Introduction

Sustainable polymers are essential for harmonizing industrial advancement with ecological conservation1,2,3,4,5,6. Developing environmentally friendly production and preparation methods is crucial for achieving comprehensive sustainability7,8,9,10,11. In this context, Waterborne polyurethane elastomers (WPUE) have emerged as a promising option for the sustainable development of the polyurethane family, as they utilize water instead of costly and harmful organic solvents12,13. Despite its advantages in minimizing environmental impact, WPUE is often viewed as lacking the mechanical strength of thermoplastic polyurethane elastomers (TPUE) produced through traditional synthetic processes using organic solvents. This perception significantly limits the applications of this sustainable alternative14. The development of high-performance WPUE has the potential to revolutionize the sourcing of sustainable elastomeric materials.

When polymers are stretched, their microstructures can change significantly, enhancing the material’s strength and toughness15,16,17,18,19. For TPUE, strain-induced crystallization (SIC) is a key mechanism that enhances mechanical properties, enabling strength increases of 40–80 MPa while maintaining toughness levels between 300 and 500 MJ/m³ 20,21,22,23,24,25. However, this process is seldom observed in WPUE due to the presence of hydrophilic ionic functional groups, which are essential for water processing but disrupts the alignment of polymer chains26 necessary for crystallization because of its non-planar nature. Additionally, this non-planar structure and the steric hindrance of the hydrophilic segments modulate the arrangement of the hard and soft segments in WPUE, resulting in a different microstructure compared to TPUE (Supplementary Fig. 1) and complicating the SIC process.

Here, we demonstrate the potential of harnessing molecular design to create high-performance WPUE (HPWPUE) with exceptional toughness. Our success lies in the use of symmetrical monomers that enhance the formation of a dynamic biphase structure, promoting the formation of hydrogen bonding during stretching and enabling crystallization at higher strains. Under tensile stress, the hard-segment domains in HPWPUE break apart into separate hard segments that mix with the soft segments. Notably, this process continues until a high stretch ratio (λ) of 20 without fracturing, attributable to the hydrogen bonds formed throughout the process. This ability to endure significant strains enhances the alignment of polymer segments during stretching. As λ exceeds 13, the hard and soft segments align and stack along the strain direction, and at λ ~ 19, co-crystallization occurs, significantly enhancing both tensile strength and toughness. We term this phenomenon delayed crystallization response (DCR) (Fig. 1a), as crystallization in SIC typically occurs at a much lower λ of ~427,28,29,30,31,32,33 (Supplementary Fig. 1e).

Fig. 1: Microstructural changes in our HPWPUE (high-performance waterborne polyurethane elastomer) during the delayed crystallization response (DCR) process and their mechanical properties.
Fig. 1: Microstructural changes in our HPWPUE (high-performance waterborne polyurethane elastomer) during the delayed crystallization response (DCR) process and their mechanical properties.
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a Schematic illustrations depicting the evolution of microstructure during the DCR process. The leftmost and rightmost photographs show the HPWPUE in its virgin and highly stretched states (λ = 30), with the crystalline regions highlighted in pink. b Photograph of 1 mm thick HPWPUE showcasing its transparency. c Scatter plot of tensile strength (σe, max) versus elongation at break (εe, max) for various commercial TPUEs. The meanings of the symbols and additional information about these elastomers can be found in Supplementary Table 1. d The synthetic route for HPWPUE is shown, with further details available in the Methods section. e An engineering stress-strain (σe-εe) curve of our HPWPUE is shown until break. Inset: Photograph of a HPWPUE sample (81.7 mg by weight) supporting a 5-kg weight. f A true tensile stress-strain (σt-εt) curve of HPWPUE. g Left: Force-displacement curves from puncture tests comparing HPWPUE with two commercial TPUEs. Right: A photograph demonstrating the puncture resistance of HPWPUE, colored with blue dye for better visibility. Scale bar in the leftmost photograph of panel (a) is 2 mm, in (b) it is 20 cm, and in the other panels, it is 2 cm.

With DCR, our HPWPUE can be processed into transparent thin films at a large scale (1.5 m × 1.5 m × 1 mm) using commercially available feedstocks (Fig. 1b). It achieves an impressive maximum stretch ratio of 31.9, a tensile strength of 81.8 MPa, and a toughness of 0.959 GJm⁻³, surpassing the performance of typical commercial TPUE (Fig. 1c; see Supplementary Table 1 for details).

