While this work demonstrates that our single-layer, gradient composite actuator achieves exceptional performance in terms of strain, low mass (~50 mg), and ultra-low input power (on the order of milliwatts), we acknowledge that its current operational framework requires a high driving voltage (~1.5 kV). This requirement for an external high-voltage source presents a recognized challenge for the development of fully untethered, autonomous soft robots, as it necessitates ancillary power conversion circuitry. The core trade-off in our present design is indeed the achievement of a simplified, monolithic fabrication process and high electromechanical performance at the cost of operational voltage. However, it is crucial to distinguish between operating voltage and input power. The low power claim central to this work refers to the minimal energy consumption (in mW) during operation, a metric in which our actuators are highly competitive, as benchmarked in Table S1. This low power consumption, combined with the exceptionally low mass of the active material itself, is a critical step toward future energy-autonomous systems. The single-material, unimorph architecture eliminates interfacial energy losses common in multi-layer laminates and is the direct cause of the low mass and simple structure that enable complex biomimetic motions. Future work will squarely focus on overcoming the voltage limitation, through strategies such as developing much thinner films or exploring multi-layer co-fabrication to reduce the required driving voltage to more practical levels without compromising the achieved benefits of low mass and power.

Methods
Materials
P(VDF-TrFE-CFE) terpolymer with the composition of VDF/TrFE/CFE = 61.2/31.3/7.5 mol% was obtained from Piezo Technologies, France. N,N-dimethylformamide (DMF) was purchased from Sigma-Aldrich. Citric acid (CA) was obtained from Aladdin, and ethylenediamine (EDA) was received from Sinopharm Chemical Reagent co., Ltd. All the chemicals were used as received.

Synthesis of polymer dots
In this work, the PDs were prepared following a classical bottom-up method, hydrothermal crosslinking24, as shown in the Supplementary Fig. S29. CA (1.0507 g) and EDA (335 µL) were initially dissolved in deionized water (10 mL). Then the mixed solution was transferred to a poly(tetrafluoroethylene) (Teflon)-lined autoclave (30 mL) and heated at 180 °C for 7 h, the reactors were naturally cooled to room temperature. At the end of the reaction, the solution was transferred to a dialysis bag for 30 h. After dialysis, the solution was freeze-dried for 72 h to obtain brown PDs. Finally, the brown-black powders PDs were obtained (Supplementary Figs. S30–S32, Supplementary Information for relevant characterization of PDs).

Polymer/polymer dots (PDs) composite films preparation
In the preparation of the composite films (terpolymer/PDs-G) shown in Supplementary Fig. S33, 0.3 g of terpolymer was dissolved in 5 mL DMF and stirred overnight at room temperature. Depending on the required mass fraction in the composites, the appropriate weight of PDs was dissolved in 3 mL of DMF and sonicated for 1 h before using. These homogeneous solutions were then dropped onto a 6 cm × 6 cm clean glass slide and placed in an oven at 60 °C for 4 h, then the temperature was slowly raised to 120 °C within 8 h. Afterward, the polymer films were placed in a vacuum oven and annealed at 120 °C for 12 h (corresponding to a controlled ramp rate of 7.5 °C h−1). The thickness of the films was in the range of 30 μm to 40 μm. The distribution of characteristic elements in terpolymers and composites was characterized by energy dispersive X-ray spectroscopy (EDS), as shown in Supplementary Fig. S32. The appearance of the O element in the EDS image of the composite indicates the successful introduction of PDs into the terpolymer, and the uniform distribution of elements suggests the uniformity of the composites. The terpolymer/PDs films with uniformly distributed all-trans conformation were treated at 60-120 °C for 12 h during the preparation process to evaporate the solvent, and the other treatment methods were the same as the above process for terpolymer/PDs-G.

The fabrication and construction of the bionic caterpillar crawling actuator and the bionic butterfly flight actuator
The prepared polymer composite film was cut into homogeneous texture according to the shape and size as shown in Supplementary Figs. S34 and S35 and carbon nanotubes (CNT, XFNANO, purity > 90%, 1–2 nm in diameter and 1–3 µm in length) were deposited on the surface of the composite film by spray-coating (Dahua, DH350E) to form the elliptical shaped active area as electrodes, and the active area was extended to the upper and lower edges of the film through elongated conductive regions, respectively. The ultra-thin copper strips (0.06 mm in thickness) were attached onto the composite film with conductive adhesive to serve as a skeleton support, and connected to the high-voltage power supply through ultrafine copper wires (0.03 mm in diameter). In the bionic caterpillar crawling actuator, we designed the crawling route as a stepped structure with a slope of 40° and a length of 90 mm to simulate the rough tree-like structure in nature. Specific details can be seen in Supplementary Fig. S19, Supplementary Movie 2. For the bionic butterfly flight actuator, in order to avoid the impact of the force generated by the wings flapping on the ground on the mechanical and electrical performance display of the composites, all flight tests were conducted on narrow glass platforms, with a 1 cm small square set in the center of the glass platform to slightly sink the active area of the flight actuator, allowing the wings to be upwardly folded (Supplementary Fig. S25 and Supplementary Movies 3 and 4).

Characterizations and measurements
The SEM experiment was captured on a JEOL JSM-7401F scanning electron microscope with an EDS analysis attachment. The TEM and HRTEM images were obtained by a JEM-2100F instrument. The films were embedded in epoxy resin and then sectioned using an EM UC7 ultramicrotome. Fourier transform infrared (FT-IR) measurement was carried out on a Varian Digilab FTS-8010 spectrophotometer. Peak Force Quantitative Nanomechanical Mapping (PFQNM) characterizations of surface elastic modulus were measured by Multimode 8. AFM-IR measurements were obtained from a Bruker Anasys nanoIR3 system. Simultaneous acquisition of topographic and chemical images by tapping mode. The HRTEM images were captured on a JEM-2100F transmission electron microscope. A UV2501-PC spectrophotometer was used to get Ultraviolet-visible (UV-vis) absorption spectra. Two-dimensional (2D) glancing incidence experiments were performed at the Xeuss 3.0. Dielectric constants and dissipation factors were measured using an Alpha-A High Performance Frequency Analyzer, and the samples (diameter 3.4 mm and thickness about 30-40 μm) were deposited on a 30 nm thick copper. The polarization-electric field (P-E) loops were collected using a Polarization Loop & Dielectric Breakdown Test System from PolyK, where AC electric fields with an increment of 10 MV m−1 were applied across polymer films with a triangular waveform at a frequency of 1 Hz and 10 Hz. Leakage current density was acquired from a Keithley Model 6517A electrometer equipped with a Trek Model 610C amplifier. Thermally stimulated depolarization current (TSDC) was measured by a Keithley 6517A electrometer. The samples were firstly polarized under 20 MV m−1 at 40 °C for 20 minutes a Trek Model 610C amplifier. Then the temperature was quickly dropped to -10 °C with the electric field on. After that, the samples were disconnected from the electric field and short-circuited for 2 minutes. Finally, the samples were heated from -10 °C to 60 °C at a rate of 1 °C min−1 and the current was recorded. Differential scanning calorimetry (DSC) data was collected by a DSC Q2000 with a heating rate of 10 °C min−1. High-speed camera images and movies were captured by the SONY DSC-RX10III digital camera. High voltage was provided by a high-voltage power supply (Dongwen, DW-P103-0). The electrostrictive strain (thickness deformation) of the films was measured using a non-contact photonic sensor system (MTI-2100 Fotonic Sensor, 2032R probe). As illustrated in Supplementary Fig. S36, the sample was placed between parallel plate electrodes, and the thickness change was measured perpendicularly by the probe. The reported strain values were recorded at 0.1 s or 1 s after field application to ensure the capture of steady-state response.

