Shape-memory polyurethanes for polar wearables with ultrasensitive multi-monitoring

Domain: physics
Published: 2025-12-11
Editorial outcome: major_revision

Abstract
Integrating environmental stability and multi-monitoring modules into flexible sensor remains a pivotal scientific challenge. This study presents a supramolecular polyurethane (PU) engineered with fluorine-rich segments that form electrostatic crosslinks with positively charged ionic groups at polymer chain terminals and establish fluorine-dipole interactions with blended ionic liquid (IL) to stabilize ion transport pathways. The resulting ionically conductive elastomer combines shape memory capacity, self-healing property, and cryogenic tolerance, retaining robust mechanical strength (~32.31 MPa), toughness (~107.05 MJ m⁻3) and substantial ionic conductivity even at −40 °C. Notably, it exhibits a high temperature coefficient of resistance (TCR = 8.05% °C⁻1) at cryogenic temperatures (−40 °C to −30 °C), making it attractive for the development of cryogenic sensing materials. Additionally, the material exhibits high-sensitivity physiological monitoring capabilities with signal fidelity, serving as ionicskins for accurate physiological signal acquisition. Such multifunctional adaptability positions it as an ideal candidate for next-generation flexible electronics requiring reliable performance in extreme environments.

Introduction
The rapid advancement of flexible electronics technology is driving the innovation of sensing systems toward stretchable, adaptive, and self-healing directions. Compared to traditional rigid sensors, flexible sensors demonstrate disruptive application potential in health monitoring, human–machine interaction, and soft robotics due to their conformal adhesion, mechanical durability, and biocompatibility1–5. The core development of flexible sensors lies in creating materials that exhibit environmental stability, self-healing capability, and high signal fidelity. Based on their transmission mechanisms, flexible sensors are categorized into two types, electronic sensors relying on electrical conductors and ionic sensors based on ion-conductive systems6. Electronic sensors face limitations, such as opacity, excessive substrate dependency, and high rigidity with low stretchability, hindering their direct application7–13. In contrast, ionic sensors inherently possess homogeneous conductive epidermal characteristics. Their emergence has advanced wearable medical technologies and soft robotics sensing fields due to significant conductivity changes under strain or pressure stimuli6,14–18. Despite progress, conventional flexible sensors primarily depend on hydrogel or elastomer matrices infused with ionic liquids (ILs) or conductive fillers, struggling to simultaneously achieve high ionic conductivity, environmental stability, and shape memory performance critical for epidermal sensor applications19,20. For instance, hydrogel-based ionic conductors suffer from poor stability, rapidly swelling in humid environments or dehydrating in arid conditions21–23. Recent efforts to address these limitations focus on dynamic polymer networks.

Among diverse materials, polyurethanes (PUs) have been widely utilized due to its chemical stability and remarkable molecular designability24–27. The development of high-performance PUs with functional applications through rational molecular design has increasingly become a research focus28–35. Benefiting from the rapid progress in supramolecular polymers, the incorporation of dynamic interactions, such as hydrogen bonds, metal coordination bonds, and ionic bonds, can enable energy dissipation through bond rupture and reconfiguration while enhancing crosslinking density, thereby improving tensile strength and fracture elongation36–41. However, integrating self-healing and shape memory capabilities with environmentally stable ionic conduction in PUs remains a significant challenge, rooted in the inherent conflict between energy-dissipative network reconfiguration and continuous charge transport pathways. A promising strategy involves leveraging fluorine-mediated interactions to enhance both ionic transport and stretchability in PU systems. Concurrently, the introduction of ILs, particularly 1-ethyl-3-methylimidazolium bis(trifluoromethylsulfonyl)imide ([EMIM]+[TFSI]−), known for its low viscosity and broad electrochemical window—exploits fluorine’s strong dipole–dipole interactions and chemical inertness, endowing PU/IL hybrid systems with high stability42–45. For instance, Chen et al. reported a fluorinated PU-based material blended with ILs, achieving an ionotronicskin with high ionic conductivity (1.04 × 10−3 S cm−1) and exceptional underwater stability46. Nevertheless, existing studies predominantly focus on fluorine’s dielectric or hydrophobic properties, overlooking its potential to synchronously optimize environmental stability, self-healing kinetics and shape memory abilities.

This study proposed a molecular engineering paradigm that leverages the dual functionality of fluorine to construct spatially correlated yet functionally independent dynamic networks within a PU matrix. By introducing 2,2′-bis(trifluoromethyl)benzidine (TFMB) as a fluorinated chain extender and malonic acid dihydrazide (MDH) as a multi-H-bonding unit, we established a hierarchical architecture that decouples energy-dissipative sacrificial bonds from ion-conductive pathways. The trifluoromethyl groups in TFMB exhibited dual roles, generating synergistic effects in the PU matrix, forming ionic bonds with the end-capping agent trimethylammonium bromide (TLB) via their high electronegativity while reinforcing the molecular network alongside hierarchical H-bonds to preserve topological integrity and stability. Concurrently, polarized fluorine centers created localized electrostatic traps that immobilize cations through directional dipole interactions. And the fluorine-mediated cation confinement constructed “anion highways” for rapid anion migration and conductive pathway establishment. Remarkably, even when these highways were narrowed under harsh conditions, including low temperatures or mechanical stretching, the fluorine-stabilized framework preserved continuous ion-conducting channels through persistent dipole effects, maintaining stable ion transport efficiency.

Experimental results demonstrated breakthrough synergistic performance. The ionically conductive elastomer combines high stretchability (>900%) and ionic conductivity (1.34 × 10−4 S cm−1) with self-healing behavior (85.4% for 1 h at room temperature). Additionally, the material exhibits shape-memory performance (shape recovery rate up to 94.8%) and damping capacity. The developed flexible sensor demonstrated cryogenic tolerance, effectively addressing the critical challenge of performance degradation under extreme thermal conditions. Notably, even at −40 °C that most polymeric sensors undergo catastrophic embrittlement, it retained mechanical resilience, achieving a tensile strength of 32.31 MPa and toughness of 107.05 MJ m−3. Furthermore, the fabricated sensor not only enabled precise temperature sensing with a high temperature coefficient of resistance (TCR) of 8.05% °C−1, but also served as a strain sensor for comprehensive health monitoring applications. Most importantly, it achieved clinically relevant sensitivity in electrocardiogram (ECG) monitoring, with T-wave to R-wave amplitude ratios within the critical diagnostic threshold range of 0.2–0.4, meeting medical-grade cardiovascular assessment standards. As illustrated in Fig. 1a, this combination of properties enables applications in polar-environment sensing and cryotherapy. Our work provides insights into the design of durable ionic conductors and offers a polymer engineering strategy that mitigates conventional performance trade-offs.

Results
Polymer design, preparation, and characterization
The high electronegativity of fluorine enables effective immobilization of cationic moieties within the material matrix, facilitating the multifunctional transformation of conventional PUs through fluorine engineering. Our design strategy focuses on constructing a synergistic dynamic network via the integration of fluorine–cation interactions (ionic bonds) and high-density hydrogen bonding units in PU architecture. This approach simultaneously enhances mechanical strength while leveraging the strong electrostatic nature of C–F bonds to amplify electron polarization effects. The resultant electronic polarization promotes cation immobilization while creating additional free volume for mobile ions, thereby improving ionic conductivity. As illustrated in Supplementary Figs. 1 and 2, a series of poly(urea-urethane) (PUU) and PUFTx were synthesized via a two- or four-step process. PUU referred to the control sample without both TFMB and TLB, whereas fluorinated counterparts were designated as PUFTx, where x denotes the millimolar ratio of TLB incorporation.

Within the PU network, MDH reacted with isocyanate groups to form acyl semicarbazide units rich in H-bond donors and acceptors. Concurrently, ionic interactions between TFMB and TLB established a collaborative dynamic crosslinking network. This dual dynamic network architecture enabled the optimized PUFT2 sample to achieve exceptional mechanical properties with a tensile strength of 86.3 MPa and toughness of 311.5 MJ m−3 (Supplementary Fig. 3). To develop multifunctional ionic elastomers, [EMIM]+ [TFSI]− was incorporated into PUU and PUFT matrices, designated as PUFTx-yIL (PUU-yIL when x = 0), where y represents the mass fraction of [EMIM]+ [TFSI]−. The [EMIM]⁺ cations preferentially interacted with fluorinated TFMB segments through fluorine–cation interactions, while [TFSI]− anions remained mobile to establish continuous conductive pathways (Fig. 1b). Systematic modulation of [EMIM]+ [TFSI]− content (y) allowed precise control over free ion concentration and ionic conductivity. Consequently, these fluorine-engineered PU ionic elastomers with synergistic dynamic networks enabled a balanced integration of mechanical robustness and electrical functionality.

Fourier-transform infrared (FT-IR) spectra of all PUFTx samples (Supplementary Fig. 4) confirmed successful PU synthesis through the absence of characteristic isocyanate peaks between 2260 and 2280 cm−147. Simultaneously, the 1H NMR results (Supplementary Figs. 5 and 6) also confirmed the successful synthesis of PU and the effective incorporation of [EMIM]+ [TFSI]−. Given the structural similarity across PUFTx variants and the optimal mechanical performance of PUFT2, we selected PUFT2-yIL as the primary model system for generalized conclusions. Thermal gravimetric analysis (TGA) curves in Supplementary Fig. 7 demonstrated thermal stability for all PUFT2-yIL elastomers, with the thermal decomposition temperatures (Td, temperature of 5% weight loss) exceeding 255 °C, attributable to the inherent thermal resistance of [EMIM]+ [TFSI]−. Differential scanning calorimetry (DSC) revealed depressed glass transition temperatures (Tg) with increasing IL content (Fig. 2a), arising from IL-induced effects that reduced intermolecular interactions and increased free volume. Notably, the lowered Tg facilitated rapid chain mobility at ambient or low temperatures, enabling dynamic non-covalent bond reconfiguration and conferring room-temperature self-healing capabilities with retained low-temperature tolerance. The absence of crystallization-related endothermic and exothermic peaks in DSC profiles, coupled with the results of X-ray diffraction (XRD) analysis (Fig. 2b), confirmed the noncrystalline nature of these ionic elastomers.

Validation of fluorine–cation interaction
The fluorine–cation dipole interaction constitutes the core of this molecular design strategy, playing dual critical roles in stabilizing the dynamic network and enabling an anion-selective conduction. However, the existence patterns and interaction intensities of such interactions within multicomponent PU systems remain debated. To clarify the fluorine-mediated cation immobilization mechanism, we employed a multimodal characterization approach integrating quantum chemical calculations, spectroscopic and energy spectrum analyses, and surface ion distribution imaging. This systematic investigation not only decouples the contributions of fluorine’s electronegativity and steric effects but also establishes quantitative correlations between interaction energy and macroscopic performance.