Results

HPWPUE material and its mechanical performance

The synthetic route of our HPWPUE is shown in Fig. 1d. We use hexamethylene diisocyanate (HDI) and methylene-bis(4-cyclohexylisocyanate) (HMDI) as isocyanates, poly (adipic acid-hexylene glycol-neopentyl glycol) (PAA-PHG-PNPG) (Number average molecular weight (Mn) = 2000 g/mol) as polyol, and 2,2-bis(hydroxymethyl) propionic acid (DMPA), ethylenediamine (EDA), and 2-(2-aminoethylamino) ethanol (AEEA) as chain extenders. Nuclear magnetic resonance data can be found in Supplementary Fig. 2. Due to its hydrophilic nature, DMPA enables the stability of HPWPUE in aqueous solutions (Supplementary Fig. 3). Using various measurements and analyses, we confirm the successful synthesis of the product (Supplementary Fig. 4) and establish that HPWPUE is thermally stable up to ~250 °C (Supplementary Fig. 5). The glass transition temperature (Tg) of HPWPUE is below room temperature (Supplementary Fig. 1b), indicating its elastomer behavior. Furthermore, it exhibits a non-crystalline, microphase-separated structure (Supplementary Fig. 1b and Supplementary Fig. 6). Stress-strain cycle measurements at various delay times (Supplementary Fig. 7) following an initial cycle confirm the presence of reversible inter-segment bonds, which we attribute to inter-segment hydrogen bonds. These bonds function as sacrificial bonds (as elaborated below), preventing molecular bond breakage during tensile stretching and enhancing toughness of the elastomer.

To evaluate the mechanical strength of HPWPUE, we first conducted uniaxial tensile tests at an extension rate of 50 mm/min. Notably, HPWPUE demonstrates impressive tensile strength (σe, max = 81.8 MPa), elongation at break (εe, max = 30.9), and toughness (τ = 0.959 GJ/m³) (Fig. 1e). We also observed that HPWPUE starts off completely transparent (leftmost photograph in Fig. 1a) but becomes opaque as the stretch ratio increases to 30 (rightmost photograph in Fig. 1a), a change attributed to strain-induced crystallization. Remarkably, the toughness of HPWPUE is 2.7 times greater than that of Darwin’s bark spider silk, the toughest known natural material34. During tensile tests, the sample’s cross-sectional area significantly shrinks, making true tensile strength a better measure of HPWPUE’s actual strength, which is determined to be 2.65 GPa. We determine that a HPWPUE specimen weighing only 81.7 mg can lift a load 61,200 times its weight (Fig. 1f). Additionally, we conducted puncture resistance tests, revealing that HPWPUE is difficult to pierce (right side of Fig. 1g). Force-distance curves from these tests show that HPWPUE outperforms commercial TPUE in puncture resistance (left side of Fig. 1g). Furthermore, HPWPUE exhibits a tear energy of 238.7 kJ/m2 and a fracture energy of 147.3 kJ/m2, which are significantly higher than those of human tendons and ligaments (20–30 kJ/m2) (Supplementary Fig. 8).

Elucidating the delayed crystallization response (DCR) process

We believe that not all WPUE are suitable for the DCR as a reinforcement strategy. To investigate the requirements for this strategy, we synthesized additional WPU polymers using different chain extenders in place of the original EDA. The new chain extenders are diethylenetriamine (DETA), N, N’-di-tert-butylethylenediamine (DBEDA), and isophorone diamine (IPDA(NH)). The chemical structures of these extenders are shown in Fig. 2a. The rationale for selecting these chain extenders is as follows: DETA has an extra −NH− group compared to EDA, which enhances cross-linking density. DBEDA contains tertiary nitrogen that may hinder hydrogen bond formation and the absence of urea groups in WPU-DBEDA may also reduce interchain bonding. Lastly, IPDA has an asymmetric structure, which allows us to explore how molecular symmetry affects the DCR process. We also synthesized new WPUE by substituting the original isocyanates, HDI and HMDI, in HPWPUE with one of the following diisocyanates: hexamethylene diisocyanate (HDI), methylene-bis(4-cyclohexyl isocyanate) (HMDI), isophorone diisocyanate (IPDI(OCN)), and diphenylmethane diisocyanate (MDI). Among these, HDI is an aliphatic diisocyanate, HMDI is a cycloaliphatic diisocyanate, MDI is an aromatic diisocyanate, and IPDI(OCN) is an asymmetrically structured diisocyanate (see Fig. 2a). These four diisocyanates allow us to examine the influence of diisocyanate type on the DCR process. Gel permeation chromatograph (GPC) shows that all the new WPUE have approximately the same Mn of ~20 kg/mol and a polydispersity index of ~2 (Supplementary Table 2), similar to HPWPUE, ensuring that variations in molecular weight will have minimal impact performance. The resulting polymers are designated as WPU-X, with X representing the specific monomer used.