Data availability
The data generated in this study are provided in the Supplementary Information/Source Data file. All data are available from the corresponding author upon request. Source data are provided with this paper.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This manuscript presents a highly innovative and well-executed study that explores a promising strategy for developing high-performance soft actuators based on a relaxor ferroelectric polymer nanocomposite. This work offers significant potential for advancing the fields of soft robotics and insect-scale actuation systems. With some minor revisions and clarifications, this manuscript would make a valuable contribution to its area. I am pleased to recommend this manuscript for acceptance. Detailed comments and suggestions are provided below:\n1. The choice of the relaxor ferroelectric polymer P(VDF-TrFE-CFE) as the matrix should be justified by explaining its relevance to the material design strategy and actuator performance.\n2. A more detailed description of the strain measurement setup is needed. The authors should specify the methods and instrumentation used to quantify strain, and it would be very helpful to include schematic diagrams and/or photographs of the measurement setup for clarity and reproducibility.\n3. The electrical measurements throughout the manuscript are conducted at 10 Hz and 1 Hz. The reason for explaining the rationale behind selecting these two specific frequencies should be given. A brief discussion of how the frequency affects actuator performance would strengthen the interpretation of the results.\n4. The exceptional actuator performance is a key strength of this work, and the manuscript would benefit from a more thorough description of the electrode design. Specifically, details such as the dimensions, materials, and placement of the ultra-thin copper strips and ultrafine copper wires should be provided to offer a more complete and vivid picture of the device architecture.\n5. The slight enhancement in the breakdown strength of the composite material is an interesting observation. Possible explanation or hypothesis for this improvement is encouraged to provide, which may further reinforce the understanding of the structure–property relationship in the composite system.\n6. The article mentions achieving a gradient distribution of PDs through gradual heating at 60-120 °C, but does not specify the heating rate.It is recommended to supplement the specific heating process to improve process repeatability.\n7. Why does the difference in solvent evaporation rates result in a higher proportion of all-trans conformations in the lower layer?\n8. The article shows that the material performs best at 0.5 wt% PDs. Please explain why the material's performance declines when the PDs content exceeds 0.5 wt% and provide relevant supporting materials.\n9. Please provide the DSC test results for pure terpolymers (terpolymer-G) and composite materials with different PDs contents (terpolymer/PDs-G).\n10. Please check the chart information in fig.2 e-f.", "reviewer_2": "This manuscript, \"Biomimetic actuator crafted from a relaxor ferroelectric polymer nanocomposite,\" presents the fabrication of a single-layer actuator based on a P(VDF-TrFE-CFE) terpolymer with a gradient distribution of polymer dots (PDs). The authors claim that this novel structure leads to a functional unimorph actuator with exceptionally high strain (14.4% at 100 MV/m) and ultra-low power consumption (3-8 mW), demonstrating its potential through caterpillar-like and butterfly-like soft robots. While the concept of creating a functional gradient within a single material is intriguing and the demonstrated actuators are visually impressive, the reviewer has several major concerns regarding the experimental design, the interpretation of the results, and the validity of the central claims. These concerns are significant enough to undermine the core conclusions of the study.\n1. The abstract prominently features an \"exceptionally high strain\" of 14.4%. This is a misleading representation of the actuator's performance. The reported value is the thickness strain (S3), which is an indirect, internal material response, not the functional output of the device. For a bending actuator, the useful performance is defined by macroscopic motion, such as bending curvature, tip deflection, or angular displacement. The high thickness strain is merely the cause that, due to the engineered gradient structure, gets converted into bending. Presenting the cause as the primary performance metric in the abstract, without this critical context, conflates an intrinsic material property with a structural amplification effect. It creates the false impression that the material itself linearly contracts or expands by this magnitude, which is not the case. This is a significant overstatement of the material's direct actuation capability.\n2. The authors claim a significant material innovation, stating that the addition of PDs doubles the electrostrictive strain of the terpolymer (from ~3% to 6.7% for uniform films at 100 MV/m, Fig. S2). However, the baseline performance of their control sample (pure terpolymer-G, ~3% strain at 100 MV/m) appears to be considerably lower than what is commonly reported in the literature for this class of terpolymers, where strains of 5% or higher at 100 MV/m are typical [1]. This raises a critical question about the validity of the baseline. By using a potentially underperforming control sample, the claimed two-fold improvement due to PDs may be an overestimation. The authors must compare their results against well-established literature values or provide a compelling reason for their low baseline performance to substantiate their claim of material innovation.\n3. The claim of \"ultra-low power consumption\" is a cornerstone of the paper, but the comparison presented in Fig. 4d and Table S1 is fundamentally flawed. The authors compare their electrostrictive actuator to actuators based on entirely different physical mechanisms (e.g., thermal, chemical, phase-change). This is not a scientifically valid. The only rigorous comparison would be against an actuator of the same design and scale fabricated from the pure terpolymer control material, which is absent from the manuscript. Without this direct comparison, the claim that their specific material/process innovation leads to an order-of-magnitude reduction in power consumption is unsubstantiated. This appears to be a case of \"cherry-picking\" data from disparate fields to make the results seem more significant than they are.\n4. The authors frame their \"single-process\" approach as a major breakthrough over conventional multi-layer, laminated actuators. However, the core trade-off they have made is to simplify the fabrication process at the cost of requiring an extremely high operating voltage. The reported operating field of ~42 MV/m for a 35 µm thick film translates to a driving voltage of approximately 1.5 kV, and the fields for maximum strain (100 MV/m) require 3.5 kV. Such high kV-level voltages are a fatal flaw for the target application of small-scale, untethered biomimetic robots. It necessitates bulky, heavy, and inefficient voltage amplification circuitry, which negates the benefits of a lightweight actuator and contradicts the practical meaning of \"low power.\" While other methods (e.g., multilayering) have more complex fabrication, they successfully address this critical voltage issue, enabling operation at much lower and more practical voltages (<100 V). The authors' failure to acknowledge and discuss this severe limitation, while heavily emphasizing the process simplification, makes the claims of practical advancement feel unbalanced and misleading. The innovation is not the simplification of the process itself, but rather the creation of a functional unimorph structure. However, the practical viability of this structure is highly questionable due to the voltage requirements.\nIn summary, the manuscript's central claims of significant material improvement, ultra-low power consumption, and practical innovation are built on a foundation of questionable baselines, inappropriate comparisons, and the omission of a critical discussion on the impracticality of the high operating voltage. While the idea is interesting, the work in its current form does not provide sufficient rigorous and contextual evidence to support its claims. Therefore, the reviewer regrets that the reviewer cannot recommend this manuscript for publication.\nReference\n[1] S. Zhang, C. Huang, Z.-Y. Cheng, F. Xia, R. Klein, Q. M. Zhang, V. Bobnar, A. Levstik, and F. Bauer, Relaxor Ferroelectric Polymers, Ferroelectrics, 339:37–45, 2006. (DOI: 10.1080/00150190600737917)", "reviewer_3": "In this manuscript, the authors reported multi-mode soft actuators with low power consumption of 3-8 mW but high loading weight up to 2000% of its own weight, which was achieved using a monolayer all-organic nanocomposite with gradually distributed polymer dots (PDs). They found the proportion of all-trans conformations in the terpolymer was increased by the hydrogen bonding between the functional groups on the PD surface and the fluorine atoms in the terpolymer chains, resulting a gradient distribution of the high-polarization moiety. Notably, the nanocomposite film exhibited ultrahigh electrostriction performance with maximum strain of 14.4% and an output mechanical energy density of 1.92 J cm-3 at 100 MV m-1. This finding represents a significant advancement in in ferroelectric polymers and soft actuators, paving the way to truly bionic soft robots. I strongly recommend the publication of this nice work in Nature Communications after some minor revisions.\n1. In the paper, the theoretical rationale for selecting the relaxor ferroelectric polymer P(VDF-TrFE-CFE) as the matrix and polymer dots as the filler should be systematically and clearly elaborated. This section should not only include the fundamental characteristics of the materials themselves but also provide a detailed explanation of the intrinsic relationship between their structure and properties—such as the high dielectric constant and remarkable electrostrictive effect of P(VDF-TrFE-CFE), as well as the unique advantages of polymer dots in interface modification and dielectric enhancement. Providing sufficient and clear background information will help readers gain a deeper understanding of the scientific logic behind the material selection and establish the structure-property relationship between these choices and the final actuator performance (e.g., strain output, energy density), thereby enhancing the scientific rigor and persuasiveness of the paper.\n2. The gradual distribution of polymer dots (PDs) in the monolayer all-organic nanocomposite is the highlight of this work. What is the driving force of this gradual distribution? Gradual heating temperature-rise periods. Or the graded distribution of temperatures in the films on the heating plate, where temperatures decrease from the bottom to the top.\n3. Is the \"gradual heating process\" discussed in the paper straightforward to standardize? How consistent are results from one batch to the next? Are there any data to back up its process stability?\n4. The authors assert that the incorporation of polymer dots leads to an enhanced breakdown strength. It is recommended that they provide potential explanations or hypotheses for this observed improvement, thereby facilitating a more profound understanding of the structure-property relationship within the composite system. Furthermore, the authors should consider performing leakage current and/or thermally stimulated depolarization current (TSDC) measurements to substantiate these findings.\n5. The outstanding performance of the actuator is a principal highlight of this research. To facilitate a deeper reader understanding of the device's structure-property relationship, a more detailed description of the actuator's design is recommended. Furthermore, more thorough elucidation of the strain measurement process, including the methodologies, pertinent parameters, and fundamental principles, is advised. This will significantly augment reader comprehension of the measurement approach and streamline the reproducibility of the experimental protocol.\n6. Please check the text in Fig. 2e, f. According to the caption, e is terpolymer/PDs，f is terpolymer. Please confirm it.", "reviewer_4": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts."}, "author_response": "Answers to the Comments\nWe would first like to express our gratitude to the reviewers for their insightful comments, which helped us further improve this manuscript. Detailed responses to the comments are as follows:\nReviewer #1 (Remarks to the Author):\nThis manuscript presents a highly innovative and well-executed study that explores a promising strategy for developing high-performance soft actuators based on a relaxor ferroelectric polymer nanocomposite. This work offers significant potential for advancing the fields of soft robotics and insect-scale actuation systems. With some minor revisions and clarifications, this manuscript would make a valuable contribution to its area. I am pleased to recommend this manuscript for acceptance. Detailed comments and suggestions are provided below:\n1. The choice of the relaxor ferroelectric polymer P(VDF-TrFE-CFE) as the matrix should be justified by explaining its relevance to the material design strategy and actuator performance.