Quantum chemical calculations were first conducted to evaluate the binding energy of [EMIM]+ [TFSI]− in the presence and absence of polymer chains with the computational details provided in Supplementary Information. As shown in Fig. 2c, the binding energy of [EMIM]+ [TFSI]− under polymer chain influence was significantly more negative compared to its isolated state. Given that binding energy values were typically negative with more negative values indicating stronger binding and structural stability, this result suggested that the cations and anions of the IL readily dissociate via fluorine–cation interactions without disrupting the ionic bonding between molecular chains. This separation mechanism ensured the integrity of the dynamic network while facilitating ion mobility.

Structural validation of fluorine–cation interactions was further pursued through spectroscopic techniques. Supplementary Fig. 8 and Fig. 2d presented the FT-IR spectra of PUFT2-yIL in the 600–3600 and 3050–3200 cm−1 regions. Within the 3050–3200 cm−1 range, two characteristic C–H stretching vibration peaks were observed, corresponding to the imidazolium ring of [EMIM]+ cations. Four distinct vibrational modes were resolved through spectral deconvolution, consisting of asymmetric C(4)-H and C(5)-H (double-bonded carbon atoms in imidazole ring) stretching (fit peek 1) as well as symmetric C–H stretching (fit peek 2), asymmetric C (2)-H stretching between nitrogen atoms (fit peek 3), and CH₃(N)HCH side chain stretching (fit peek 4)46,48. Notably, fit peaks 3 and 4 in PUFT2-50% IL exhibited blue shifts compared to PUU-50% IL, attributable to enhanced electron-withdrawing effects at C (2)-H sites through fluorine’s induction effect. Conversely, fit peaks 1 and 2 displayed red shifts, indicating reduced constraints on C (4) and C (5) positions distal to polymer chains. Additionally, the C–F stretching vibration peaks in the 1180–1200 cm−1 range exhibited pronounced intensity enhancement and broadening upon IL incorporation, confirming the dominance of [TFSI]−-derived C–F vibrations and verifying successful IL integration. Complementing FT-IR analysis, Raman spectroscopy, with its sensitivity to symmetric vibrational modes and polarizability changes, provides further insights into fluorine-mediated structural reorganization. Supplementary Fig. 9 illustrated the characteristic [EMIM]⁺ bands in the 1300–1600 cm−1 region. For PUFT2-50% IL, the H–C–H rock vibration at 1458 cm−1 and scissor vibration at 1482 cm−1 shifted to 1473 and 1495 cm−1, respectively. These blue shifts reflected restricted bond vibrations of [EMIM]⁺ cations, directly evidencing their attraction to the fluorine-rich PUFT2 matrix46,48.

X-ray photoelectron spectroscopy (XPS) was employed to resolve fluorine’s chemical states and quantify its binding with cationic species. The F 1s spectra of PUFT, PUU-50% IL, and PUFT2-50% IL displayed a single peak corresponding to fluorine in either polymer backbones or [TFSI]− anions (Supplementary Fig. 10). Notably, the C (–CF3) binding energy in PUFT2 increased by 0.3 eV upon IL incorporation (Fig. 2e), indicating reduced electron cloud density on carbon atoms within trifluoromethyl groups. This reduction stemmed from fluorine’s strong electronegativity and its interactions with [EMIM]⁺ cations. Concurrently, F 1s peak’s binding energy shifted from 687.3 to 688.3 eV showing a 1.0 eV energy gap (Fig. 2f), further corroborating interactions between polymer chains and [EMIM]⁺. To spatially resolve fluorine distribution and track cation-anion coordination at the molecular scale, time-of-flight secondary ion mass spectrometry (ToF-SIMS) mapping was performed. Line-scan profiles of [TFSI]−-related fragments (selected ion: CNO₄F₃S₂−) revealed stark contrasts between matrices (Fig. 2g). PUFT2-50% IL exhibited uniform brightness with minimal dark regions, whereas PUU-50% IL displayed pronounced brightness variations and extensive dark areas. This disparity confirmed that fluorine-mediated interactions promote homogeneous IL distribution within the PUFT2 matrix, directly validating the spatial uniformity of fluorine–cation interactions.

Collectively, this multimodal characterization, including spanning quantum calculations, vibrational spectroscopy, surface chemistry analysis, and ion mapping unambiguously confirms the existence, intensity, and functional dominance of fluorine–cation interactions. These interactions not only stabilize the dynamic network but also govern anion-selective ion transport by spatially confining [EMIM]⁺ cations while liberating [TFSI]− anions for conductive pathways. The decoupling of fluorine’s electronegativity and steric effects provides a foundational framework for rational design of fluorine-engineered PUs with tunable mechanical and electrical properties.

Mechanical and electrical properties
The mechanical properties of PUFT2-yIL, including tensile strength, elongation at break, and toughness, were summarized in Supplementary Table 1, with corresponding stress–strain curves illustrated in Fig. 3a. The introduction of IL immediately reduced material strength and toughness, likely due to IL–fluorine interactions disrupting original ionic bonds and weakening interchain cohesion. As IL content increased, mechanical strength and stretchability further decreased (Fig. 3b). Although PUFT2 exhibited 86.3 MPa strength, the incorporation of 50% IL altered its strength and elongation to 1.32 MPa and 948%, respectively, achieving flexibility and mechanical compliance suitable for epidermal electronics. Small-angle X-ray scattering (SAXS) analysis revealed distinct scattering peaks for PUU (microphase separation period d=7.46 nm) and PUFT2 (d=8.62 nm), confirming enhanced microphase separation due to ionic bonding between TFMB and TLB (Fig. 3c). The increased d-spacing in PUFT2 indicated that ionic bonds amplify microphase-separated domains, positively influencing mechanical performance, as further corroborated by stronger halos in 2D-SAXS patterns compared to PUU (Supplementary Fig. 11). However, no scattering peaks were observed for PUU-50% IL or PUFT2-50% IL, suggesting IL-induced disruption of microphase separation. Atomic force microscopy (AFM) images revealed progressively blurred boundaries between hard segments (bright regions) and soft segments (dark regions) in PUFT2-yIL with increasing IL content (Supplementary Fig. 12), further evidencing weakened microphase separation.

Building on the mechanical adaptability of the fluorine-engineered dynamic network, we further investigated its electrical properties to elucidate how fluorine–cation interactions synergistically regulate ion transport mechanisms. Electrochemical impedance spectroscopy (EIS) of PUU-50% IL and PUFTx-50% IL (Fig. 3d) showed that bulk resistance initially increased and then decreased with the x value changed. As shown in Fig. 3e, PUFT2-50% IL achieved the highest ionic conductivity (1.34 × 10−4 S cm−1), which was attributed to the fluorine–cation interactions that preferentially immobilized the [EMIM]⁺ cations while weakening the anion–cation interaction, thereby promoting free ion release. Increasing IL concentration enhanced ionic conductivity (Fig. 3f, g and Supplementary Table 2), though concentrations exceeding 50% compromised mechanical stability and self-supporting capability, rendering them unsuitable for practical polymer electrolyte applications. Combined with mechanical data, this suggested that IL-induced destruction of polymer crystalline structures enhances amorphous domains, leading to mechanical degradation. The enhanced effect of fluorine engineering on ionic conductivity was further investigated by all-atom molecular dynamics simulations. Supplementary Figs. 13 and 14 illustrated the modeling of the polymer chain in the PUU-50% IL and PUFT-50% IL systems, with modeling and simulation details provided in the Supplementary Information. The results revealed superior conductivity in the fluorine-engineered PUFT-IL system compared to PUU during tensile-compression cycles (Fig. 3h). Notably, the PUFT-IL system maintained stable conductivity during simulated epidermal sensor stretching-recovery processes, whereas PUU exhibited significant conductivity fluctuations, confirming fluorine engineering’s critical role in ensuring sensing stability. Linear sweep voltammetry demonstrated electrochemical stability windows of 3.8–3.6 V for PUFT2-yIL with 10–50% IL content, consistently exceeding the 2.7 V window of PUU-50% IL (Fig. 3i). The preserved high electrochemical stability, coupled with tunable ionic conductivity and mechanical compliance, establishes PUFT2-yIL as a viable candidate for epidermal electronic devices.

Self-healing, damping and shape memory properties
Driven by reversible intermolecular multi-hydrogen bonding and ionic interactions, the PUFT2-50% IL ionic elastomer exhibits self-healing behavior. Real-time resistance measurements during multiple surgical blade cuts showed that the material restored its original resistance within 0.5 s, confirming rapid electrical recovery (Fig. 4a). To quantify self-healing efficiency, PUFT2-50% IL samples were cut into two pieces, then the two completely separated ruptured surfaces were brought into full contact at room temperature. As shown in Fig. 4b, tensile strength and strain gradually increased with healing time. After 1 h, the tensile strength and strain recovered to 0.74 MPa and 760%, respectively. After 3 h, they reached 1.11 MPa and 912%, similar to the original sample. For strain sensors, stretchability is a critical parameter. Based on strain recovery, the healing efficiency reached 85.4% after 1 h and 97.4% after 3 h (Fig. 4c), demonstrating its efficient self-healing capability.

For sensor materials, exceptional damping performance is essential. Tan δ is a key factor for evaluating damping properties. A temperature range with tan δ greater than 0.3 is considered an effective damping temperature range, while tan δ above 0.5 indicates a broad damping temperature range49–52. The temperature-dependent damping properties of PUFT2-50% IL were measured via dynamic mechanical analysis (DMA) at 1 Hz (Fig. 4a). The results showed that tan δ remained above 0.3 from −50 to −15 °C, demonstrating the material’s low-temperature stability. However, for sensing materials in polar environments, a broad damping temperature range was more critical. PUFT2-50% IL exhibited good damping performance (tan δ > 0.5) between −45 and −25 °C, which significantly increased its potential for polar applications. Due to its observed damping behavior, PUFT2-50% IL was also suitable for protecting fragile objects. For example, as shown in Fig. 4b, when a 100 g ball was dropped from a height of 50 cm, unprotected glass slides shattered completely. In contrast, a 0.7 mm thick PUFT2-50% IL film effectively prevented glass breakage under both room temperature and cryogenic conditions (Supplementary Movie 1).

Furthermore, considering that PUFT2-50% IL possessed a low Tg, this material was expected to exhibit temperature-induced shape memory behavior. The shape memory properties were quantitatively investigated using DMA (Fig. 4c). The results demonstrated that the shape recovery rate of PUFT2-50% IL increased with cycling, rising from 86.5 to 94.8%, while its fixation rate remained stable at approximately 98.3%. This increase in recovery rate was likely attributed to the gradual alignment of molecular chains along a consistent orientation and reduced chain entanglements during cyclic loading. These quantitative results verified the shape memory capability of PUFT2-50% IL. Here, the focus was primarily on the material’s unidirectional shape memory effect and its potential for smart wearable applications. Specifically, the ionically conductive elastomer could be coiled on flat surfaces or fingers and fixed into a temporary shape at low temperatures. This temporary shape then recovered to nearly indistinguishable original configurations within a short time at room temperature (Fig. 4d; Supplementary Movies 2 and 3). Notably, simulated finger-mounted demonstrations revealed that the material maintained its coiled state for temperature monitoring under extreme cryogenic conditions without restricting normal finger flexibility. Upon returning to room temperature, the material reverted to its original shape and conformed to the finger, enabling subsequent human motion detection. These capabilities highlighted the material’s promising potential for intelligent wearable devices and signal transmission in polar environments.