Fig. 2: Prerequisites for reinforcement by DCR and HPWPUE phase deformation.
Fig. 2: Prerequisites for reinforcement by DCR and HPWPUE phase deformation.
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a Chemical structure of HPWPUE and its variants, highlighting the segments derived from the varying chain extenders (left) and isocyanates (right) investigated. b Stress-strain curves for HPWPUE and its seven variants: WPU-DBEDA, WPU-DETA, WPU-IPDA(NH), WPU-HMDI, WPU-HDI, WPU-MDI, and WPU-IPDI(OCN). c Bar chart comparing the σe, max and toughness values for HPWPUE and its variants. The error bars represent the standard deviation calculated from three independent measurements. d Polarized optical microscopic images of HPWPUE under varying horizontal stress ratios, with the analyzer oriented vertically (scale bar = 200 μm). e Differential modulus (dσe /dεe) plotted against εe. f Haward-Thackray plot of true stress versus (λ2 −1/λ), with an inset magnifying the range 0 ≤ λ2 -1/λ ≤ 100. The blue region represents the HPWPUE in the amorphous phase, while the red region indicates crystallization of the HPWPUE.

Tensile testing of these WPUE shows that HPWPUE exhibits significantly higher toughness and strength compared to the new WPUEs (Fig. 2b, c). In WPU-DETA, high cross-linking causes the material to fracture at small elongations, because the cross-linked network restricts the movement of the polymer chain segments. In WPU-DBEDA, insufficient hydrogen bonding leads to premature fracture of the material. In WPU-IPDA(NH), reinforcement does not occur at the stretch ratios where HPWPUE shows crystallization (Fig. 2b), suggesting that crystallization may be hindered in WPU-IPDA(NH). Among the WUP-X with varying diisocyanates, WPU-HMDI exhibits mechanical enhancements similar to those from SIC. This behavior can be attributed to its higher chain-segment rigidity, which enhances microphase separation and thereby impedes fragmentation of hard domains into separate segments required for DCR. WPU-HDI, which incorporates a flexible aliphatic diisocyanate, forms tightly bonded hard domains that do not effectively dissipate stress, resulting in low tensile strength and an absence of delayed crystallization. WPU-MDI exhibits an elevated elastic modulus as well as superior tensile strength and elongation at break, which can be ascribed to π–π stacking of the aromatic segments. Notably, both WPU-MDI and WPU-HDI exhibit the same qualitative trend during stretching and do not undergo self-enhancement before fracture. This observation indicates a physical network structure with strong hydrogen bonding, likely corresponding to tightly stacked hard domains within the polyurethane matrix before stretching. The strength of WPU-IPDI(OCN) is notably lower than that of WPU-HMDI due to its asymmetric structure, which hinders orderly stacking of hard domains. Additionally, the steric hindrance of the methyl side groups inhibits the formation of interchain hydrogen bonding, resulting in a lower elastic modulus compared to WPU-HMDI (Fig. 2b).

The structures of the seven WPUEs were analyzed using atomic force microscopy (AFM), Fourier Transform Infrared (FT-IR) spectroscopy, small-angle X-ray scattering (SAXS), and other characterization techniques (details provided in the discussion beneath Supplementary Figs. 9–11). Based on these analyses, we propose the following prerequisites for effective mechanical reinforcement through DCR: First, interchain hydrogen bonding must provide sufficient energy dissipation to prevent premature chain breakage during stretching. Second, the elastomer should not have a highly cross-linking structure, as this can hinder molecular rearrangements and lead to stress concentration. Third, the hard-domain structures within the polyurethane network cannot pack tightly, as this would impede their dissociation and reorganization during stretching. Finally, the elastomer chain segments should exhibit a strong tendency to crystallize upon alignment.