\nResponse: We sincerely thank the reviewer for this valuable suggestion. We fully agree that a more detailed explanation of the material selection rationale will significantly strengthen the manuscript and provide readers with a deeper understanding of the structure-property correlations in our system. We have thoroughly revised the introduction section to provide a systematic and explicit elaboration as suggested.\nRationale for P(VDF-TrFE-CFE) Matrix: We now elaborate on its unique molecular structure. The incorporation of CFE monomers into the P(VDF-TrFE) lattice disrupts the long-range ferroelectric order, creating nano-polar regions that are highly responsive to electric fields. This structure is responsible for the material's high dielectric constant and relaxor ferroelectric behavior, characterized by a large electric-field-induced polarization with minimal hysteresis – a fundamental requirement for high-performance electrostrictive actuators.\nTo address the Reviewer’s suggestion,\nThese additions have been integrated into the introduction of the revised manuscript (Page 3).\n“The incorporation of CFE monomers into the P(VDF-TrFE) lattice disrupts the long-range ferroelectric order, creating nano-polar regions that are highly responsive to electric fields. This structure is responsible for the high dielectric constant and relaxor ferroelectric behavior of the material, characterized by a large electric-field-induced polarization with minimal hysteresis, which is a fundamental requirement for high-performance electrostrictive actuators.”\n2. A more detailed description of the strain measurement setup is needed. The authors should specify the methods and instrumentation used to quantify strain, and it would be very helpful to include schematic diagrams and/or photographs of the measurement setup for clarity and reproducibility.\nResponse: We thank the reviewer for this suggestion to improve the clarity of our experimental methods. We have thoroughly revised the \"Experimental Section\" to provide a more detailed description of the strain measurement setup.\nWe characterized the electrostrictive strain (deformation in the thickness direction) of the polymer films using a non-contact photonic sensor system (MTI-2100 Fotonic Sensor with a 2032R probe). This method offers high-resolution displacement measurement, which is critical for accurately quantifying the small deformations in thin films. A schematic diagram of the measurement setup is provided in Fig. R1, and a photograph of the actual instrument is shown in Fig. R2. The film sample was placed between two parallel plate electrodes, with the photonic sensor probe aligned perpendicularly to measure the thickness change. Upon application of the electric field, the temporal evolution of the displacement was recorded. The reported electrostrictive strain values were determined after the signal had stabilized, typically at 0.1 s or 1 s after the field was applied, ensuring the measurement captured the steady-state material response rather than any transient effects.\nFig. R1 Schematic of a photonic sensor setup for the longitudinal strain measurement.\nFig. R2 The physical photo of the photonic sensor setup for the longitudinal strain measurement.\nTo address the Reviewer’s suggestion,\nThe measure instrument, measure method and related test details of electroactuation strains were supplemented in the Materials and Methods section of the revised manuscript. And we have supplemented Fig. S36 in the revised Supplementary Information and marked in red.\n“The electrostrictive strain (thickness deformation) of the films was measured using a non-contact photonic sensor system (MTI-2100 Fotonic Sensor, 2032R probe). As illustrated in Fig. S36, the sample was placed between parallel plate electrodes, and thickness change was measured perpendicularly by the probe. The reported strain values were recorded at 0.1 s or 1 s after field application to ensure the capture of steady-state response.”\nFig. S36 Schematic of a photonic sensor setup for the longitudinal strain measurement.\n3. The electrical measurements throughout the manuscript are conducted at 10 Hz and 1 Hz. The reason for explaining the rationale behind selecting these two specific frequencies should be given. A brief discussion of how the frequency affects actuator performance would strengthen the interpretation of the results.\nResponse: We thank the reviewer for this valuable suggestion. Our reasons for selecting 1 Hz and 10 Hz are twofold:\nEnsuring Broad Comparability: These frequencies are widely adopted standards in the field of relaxor ferroelectric polymers. Some key studies focus on 1 Hz1, while others, including seminal works, utilize 10 Hz2. By presenting data at both frequencies, we provide a direct and meaningful basis for comparison with a broader body of literature, allowing readers to better contextualize the performance of our material.\nProbing Material Dynamics: The choice of frequency allows us to assess the dynamic electromechanical response of our material. As expected for relaxor ferroelectric polymers, a lower driving frequency (1 Hz) provides a longer period for dipole switching and domain wall motion, often resulting in a marginally higher measured strain compared to a higher frequency (10 Hz). The slight decrease in strain observed at 10 Hz (as seen in our Fig. S3) can be attributed to the inertia of the polar nanoregions and increased viscous losses within the polymer matrix at shorter time scales.\nThis frequency-dependent behavior is not just a metric but has practical implications for actuator design. The performance at 1 Hz is highly relevant for applications requiring larger strains in quasi-static or slow-cycle operations (e.g., robotic gripping, crawling). Conversely, the performance maintained at 10 Hz demonstrates the material's capability for faster, more dynamic actuation (e.g., flapping, vibration control), highlighting its versatility. We agree with the reviewer that this discussion strengthens the paper, and we are grateful for the suggestion to include it.\nTo address the Reviewer’s suggestion,\nWe have now added the rationale for frequency selection and a discussion on frequency dependence to the revised manuscript (Page 7).\n“The electrostrictive strain was characterized at 1 Hz and 10 Hz to ensure direct comparability with established literature and to assess the frequency-dependent response 1,2. The marginally higher strain at 1 Hz is typical of relaxor ferroelectrics, reflecting the enhanced polarization alignment at lower frequencies. The maintained performance at 10 Hz confirms the material's capability for faster actuation, highlighting its versatility across different operational regimes.”\n4. The exceptional actuator performance is a key strength of this work, and the manuscript would benefit from a more thorough description of the electrode design. Specifically, details such as the dimensions, materials, and placement of the ultra-thin copper strips and ultrafine copper wires should be provided to offer a more complete and vivid picture of the device architecture.\nWe thank the reviewer for this valuable suggestion to improve the clarity of our device fabrication. In direct response, we have now revised the \" Materials and Methods Section\" in the manuscript to include the specific details regarding the electrode and structural support components.\nThe prepared polymer composite film was cut into the desired shapes (Fig. S34 and S35). Carbon nanotubes (CNT, XFNANO, purity > 90%, 1-2 nm in diameter and 1-3 µm in length) were spray-coated (Dahua, DH-350E) onto the film surface to form the elliptical active electrode area, with elongated conductive regions extending to the upper and lower edges. Ultra-thin copper strips (0.06 mm in thickness) were attached using conductive adhesive to serve as a skeletal support and current collector. These strips were then connected to the high-voltage power supply via ultrafine copper wires (0.03 mm in diameter).\nThis revision explicitly includes the materials (copper), key dimensions (thickness and diameter), and the specific roles (skeletal support and electrical connection) of these components. We believe this addition offers a vivid and complete depiction of our actuator design, directly addressing your comment.\nTo address the Reviewer’s suggestion,\nThe modified text can be found on Page 19 of the revised manuscript and has been highlighted for your convenience.\n“The prepared polymer composite film was cut into homogeneous texture according to the shape and size as shown in Fig. S34 and Fig. S35. And carbon nanotubes (CNT, XFNANO, purity > 90%, 1-2 nm in diameter and 1-3 µm in length) were deposited on the surface of the composite film by spray-coating (Dahua, DH350E) to form the elliptical shaped active area as electrodes, and the active area was extended to the upper and lower edges of the film through elongated conductive regions, respectively. The ultra-thin copper strips (0.06 mm in thickness) were attached onto the composite film with conductive adhesive to serve as a skeleton support, and connected to the high-voltage power supply through ultrafine copper wires (0.03 mm in diameter).”\n5. The slight enhancement in the breakdown strength of the composite material is an interesting observation. Possible explanation or hypothesis for this improvement is encouraged to provide, which may further reinforce the understanding of the structure–property relationship in the composite system.\nResponse: We sincerely thank the reviewer for this valuable suggestion. Following the reviewer's advice, we have conducted both thermal-stimulated depolarization current (TSDC) and leakage current measurements on the pure terpolymer-G and the terpolymer/PDs-G composite. The results provide profound insights into the underlying mechanism of breakdown strength enhancement, which we have elaborated in the revised manuscript. Our findings and proposed explanations are summarized as follows:\n1) TSDC Analysis: Revealing Deepened Charge Traps and Interface Polarization\nThe TSDC curves offer direct evidence of how PDs modify the charge trapping landscape and polarization behavior within the composite (Fig. R3):\nShift of Main Relaxation Peak: The main relaxation peak of the composite (terpolymer/PDs-G) shifts to a higher temperature (24.9 °C) compared to that of the pure terpolymer-G (23.8 °C). In TSDC, a shift to a higher temperature indicates that a greater amount of energy is required to release the trapped charges. We attribute this to the strong hydrogen bonding between the PDs and the terpolymer matrix. These interfacial hydrogen bonds act as deep, stable traps for space charges, immobilizing them and preventing their early activation and long-range migration.\nEmergence of a New Low-Temperature Peak: A distinct new peak appears at ~2.6 °C in the terpolymer/PDs-G composite, which is absent in the pure polymer. This new peak is attributed to the interface-induced polarization. The abundant PDs/terpolymer interfaces, coupled with the hydrogen bonding, create new interfacial polarization modes. These modes are activated at lower temperatures and contribute to the overall polarizability of the material without leading to destructive charge conduction.\nIncrease in Total Polarization (Peak Area): The total depolarization charge, which is proportional to the area under the TSDC curve, is larger for the terpolymer/PDs-G composite. This confirms that the composite can store a greater amount of polarized charge under an electric field. The synergistic effect of the intrinsic dipolar relaxation of the terpolymer and the newly introduced interfacial polarization leads to this enhanced charge storage capability.\nFig. R3 TSDC curves of (a) terpolymer-G and (b) terpolymer/PDs-G.\n2) Leakage Current Analysis: Direct Evidence of Suppressed Charge Transport\nThe leakage current measurement provides a direct link between the TSDC findings and the macroscopic dielectric strength (Fig. R4):\nThe terpolymer/PDs-G composite exhibits a slightly reduced leakage current density compared to the pure terpolymer-G. This result directly validates the role of PDs as effective charge traps. The deep traps revealed by TSDC effectively capture and immobilize charge carriers that are injected under a high electric field. By suppressing the formation of a conductive path, these trapped charges prevent the rapid increase of leakage current, thereby delaying the onset of breakdown strength and enhancing the overall dielectric strength of the material.\nFig. R4 The leakage current density of the terpolymer-G and the terpolymer/PDs-G.