Low temperature tolerance and temperature sensitivity
Existing flexible sensors commonly suffer from instability under varying environmental conditions. As previously mentioned, PUFT2-50% IL exhibited good damping performance, strongly suggesting its potential for cryogenic stability. Figure 5a demonstrated the typical stress–strain curves of PUFT2-50% IL at 0, −20, and −40 °C. Notably, with continuous temperature reduction, the material’s strength increased while its elongation at break decreased. Compared with room temperature conditions, the mechanical properties showed varying degrees of enhancement (Supplementary Fig. 15). Remarkably, at −40 °C, the strength and toughness reached 32.31 MPa and 107.05 MJ m−3, representing 2348 and 1515% improvements, respectively, while still maintaining high stretchability (Supplementary Table 3). This phenomenon could be attributed to the restricted molecular chain mobility in polyurethane elastomers at lower temperatures, which intensified intermolecular interactions and facilitated the formation of ordered structures. Consequently, the polymer chains exhibited reduced sliding capacity and deformation capability, thereby increasing the resistance to plastic deformation in the elastomer (the serrated stress–strain patterns observed were attributed to the liquid nitrogen cooling process employed during tensile testing). Similarly, ionic conductivity depends on ion mobility, which was intrinsically linked to polymer chain conformational changes and thermal motion. To investigate the electrical performance of PUFT2-50% IL under cryogenic conditions, EIS measurements were conducted within the temperature range of −40 to 0 °C (Supplementary Fig. 16). As shown in Fig. 5b, the ionic conductivity of PUFT2-50% IL demonstrated a progressive decline with decreasing temperature. This reduction stemmed from strengthened intermolecular interactions at lower temperatures, leading to denser polymer networks that restricted ion transport channels and impeded ionic migration. Notably, even at −40 °C, PUFT2-50% IL maintained a relatively high ionic conductivity of 1.46 × 10−6 S cm−1, significantly expanding its operational temperature range as a functional sensor material.

Benefiting from its favorable combination of low-temperature toughness, ionic conductivity, and stretchability, PUFT2-50% IL demonstrated promising potential as a resistive temperature sensor for cryogenic applications. Figure 5c illustrated the temperature-dependent relative resistance (ΔR/R₀) curve of PUFT2-50% IL. Notably, the relative resistance values at all tested temperatures were negative compared to the initial resistance at −40 °C, indicating a significant decrease in resistance with increasing temperature. This observation confirmed that PUFT2-50% IL exhibited negative temperature coefficient behavior. To quantitatively evaluate its sensitivity and stability as a temperature sensor, we introduced the TCR, defined as the relative resistance change per unit temperature variation. The TCR value, derived from the slope of the linear fitting curve in Fig. 5c, reflected the sensor’s sensitivity, with higher absolute TCR values corresponding to greater sensitivity. As shown in Fig. 5d, the absolute TCR of PUFT2-50% IL reached a maximum of 8.05% °C−1 in the −40 to −30 °C range. Remarkably, as shown in Fig. 5e, this material operated at a temperature of −40 °C while maintaining a high TCR, a performance that is notable among reported flexible temperature sensors49,53–67. Furthermore, the temperature-resistance relationship in viscoelastic ionic elastomers adhered to the Vogel–Tamman–Fulcher equation, which governed temperature-dependent ionic conductivity. The coefficient of determination (R²) of the linear fitting curve determined whether temperature could be accurately represented by resistance variations. As annotated in Fig. 5d, all R² values exceeded 0.96 across tested temperature ranges, demonstrating the precise quantifiability of temperature through resistance changes. This combination of high TCR sensitivity and robust linear correlation underscored PUFT2-50% IL’s viability as a high-performance temperature-sensing material in extreme cold environments. Notably, the PUFT2-50% IL sensor demonstrated good temperature-response cyclability, as evidenced by its stable relative resistance variation during five consecutive thermal cycling tests between −40 and 20 °C (Fig. 5f). Furthermore, its low-temperature stability was confirmed by a minimal resistance change of less than 4% over a continuous 120-min test at −40 °C (Supplementary Fig. 17), underscoring its strong potential for reliable applications in cryogenic environments. These robust performances under periodic thermal shocks confirmed the material’s structural integrity and temperature-tolerance capability, critical for reliable operation in extreme environments.

Highly sensitive and healing ionic skin
Strain sensors represent an ideal direction for future development of flexible wearable sensors. As shown in Supplementary Fig. 18, the PUFT2-50% IL demonstrated the capability to detect distinct signal variations induced by minor strain changes (20%) within a 20–100% strain range. Meanwhile, the sensitivity of PUFT2-50% IL was evaluated under strain conditions ranging from 0 to 500% (Fig. 6a). The gauge factor was determined to be 1.93 within the 0–100% strain range and was observed to increase with increasing strain, reaching a value of 2.41 at 200–300% strain. Notably, the sensor exhibited rapid response characteristics with measured response and recovery times of 72 and 82 ms respectively, indicating high stability and sensitivity (Fig. 6b). To evaluate its potential as an epidermal biosensor, PUFT2-50% IL was employed for real-time monitoring of diverse human movements and physiological signals (Fig. 6c). Figure 6d demonstrated that periodic finger flexion-extension motions induced corresponding cyclic resistance variations, through which the sensor accurately identified different bending angles while maintaining stable signal output across multiple cycles, thereby confirming its reliability and sensitivity. Furthermore, as illustrated in Fig. 6e, f PUFT2-50% IL effectively detected wrist and elbow flexion with consistent resistance signal outputs. These quantifiable electrical signals enabled precise recording and analysis of athletes’ motion patterns during training sessions, facilitating timely evaluation of movement accuracy and subsequent improvement of training efficiency.

In addition to the strong resistance signals described above, PUFT2-50% IL successfully detected subtle physiological activity signals. Respiratory monitoring, which serves as an indicator of exercise intensity or health status, constitutes an essential component of comprehensive health monitoring systems. As shown in Fig. 6g, thoracic expansion and contraction during breathing generated strain variations in the attached PUFT2-50% IL, which induced measurable resistance changes. The sensor differentiated between deep and shallow breathing patterns with precision, further confirming its stability and sensitivity. Although respiratory rhythm reflects autonomic nervous system activity and its detection accuracy directly influences metabolic state assessments, the limitations of single-parameter respiratory monitoring necessitated the application of PUFT2-50% IL as an electrode for ECG detection. Figure 6h illustrated the stable ECG signals acquired using PUFT2-50% IL. Signal quality was evaluated through characteristic waveforms, including the P-wave, QRS complex, and T-wave, which correlated with cardiac electrical activity and were critical for clinical diagnosis (Fig. 6i). The T-wave amplitude, representing ventricular repolarization, typically ranges between 20 and 35% of the R-wave amplitude associated with ventricular depolarization. Clinically acceptable ECG signal quality requires a T/R peak amplitude ratio within 0.2–0.446,68. As demonstrated in Fig. 6j, the T/R ratios derived from PUFT2-50% IL-detected ECG signals consistently fell within this range, confirming reliable signal acquisition. These results collectively demonstrated the high sensitivity and stability of PUFT2-50% IL in sensing applications.

Previously, the self-healing capability of PUFT2-50% IL had already been confirmed. Therefore, it could be inferred that PUFT2-50% IL would still be able to maintain its sensing ability after healing. Healed samples (1 h healing) successfully monitored repetitive finger and wrist flexion cycles, with ΔR/R0 amplitude consistency confirming stable real-time motion detection (Supplementary Figs. 19 and 20). Furthermore, the healed material maintained stable ECG signal acquisition capability post-repair, as evidenced by unaltered signal fidelity in Supplementary Fig. 21. These results collectively demonstrated that the self-healing mechanism preserved both mechanical integrity and sensing functionality, significantly extending the material’s applicability in demanding environments. This work substantiated that hierarchical dynamic bonding networks not only enabled robust self-healing but also ensured performance retention across mechanical, electrical, and bio-signal sensing modalities.

Methods
Materials
Polycarbonate diol (PCDL, Mn = 2000 Da, 99%) was purchased from Liduo Chemical Co., Ltd. (Jining, China). Isophorone diisocyanate (IPDI, 99%), MDH (97%) and dibutyltin dilaurate (DBTDL, 98%) were purchased from Energy Chemical Co., Ltd. (China). 2,2′-bis(trifluoromethyl) benzidine (TFMB, 98%), (2-hydroxyethyl) trimethylammonium bromide (TLB, 98%) and 1-ethyl-3-methylimidazolium bis(trifluoromethanesulfonyl)imide ([EMIM]+[TFSI]−, 99%) were purchased from TCI (Shanghai) Development Co., Ltd. N,N-dimethylformamide (DMF, 99.5%) was purchased from Rionlon Bohua Pharmaceutical and Chemical Co., Ltd. (Tianjin, China). All other chemicals and solvents were of analytical grade and used without further purification.

Synthesis of the PUFTx elastomer
For the synthesis of the PUFT: PCDL (Mn = 2000 Da, 5.0 mmol, 10.00 g) was mechanically stirred at 110 °C for 2 h under nitrogen (the following operations are carried out under nitrogen atmosphere if not otherwise specified.) in a three-necked round-bottom flask to remove residual moisture, cooled to 80 °C, and supplemented with a solution of IPDI (10.0 mmol, 2.25 g) and DBTDL (0.02 g) in DMF (40 mL). The mixture was stirred at 80 °C for 3 h, and the resulting prepolymer solution was supplemented with MDH (4.0 mmol, 0.53 g) and TFMB (0.5 mmol, 0.16 g). The mixture was heated at 80 °C for 2 h to a chain extension reaction, then supplemented with TLB (1 mmol, 0.18 g), further heated at 80 °C for 5 h, until the reaction was completed. The resulting liquid was poured into clean Teflon molds, and the samples were vacuum-dried at 80 °C for 48 h. The product was named PUFT. Following the same procedure for PUFTx were synthesized by changing the feed mass of MDH, TFMB and TLB. The mixing ratios used to prepare PUFTx are given in Supplementary Table 4.

Synthesis of the PUFTx-yIL elastomer
For the synthesis of the PUFT2-50% IL: PUFT2 (2.00 g) was dissolved in DMF (40 mL) and [EMIM]+ [TFSI]− (2.00 g) was added after the dissolution was completed. After 1 h of stirring to mix well, the mixture solution was poured into clean Teflon molds, and the samples were vacuum-dried at 80 °C for 48 h. The product was named PUFT2-50% IL.