Additional results from stress–strain tests on HPWPUE at various temperatures reveal that temperature influences the effective implementation of the DCR process (Supplementary Fig. 12). As temperature increases to 40 and 50 °C, the stretch ratio at which strain hardening begins to occur decreases. We believe that the additional energy provided by the temperature increase disrupts some of the hard domains and disentangles the soft segments, allowing for an earlier crystallization process. Furthermore, the concave upward trend in the stress-strain curves gradually disappears at temperatures above 60 °C, preventing the realization of the delayed self-enhancement process. This phenomenon occurs because elevated temperature weakens the interactions between polymer chain segments and break interchain hydrogen bonds, leading to reductions in both tensile strength and elastic modulus. Elevated temperature also promotes chain mobility and thereby chain disentanglement, explaining decreases in both stress and strain at break (Supplementary Fig. 12b).

To understand the mechanical self-reinforcement of HPWPUE during the DCR process, we analyze the microphase behavior of HPWPUE at various stretch ratios (λ) or engineering strains (εe) where λ = 1 + εe. The polarizing optical microscopic (POM) images in Fig. 2d show the development of a notable alignment signature, indicating crystallization, when λ reaches 19. This observation is supported by AFM imaging at different stretching ratios (Supplementary Fig. 13). For 1 < λ < 12, blurred phase boundaries, smaller hard domain sizes, and weakened phase contrast are observed, indicating disruption of hard domains and mixing with soft segments. As deformation continues, the two phases start to blend. Beyond a stretch ratio of 20, a crystalline phase with distinct boundaries emerges. Figure 2e displays the differential modulus (dσe/dεe) against εe. As εe increases from zero, dσe/dεe initially decreases rapidly, indicating softening behavior. Between εe = 1.25 and 18, dσe /dεe stabilizes. Beyond this range, it increases with εe, demonstrating strain-hardening behavior, similar to how a hard solid transitions to a soft rubber after yielding35,36. We also calculated the true stress (σt) to analyze changes in the microphase structure. Treating HPWPUE as a cross-linked Gaussian-chain network, σt can be expressed as σt = Gp (λ2 −1/λ), where Gp represents the rubber elastic modulus, calculated as ρRT /<Mc>. Here, ρ is the density (1.12 g cm-3 for HPWPUE), R is the gas constant, T is the absolute temperature, and <Mc> is the average molecular weight between cross-links28,37,38,39. To relate the mechanical properties of HPWPUE to its microstructure, we plot σt against (λ2−1/λ) in Fig. 2f. The inset reveals that HPWPUE quickly yields into a softer elastomer when λ2 −1/λ ≈ 5 or εe ≈ 1, consistent with Fig. 2e. Beyond this yield point, the data follows the relationship35, σt = Y + Gp (λ2−1/λ), which is also exhibited by TPUEs28,35,40. Here, Y represents the yield stress, and Gp is the rubber elastic modulus during strain-hardening30,35,37,41. Using this formula along with the Gp value obtained from the inset of Fig. 2f, we calculate <Mc> to be 2216 g/mol. This value is close to the molecular weight of the soft segments (PAA-PHG-PNPG) in our HPWPUE, suggesting that the hard domains act as network junctions.

The distinctly higher modulus for small λ2−1/λ values below 5 is attributable to the soft segments in the soft domains wrapping around the hard domains, effectively shortening the length of the soft segments between junctions. However, this wrapping can loosen under tension. When λ2 −1/λ exceeds 300 (or λ > ~19), the slope of the σt versus (λ2 −1/λ) curve increases significantly beyond Gp, indicating the onset of crystallization, which aligns with the emergence of macroscopic alignment observed by POM at the same λ in Fig. 2d.