\nTo address the Reviewer’s suggestion,\nWe have incorporated these new TSDC and leakage current data, along with the above discussion, into the revised manuscript (Page 13) to strengthen our claim. The relevant test instructions for TSDC and leakage current have been supplemented in the Materials and Methods section and marked in red. The TSDC curves and the leakage current test data have been added to the revised Supplementary Fig. S21 and Fig. S22.\n“To gain deeper insight into the mechanism behind the enhanced dielectric breakdown strength, we employed thermal-stimulated depolarization current (TSDC) and leakage current measurements (Fig. S21 and Fig. S22). The TSDC spectrum of the terpolymer/PDs-G composite revealed a main relaxation peak shifted to a higher temperature (24.9 °C) compared to that of the pure terpolymer-G (23.8 °C), indicating the presence of deepened energy traps for charge carriers, which we attribute to the strong hydrogen-bonded interfaces. Furthermore, the composite exhibited a lower leakage current density than the pure polymer under high electric fields (Fig. S22). These findings collectively demonstrate that the PDs, via hydrogen bonding, create effective charge traps that suppress charge injection and transport. This suppressed conduction is a key factor in delaying the onset of dielectric breakdown, thereby contributing to the breakdown strength and reliability of the composite.”\nMaterials and Methods:\n“Leakage current density was acquired from Keithley Model 6517A electrometer equipped with a Trek Model 610C amplifier. Thermally stimulated depolarization current (TSDC) was measured by a Keithley 6517A electrometer. The samples were firstly polarized under 20 MV m-1 at 40 °C for 20 minutes a Trek Model 610C amplifier. Then the temperature was quickly dropped to -10 °C with the electric field on. After that, the samples were disconnected from the electric field and short-circuited for 2 minutes. Finally, the samples were heated from -10 °C to 60 °C at a rate of 1 °C min-1 and the current was recorded.”\nFig. S21 TSDC curves of (a) terpolymer-G and (b) terpolymer/PDs-G.\nFig. S22 The leakage current density of the terpolymer-G and the terpolymer/PDs-G.\n6. The article mentions achieving a gradient distribution of PDs through gradual heating at 60-120 °C, but does not specify the heating rate. It is recommended to supplement the specific heating process to improve process repeatability.\nResponse: We thank the reviewer for this valuable suggestion to improve the reproducibility of our method. We have now revised the corresponding section in the \"Experimental Section\" of the manuscript to explicitly state the heating rate.\nThe original description stated the film was placed in an oven at 60 °C for 4 h, and then the temperature was \"slowly raised to 120 °C within 8 h.\" To eliminate any ambiguity, we have revised this sentence for precision. The updated text now reads:\n\"...placed in an oven at 60 °C for 4 h, and then the temperature was raised to 120 °C over 8 h (corresponding to a controlled ramp rate of 7.5 °C h-1)...\"\nThis clarification provides the specific heating rate used in our process, which we agree is essential for ensuring the repeatability of gradient formation. The modification has been made on Page 19 of the revised manuscript.\n7. Why does the difference in solvent evaporation rates result in a higher proportion of all-trans conformations in the lower layer ?\nResponse: We thank the reviewer for this profound question regarding the core physical mechanism. The higher proportion of all-trans conformations at the bottom layer is a direct consequence of the slower solvent evaporation rate there, which is governed by the principles of polymer crystallization kinetics. The key lies in the competition between the timescales of chain mobility and solvent removal:\nAt the Bottom Surface (Slower Evaporation): The film's bottom is in direct contact with the hot substrate. While the absolute temperature is higher, the solvent concentration remains high for a longer duration, plasticizing the polymer chains and granting them high mobility. This prolonged \"wet\" state provides ample time for the polymer chains to undergo conformational rearrangement, align, and pack into the thermodynamically more stable, extended all-trans conformation.\nAt the Top Surface (Faster Evaporation): The top surface loses solvent rapidly into the air, especially as the oven temperature increases. This quickly increases the polymer concentration and viscosity, effectively \"freezing\" the polymer chains in place before they have sufficient time to fully unwind and organize into the optimal all-trans structure. This often results in a higher fraction of random coils or less ordered phases (like the α-phase).\nThis phenomenon is supported by literature on PVDF-based systems (J. Appl. Polym. Sci., 2010, 116, 2, 785-791) 3, which reports that slower evaporation favors the polar β-phase. Our gradual heating process deliberately creates this evaporation rate gradient, thereby programming the gradient in chain conformation. The role of the PDs is to further promote and stabilize the all-trans conformation locally via hydrogen bonding, amplifying the final effect of this physically templated gradient.\nTo address the Reviewer’s suggestion,\nWe have clarified this mechanism in the revised manuscript (Page 5).\n“The gradient distribution of the all-trans conformation, rather than the spatial distribution of PDs, is the primary outcome of our gradual-heating process. While the PDs are uniformly dispersed due to good compatibility with the terpolymer matrix, the thermal gradient during film formation creates a corresponding solvent evaporation rate gradient. The bottom surface of the film, experiencing a slower evaporation rate, facilitates a more orderly arrangement of polymer chains into the all-trans conformation. The PDs, rich in hydrogen-bonding sites, locally amplify this effect by stabilizing the all-trans conformation around them. Thus, the driving force is the evaporation rate gradient templated by the thermal gradient, which directly programs the gradient in polymer chain conformation.”\n8. The article shows that the material performs best at 0.5 wt% PDs. Please explain why the material's performance declines when the PDs content exceeds 0.5 wt% and provide relevant supporting materials.\nResponse: We sincerely thank the reviewer for this insightful question regarding the optimal PDs content. We have conducted a thorough analysis and propose the following mechanisms, supported by relevant literature, to explain the performance decline beyond 0.5 wt% PDs. The decline in electrostrictive strain at higher PDs loadings is likely attributable to two interrelated factors:\nIncreased Material Modulus and Restricted Chain Mobility: While the hydrogen bonds between PDs and the terpolymer matrix are crucial for promoting the electroactive all-trans conformation, an excessive density of these hydrogen bonds can form a rigid network that acts as physical cross-links. This network significantly restricts the mobility of the polymer chains. Since the large electrostrictive strain in relaxor ferroelectrics relies on the reorientation of polar segments under an electric field, this increased modulus directly impedes the achievable strain.\nInhibition of Dipole Switching under Electric Field: The hydrogen bonds effectively stabilize the all-trans conformations. However, beyond an optimal concentration, these bonds may become too numerous and stable, effectively \"locking\" the dipoles in place. In an alternating electric field, these overly stabilized dipoles cannot switch orientation freely, leading to a decrease in polarization response and thus a lower electrostrictive strain.\nIn summary, we posit that at 0.5 wt% PDs, a balance is achieved where hydrogen bonds optimally promote the all-trans conformation without excessively restricting chain mobility or dipole switching. At higher loadings, the detrimental effects of increased modulus and inhibited dipole reversal outweigh the benefits of further increasing the all-trans conformation content.\nTo address the Reviewer’s suggestion,\nWe have added the following content on Page 5 of the revised manuscript.\n\" Figs. 1e and S1a illustrate the thickness strain (S3) of nanocomposites with different concentrations of PDs as a function of electric field. As PD content increases, the strain initially rises and then decreases, and the composite with 0.5 wt% PDs content has the highest strain. This is due to the best hydrogen bonding interactions between PDs and terpolymer in the 0.5 wt% PDs doped composites.\"\n9. Please provide the DSC test results for pure terpolymers (terpolymer-G) and composite materials with different PDs contents (terpolymer/PDs-G).\nResponse: We thank the reviewer for this suggestion. We have conducted DSC measurements (Fig. R5). The DSC data from the second heating cycle show a clear trend: both the melting temperature and the melting enthalpy of the composite films gradually decrease as the PDs content increases. This phenomenon indicates that the incorporation of PDs influences the thermal behavior and microstructure of the polymer matrix.\nWe interpret these observations in the context of the all-trans conformation. The formation of hydrogen bonds between PDs and the terpolymer chains is known to promote the all-trans conformation. The changes in thermal properties observed by DSC suggest that this interaction also alters the packing and organization of the polymer chains on a broader scale. The specific nature of this microstructural evolution, as reflected in the thermal data, is a subject of our ongoing investigation.\nFig. R5 The second heating curves of the terpolymer-G and composite materials with different PDs contents.\nTo address the Reviewer’s suggestion,\nWe have supplemented the relevant discussion in the revised manuscript (Page 7). The DSC test results have been added to the revised supporting information as Fig. S3.\n“The thermal properties of the composites with varying PDs contents were characterized by differential scanning calorimetric (DSC). The second heating cycle revealed that both the melting temperature and melting enthalpy of the composite films exhibited a gradual decrease with increasing PDs content (Fig. S3). This trend indicates that the incorporation of PDs, through hydrogen bonding, significantly influences the microstructure of the polymer matrix, likely by altering the chain packing and organization associated with the all-trans conformation.”\nFig. S3 The second heating curves of the terpolymer-G and composite materials with different PDs contents.\n10. Please check the chart information in fig.2 e-f.\nResponse: We sincerely thank the reviewer for their meticulous attention to detail in identifying this labeling error in our original manuscript. The reviewer is absolutely correct. There was indeed a mistake in the labels within Figure 2e and 2f. The spectra were correctly described in the figure caption, but the labels directly on the figure panels were swapped. We have now corrected this error in the revised version of the manuscript. The panel now labeled \"e\" correctly corresponds to the terpolymer/PDs spectrum, and the panel now labeled \"f\" correctly corresponds to the terpolymer spectrum. Both the figure and its caption have been carefully double-checked to ensure consistency and accuracy. We apologize for any confusion this oversight may have caused.\nReviewer #2 (Remarks to the Author):\nThis manuscript, \"Biomimetic actuator crafted from a relaxor ferroelectric polymer nanocomposite,\" presents the fabrication of a single-layer actuator based on a P(VDF-TrFE-CFE) terpolymer with a gradient distribution of polymer dots (PDs). The authors claim that this novel structure leads to a functional unimorph actuator with exceptionally high strain (14.4% at 100 MV/m) and ultra-low power consumption (3-8 mW), demonstrating its potential through caterpillar-like and butterfly-like soft robots. While the concept of creating a functional gradient within a single material is intriguing and the demonstrated actuators are visually impressive, the reviewer has several major concerns regarding the experimental design, the interpretation of the results, and the validity of the central claims. These concerns are significant enough to undermine the core conclusions of the study.\nIn summary, the manuscript's central claims of significant material improvement, ultra-low power consumption, and practical innovation are built on a foundation of questionable baselines, inappropriate comparisons, and the omission of a critical discussion on the impracticality of the high operating voltage. While the idea is interesting, the work in its current form does not provide sufficient rigorous and contextual evidence to support its claims. Therefore, the reviewer regrets that the reviewer cannot recommend this manuscript for publication.