General characterizations
TGA was evaluated in an SDT650 synchronous thermal analyzer (TA, America) in the range of 25–800 °C under a nitrogen flow (100 mL min−1) at a heating rate of 10 °C min−1. DSC measurements (DSC STA449F3, Netzsch, Germany) were performed in a flow of nitrogen. The samples (5–6 mg) were heated to 150 °C, cooled to −100 °C, and reheated to 150 °C at a rate of 10 °C min−1. XRD measurements were characterized on a Bruker D8 Advance (BRUKER, Germany) using Cu-Kα radiation with λ = 1.5418 Å, and the data were recorded in a range from 5° to 90° at a scanning speed of 5° min−1. FT-IR spectra was recorded on a Nicolet iS20 (Thermo Scientific, America) spectrometer from 500 to 4000 cm−1 in attenuated total reflectance mode. Hydrogen proton nuclear magnetic resonance (1H NMR) spectra were measured on an AVANCE NEO 400 MHz spectrometer (Bruker, Germany), using CDCl3 as solvent. Raman spectroscopy was recorded using a Raman spectrometer (LabRAM HR Evolution, HORIBA), with the laser power of 25 mW for 532 nm laser excitation. XPS was carried out on K-Alpha (Thermo Scientific, America), excited by an Al Kα (hν = 1486.6 eV) Mono X-ray with a beam spot of 400 µm. ToF-SIMS measurements were conducted utilizing a PHI nanoTOF II Time-of-Flight SIMS (ULVAC-PHI.INC, JAPAN) instrument. Ion mass spectrometry images were acquired by sputtering the sample surface using a 30 keV Bi ++ ion beam in high mass resolution mode over a surface area of 200 × 200 μm of the test sample. The microphase structure of the samples was detected by an atomic force microscope (AFM, Bruker Dimension Icon, Germany) in tapping mode with a scan area of 0.6 × 0.6 μm. Thermomechanical properties were characterized using a dynamic mechanical analyzer (TA Q850 DMA). To that end, all samples were cut into rectangles (20 × 4 × (0.2–0.5) mm) and heated from −120 to 100 °C with a heating rate of 5 °C min−1 and a frequency of 1 Hz.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "Accept as is.", "reviewer_2": "This manuscript describes a multifunctional polyurethane elastomer constructed through fluorine-cation interactions, combining shape memory capability, self-repairing properties, and excellent low-temperature resistance. A novel methodological strategy is proposed that is applicable to ultra-low-temperature sensors, enabling simultaneous improvements in stability and sensitivity. This approach is expected to be instructive for applications in wearable devices operating in extreme environments and will likely attract broad interest from researchers in the fields of smart wearable materials, smart sensing, and extreme-environment technologies. Overall, the manuscript presents a promising contribution that could be suitable for publication in Nature Communications after minor revisions are addressed. The following points should be considered for improvement:\n1. Many figures are not well organized. It is recommended that some of the diagrams be combined for clarity and that the overall figure layout be polished.\n2. The manuscript lacks information on the variability and reproducibility of the presented data. Without this, readers must assume that the results are based on single measurements. The reliability of the findings would be enhanced by including data from multiple samples and statistical analyses.\n3. Given that the elastomer’s properties are closely related to its molecular structure, the structural characterization appears insufficient. Additional data such as NMR spectra should be provided to confirm the chemical structure.\n4. On page 12, line 219, the FTIR peaks are attributed to different carbon atoms on the imidazole ring. It would be helpful to include the molecular structure of the [EMIM]+ cation and label the relevant carbon atoms for clarity.\n5. On page 21, line 404, ionic conductivity at –40 °C is mentioned. To better contextualize this result, a comparison with previously reported materials would be beneficial. In Figure 5e, the lowest working temperature and TCR index are compared; adding a comparison of ionic conductivity across different materials would further strengthen the discussion.\n6. In Figure 7c, the self-healing efficiency of elongation at break is significantly higher than that of tensile strength. An explanation for this discrepancy should be provided, along with data or discussion regarding changes in toughness after healing.", "reviewer_3": "This review evaluates the molecular modeling aspects of the manuscript. Overall the results are contradictory (see my comment about binding energy below), not sufficiently detailed, and (probably, but hard to judge since so much information is lacking) not run long enough. It is not acceptable in the current form. My specific comments are:\nI have some questions about the following text:\n\"Quantum chemical calculations were first conducted to evaluate the binding energy of [EMIM]+ [TFSI]- in the presence and absence of polymer chains with the computational details provided in Supporting Information. As shown in Fig. 2c, the binding energy of [EMIM]+ [TFSI]- under polymer chain influence was significantly less negative compared to its isolated state. Given that binding energy values were typically negative with more negative values indicating stronger binding...\"\nMy questions about this text are: (1) The binding energy of [EMIM]+ [TFSI]- under polymer chain influence is given as -142.4 which is more negative, not less negative, than the -83.2 value for the isolated state. This would seem to negate the argument given in the text. (2) Under the polymer chain influence, I don't know what was computed using equation 10 of the supplementary file? Be more specific here since there are more than two species.\nThe authors say: \"Initially, the simulation systems underwent structural optimization.\" (1) What does this mean? (2) How were the 24 or 36 polymers arranged? (3) What does Figure 12 show? Is it a snapshot of a single polymer or is this the 36 polymer system? I'm guessing it is just a single polymer. (4) We are told it is \"optimized\" but what does this mean? Some parts of the polymer chain are extended and some are coiled. How was this structure created? (5) More importantly, how are the 36 polymer chains arranged? Was this single polymer configuration replicated 36 times? How were the 36 polymers packed together? It is impossible to judge the realism and quality of the simulations without these details.\nThe authors state that \"The diffusion coefficient is obtained from the long-time limit of the MSD\". But the simulation was only run for 5 ns. It is doubtful that there is a long-time limit with this short a simulation. The authors need to show the MSD vs. time plot and show the fitted line to the data that they used to compute the diffusion coefficient.\nI have some questions about the following text:\n\"...relaxation for 1 nanosecond until the systems reached a steady state. Following this, over a period of 5 nanoseconds, the systems were uniformly stretched in the x direction to twice their original length, while maintaining a pressure of 1 atmosphere in the y and z directions. Then, over another 5 nanoseconds, the systems were compressed back to their original length in the x-direction, again maintaining a pressure of 1 atmosphere in the y and z directions. For the structures before stretching, after stretching, and after compression, we conducted 5-nanosecond simulations under the NVT ensemble\"\nMy questions about this text are: (1) what does \"steady state\" mean? 1 ns is very short. Again this goes back to my previous comment about how the polymer chains were arranged relative to one another. Steady state could mean the box volume stops changing, or the total energy converges to a steady value, or that the RMSD of the polymers stops changing, or any number of other criteria. It is hard to believe that 1 ns is enough to reach equilibrium. (2) how do the y and z box sizes compare before and after the stretching and compression? (3) After the 5 ns stretching, was the system immediately compressed or was the 5 ns NVT run done in-between with the compression done from after the NVT run? Or is the NVT run done on a pull-out of the NPT configuration without interrupting the stretch/compress run? This isn't clear. (4) If the NVT runs to measure conductivity are done immediately following the stretching run (and the compression run), presumably the system is not at equilibrium yet (the system would relax to equilibrium after the stretching is stopped over some time scale) so this would be a non-equilibrium simulation which is not appropriate.", "reviewer_4": "This study presents a supramolecular polyurethane engineered with fluorine-rich segments that form electrostatic crosslinks with positively charged ionic groups at polymer chain terminals and establish fluorine-dipole interactions with blended ionic liquid to stabilize ion transport pathways. Many interesting results are presented. However, the reviewer suggests the following points to be seriously addressed:\n1. In Abstract, The author points out: “The resulting ionically conductive elastomer combines shape memory capacity, self-healing property, and exceptional cryogenic tolerance, retaining robust mechanical strength (~32.31 MPa), toughness (~107.05 MJ m⁻³) and substantial ionic conductivity even at -40°C. Notably, it achieves an unprecedented combination of record-low operational temperatures (-40°C to -30°C) and an ultrahigh temperature coefficient of resistance (TCR = 8.05% °C⁻¹), setting a new benchmark for cryogenic sensing materials.” However, the author did not focus on testing the various characteristics and application under low-temperature conditions.\n2. In Section 2.6 (Highly sensitive ionic skin), there are too few tests on static characteristics, and basic characteristic indicators such as sensitivity, linearity and range are missing.\n3. In Section 2.7 (Self-healing properties), Fig. 6d-e and Fig. 6f should be marked as Fig. 7d-e and Fig. 7f. This mistake should be due to the author's carelessness.\n4. As shown in Fig.5 and Fig.6, compared with the gauge factor, the resistance temperature coefficient is too strong. In the actual application process, it will affect its application in strain or pressure.\n5. It is recommended to describe the material properties first and then the application properties. For example, Section 2.7 should be described earlier.\n6. Generally, ionic conductors are commonly used in capacitive sensors, demonstrating high sensitivity, but are less applied in resistive principles. It can also be roughly seen from the results in Section 2.6 (Highly sensitive ionic skin) that the gauge factor of the designed sensor is relatively low.\n7. The application part of the material is too simple to support its application value."}, "author_response": "Response to Previous Reviewers’ Comments\nDear reviewers,\nWe would like to thank all reviewers for their constructive comments and\nsuggestions. After carefully reading and thinking about the reviewers' comments, we\nhave supplemented and revised this manuscript to reflect most of their suggestions.\nThese comments are valuable and helpful for the improvement of our paper, and they\nalso have important guiding significance for research. Therefore, we have studied the\ncomments carefully and have made further revisions, hoping that the reviewers and\neditors will be satisfied with the revised manuscript.\nOur point-by-point responses to these comments are provided in another document,\nand the main changes in the revised manuscript are marked in yellow (revised contents)\nand blue (supplementary contents). In any case, we are open to consideration of any\nfurther comments and suggestions. We look forward to hearing from you on the status of our submission. Thanks again to the four reviewers for their constructive comments,\nand to the editors for their affirmation and support of this work.\nSincerely,\nJing Xu and Xinrui Zhang\n Reviewer #1 (Remarks to the author):\nAccept as is.\n Response #1:\nWe are very grateful to the reviewer for reviewing our manuscript and we sincerely\nappreciate your recognition of our work. Your unequivocal endorsement of our work is\nboth deeply encouraging and a significant validation of our research efforts. We\nsincerely appreciate your insightful recognition of the manuscript's scientific rigor and\ncontribution to the field. Such support from a respected expert is invaluable to our team.