The presence of C = O, ether, alkoxy (acceptor), and −NH− (donor) in WPU leads to abundant hydrogen bonding in HPWPUE. These bonds form reversible cross-links during stretching, which we believe significantly influences the DCR process. To investigate this, we collect stretching-dependent FT-IR spectra of HPWPUE from λ = 1 to 31, as shown in Fig. 3a for the C = O and Fig. 3b −NH− spectral bands. The relative spectral intensities in these bands change noticeably with stretching, indicating dynamic alterations in hydrogen bonding between the C = O and −NH− groups. We conduct further analysis by calculating synchronous and asynchronous two-dimensional correlation spectra (2DCOS)42,43,44,45, with the results presented in Fig. 3c−f.

Fig. 3: Change in hydrogen bonding during stretching.
Fig. 3: Change in hydrogen bonding during stretching.
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a FT-IR spectra showing the C = O and b −NH− stretching bands at different stretch ratios (λ). c, d 2DCOS synchronous and asynchronous spectra calculated for the 1800−1600 cm-1 band of HPWPUE. e, f 2DCOS synchronous and asynchronous spectra calculated for the 3600−3150 cm-1 band of HPWPUE. Red and blue areas indicate positive and negative values, respectively, and signal intensity increases with color intensity. g Schematic diagram illustrating the changes in hydrogen bonding and binding energy, as inferred from the 2D correlation spectra during stretching.

By applying Noda’s judging rules46 to the 2DCOS of the carbonyl (C = O) band and analyzing the signs of cross-peaks in the synchronous and asynchronous spectra (Fig. 3c, d), as discussed further below Supplementary Fig. 14, we determine the response order of four sub-peaks in the C = O band with stretching: 1730/1740 cm−1 → 1693 cm−1 → 1710 cm−1 → 1674 cm−1. This sequence represents the progression from free or disordered C = O in urethane groups to free C = O in urea groups, followed by ordered C = O in urethane groups, and finally ordered C = O in urea groups (see Supplementary Fig. 14e for the specific chemical environments of these C = O groups).

From this sequence, we infer that upon stretching, the hydrogen bonds in the disordered state with the lowest bond energy respond first, followed by those of C = O in the hard domains. We also deduce the order of response for the −NH− sub-peaks as follows: 3450 cm−1 → 3260 cm−1 → 3340 cm−1 → 3320 cm−1 (Fig. 3e, f). This sequence indicates the progression from free (non-hydrogen-bonded) −NH− groups (3450 cm-1) to −NH− groups hydrogen-bonded to ether oxygen (3260 cm-1), then to disordered (hydrogen-bonded) −NH− groups (3340 cm-1), and finally to ordered (hydrogen-bonded) −NH− groups (3320 cm-1) (refer to Supplementary Fig. 14e for details on the chemical environments of these −NH− groups). Using similar reasoning, this sequence indicates that changes in hydrogen bonding between the soft and hard domains occur first, followed by changes within the hard domain.

A large number of hydrogen bonds increase the energy dissipation, allowing HPWPUE to undergo interchain breakage and morphological change during the initial stretching stage, providing the basis for delayed crystallization. As the structure of the stretched HPWPUE becomes more compact and uniform, it is evident upon the initiation of crystallization (at λ > ∼21), the hydrogen bonding transitions to a stable structure. To further elucidate the role of hydrogen bonds in the DCR process, we employed density functional theory (DFT) to calculate the hydrogen bond binding energies between various donor and acceptor groups in the elastomer. Please refer to the supplementary information for details.

The calculation results show that the binding energies of the hydrogen bonds formed by different donor/acceptor groups vary (Supplementary Table 3). In the initial state, the bonding energy of hydrogen bonds in HPWPUE are calculated as follows: -9.000 kcal/mol between urea-urea groups, -7.949 kcal/mol between urethane-urea groups, and -8.096 kcal/mol between urethane-urethane groups. After stretching, these values increase to -14.879, -13.942 and -13.490 kcal/mol, respectively, indicating transformation of the hydrogen bonds from monodentate to bidentate (Fig. 3g and Supplementary Table 3) as the polymer chains are arranged along the stress direction. It is generally agreed that the primary driving force for hard domain aggregation is the strong intermolecular interaction between the urethane units47,48, which facilitate the formation of inter-urethane hydrogen bonds and microphase separation. During the reorganization process of hydrogen bonds, the formation of additional bidentate hydrogen bonds enhances the structural ordering of the hard domains. This increase in binding energy from the bidentate hydrogen bonds can help delay early fracture in HPWPUE. As we will discuss further, this characteristic is essential for the high strength of HPWPUE.