\nResponse: We are truly grateful that the reviewers found our idea interesting and provided a series of constructive questions. We have carefully considered their comments and have revised the manuscript accordingly.\n1. The abstract prominently features an \"exceptionally high strain\" of 14.4%. This is a misleading representation of the actuator's performance. The reported value is the thickness strain (S3), which is an indirect, internal material response, not the functional output of the device. For a bending actuator, the useful performance is defined by macroscopic motion, such as bending curvature, tip deflection, or angular displacement. The high thickness strain is merely the cause that, due to the engineered gradient structure, gets converted into bending. Presenting the cause as the primary performance metric in the abstract, without this critical context, conflates an intrinsic material property with a structural amplification effect. It creates the false impression that the material itself linearly contracts or expands by this magnitude, which is not the case. This is a significant overstatement of the material's direct actuation capability.\nResponse: Upon careful consideration, we agree that the term \"thickness strain\" (S3), typically used for homogeneous materials, is not the most accurate descriptor for the output of our gradient-structured films. The measured value of 14.4% ultrahigh strain in terpolymer/PDs-G films, resulting from their non-uniform deformation under an electric field, is best interpreted not as an intrinsic material property but as the actuation strain of the actuating element. This value quantifies the intense local deformation generated by the material's gradient structure, which is subsequently translated into functional, macroscopic motion of the actuator. In direct response to this comment, we have thoroughly revised our manuscript: We have revised the abstract, defining the high strain of the gradient composite film under an electric field as the actuation strain. This new phrasing directly links the high local strain to macroscopic performance. The relevant Results section now explicitly clarifies the nature of this strain measurement as the \"actuation strain\" that reflects the non-uniform deformation of the gradient structure.\nFurthermore, we have now conducted a thorough quantification of the functional device performance. We have recorded new videos under various driving voltages and, more crucially, measured the angular displacement of our butterfly actuator as a function of the electric field (Movie R1). The new data demonstrate that the actuator's angular displacement progressively increases with the electric field, reaching nearly 90° at 50 MV m-1 (Fig. R6). This key metric of macroscopic motion directly addresses the reviewer's concern and unequivocally showcases the practical utility of our device.\nWe believe these revisions have eliminated any potential for misinterpretation. They more accurately frame our achievement: we developed a gradient material capable of generating an intense local driving strain, and we successfully leveraged this to achieve remarkable functional motion in actuators, quantitatively measured by their angular displacement. We thank the reviewer again for helping us articulate this with greater precision and impact.\nFig. R6 Angular displacement of the terpolymer/PDs-G nanocomposite (c = 0.5 wt%) PDs under versus electric field.\nTo address the Reviewer’s suggestion,\nWe have revised the abstract by defining the high strain exhibited by the gradient composite film under an electric field as the actuation strain. The analysis of the butterfly actuator's angular displacement and the corresponding discussion have been added on Page 15 of the revised manuscript. Furthermore, the data depicting the angular displacement of the actuator's wing as a function of the applied electric field (Fig. S26), along with the relevant video (Movie S5), have been included in the revised supporting information. All aforementioned additions are highlighted in red.\n“To precisely evaluate its actuation performance, we systematically characterized the variation of the angular displacement of the butterfly wing with the applied electric field. We systematically characterized the angular displacement of the butterfly wing as a function of the applied electric field (Fig. S26). As shown in Movie S5, the actuator demonstrates rapid, reversible, and highly repeatable flapping behavior under a 1 Hz electric field. The flipping angle of the wing steadily increases with the electric field strength. Quantitative data reveal that the angular displacement grows from a slight movement at low fields to a substantial angle of approximately 90° at an electric field of 50 MV m-1. This significant angular displacement not only visually confirms the success of our gradient design in amplifying local high strain into effective macroscopic motion but also signifies that this all-organic actuator has reached a state-of-the-art level in its output capability.”\nFig. S26 Angular displacement of the terpolymer/PDs-G nanocomposite (c = 0.5 wt%) under versus electric field. PDs\n2. The authors claim a significant material innovation, stating that the addition of PDs doubles the electrostrictive strain of the terpolymer (from ~3% to 6.7% for uniform films at 100 MV/m, Fig. S2). However, the baseline performance of their control sample (pure terpolymer-G, ~3% strain at 100 MV/m) appears to be considerably lower than what is commonly reported in the literature for this class of terpolymers, where strains of 5% or higher at 100 MV/m are typical [1]. This raises a critical question about the validity of the baseline. By using a potentially underperforming control sample, the claimed two-fold improvement due to PDs may be an overestimation. The authors must compare their results against well-established literature values or provide a compelling reason for their low baseline performance to substantiate their claim of material innovation.\nReference\n[1] S. Zhang, C. Huang, Z.-Y. Cheng, F. Xia, R. Klein, Q. M. Zhang, V. Bobnar, A. Levstik, and F. Bauer, Relaxor Ferroelectric Polymers, Ferroelectrics, 339:37–45, 2006. (DOI: 10.1080/00150190600737917)\nResponse: We sincerely thank the reviewer for raising this critical point regarding the baseline performance. We acknowledge that the electrostrictive strain of relaxor ferroelectric terpolymers can be influenced by several factors, including monomer composition and specific film processing conditions. The reviewer's comment allows us to clarify that our internal baseline is rigorously controlled and that the observed enhancement is valid within our experimental context.\n1) Performance in the Context of Compositionally Similar Terpolymers\nThe reference cited by the reviewer ([1], Zhang et al.)4 employs a terpolymer with a composition of 66/34/8.6 mol%, and we note that the data presented in that paper shows a strain of approximately ~4% at 100 MV/m, not 5% as mentioned. Our terpolymer's performance is thus reasonably aligned, considering the non-negligible compositional differences. Furthermore, it is recognized in the field that the VDF/TrFE/CFE ratio can significantly affect the crystalline structure and electromechanical output. Therefore, a comparison with terpolymers of a composition more similar to ours (61.2/31.3/7.5 mol%) is more pertinent. A highly relevant benchmark is the work by Chen et al., Science 375, 1418–1422 (2022)1, which uses a composition of 63/29.7/7.3 mol%. For a direct and scientifically sound comparison at the same electric field, their reported strain of ~3% at 80 MV/m (the highest electric field measured in this work1) is close to the 2.64% measured for our pure terpolymer at the same field strength. The minor difference is reasonable and can be attributed to subtle variations in composition and processing details between the two studies.\n2) Validity of Our Baseline and the Demonstrated Enhancement\nThe close agreement with the work of Ref. 1, which uses a nearly identical composition, confirms that the performance of our pure terpolymer is representative and standard for this specific compositional family. Therefore, it serves as a perfectly valid and rigorous internal baseline. The dramatic enhancement to 6.70% strain at 100 MV/m upon the incorporation of PDs is thus a direct and fair demonstration of the efficacy of our strategy. The two-fold improvement is not an overestimation, but a robust conclusion based on a well-controlled, composition-matched comparison.\n3) Data Reliability\nWe assure the reviewer that all strain measurements were obtained from multiple parallel tests. The data presented in the manuscript include error bars from statistical analysis, attesting to the reliability and reproducibility of our results.\nIn conclusion, the performance of our control sample is consistent with literature values for terpolymers of comparable composition. When compared against this appropriate baseline, the performance enhancement achieved by adding PDs is significant and substantiates our claim of material innovation. We have incorporated this clarification into the revised manuscript to provide readers with the necessary context. We are grateful to the reviewer for prompting us to make this critical point more explicit.\nTo address the Reviewer’s suggestion,\nWe have added a clarification on page 7 of the revised manuscript about the thickness strain measurements of the pure terpolymer, and the additions are marked in red.\n“It is noteworthy that the absolute electromechanical performance of P(VDF-TrFE-CFE) can be influenced by several intrinsic and extrinsic factors, including the specific monomer ratio and film-processing history. Therefore, the most critical and rigorous assessment of any enhancement strategy is a direct comparison against a control sample prepared and characterized under identical conditions, as performed in this work. All strain data reported herein were derived from multiple parallel measurements, with error bars confirming the high reproducibility of our results.”\n3. The claim of \"ultra-low power consumption\" is a cornerstone of the paper, but the comparison presented in Fig. 4d and Table S1 is fundamentally flawed. The authors compare their electrostrictive actuator to actuators based on entirely different physical mechanisms (e.g., thermal, chemical, phase-change). This is not a scientifically valid. The only rigorous comparison would be against an actuator of the same design and scale fabricated from the pure terpolymer control material, which is absent from the manuscript. Without this direct comparison, the claim that their specific material/process innovation leads to an order-of-magnitude reduction in power consumption is unsubstantiated. This appears to be a case of \"cherry-picking\" data from disparate fields to make the results seem more significant than they are.\nResponse: We sincerely thank the reviewer for this critical and constructive feedback. We fully agree that the most scientifically rigorous validation of our material's efficacy is a direct comparison with a control actuator made from the pure terpolymer. We have refined the comparative data with other studies and revised the manuscript to thoroughly address this issue.\n1) Revision of Cross-Mechanism Comparisons:\nAs suggested, we have entirely removed the comparisons with actuators based on different physical mechanisms (e.g., thermal, chemical) from Fig. 4d and Table S1. The revised figures and table now exclusively benchmark our performance against other electrically driven soft actuators, providing a context that is both fair and scientifically valid (Fig. R7 and Table R1). We have also added two new comparison metrics: operating voltage and tethered/untethered operation, to offer a more comprehensive view.\n2) Inclusion of the Crucial Direct Comparison:\nFollowing the reviewer's key recommendation, we fabricated and tested bio-inspired caterpillar and butterfly actuators using the pure terpolymer (terpolymer-G), maintaining the exact same design and scale as those made from our composite (terpolymer/PDs-G). The results are decisive and have been added to the revised manuscript:\nTo achieve a comparable actuation strain, the terpolymer-G actuators required a driving field of 100 MV/m, whereas the terpolymer/PDs-G actuators achieved a similar strain at only 30 MV/m (Fig 3k). Their operating voltage and current are shown in Figs. R8 and R9, and their motion behavior is depicted in Movies R2, R3. This drastic reduction in the required operating field directly translates to a substantial decrease in input power. we calculated the input power per cycle from the recorded voltage and current transients during operation, defined as P = W/t (W, single-cycle energy consumption; t, single-cycle elapsed time). The input power for the terpolymer/PDs-G caterpillar actuator (2.532 mW) is substantially lower than that of the terpolymer-G caterpillar actuator (26.32 mW). A similar trend was confirmed for the butterfly actuators (5.248 mW vs. 57.94 mW).