\n Reviewer #2 (Remarks to the author):\nThis manuscript describes a multifunctional polyurethane elastomer constructed\nthrough fluorine-cation interactions, combining shape memory capability, self-\nrepairing properties, and excellent low-temperature resistance. A novel methodological\nstrategy is proposed that is applicable to ultra-low-temperature sensors, enabling\nsimultaneous improvements in stability and sensitivity. This approach is expected to be\ninstructive for applications in wearable devices operating in extreme environments and\nwill likely attract broad interest from researchers in the fields of smart wearable\nmaterials, smart sensing, and extreme-environment technologies. Overall, the\nmanuscript presents a promising contribution that could be suitable for publication in\nNature Communications after minor revisions are addressed. The following points\nshould be considered for improvement.\n Response #2:\nWe are very grateful to the reviewer for reviewing our manuscript and providing\nconstructive comments, which are immensely helpful in improving the quality of our\nmanuscript. To fully address your comments and concerns, we have updated the\nmanuscript and provided detailed explanations. Besides, we also have supplemented or\nrevised figures, expanded discussions, and improved the language in the revised\nmanuscript. The main changes in the revised manuscript or supporting information are\nin yellow (revised contents) and blue (supplementary contents). In the following text,\nwe will address each comment point by point.\n• Comment #2-1\nMany figures are not well organized. It is recommended that some of the diagrams\nbe combined for clarity and that the overall figure layout be polished.\n Response #2-1:\nWe sincerely thank you for your insightful comments and have thoughtfully\nrevised the manuscript based on the reviewers' comments. We have reconsidered the\ncombination of the figures and have revised this in the re-submission.\n Comment #2-2\nThe manuscript lacks information on the variability and reproducibility of the\npresented data. Without this, readers must assume that the results are based on single\nmeasurements. The reliability of the findings would be enhanced by including data from\nmultiple samples and statistical analyses.\n Response #2-2:\nThanks for your careful review and constructive suggestions regarding our\nmanuscript. In this work, we have conducted repeated experiments to prove our\nexperimental results, and the test data obtained are those we have obtained after\nrepeated proof. The relevant performance data of the samples are fully reliable. We also\nprovide detailed test descriptions at the end of the paper, and all the results obtained\nfrom the tests and experiments support each other. In addition, we have uploaded the\nraw data for experimental replication. Overall, the research results are highly reliable.\n Comment #2-3\nGiven that the elastomer’s properties are closely related to its molecular structure,\nthe structural characterization appears insufficient. Additional data such as NMR\nspectra should be provided to confirm the chemical structure.\n Response #2-3:\nThanks again for your careful review and constructive suggestions on our\nmanuscript, which greatly improved its quality. According to your advice, we have\nperformed 1H NMR tests to confirm the structure of the as-synthesized polymers.The\nresults have been added in the revised Supplementary Information as shown in Figs.\nR1-R2 (Figs. S2-S6 in Supplementary Information). From the results of NMR further\ndemonstrated that the materials have been successfully synthesized.\nFig.R1 (Fig. 5 in Supplementary Information) 1H NMR spectrum of PUFT in CDCl .\n2 3\nThe specific 1H NMR analysis is as follows:\n1H NMR (400 MHz, CDCl ) δ 6.99 (m, H-h), 6.79 (d, J = 2.5 Hz, H-f), 6.77 (m,\n1.46 – 1.22 (s, H-i), 1.16 – 0.70 (m, H-j).\nFig.R2 (Fig. 6 in Supplementary Information) 1H NMR spectrum of PUFT -50%IL in\n1H NMR (400 MHz, CDCl ) δ 8.94 (s, H-k), 7.30 (d, J = 1.8 Hz, H-l), 7.28 (d, J\n3\n= 1.8 Hz, H-m), 7.09 – 6.94 (m, H-h), 6.79 (d, J = 2.4 Hz, H-f), 6.77 (d, J = 2.6 Hz,\nH-c), 4.27 (q, J = 7.4 Hz, H-n), 4.12 (m, H-b), 3.97 (s, H-p), 3.26 (s, H-a), 3.16 (m, H-\nd), 2.97 – 2.86 (m, H-g), 1.67 (m, H-e), 1.57 (t, J = 7.4 Hz, H-o), 1.41 (s, H-i),\n1.11 – 0.77 (m, H-j).\nThe detailed descriptions added to the revised manuscript are as follows:\nSimultaneously, the 1H NMR results also confirmed the successful synthesis of\nPU and the effective incorporation of [EMIM]+ [TFSI]-.\n Comment #2-4\nOn page 12, line 219, the FTIR peaks are attributed to different carbon atoms on\nthe imidazole ring. It would be helpful to include the molecular structure of the\n[EMIM]+ cation and label the relevant carbon atoms for clarity.\n Response #2-4:\nWe sincerely thank you for your kind suggestions, as well as apologize for the\nconfusion caused to the reviewers due to unclear description. We have revised this in\nthe re-submission, provided a more clearly figure (Fig R3).\nFig. R3 (Fig. 2 in revised manuscript)\n Comment #2-5:\nOn page 21, line 404, ionic conductivity at –40 °C is mentioned. To better\ncontextualize this result, a comparison with previously reported materials would be\nbeneficial. In Figure 5e, the lowest working temperature and TCR index are compared;\nadding a comparison of ionic conductivity across different materials would further\nstrengthen the discussion.\n Response #2-5:\nThank you very much for your careful review and constructive suggestions\nregarding our manuscript. According to your advice, we searched relevant works for\ncomparison of ionic conductivity. Nevertheless, we noted that numerous cryogenic\nmaterials exhibited superior electrical conductivity at low temperatures compared to\nour material, yet they had not been utilized in sensors, nor had their TCR indices been\ncharacterized under cryogenic conditions. Concurrently, some literature on sensor\nmaterials either omitted low-temperature ionic conductivity testing or reported TCR\nindices solely above room temperature. These limitations hinder meaningful\ncomparative analysis. Crucially, this work prioritized the ultra-low-temperature sensing\nperformance of the material. Expanding the scope to include comprehensive low-\ntemperature conductivity studies risked compromising the manuscript’s conciseness.\nHerein, we enthusiastically provide detailed comparative data where feasible.\nSpecifically, Fig. R4 now directly contrasts the cryogenic ionic conductivity of our\nmaterial with literature values cited in the original Fig. 5e. This comparison\nconclusively demonstrates that our material outperforms these referenced materials in\nlow-temperature ionic conductivity.\nFig. R4 Superiority of this materials over similar previously reported samples.\n Comment #2-6:\nIn Figure 7c, the self-healing efficiency of elongation at break is significantly\nhigher than that of tensile strength. An explanation for this discrepancy should be\nprovided, along with data or discussion regarding changes in toughness after healing.\n Response #2-6:\nThank you again for your valuable advice to improve the quality of our manuscript.\nBased on your advice, changes in toughness after healing have been added in the revised\nmanuscript, as shown in Fig. R5 (Fig. 4c in revised manuscript).\nFig. R5 (Fig. 7c in manuscript) Healing efficiency at different times.\nIn our system, the significantly higher self-healing efficiency of elongation at\nbreak compared to tensile strength can be attributed to the distinct structural\ndependencies of these two mechanical parameters. Elongation at break is mainly\ngoverned by the flexibility and mobility of polymer chains, as well as their ability to\nre-entangle and rearrange at the fracture interface. Under the influence of reversible\nnon-covalent interactions (e.g., H-bonding and ionic interactions) and the plasticizing\neffect of the incorporated ionic liquid, these flexible chain segments can rapidly\nreconfigure during the healing process, leading to a substantial recovery of extensibility.\nIn contrast, tensile strength relies more on the integrity and spatial continuity of\nload-bearing permanent cross-linking nodes (chemical cross-links or highly ordered\nphysical cross-linking domains). Microcracks, interfacial defects, and irreversible chain\nscission generated during fracture are difficult to completely eliminate during self-\nhealing, resulting in stress concentration and premature failure under high loads,\nthereby limiting the recovery of tensile strength.\nIt should be noted that the healing efficiency of toughness in our system is often\neven lower than that of tensile strength. Toughness is essentially the integrated product\nof strength and elongation, making it highly sensitive to both the network integrity and\nchain extensibility. During healing, even if extensibility is partially restored, any\nresidual defects that compromise the load-bearing capability at high stress will\nsignificantly reduce toughness, and vice versa. Therefore, complete recovery of\ntoughness requires simultaneous high recovery of both strength and elongation, which\nis generally more challenging to achieve in practical self-healing processes.\n Reviewer #3 (Remarks to the author):\nThis review evaluates the molecular modeling aspects of the manuscript. Overall\nthe results are contradictory (see my comment about binding energy below), not\nsufficiently detailed, and (probably, but hard to judge since so much information is\nlacking) not run long enough. It is not acceptable in the current form.\n Response #3:\nWe sincerely thank the reviewer for carefully evaluating the molecular modeling\naspects of our manuscript and for providing constructive comments. We acknowledge\nthat some of our modeling results were insufficiently detailed and that certain\nexplanations were unclear, which may have led to apparent contradictions. We\napologize for the lack of clarity and detail in the initial submission and have\nsubstantially revised the relevant sections to address these concerns, as detailed in our\npoint-by-point responses below.\n• Comment #3-1:\n\"Quantum chemical calculations were first conducted to evaluate the binding\nenergy of [EMIM]+ [TFSI]- in the presence and absence of polymer chains with the\ncomputational details provided in Supporting Information. As shown in Fig. 2c, the\nbinding energy of [EMIM]+ [TFSI]- under polymer chain influence was significantly\nless negative compared to its isolated state. Given that binding energy values were\ntypically negative with more negative values indicating stronger binding...\"\nMy questions about this text are: (1) The binding energy of [EMIM]+ [TFSI]-\nunder polymer chain influence is given as -142.4 which is more negative, not less\nnegative, than the -83.2 value for the isolated state. This would seem to negate the\nargument given in the text. (2) Under the polymer chain influence, I don't know what\nwas computed using equation 10 of the supplementary file? Be more specific here since\nthere are more than two species.\n Response #3-1:\n(1) We apologize for the inconvenience caused to the reviewer by the incorrect\ndescription. Based on the quantitative calculation results, the binding effect of\n[EMIM]+[TFSI]- is weaker in the absence of polymer chain influence, whereas it\nbecomes stronger in the composite system when polymer chains are present. In the\npresence of polymer chain segments, the binding energy is not “less negative” but rather\n“more negative”. What we intended to convey is that a larger absolute value of the\nbinding energy corresponds to a stronger binding effect. This led to our previous\nincorrect description, which we have now corrected in the revised manuscript. We\nsincerely thank the reviewer again for this valuable suggestion.\nThe specific changes in the manuscript are as follows:\n“Quantum chemical calculations were first conducted to evaluate the binding\nenergy of [EMIM]+ [TFSI]- in the presence and absence of polymer chains with the\ncomputational details provided in Supporting Information. As shown in Fig. 2c, the\nbinding energy of [EMIM]+ [TFSI]- under polymer chain influence was significantly\nmore negative compared to its isolated state. Given that binding energy values were\ntypically negative with more negative values indicating stronger binding and structural\nstability, this result suggested that the cations and anions of the IL readily dissociate via\nfluorine-cation interactions without disrupting the ionic bonding between molecular\nchains. This separation mechanism ensured the integrity of the dynamic network while\nfacilitating ion mobility.”\n(2) Thank you for your careful review and constructive suggestions on our\nmanuscript, which greatly improved its quality. Our binding energy calculation formula\nis defined as the total energy of the composite system after binding minus the energies\nof each component in their isolated states. For multi-component systems, the resulting\nbinding energy refers to the overall binding energy of the system when all components\nare aggregated together. Our original intention was to use Part A and Part B to represent\nindividual components, indicating that the total energy should be subtracted by the\nenergies of all such components. However, we did not take into account the potential\nambiguity that may arise when more components are involved, and we sincerely\napologize for this oversight. We have now provided a more detailed explanation of the\nbinding energy calculation formula in the Supplementary Information.\nThe specific changes in the Supplementary Information are as follows:\nThe interaction energy of [EMIM]+[TFSI]- were calculated by the following\nEquation 1(Equation 10 in the Supplementary Information):\n𝐸 = 𝐸 −(𝐸 +𝐸 ) (1)\n𝐵𝑖𝑛𝑑(𝐼𝐿) 𝐼𝐿 + −\nwhere 𝐸 denotes the interaction energy of [EMIM]+[TFSI]- in the absence of\n𝐵𝑖𝑛𝑑(𝐼𝐿)\npolymer chains, 𝐸 , 𝐸 and 𝐸 correspond to the energies of [EMIM]+[TFSI]-,\n𝐼𝐿 + −\nisolated [EMIM]+ and isolated [TFSI]-.\nThe interaction energy of [EMIM]+[TFSI]- under polymer chain influence were\ncalculated by the following Equation 2 (Equation 11 in the Supplementary Information):\n𝐸 = 𝐸 −(𝐸 +𝐸 +𝐸 +𝐸 ) (11)\n𝐵𝑖𝑛𝑑(𝐶𝑜𝑚𝑝𝑙𝑒𝑥) 𝐶𝑜𝑚𝑝𝑙𝑒𝑥 + − 𝑇𝐹𝑀𝐵 𝑇𝐿𝐵\nwhere 𝐸 denotes the interaction energy of [EMIM]+[TFSI]- in the\n𝐵𝑖𝑛𝑑(𝐶𝑜𝑚𝑝𝑙𝑒𝑠)\npresence of polymer chains, 𝐸 , 𝐸 , 𝐸 , 𝐸 and 𝐸 correspond to the\n𝐶𝑜𝑚𝑝𝑙𝑒𝑥 + − 𝑇𝐹𝑀𝐵 𝑇𝐿𝐵\nenergies of this system, isolated [EMIM]+, isolated [TFSI]-, isolated TFMB part and\nisolated TLB part.\n• Comment #3-2:\nThe authors say: \"Initially, the simulation systems underwent structural\noptimization.\" (1) What does this mean? (2) How were the 24 or 36 polymers arranged?\n(3) What does Figure 12 show? Is it a snapshot of a single polymer or is this the 36\npolymer system? I'm guessing it is just a single polymer. (4) We are told it is\n\"optimized\" but what does this mean? Some parts of the polymer chain are extended\nand some are coiled. How was this structure created? (5) More importantly, how are\nthe 36 polymer chains arranged? Was this single polymer configuration replicated 36\ntimes? How were the 36 polymers packed together? It is impossible to judge the realism\nand quality of the simulations without these details.\n Response #3-2:\nMany thanks to the reviewer for your careful reading of our manuscript and\nrelevant remarks. Additionally, we sincerely apologize for any confusion caused by\nunclear descriptions. Below are our full responses to these issues, and we are open to\nconsideration of any further comments and suggestions\n(1) Here, we perform structural optimization for two systems that are undergoing\nsimulation calculations. Structural optimization is initially carried out by adjusting the\natomic positions based on the direction of atomic forces, moving them toward a lower\ntotal potential energy. This process is done to avoid excessively high initial energy of\nthe structure.\n(2) The polymers within the system are randomly arranged.\n(3) Figure 12 shows the snapshot of a single polymer molecular chain. Snapshots\nof the specific polymer systems are shown in the Fig. R6 below. Since the polymer\nsystems are confined within the simulation box, the box boundaries limit the visibility\nof the detailed modeling of the polymer molecular chains. Therefore, we chose to\npresent the snapshot of a single polymer molecular chain rather than that of the entire\nsystem.\nFig. R6 The snapshot of the two systems.\n(4) We reduced the total energy of the initial structure through structural\noptimization. By using Packmol, we packed a specified number of polymer chains into\na cubic simulation box. The system was then relaxed under the NPT ensemble at a\ntemperature of 300 K and a pressure of 1 atm. Driven by thermal motion, the polymer\nchains spontaneously bent and entangled, forming the structure corresponding to\nambient temperature and pressure.\n(5) The 36 polymer chains within the box are randomly arranged and generated by\nreplicating a single polymer configuration 36 times. However, due to the random\nmotion of molecules during molecular simulation, the molecular configurations\nundergo certain changes, resulting in each polymer having a distinct and completely\nrandom configuration. Subsequently, we used the Packmol software to pack these\npolymers, together with the corresponding number of ions, into a cubic box, and then\nobtained a stable initial structure through relaxation.\n• Comment #3-3:\nThe authors state that \"The diffusion coefficient is obtained from the long-time\nlimit of the MSD\". But the simulation was only run for 5 ns. It is doubtful that there is\na long-time limit with this short a simulation. The authors need to show the MSD vs.\ntime plot and show the fitted line to the data that they used to compute the diffusion\ncoefficient.\n Response #3-3:\nWe thank the reviewer for your insightful and thoughtful comments. And we\napologize for the omission of the MSD curve. We have now examined the MSD data\nand the corresponding fitted curve, as shown in the Fig. R7 below. The data fitting was\nperformed using the last one-third of the data for linear regression. The high degree of\nlinearity in the fitting results demonstrates that the simulation has essentially reached\nconvergence.\nFig. R7 The MSD curves of system 1 and system 2.\n• Comment #3-4:\n\"...relaxation for 1 nanosecond until the systems reached a steady state. Following\nthis, over a period of 5 nanoseconds, the systems were uniformly stretched in the x\ndirection to twice their original length, while maintaining a pressure of 1 atmosphere in\nthe y and z directions. Then, over another 5 nanoseconds, the systems were compressed\nback to their original length in the x-direction, again maintaining a pressure of 1\natmosphere in the y and z directions. For the structures before stretching, after\nstretching, and after compression, we conducted 5-nanosecond simulations under the\nNVT ensemble\".\nMy questions about this text are: (1) what does \"steady state\" mean? 1 ns is very\nshort. Again this goes back to my previous comment about how the polymer chains\nwere arranged relative to one another. Steady state could mean the box volume stops\nchanging, or the total energy converges to a steady value, or that the RMSD of the\npolymers stops changing, or any number of other criteria. It is hard to believe that 1 ns\nis enough to reach equilibrium. (2) how do the y and z box sizes compare before and\nafter the stretching and compression? (3) After the 5 ns stretching, was the system\nimmediately compressed or was the 5 ns NVT run done in-between with the\ncompression done from after the NVT run? Or is the NVT run done on a pull-out of the\nNPT configuration without interrupting the stretch/compress run? This isn't clear. (4) If\nthe NVT runs to measure conductivity are done immediately following the stretching\nrun (and the compression run), presumably the system is not at equilibrium yet (the\nsystem would relax to equilibrium after the stretching is stopped over some time scale)\nso this would be a non-equilibrium simulation which is not appropriate.\n Response #3-4:\nWe thank the reviewer for your insightful and thoughtful comments. We have\ncarefully checked all the details of the simulation calculations, and we apologize for\nany confusion caused by our unclear description. We have refined the simulation details,\nand below are our point-by-point responses.\n(1) We define the steady state as the condition in which the volume of the\nsimulation box no longer changes and the total energy has converged. The convergence\nprocess is illustrated in the Fig. R8 below, from which it can be observed that the system\nreaches a steady state after approximately 1 ns.\nFig. R8 The changes of energy (a) and volume (b) of system 1 as well as energy (c) and\nvolume (d) of system 2.\n(2) During the stretching process, the lengths in the y and z directions decrease.\nDuring the compression process, the lengths in the y and z directions increase. After\ncompression, the lengths in the y and z directions return to the values observed before\nthe stretching began. The changes in the lengths of the box before and after stretching\nand compression are shown in the Fig. R9 below.\nFig. R9 The changes of the box sizes of system 1 (a) and system 2 (b).\n(3) After 5 ns of stretching, the stretched system configuration was saved. This\nconfiguration was then used for both performing the MSD calculation and carrying out\nthe compression simulation. In this sense, the NVT simulations can be considered as\nextracted from the NPT configuration without interrupting the ongoing stretching and\ncompression process.\n(4) Due to the presence of numerous small ionic liquid molecules in the simulated\nsystem, the response to stretching and compression is very fast. In fact, after the\nstretching is stopped, the system relaxes to equilibrium within a very short time.\nFurthermore, the high degree of linearity confirms that the simulation has essentially\nreached convergence (Fig. R10).\nFig. R10 The MSD curves of system 1 and system 2.\n Reviewer #4 (Remarks to the author):\nThis study presents a supramolecular polyurethane engineered with fluorine-rich\nsegments that form electrostatic crosslinks with positively charged ionic groups at\npolymer chain terminals and establish fluorine-dipole interactions with blended ionic\nliquid to stabilize ion transport pathways. Many interesting results are presented.\nHowever, the reviewer suggests the following points to be seriously addressed.\n Response #4:\nWe greatly appreciate your positive assessments and constructive comments on\nour work. To fully address your comments and concerns, we have updated the\nmanuscript and provided detailed explanations. Moreover, we are sincerely honored by\nthe positive comments from the reviewers regarding our work. The main changes in the\nrevised manuscript or supporting information are in yellow (revised contents) and blue\n(supplementary contents). In the following content, we will provide in-depth\ndiscussions and responses to each of your questions.\n• Comment #4-1:\nIn Abstract, The author points out: “The resulting ionically conductive elastomer\ncombines shape memory capacity, self-healing property, and exceptional cryogenic\ntolerance, retaining robust mechanical strength (~32.31 MPa), toughness (~107.05 MJ\nm-3) and substantial ionic conductivity even at -40 ℃. Notably, it achieves an\nunprecedented combination of record-low operational temperatures (-40 ℃ to -30 ℃)\nand an ultrahigh temperature coefficient of resistance (TCR = 8.05 % ℃-1), setting a\nnew benchmark for cryogenic sensing materials.” However, the author did not focus on\ntesting the various characteristics and application under low-temperature conditions.\n Response #4-1:\nThank you for your careful review and constructive suggestions on our manuscript,\nwhich greatly improved its quality. We agree that a comprehensive evaluation of the\nmaterial's performance under cryogenic conditions is crucial to support our claims in\nthe abstract. In response to this valuable feedback, we have now significantly expanded\nthe low-temperature testing section to provide a more robust demonstration of its\ncryogenic stability and applicability.