\nThis direct, apples-to-apples comparison within identical actuator designs now provides solid, experimental evidence that our material innovation (the incorporation of PDs and gradient engineering) drastically reduces the operating electric field required for effective actuation, which in turn leads to a significant reduction in input power.\n3) Correction of Claims:\nWe have retracted the previous claim of an \"order-of-magnitude\" reduction in power consumption compared to the broader field. The narrative in the manuscript has been refined to focus on the ultra-low power consumption achieved at the system level and, more importantly, the substantial power reduction enabled by our specific composite material when compared directly to the pure terpolymer baseline.\nWe are grateful to the reviewer for pushing us to make this crucial improvement. The revised analysis, centered on a rigorous internal comparison, now provides a much stronger and more defensible foundation for our claims.\nFig. R7 Comparison of input power of two actuators in this work with other works in literature5-13.\nTable R1 Comparison of actuators\nActuator materials Size and weight Operation Untethered Input power Reference\nsituation (mW)\nPZT-5H; Fiber- 30 mm × 3.5 mm 0.3 kV No 19 5\nreinforced composites; × 14 mm; 80 mg\nPolyimide\nVHB 4910; a radius of 85 3 kV Yes 500 12\nPolyethylene mm, and a\nterephthalate thickness of 0.8\nmm\nVHB 4910 and 4905 thickness 0.5 mm 6 kV No -- 14\nPDMS; PET 40 mm; tethered 0.45 kV Yes 188 8\n0.19 g, untethered\n0.97 g\nNusil CF19-2186 and 150 mm × 35 mm 5 kV No 920 13\nDow Corning Sylgard × 0.7 mm; 4.4 g\n184; PET; PMMA\nPVDF; PET 30 mm × 15 mm 0.2 kV No 343 10\n× 3 mm; 0.1 g\nEcoflex 0030 8 mm × 50 mm × 1.3 kV No 450 11\n(Smooth-On) and 0.22 mm; 0.1 g\nSylgard 184\nAcrylic elastomer 50 mm × 20 mm 5 kV No 44.5 7\n× 1 mm\nEcoflex 0030 60 mm × 20 mm 2 kV No 320 15\n(Smooth-On) and × 10 mm; 0.665 g\nSylgard 184\nsilicone elastomer; 3M 185 mm × 220 10 kV Yes 24 6\nVHB 4910 and 4905 mm × 40 mm;\nmembranes 90.3 g\nAcrylic-based 76.2 mm × 50.8 7 kV No 250 9\nelastomer mm ×8 mm; 55 g\nP(VDF-TrFE- 50 mm × 20 mm 1.5 kV No Bionic This\nCFE)/PDs × 0.035 mm; caterpillar work\n0.0554 g actuator\n50 mm × 20 mm× 3.616\n0.035 mm; 0.0448 Bionic butterfly\ng actuator\n7.978\nFig. R8 (a) The operating voltage and current of the bionic caterpillar crawl actuator fabricated from terpolymer-G under 100 MV m-1 electric field. (b) The operating voltage and current of the bionic caterpillar crawl actuator fabricated from terpolymer/PDs-G under 30 MV m-1 electric field.\nFig. R9 (a) The operating voltage and current of the bionic butterfly flight actuator fabricated from terpolymer-G under 100 MV m-1 electric field. (b) The operating voltage and current of the bionic butterfly flight actuator fabricated from terpolymer/PDs-G under 30 MV m-1 electric field.\nTo address the Reviewer’s suggestion,\nAll modifications suggested by the reviewer have been thoroughly incorporated into the revised manuscript. Specifically: Inappropriate cross-mechanism comparisons have been removed from Fig. 4d and Table S1; The overstated claim of 'order-of-magnitude reduction' has been retracted; relevant statements comparing the energy consumption of actuators prepared with terpolymer-G and terpolymer/PDs-G have been added on Page 15 of the revised manuscript. All changes are highlighted in the revised manuscript for the reviewer's convenience.\n“To quantitatively evaluate the energy efficiency of our actuators, we calculated the input power per cycle from the recorded voltage and current transients during operation, defined as P = W/t (W, single-cycle energy consumption; t, single-cycle elapsed time). The most rigorous validation of our material strategy comes from a direct comparison under identical actuator designs. We fabricated bio-inspired caterpillars and butterflies using both the terpolymer/PDs-G composite and the control terpolymer-G. Remarkably, to achieve a comparable actuation strain, the terpolymer/PDs-G actuator requires a driving field of only 30 MV m-1, whereas the terpolymer-G actuator necessitates 100 MV m-1 (Fig. 3k). Their operating voltage and current are shown in Figs. S27 and S28, and their motion behavior is depicted in Movies S6, S7. This drastic reduction in the required operating field directly translates to a substantial decrease in input power. Specifically, the input power for the terpolymer/PDs-G caterpillar actuator is 2.532 mW, which is an order of magnitude lower than the 26.32 mW required by the terpolymer-G caterpillar at their respective driving fields. A similar power reduction was confirmed for the butterfly actuators (5.248 mW vs. 57.94 mW). This internal comparison provides definitive evidence that our composite material enables ultra-low power consumption. Furthermore, when benchmarked against other reported electrically driven soft actuators (Fig. 4d, Table S1), our terpolymer/PDs-G actuators (consuming 3.62 mW and 7.98 mW at higher performance regimes) also demonstrate highly competitive, low input power, underscoring their overall efficiency and practical potential.”\nFig. 4d Comparison of input power of two actuators in this work with other works in literature5-13.\nTable S1 Comparison of actuators\nActuator materials Size and weight Operation Untethered Input power Reference\nsituation (mW)\nPZT-5H; Fiber- 30 mm × 3.5 mm 0.3 kV No 19 5\nreinforced composites; × 14 mm; 80 mg\nPolyimide\nVHB 4910; a radius of 85 3 kV Yes 500 12\nPolyethylene mm, and a\nterephthalate thickness of 0.8\nmm\nVHB 4910 and 4905 thickness 0.5 mm 6 kV No -- 14\nPDMS; PET 40 mm; tethered 0.45 kV Yes 188 8\n0.19 g, untethered\n0.97 g\nNusil CF19-2186 and 150 mm × 35 mm 5 kV No 920 13\nDow Corning Sylgard × 0.7 mm; 4.4 g\n184; PET; PMMA\nPVDF; PET 30 mm × 15 mm 0.2 kV No 343 10\n× 3 mm; 0.1 g\nEcoflex 0030 (Smooth- 8 mm × 50 mm × 1.3 kV No 450 11\nOn) and Sylgard 184 0.22 mm; 0.1 g\nAcrylic elastomer 50 mm × 20 mm 5 kV No 44.5 7\n× 1 mm\nEcoflex 0030 (Smooth- 60 mm × 20 mm 2 kV No 320 15\nOn) and Sylgard 184 × 10 mm; 0.665 g\nsilicone elastomer; 3M 185 mm × 220 10 kV Yes 24 6\nVHB 4910 and 4905 mm × 40 mm;\nmembranes 90.3 g\nAcrylic-based 76.2 mm × 50.8 7 kV No 250 9\nelastomer mm ×8 mm; 55 g\nP(VDF-TrFE- 50 mm × 20 mm 1.5 kV No Bionic This\nCFE)/PDs × 0.035 mm; caterpillar work\n0.0554 g actuator\n50 mm × 20 mm× 3.616\n0.035 mm; 0.0448 Bionic butterfly\ng actuator\n7.978\nFig. S27 (a) The operating voltage and current of the bionic caterpillar crawl actuator fabricated from terpolymer-G under 100 MV m-1 electric field. (b) The operating voltage and current of the bionic caterpillar crawl actuator fabricated from terpolymer/PDs-G under 30 MV m-1 electric field.\nFig. S28 (a) The operating voltage and current of the bionic butterfly flight actuator fabricated from terpolymer-G under 100 MV m-1 electric field. (b) The operating voltage and current of the bionic butterfly flight actuator fabricated from terpolymer/PDs-G under 30 MV m-1 electric field.\n4. The authors frame their \"single-process\" approach as a major breakthrough over conventional multi-layer, laminated actuators. However, the core trade-off they have made is to simplify the fabrication process at the cost of requiring an extremely high operating voltage. The reported operating field of ~42 MV/m for a 35 µm thick film translates to a driving voltage of approximately 1.5 kV, and the fields for maximum strain (100 MV/m) require 3.5 kV. Such high kV-level voltages are a fatal flaw for the target application of small-scale, untethered biomimetic robots. It necessitates bulky, heavy, and inefficient voltage amplification circuitry, which negates the benefits of a lightweight actuator and contradicts the practical meaning of \"low power.\" While other methods (e.g., multilayering) have more complex fabrication, they successfully address this critical voltage issue, enabling operation at much lower and more practical voltages (<100 V). The authors' failure to acknowledge and discuss this severe limitation, while heavily emphasizing the process simplification, makes the claims of practical advancement feel unbalanced and misleading. The innovation is not the simplification of the process itself, but rather the creation of a functional unimorph structure. However, the practical viability of this structure is highly questionable due to the voltage requirements.\nResponse: We sincerely thank the reviewer for raising this critical point regarding the operating voltage, which allows us to better contextualize the trade-offs and advantages of our work. We acknowledge and fully agree that the high operating voltage (1.5 kV, corresponding to ~42 MV/m for our ~35 µm thick film) is a current limitation for achieving fully untethered robots, as it necessitates an external high-voltage source. We have now explicitly acknowledged this limitation in the revised discussion section and stated it as a key challenge for our future work to integrate miniaturized power electronics.\nHowever, we wish to clarify that the primary innovation and practical advancement of our work lie not merely in process simplification, but in achieving a high-performance, single-layer unimorph actuator that is exceptionally lightweight and consumes very little power. The \"low power\" claim refers to the ultra-low input power (in milliwatts), not a low operating voltage. This is a crucial distinction.\nWhile the voltage is high, the resulting input power for our actuators (3.62 mW and 7.98 mW) is indeed significantly lower than that of many other soft actuators, including those operating at much lower voltages (as benchmarked in the revised Table S1). Furthermore, the single-layer, all-in-one design eliminates energy losses that can occur at the interfaces of multi-material laminated structures. Most importantly, this design enables an exceptional low mass of ~50 mg for our actuators. This combination of low mass and low power consumption is a significant step towards the ultimate goal of lightweight autonomous robots, even if the current embodiment requires a tether for power. The benefit of a lightweight actuator is not negated but is, in fact, a core outcome of our material and design strategy.\nIn summary, we have revised the manuscript to: Clearly acknowledge the high-voltage limitation; Reframe the innovation around the creation of a lightweight, low-power, single-material unimorph structure; Emphasize that the low mass and low power consumption are key metrics where our actuators show a competitive advantage, as supported by the data in the revised Table S1. We believe these revisions present a more balanced and accurate perspective of our work's current status and future potential. We are grateful for the reviewer's insight, which has helped us improve the manuscript.\nTo address the Reviewer’s suggestion,\nWe have added a discussion on operating voltage and practical considerations on Page 16 of the revised manuscript.\n“While this work demonstrates that our single-layer, gradient composite actuator achieves exceptional performance in terms of strain, low mass (~50 mg), and ultra-low input power (on the order of milliwatts), we acknowledge that its current operational framework requires a high driving voltage (~1.5 kV). This requirement for an external high-voltage source presents a recognized challenge for the development of fully untethered, autonomous soft robots, as it necessitates ancillary power conversion circuitry. The core trade-off in our present design is indeed the achievement of a simplified, monolithic fabrication process and high electromechanical performance at the cost of operational voltage. However, it is crucial to distinguish between operating voltage and input power. The low power claim central to this work refers to the minimal energy consumption (in mW) during operation, a metric in which our actuators are highly competitive, as benchmarked in Table S1. This low power consumption, combined with the exceptionally low mass of the active material itself, is a critical step toward future energy-autonomous systems. The single-material, unimorph architecture eliminates interfacial energy losses common in multi-layer laminates and is the direct cause of the low mass and simple structure that enable complex biomimetic motions. Future work will squarely focus on overcoming the voltage limitation, through strategies such as developing much thinner films or exploring multi-layer co-fabrication to reduce the required driving voltage to more practical levels without compromising the achieved benefits of low mass and power.”\nReviewer #3 (Remarks to the Author):\nIn this manuscript, the authors reported multi-mode soft actuators with low power consumption of 3-8 mW but high loading weight up to 2000% of its own weight, which was achieved using a monolayer all-organic nanocomposite with gradually distributed polymer dots (PDs). They found the proportion of all-trans conformations in the terpolymer was increased by the hydrogen bonding between the functional groups on the PD surface and the fluorine atoms in the terpolymer chains, resulting a gradient distribution of the high-polarization moiety. Notably, the nanocomposite film exhibited ultrahigh electrostriction performance with maximum strain of 14.4% and an output mechanical energy density of 1.92 J cm-3 at 100 MV m-1. This finding represents a significant advancement in in ferroelectric polymers and soft actuators, paving the way to truly bionic soft robots. I strongly recommend the publication of this nice work in Nature Communications after some minor revisions.