\nSpecifically, we have supplemented the manuscript with the following key\nexperiment:\nLong-term Stability Test at -40 ℃: To directly prove the exceptional cryogenic\ntolerance and operational stability, we conducted a continuous 120-minute resistance\nmonitoring test at -40 ℃. The results show an outstandingly minimal resistance drift\nof less than 4% over this extended period, which strongly confirms the material's\nrobustness for long-term use in ultra-low temperature environments. This new data has\nbeen added to Fig. R11 (Fig. 17 in the revised Supplementary Information).\nFig. R11 (Fig. 17 in the revised Supplementary Information) ΔR/R of PUFT2-\n0\n50%IL at -40 ℃ for 120 minutes.\nThis new data, combined with the previously presented:\n(1) High ionic conductivity at -40 ℃ (Fig. 5b in the manuscript).\n(2) Ultrahigh TCR value (8.05 % ℃-1) in the cryogenic range (-40 to -30 ℃) (Figs.\n5c-d in the manuscript).\n(3) Stable performance over five rapid thermal cycles (Figs. 5f in the\nmanuscript).\nThese collectively forms a comprehensive suite of evidence that fully supports our\nclaim of \"exceptional cryogenic tolerance\" and its suitability as a benchmark material\nfor cryogenic sensing applications. The long-term stability test specifically addresses\nthe concern about durability under sustained low-temperature operation, moving\nbeyond mere functionality to proven reliability.\nWe have revised the relevant results section and discussion to incorporate the\nanalysis of this new finding.\nThe specific changes in the manuscript are as follows:\n“Furthermore, its outstanding low-temperature stability was confirmed by a\nminimal resistance change of less than 4% over a continuous 120-minute test at -40°C\n(Supplementary Fig. 17), underscoring its strong potential for reliable applications in\ncryogenic environments.”\n• Comment #4-2:\nIn Section 2.6 (Highly sensitive ionic skin), there are too few tests on static\ncharacteristics, and basic characteristic indicators such as sensitivity, linearity and range\nare missing.\n Response #4-2:\nMany thanks to the reviewer for your careful reading of our manuscript and\nrelevant remarks. We apologize for missing these important parameters. Based on your\nadvice, we have added experiments and descriptions in the revised manuscript. The\nsensitivity, linearity and range are shown in Fig. R12 (Fig. 6a in the revised\nmanuscript).\nFig. R12 (Fig. 6a in the revised manuscript) Strain sensitivity curve of PUFT -50%IL.\n“Meanwhile, the sensitivity of PUFT2 was evaluated under strain conditions\nranging from 0 to 500%. The gauge factor (GF) was determined to be 1.93 within the\n0-100% strain range and was observed to increase with increasing strain, reaching a\nvalue of 2.41 at 200-300% strain.”\n• Comment #4-3:\nIn Section 2.7 (Self-healing properties), Fig. 6d-e and Fig. 6f should be marked as\nFig. 7d-e and Fig. 7f. This mistake should be due to the author's carelessness.\n Response #4-3:\nMany thanks to the reviewer for your careful reading of our manuscript and\nrelevant remarks. We apologize for the inconvenience caused to the reviewer by the\nincorrect numbering. We have corrected it in the revised manuscript.\nThe specific changes in the manuscript are as follows:\n“Healed samples (1 hour healing) successfully monitored repetitive finger and\nwrist flexion cycles, with ΔR/R0 amplitude consistency confirming stable real-time\nmotion detection (Supplementary Figs. 19-20). Furthermore, the healed material\nmaintained stable ECG signal acquisition capability post-repair, as evidenced by\nunaltered signal fidelity in Supplementary Fig. 21.”\n• Comment #4-4:\nAs shown in Fig.5 and Fig.6, compared with the gauge factor, the resistance\ntemperature coefficient is too strong. In the actual application process, it will affect its\napplication in strain or pressure.\n Response #4-4:\nWe appreciate the reviewer’s insightful comment. Indeed, the high temperature\ncoefficient of resistance (TCR = 8.05 % ℃-1 in the range of –40 to –30 ℃) indicates\nthat the material is extremely sensitive to temperature changes, which is beneficial for\ncryogenic temperature sensing but may introduce thermal interference in strain or\npressure sensing under fluctuating temperatures. Based on your advice, we would like\nto clarify the scope and operating modes evaluated in this work.\nWe would like to clarify that a primary objective of this work was to develop a\nfunctional material capable of reliable operation under cryogenic conditions. Therefore,\nits exceptional cryogenic temperature-sensing performance (exemplified by the ultra-\nhigh TCR and record-low operational temperature) is a core innovation and a\ndeliberately pursued feature of our work, rather than an unintended side effect.\nThat said, we also value its strain-sensing capability. As the reviewer pointed out,\nthe strong TCR implies that this sensor is more suitable for strain detection\nin environments with minimal temperature fluctuations. For instance, when attached to\nhuman skin for motion monitoring (where temperature is stable around ~25-35 °C),\nthe temperature interference is negligible. In response, we have supplemented the\nmanuscript with a new linear fitting analysis of the resistance under stretching at room\ntemperature (Fig. 6a in the revised manuscript), which confirms its reliable and linear\nstrain response under such conditions, making it fully adequate for biomechanical\nmonitoring like human motion detection.\nIn conclusion, we believe this material opens a unique application niche:\nits primary application is as a high-performance cryogenic temperature sensor,\nwhile secondarily, it can function as a reliable strain sensor in thermally stable\nenvironments. Future work could focus on decoupling the signals through integrated\nreference units or compensation algorithms, but that extends beyond the scope of this\nmanuscript as a materials study. We are grateful to the reviewer for prompting us to\nprovide a clearer explanation of this important aspect.\n• Comment #4-5:\nIt is recommended to describe the material properties first and then the application\nproperties. For example, Section 2.7 should be described earlier.\n Response #4-5:\nThank you very much for your careful review and constructive suggestions\nregarding our manuscript. In the revision, we have reorganized the “Section Results”\nso that material properties precede application demonstrations. Specifically, the\nsections now appear in the following order:\n(1) Self-Healing, damping and shape memory properties (previously Sec. 2.4)\nThe specific changes in the manuscript are as follows:\n“Driven by reversible intermolecular multi-hydrogen bonding and ionic\ninteractions, the PUFT2-50%IL ionic elastomer exhibits exceptional self-healing. Real-\ntime resistance measurements during multiple surgical blade cuts showed that the\nmaterial restored its original resistance within 0.5 seconds, confirming rapid electrical\nrecovery (Fig. 4a). To quantify self-healing efficiency, PUFT -50%IL samples were cut\n2\ninto two pieces, then the two completely separated ruptured surfaces were brought into\nfull contact at room temperature. As shown in Fig. 4b, tensile strength and strain\ngradually increased with healing time. After 1 hour, the tensile strength and strain\nrecovered to 0.74 MPa and 760%, respectively. After 3 hours, they reached 1.11 MPa\nand 912%, similar to the original sample. For strain sensors, stretchability is a critical\nparameter. Based on strain recovery, the healing efficiency reached 85.4% after 1 hour\nand 97.4% after 3 hours (Fig. 4c), indicating nearly ideal self-healing performance.”\n(2) Low temperature tolerance and temperature sensitivity (previously Sec. 2.5)\n(3) Highly sensitive and healing ionic skin (previously Sec. 2.6)\nThe specific changes in the manuscript are as follows:\n“Previously, the excellent self-healing performance of PUFT -50%IL had already\n2\nbeen confirmed. Therefore, it could be inferred that PUFT -50%IL would still be able\n2\nto maintain its sensing ability after healing. Healed samples (1 hour healing)\nsuccessfully monitored repetitive finger and wrist flexion cycles, with ΔR/R amplitude\n0\nconsistency confirming stable real-time motion detection (Fig. 7d–e). Furthermore, the\nhealed material maintained stable ECG signal acquisition capability post-repair, as\nevidenced by unaltered signal fidelity in Fig. 7f. These results collectively demonstrated\nthat the self-healing mechanism preserved both mechanical integrity and sensing\nfunctionality, significantly extending the material’s applicability in demanding\nenvironments. This work substantiated that hierarchical dynamic bonding networks not\nonly enabled robust self-healing but also ensured performance retention across\nmechanical, electrical, and bio-signal sensing modalities, establishing PUFT -50%IL\n2\nas a durable platform for next-generation adaptive sensors.”\nThis reordering separates intrinsic material behaviors (temperature response,\nhealing) from application figures, while preserving the key message that healed devices\nremain fully functional for sensing. We hope this addresses the reviewer’s concern.\n• Comment #4-6:\nGenerally, ionic conductors are commonly used in capacitive sensors,\ndemonstrating high sensitivity, but are less applied in resistive principles. It can also be\nroughly seen from the results in Section 2.6 (Highly sensitive ionic skin) that the gauge\nfactor of the designed sensor is relatively low.\n Response #4-6:\nThanks for your careful review and constructive suggestions regarding our\nmanuscript. We fully agree that ionic conductors are predominantly employed in\ncapacitive sensing modes due to their ability to achieve high sensitivity, particularly for\ndetecting minute pressures or strains.\nHowever, the primary objective of this part of our work was to explore the often-\noverlooked potential of ionic conductors in resistive sensing schemes. We intentionally\ndesigned our sensor for applications that require large-scale deformations (e.g., >200%\nstrain, such as in joint movement monitoring or soft robotics). To thoroughly\ndemonstrate this advantage, we have now supplemented the sensitivity data under 0-\n500% strain in the revised manuscript, as kindly suggested by the reviewer (Fig. R13).\nThe new results clearly show that our sensor exhibits a stable and monotonic increase\nin resistance with strain, with a GF of 1.93 in the 0-100% range, increasing to 2.41 at\n200-300% strain. Most importantly, it maintains excellent linearity and reliability even\nup to 500% strain, a performance that is challenging to achieve with many capacitive\nsensors requiring fixed parallel plates.\nTherefore, we posit that the value of our work lies not in competing with capacitive\nionic sensors for high sensitivity at small strains, but in demonstrating a viable and\nrobust resistive-type ionic skin that is uniquely suited for high-strain applications. This\nexpands the application repertoire of ionic conductors beyond the conventional\ncapacitive paradigm.\nFig. R13 (Fig. 6a in the revised manuscript) Strain sensitivity curve of PUFT -50%IL.\n2\n• Comment #4-7:\nThe application part of the material is too simple to support its application value.\n Response #4-7:\nFirst of all, I would like to thank the reviewer for pointing out this problem. We\nagree that demonstrating tangible application value is crucial. In our study, we aimed\nto highlight the application potential not through a single complex device, but by\ndemonstrating a suite of advanced material properties that, when combined, directly\naddress critical limitations in the field of soft electronics and wearable technology.\nSpecifically, the value of our material platform lies in its multifunctionality and\nenhanced durability, which are key for practical applications:\n(1) Its core functionality as a highly sensitive ionic skin for wearable motion\nsensing (Figs. 4d-h in the revised manuscript) and temperature sensing (Fig. 4c in\nthe revised manuscript) establishes its primary application domain.\n(2) The exceptional low-temperature tolerance (Fig. 4f in the revised manuscript)\nsignificantly expands its application scenarios from room condition to harsh\nenvironments (e.g., winter sports, polar exploration), a domain where most hydrogel-\nbased devices fail.