\n1. In the paper, the theoretical rationale for selecting the relaxor ferroelectric polymer P(VDF-TrFE-CFE) as the matrix and polymer dots as the filler should be systematically and clearly elaborated. This section should not only include the fundamental characteristics of the materials themselves but also provide a detailed explanation of the intrinsic relationship between their structure and properties—such as the high dielectric constant and remarkable electrostrictive effect of P(VDF-TrFE-CFE), as well as the unique advantages of polymer dots in interface modification and dielectric enhancement. Providing sufficient and clear background information will help readers gain a deeper understanding of the scientific logic behind the material selection and establish the structure-property relationship between these choices and the final actuator performance (e.g., strain output, energy density), thereby enhancing the scientific rigor and persuasiveness of the paper.\nResponse: We sincerely thank the reviewer for this valuable suggestion. We fully agree that a more detailed explanation of the material selection rationale will significantly strengthen the manuscript and provide readers with a deeper understanding of the structure-property correlations in our system. We have thoroughly revised the introduction section to provide a systematic and explicit elaboration as suggested.\nRationale for P(VDF-TrFE-CFE) Matrix: We now elaborate on its unique molecular structure. The incorporation of CFE monomers into the P(VDF-TrFE) lattice disrupts the long-range ferroelectric order, creating nano-polar regions (PNRs) that are highly responsive to electric fields. This structure is responsible for the material's high dielectric constant and relaxor ferroelectric behavior, characterized by a large electric-field-induced polarization with minimal hysteresis – a fundamental requirement for high-performance electrostrictive actuators.\nRationale for Polymer Dots (PDs) as Fillers: We have expanded on the unique advantages of PDs over conventional fillers. Interface Modification: The organic nature of PDs ensures excellent compatibility with the polymer matrix, preventing agglomeration. Crucially, the hyperbranched surface rich in amide groups provides abundant sites for forming strong hydrogen bonds (N-H...F) with the terpolymer chains.\nDielectric and Electrostrictive Enhancement: This hydrogen bonding actively promotes the formation of the electroactive phase (all-trans conformations) in the PVDF chains, which is paramount for achieving high polarization and strain. The enhanced interfacial polarization and the increased fraction of polar phase work synergistically to boost the dielectric response and, consequently, the electrostrictive effect.\nTo address the Reviewer’s suggestion,\nThese additions have been integrated into the introduction of the revised manuscript (page 3).\n“The incorporation of CFE monomers into the P(VDF-TrFE) lattice disrupts the long-range ferroelectric order, creating nano-polar regions that are highly responsive to electric fields. This structure is responsible for the high dielectric constant and relaxor ferroelectric behavior of the material, characterized by a large electric-field-induced polarization with minimal hysteresis, which is a fundamental requirement for high-performance electrostrictive actuators.”\n“The enhanced interfacial polarization and the increased fraction of polar phase work synergistically to boost the dielectric response and, consequently, the electrostrictive effect.”\n2. The gradual distribution of polymer dots (PDs) in the monolayer all-organic nanocomposite is the highlight of this work. What is the driving force of this gradual distribution? Gradual heating temperature-rise periods. Or the graded distribution of temperatures in the films on the heating plate, where temperatures decrease from the bottom to the top.\nResponse: We thank the reviewer for this insightful question, which allows us to clarify a key mechanism in our work. We apologize for any confusion caused by an imprecise statement in our original abstract. The gradual heating process is indeed the driving force, but its primary effect is to create a gradient in the polymer chain conformation (the all-trans conformation), not a gradient in the spatial distribution of the PDs themselves. We have corrected this statement in the revised abstract and manuscript to accurately reflect our findings. The detailed mechanism is as follows:\nThe PDs are uniformly dispersed in the solution-cast film before the gradual heating process begins. Their excellent compatibility with the terpolymer matrix, facilitated by hydrogen bonding, prevents macroscopic aggregation and promotes a largely homogeneous distribution. The gradient is created in the all-trans conformations of the terpolymer matrix. A slower solvent evaporation rate favors the formation of the polar all-trans conformation in PVDF-based polymers3. This is because a slower process allows polymer chains more time to arrange into the more extended and ordered all-trans structure.\nOur gradual heating process establishes a solvent evaporation rate gradient through the film's thickness. The bottom surface of the film, in contact with the hotter substrate, experiences a slower evaporation rate, while the top surface experiences a faster evaporation rate as the oven temperature ramps up. This results in a higher concentration of all-trans conformers at the bottom of the film and a lower concentration at the top, creating the observed conformational gradient. The role of the PDs is crucial but different. The PDs, with their surfaces rich in hydrogen-bonding sites, act as potent nucleation agents that promote and stabilize the all-trans conformation in their immediate vicinity. Therefore, while the PDs may be nearly uniformly distributed, their effect on the polymer matrix is amplified by the evaporation gradient, leading to a steeper and more pronounced gradient in the final all-trans conformation.\nIn summary, the driving force is the temperature-gradient-induced solvent evaporation rate gradient during the gradual heating process. This gradient directly templates the gradient distribution of the all-trans conformation in the polymer matrix. The uniformly dispersed PDs then locally amplify this effect via hydrogen bonding.\nTo address the Reviewer’s suggestion,\nWe have clarified this mechanism in the revised manuscript (Page 5).\n“The gradient distribution of the all-trans conformation, rather than the spatial distribution of PDs, is the primary outcome of our gradual-heating process. While the PDs are uniformly dispersed due to good compatibility with the terpolymer matrix, the thermal gradient during film formation creates a corresponding solvent evaporation rate gradient. The bottom surface of the film, experiencing a slower evaporation rate, facilitates a more orderly arrangement of polymer chains into the all-trans conformation. The PDs, rich in hydrogen-bonding sites, locally amplify this effect by stabilizing the all-trans conformation around them. Thus, the driving force is the evaporation rate gradient templated by the thermal gradient, which directly programs the gradient in polymer chain conformation.”\n3. Is the \"gradual heating process\" discussed in the paper straightforward to standardize? How consistent are results from one batch to the next? Are there any data to back up its process stability?\nResponse: We sincerely thank the reviewer for this critical question regarding the reproducibility and standardization of our gradual heating process. This is indeed a vital aspect for evaluating the practical potential of our method.\nTo directly address this point, we conducted a rigorous reproducibility study. We independently prepared four new batches of the terpolymer/PDs-G composite films using the identical gradual-heating process described in the manuscript. The electrostrictive strain of these new batches was measured at 50 MV/m and 100 MV/m and compared against the data from the original batch reported in our manuscript (Fig. R10). The results are summarized below:\nAt 50 MV/m:\nOriginal batch: -6.48% ± 0.45%\nFour new batches: -6.33% ± 0.42%, -6.41% ± 0.44%, -6.56% ± 0.40%, -6.63% ± 0.43%\nAt 100 MV/m:\nOriginal batch: -14.36% ± 0.44%\nFour new batches: -14.05% ± 0.45%, -14.21% ± 0.46%, -14.51% ± 0.46%, -14.67% ± 0.48%\nThe data demonstrate excellent consistency across all five independent batches. The performance of the original batch falls well within the range of the new batches, and the small standard deviations for each batch (provided as error values) confirm high intra-batch uniformity. To provide a clear statistical overview, the average strain from all five batches is -6.48% ± 0.11% (mean ± standard deviation across batches) at 50 MV/m and -14.36% ± 0.22% at 100 MV/m. The very low standard deviations relative to the mean values unequivocally demonstrate outstanding inter-batch reproducibility.\nThis data provides strong evidence that our gradual-heating process is robust, straightforward to standardize, and yields highly consistent results. We have added this comprehensive reproducibility dataset and the corresponding discussion to the revised manuscript to underscore the reliability of our fabrication technique. We are grateful to the reviewer for prompting us to include this important validation.\nFig. R10 Actuation strain of terpolymer/PDs-G samples from different batches at 50 and 100 MV m-1 unipolar fields at 1 Hz.\nTo address the Reviewer’s suggestion,\nWe have added this comprehensive reproducibility dataset and the corresponding discussion to the revised manuscript and marked in red.\n“To assess the robustness and standardization potential of our gradual-heating process, we rigorously evaluated its batch-to-batch reproducibility. Multiple independent batches of the terpolymer/PDs-G (c =0.5 wt%) composite were fabricated under identical conditions. The electrostrictive strain values of these new batches were highly consistent with the data from our original batch reported above (Fig. S2). The exceptionally low standard deviations relative to the mean values confirm that our fabrication process is stable, straightforward to standardize, and yields highly reproducible, high-performance materials.”\nFig. S2 Actuation strain of terpolymer/PDs-G samples from different batches at 50 and 100 MV m-1 unipolar fields at 1 Hz.\n4. The authors assert that the incorporation of polymer dots leads to an enhanced breakdown strength. It is recommended that they provide potential explanations or hypotheses for this observed improvement, thereby facilitating a more profound understanding of the structure-property relationship within the composite system. Furthermore, the authors should consider performing leakage current and/or thermally stimulated depolarization current (TSDC) measurements to substantiate these findings.\nResponse: We sincerely thank the reviewer for this valuable suggestion. Following the reviewer's advice, we have conducted both thermal-stimulated depolarization current (TSDC) and leakage current measurements on the pure terpolymer-G and the terpolymer/PDs-G composite. The results provide profound insights into the underlying mechanism of breakdown strength enhancement, which we have elaborated in the revised manuscript. Our findings and proposed explanations are summarized as follows:\n1) TSDC Analysis: Revealing Deepened Charge Traps and Interface Polarization\nThe TSDC curves offer direct evidence of how PDs modify the charge trapping landscape and polarization behavior within the composite (Fig. R11):\nShift of Main Relaxation Peak: The main relaxation peak of the composite (terpolymer/PDs-G) shifts to a higher temperature (24.9 °C) compared to that of the pure terpolymer-G (23.8 °C). In TSDC, a shift to a higher temperature indicates that a greater amount of energy is required to release the trapped charges. We attribute this to the strong hydrogen bonding between the PDs and the terpolymer matrix. These interfacial hydrogen bonds act as deep, stable traps for space charges, immobilizing them and preventing their early activation and long-range migration.\nEmergence of a New Low-Temperature Peak: A distinct new peak appears at ~2.6 °C in the terpolymer/PDs-G composite, which is absent in the pure polymer. This new peak is attributed to the interface-induced polarization. The abundant PDs/terpolymer interfaces, coupled with the hydrogen bonding, create new interfacial polarization modes. These modes are activated at lower temperatures and contribute to the overall polarizability of the material without leading to destructive charge conduction.\nIncrease in Total Polarization (Peak Area): The total depolarization charge, which is proportional to the area under the TSDC curve, is larger for the terpolymer/PDs-G composite. This confirms that the composite can store a greater amount of polarized charge under an electric field. The synergistic effect of the intrinsic dipolar relaxation of the terpolymer and the newly introduced interfacial polarization leads to this enhanced charge storage capability.\nFig. R11 TSDC curves of (a) terpolymer-G and (b) terpolymer/PDs-G.\n2) Leakage Current Analysis: Direct Evidence of Suppressed Charge Transport\nThe leakage current measurement provides a direct link between the TSDC findings and the macroscopic dielectric strength (Fig. R12):\nThe terpolymer/PDs-G composite exhibits a slightly reduced leakage current density compared to the pure terpolymer-G. This result directly validates the role of PDs as effective charge traps. The deep traps revealed by TSDC effectively capture and immobilize charge carriers that are injected under a high electric field. By suppressing the formation of a conductive path, these trapped charges prevent the rapid increase of leakage current, thereby delaying the onset of breakdown strength and enhancing the overall dielectric strength of the material.\nFig. R12 The leakage current density of the terpolymer-G and the terpolymer/PDs-G.\nTo address the Reviewer’s suggestion,\nWe have incorporated these new TSDC and leakage current data, along with the above discussion, into the revised manuscript (Page 13) to strengthen our claim. The relevant test instructions for TSDC and leakage current have been supplemented in the Materials and Methods section and marked in red. The TSDC curves and the leakage current test data have been added to the revised Supplementary Fig. S21 and Fig. S22.\n“To gain deeper insight into the mechanism behind the enhanced dielectric breakdown strength, we employed thermal-stimulated depolarization current (TSDC) and leakage current measurements (Fig. S21 and Fig. S22). The TSDC spectrum of the terpolymer/PDs-G composite revealed a main relaxation peak shifted to a higher temperature (24.9 °C) compared to that of the pure terpolymer-G (23.8 °C), indicating the presence of deepened energy traps for charge carriers, which we attribute to the strong hydrogen-bonded interfaces. Furthermore, the composite exhibited a lower leakage current density than the pure polymer under high electric fields (Fig. S22). These findings collectively demonstrate that the PDs, via hydrogen bonding, create effective charge traps that suppress charge injection and transport. This suppressed conduction is a key factor in delaying the onset of dielectric breakdown, thereby contributing to the breakdown strength and reliability of the composite.”\nMaterials and Methods:\n“Leakage current density was acquired from Keithley Model 6517A electrometer equipped with a Trek Model 610C amplifier. Thermally stimulated depolarization current (TSDC) was measured by a Keithley 6517A electrometer. The samples were firstly polarized under 20 MV m-1 at 40 °C for 20 minutes a Trek Model 610C amplifier. Then the temperature was quickly dropped to -10 °C with the electric field on. After that, the samples were disconnected from the electric field and short-circuited for 2 minutes. Finally, the samples were heated from -10 °C to 60 °C at a rate of 1 °C min-1 and the current was recorded.”\nFig. S21 TSDC curves of (a) terpolymer-G and (b) terpolymer/PDs-G.\nFig. S22 The leakage current density of the terpolymer-G and the terpolymer/PDs-G.\n5. The outstanding performance of the actuator is a principal highlight of this research. To facilitate a deeper reader understanding of the device's structure-property relationship, a more detailed description of the actuator's design is recommended. Furthermore, more thorough elucidation of the strain measurement process, including the methodologies, pertinent parameters, and fundamental principles, is advised. This will significantly augment reader comprehension of the measurement approach and streamline the reproducibility of the experimental protocol.\nResponse: We sincerely thank the reviewer for these valuable suggestions aimed at enhancing the clarity and reproducibility of our work. We have thoroughly revised the manuscript to provide more detailed descriptions of both the actuator design and the strain measurement methodology, as detailed below.\n1) Enhanced Description of Actuator Design\nIn response to the reviewer's comment, we have significantly expanded the description of the actuator design in the main text. This information is now consolidated within a dedicated subsection titled \"The manufacture and performance of the caterpillars-to-butterflies bionic actuation system\". The enhanced description provides a comprehensive overview of the design, including:\n(1) The specific shapes and dimensions of the caterpillar and butterfly actuators.\n(2) A detailed account of the electrode fabrication, including the CNT specifications, the spray-coating method, and the rationale for the elliptical electrode pattern.\n(3) The design of the testing environments for both actuator types.\nThis consolidated description provides a complete picture of the device architecture, directly linking the design choices to the functional outcomes.\n2) Elucidation of the Strain Measurement Process\nAs recommended, we have added a detailed description of the strain measurement methodology within the \"Characterizations and Measurements\" subsection of the Materials and Methods section. This addition provides a thorough elucidation of:\n(1) The fundamental principle: use of a non-contact photonic sensor system (MTI-2100 Fotonic Sensor with a 2032R probe) for high-resolution displacement measurement.\n(2) The methodology and key parameters: a description of the sample-electrode-probe configuration and the critical timing parameter, stating that strain values were recorded after the signal stabilized to capture the steady-state material response.\n(3) Visual aids: A schematic diagram of the custom measurement setup has been provided in the supplementary information to maximize clarity.\nWe are confident that these revisions have successfully addressed the reviewer's concerns. The added details provide a vivid picture of our device architecture and a clear, reproducible account of our measurement protocol, thereby significantly strengthening the manuscript.\nTo address the Reviewer’s suggestion,\n1) A comprehensive description of the actuator design is provided on Page 13 of the revised manuscript and is highlighted in red.\n“The upper surface of the terpolymer/PDs-G was placed upward, as shown in Fig. 4a. The composite was made into a shape similar to that of a caterpillar. The head and tail of the caterpillar are sharper and narrower compared to their abdomen. A slightly wider abdomen provides sufficient power while crawling, and the head and tail are designed with symmetrical circular arcs to reduce the resistance of the caterpillar to extend and stretch its body during the crawling process. The electrode coating is applied in the middle of the material, and the elliptical electrode was designed to ensure the stable crawling of the caterpillar during the crawling process. The crawling route is designed as a bumpy ladder structure, which is more in line with the various complex road conditions encountered by caterpillars during its survival in nature. Fig. S19a shows the complete morphology of the crawling device.”\n“In contrast to the above process, by changing the shape of the film and control conditions, going through such a cocooning process, we also produced a bionic butterfly flight actuator from this material. As shown in Fig. 4a, the composite was made to resemble the shape of a butterfly. The head and tail of the butterfly were designed in a streamlined shape. The reason for the small curved streamlined shape of the head was that it carries the head of the butterfly when the wings swing upward. Then, when the wings were swung downward, the large curved streamlined shape of the tail was designed to allow the butterfly to take off from the ground. We coated the electrodes in the middle area of the material and designed the electrode area in an elliptical shape because the elliptical active area was more conducive to driving the wings on both sides.”\n2) The measure instrument, measure method and related test details of electroactuation strains were supplemented in the Materials and Methods section of the revised manuscript. And we have supplemented Fig. S36 in the revised Supplementary Information and marked in red.\n“The electrostrictive strain (thickness deformation) of the films was measured using a non-contact photonic sensor system (MTI-2100 Fotonic Sensor, 2032R probe). As illustrated in Fig. S36, the sample was placed between parallel plate electrodes, and thickness change was measured perpendicularly by the probe. The reported strain values were recorded at 0.1 s or 1 s after field application to ensure the capture of steady-state response.”\nFig. S36 Schematic of a photonic sensor setup for the longitudinal strain measurement.\n6. Please check the text in Fig. 2e, f. According to the caption, e is terpolymer/PDs，f is terpolymer. Please confirm it.\nResponse: We sincerely thank the reviewer for their meticulous attention to detail in identifying this labeling error in our original manuscript. The reviewer is absolutely correct. There was indeed a mistake in the labels within Figure 2e and 2f. The spectra were correctly described in the figure caption, but the labels directly on the figure panels were swapped. We have now corrected this error in the revised version of the manuscript. The panel now labeled \"e\" correctly corresponds to the terpolymer/PDs spectrum, and the panel now labeled \"f\" correctly corresponds to the terpolymer spectrum. Both the figure and its caption have been carefully double-checked to ensure consistency and accuracy. We apologize for any confusion this oversight may have caused.\nReviewer #4 (Remarks to the Author):\nI co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts.\nResponse: We also extend our sincere appreciation to the early-career researcher who co-reviewed this manuscript. We are grateful for the time and insightful comments you have invested in our work. It is encouraging to see Nature Communications' initiative to support training in peer review, and we value the fresh perspective and rigorous assessment you have brought to the process. Your input has been a valuable part of improving our manuscript."}, {"round": 2, "reviewer_comments": {"reviewer_1": "All the comments have been well addressed. No more question and agree to accept it ASAP.", "reviewer_3": "The authors have well addressed all the questions and technical concerns raised by all the reviewers. They have added new experimental results, and cited new references to support their findings.\nI have also thoroughly examined the authors' response to the comments by other peer reviewers, and I find their responses to be both logical and well-supported. While certain inquiries were undoubtedly pointed, and some might even be perceived as exhibiting a degree of bias, the authors' replies consistently struck a fitting and appropriate tone.\nA notable example is their handling of the comparison of device performance metrics. It is a common challenge in this research area that different publications tend to prioritize and test distinct device performance characteristics. This disparity in measurement focus and parameters inevitably leads to comparisons that may not be entirely comprehensive, a point that is readily understandable. Despite this inherent complexity, the authors offered a well-considered and fitting response to this particular concern.\nFurthermore, their rationale for attributing significant performance discrepancies across devices to variations in material composition, as observed in different studies, is also quite sound. Within the realm of ferroelectric terpolymer systems, even minute fluctuations in the content of the third comonomer can profoundly influence the material's dielectric constant, polarization strength, and piezoelectric coefficient. These intrinsic material property changes, in turn, significantly magnify the observed differences in device performance. This phenomenon is a recognized and established aspect of the ferroelectric polymer field. The authors meticulously and patiently addressed this point, providing a detailed explanation of the subtle compositional variations and their direct correlation with specific material and device performance outcomes.\nTherefore, I think the revision by the authors are suitable and sufficient. I support the publication of this version as it is.\nNote from the Editor: Reviewer #3 assessed the response to reviewer #2.", "reviewer_4": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts."}, "author_response": ""}]

Decision letter