\n(3) The self-healing capability and full recovery of sensing performance post-\nhealing (Figs. 4a-c and Supplementary Figs. 17-19) are not merely demonstrations\nbut are central to its application value. They directly solve the problem of mechanical\nfailure and short lifetime, promising reduced maintenance costs and longer service\nhours for wearable devices.\n(4) The shape memory property (Fig. 4g) further adds a dimension for intelligent\napplications, such as pre-programmed fitting or actuation.\nTo better integrate these points and address the reviewer's concern, we have now\nadded the resistance changes at -40℃ for two hours to prove stability (Fig. R14).\nFig. R14 (Fig.17 in the revised Supplementary Information) ΔR/R of PUFT2-\n0\n50%IL at -40 ℃ for 120 minutes.\nIn the end, thank you again for prompting us to better articulate this point."}, {"round": 2, "reviewer_comments": {"reviewer_2": "The authors have made appropriate modifications to the paper and given reasonable explanations to the questions raised by reviewers, which resolved the concerns of the reviewers. Therefore, it should be suitable for publication in Nature Communications.\n[Editor's Note: This Reviewer was asked to evaluate the the response to Reviewer #4's comments, and provided the following:]\nIn my opinion, the authors have conducted a thorough and proper revision in response to the reviewers’ comments. The incorporation of new experimental data, expanded discussions, and structural refinements has substantially elevated the overall quality and clarity of the work. In particular, the concerns raised by Reviewer #4 have been addressed comprehensively and convincingly.\nThe manuscript now includes critical low-temperature stability tests and essential sensing metrics, such as sensitivity and linear sensing range, that firmly establish the material's practical applicability. These additions not only validate the authors’ initial claims but also significantly bolster the work's completeness and scientific rigor. The revised narrative more logically progresses from intrinsic material properties to functional applications, greatly improving readability and impact.\nTherefore, I believe this revised manuscript is suitable for publication in Nature Communications.", "reviewer_3": "This review evaluates the molecular modeling aspects of the revised manuscript. I now consider the manuscript acceptable for publication but I would appreciate if the authors could address the following comments:\nIn the rebuttal the authors say \"Structural optimization is initially carried out by adjusting the atomic positions based on the direction of atomic forces, moving them toward a lower total potential energy. This process is done to avoid excessively high initial energy of the structure.\" This is fine if the sole purpose is to avoid high initial energy structures. But this local relaxation will not cause an extended chain to become coiled or vice versa. Probably \"local relaxation\" is a better term here than optimization.\nFig. R6 is helpful so I can see what the full system is. Thank you for including this. However, in the rebuttal the authors say \" Driven by thermal motion, the polymer chains spontaneously bent and entangled, forming the structure corresponding to ambient temperature and pressure.\" It is far from obvious that any entanglement occurs. Can the authors substantiate this with a measurable quantity, in other words can they quantify entanglement?\nIn Fig. R7 the authors show the MSD vs. time plot which I asked for. This is appreciated. However, since the diffusion constant is quite low (4 orders of magnitude lower than water) it is not convincing to stop the plots at 2 ns (= 2000 ps). The data sort of looks linear but it would be much more convincing if the time scale was extended.\nIn the rebuttal the authors say \"We define the steady state as the condition in which the volume of the simulation box no longer changes and the total energy has converged.\" These are very weak criteria for equilibrium. More convincing would be to track the radius of gyration of the polymer chains. This is because changes in the radius of gyration or other polymer conformational properties might not show up in the overall system volume and would be a more careful test of steady state.\nIn the rebuttal the authors say \"After compression, the lengths in the y and z directions return to the values observed before the stretching began\". Were the y and z box lengths constrained to have the same value as each other? Otherwise it would be surprising for the y and z values to return exactly to their starting values. If such a constraint was used it should be stated."}, "author_response": "Response to Previous Reviewers’ Comments\nDear reviewers,\nWe would like to thank all reviewers for their constructive comments and\nsuggestions. These comments are valuable and helpful for the improvement of our paper,\nand they also have important guiding significance for research. Therefore, we have\nstudied the comments carefully and have made further revisions, hoping that the\nreviewers and editors will be satisfied with the revised manuscript.\nOur point-by-point responses to these comments are provided below, and the main\nchanges in the revised manuscript are marked in yellow (revised contents) and blue\n(supplementary contents). In any case, we are open to consideration of any further\ncomments and suggestions. We look forward to hearing from you on the status of our\nsubmission.\nSincerely,\nXinrui Zhang and Jing Xu\n Reviewer #3 (Remarks to the author):\nThis review evaluates the molecular modeling aspects of the revised manuscript.\nI now consider the manuscript acceptable for publication but I would appreciate if the\nauthors could address the following comments:\n Response #3:\nThank you for your affirmation and reviewing the manuscript in your busy\nschedule and provided us with your professional comments and suggestions, the\nquestions and the certain responses were presented in following part orderly.\n• Comment #3-1：\nIn the rebuttal the authors say \"Structural optimization is initially carried out by\nadjusting the atomic positions based on the direction of atomic forces, moving them\ntoward a lower total potential energy. This process is done to avoid excessively high\ninitial energy of the structure.\" This is fine if the sole purpose is to avoid high initial\nenergy structures. But this local relaxation will not cause an extended chain to become\ncoiled or vice versa. Probably \"local relaxation\" is a better term here than optimization.\n Response #3-1:\nWe sincerely thank you for your insightful comments and have thoughtfully\nrevised the manuscript based on the reviewers' comments. We have corrected it in the\nrevised Supplementary Information.\nThe specific changes in the Supplementary Information are as follows:\n“Initially, the simulation systems underwent structural local relaxation.”\n• Comment #3-2：\nFig. R6 is helpful so I can see what the full system is. Thank you for including this.\nHowever, in the rebuttal the authors say \" Driven by thermal motion, the polymer chains\nspontaneously bent and entangled, forming the structure corresponding to ambient\ntemperature and pressure.\" It is far from obvious that any entanglement occurs. Can the\nauthors substantiate this with a measurable quantity, in other words can they quantify\nentanglement?\n Response #3-2:\nWe thank the reviewer for raising this critical point. We agree that the term\n\"entanglement\" may seem qualitative without rigorous quantitative support. In our\ncontext, we used it to describe the process whereby polymer chains form a stable and\nnatural network structure. We tracked the evolution of the radius of gyration of the\npolymer chains. The observed decrease and eventual stabilization of the radius of\ngyration indicate a transition from a relatively extended initial conformation to a more\ncompact one, consistent with the chains rearranging through interactions to form a\nstable structure.\nTo enhance accuracy, we have modified the description to specify that \"the chains\nunderwent relaxation, reaching a stable radius of gyration,\" as this more precisely\nconforms to the conformational state we measured.\nThe specific changes in the manuscript are as follows:\n“The time evolution of the average radius of gyration for the polymer molecules,\n[EMIM]+, and [TFSI]- was analyzed during the relaxation process to confirm that the\nsystem had reached equilibrium. ”\n• Comment #3-3：\nIn Fig. R7 the authors show the MSD vs. time plot which I asked for. This is\nappreciated. However, since the diffusion constant is quite low (4 orders of magnitude\nlower than water) it is not convincing to stop the plots at 2 ns (= 2000 ps). The data sort\nof looks linear but it would be much more convincing if the time scale was extended.\n Response #3-3:\nThanks again for your careful review and constructive suggestions on our\nmanuscript, which greatly improved its quality. In accordance with the reviewer's\nsuggestion, we have extended the MSD simulation time to 15 ns. Concurrently, the time\nspan for MSD calculation was increased to 10 ns. The resulting MSD curves for the\ndifferent ion types are shown in Fig. R1, which demonstrate good linearity.\nFig. R1 The MSD curves of system 1 and system 2.\n• Comment #3-4：\nIn the rebuttal the authors say \"We define the steady state as the condition in which\nthe volume of the simulation box no longer changes and the total energy has\nconverged.\" These are very weak criteria for equilibrium. More convincing would be\nto track the radius of gyration of the polymer chains. This is because changes in the\nradius of gyration or other polymer conformational properties might not show up in the\noverall system volume and would be a more careful test of steady state.\n Response #3-4:\nWe sincerely thank you for your kind suggestions. Following the reviewer's\nsuggestion, we tracked the radius of gyration of the polymer chains, which rapidly\nconverged to a stable value (Fig. R2), confirming the system had reached equilibrium.\nThis new stability criterion has been detailed accordingly in the revised Supplementary\nInformation.\nFig. R2 The radius of gyration of system 1 and system 2.\nThe specific changes in the manuscript are as follows:\n“The time evolution of the average radius of gyration for the polymer molecules,\n[EMIM]+, and [TFSI]- was analyzed during the relaxation process to confirm that the\nsystem had reached equilibrium. ”\n• Comment #3-5：\nIn the rebuttal the authors say \"After compression, the lengths in the y and z\ndirections return to the values observed before the stretching began\". Were the y and z\nbox lengths constrained to have the same value as each other? Otherwise it would be\nsurprising for the y and z values to return exactly to their starting values. If such a\nconstraint was used it should be stated.\n Response #3-5:\nThank you very much for your careful review and constructive suggestions\nregarding our manuscript. We apologize for the lack of this information. Yes, this\nconstraint is indeed in place because we aim to simulate uniaxial tensile loading. To\nmaintain stable box dimensions, we constrained the dimensions of y and z directions to\nremain consistent. We have now added this detail in the revised Supplementary\nInformation.\nThe specific changes in the Supplementary Information are as follows:\n“Following this, over a period of 5 nanoseconds, the systems were uniformly\nstretched in the x direction to twice their original length, while maintaining a pressure\nof 1 atmosphere in the y and z directions. To maintain stable box dimensions, we\nconstrained the dimensions of y and z directions to remain consistent. Then, over\nanother 5 nanoseconds, the systems were compressed back to their original length in\nthe x-direction…”"}, {"round": 3, "reviewer_comments": {"reviewer_3": "This review evaluates the molecular modeling aspects of the revised manuscript. The authors have comprehensively addressed all of my comments which is greatly appreciated. I have no further comments or suggestions. I recommend publication."}, "author_response": ""}]

Decision letter