Phase structure evolution during stretching

To further elucidate how DCR enhances mechanical properties, we examine the strain-induced microstructure evolution of HPWPUE using small-angle and wide-angle X-ray scattering (SAXS and WAXS). Figure 4a displays the 2D-SAXS patterns of HPWPUE at various λ, along with their 1D intensity profiles, with the equatorial direction corresponding to the stretching direction. These patterns reveal changes in the microphase-separated structure, consistent with prior findings (see Supplementary Fig. 15). The mean inter-domain spacing, d, is calculated using Bragg’s law d = 2π/qmax, representing the distance between hard domains before stretching. Initially, the SAXS pattern exhibited a ring pattern with maximum intensity at qmax = 0.58 nm-1 (d = 10.83 nm, λ = 1), indicating a microphase-separated structure of hard domains imbedded within a soft-segment matrix. The circular and hence isotropic nature of the 2D SAXS pattern suggests random dispersion of the hard domains.

Fig. 4: The delayed crystallization process of HPWPUE illustrated through SAXS and WAXS analyses and mechanical performance comparison with commercial TPUE and other WPUE.
Fig. 4: The delayed crystallization process of HPWPUE illustrated through SAXS and WAXS analyses and mechanical performance comparison with commercial TPUE and other WPUE.
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a 2D-SAXS patterns of HPWPUE acquired for q ≤ 2 nm-1 during stretching. b Azimuthal profiles averaged for q = 0.10–1.11 nm-1 from the patterns in panel (a) shown for varying stretch ratios (λ) indicated. c 1D SAXS profiles in the equatorial (φ = 90°) or stretching direction. d 1D SAXS profiles in the meridional (φ = 0°) (e) A subset of the 2D–WAXS patterns of HPWPUE acquired during stretching. Due to instrument occlusion, most part of the patterns in the equatorial direction is not visible. The dashed arrows in panels a and e indicate the azimuthal direction. f 1D–WAXS profiles in the meridional direction and (g) φ = 70° direction. h Scatter plot of tensile strength (σe,max) versus elongation at break (εe,max) for various stretchable TPUEs. The meanings of the symbols and details about these elastomers can be found in Supplementary Table 4. i Scatter plot of toughness versus tensile strength (σe,max) for other WPUE. Details and references for these elastomers are provided in Supplementary Table 5.

As strain increases, the 2D SAXS pattern changes from circular to oval and then to rhombus-shaped. Figure 4b presents the azimuthal 1-D profiles of the 2D pattern at varying λ, averaged over q = 0.10–1.11 nm-1. These profiles transition from no peaks at λ = 1 to two distinct peaks at 0° and 180° at larger λ signifying a shift from isotropic to anisotropic behavior. As λ increases from 0 to 13, peaks in the intensity profiles—along both the equatorial and meridional directions (Fig. 4c, d, respectively)—gradually disappear. This is attributed to the dissociation of hydrogen bonds and the breakdown of hard domains (Supplementary Fig. 1f). At the same time, the peak in Fig. 4c shifts to smaller q, indicating that the spacing between hard domains parallel to the stretching direction increases.

As λ increases from 14 to 31, the rhombus-shaped 2D-SAXS pattern enlarges and elongates, sharpening at the vertices perpendicular to the stretching direction. The final pattern resembles highly oriented lamellae. Along the stretching direction (Fig. 4c), a new peak appears around q = 0.44 nm-1 (d = 14.27 nm) and shifts to larger q values with increasing λ, indicating the emergence of new domains with a decreasing interdomain distance. We attribute these domains to newly formed crystals in HPWPUE. The elongation of the 2D-SAXS pattern in the meridional direction with λ shows that these crystalline domains aligned increasingly with the stretching direction. Additionally, the scattering intensity in the meridional direction increases, with a faint peak appearing at around q = 0.43 nm-1 (Fig. 4d). These results suggest that the hard chains rearranged due to the breakage of hard domains and concomitant formation of lamellar microdomains perpendicular to the chain axis49.

Figure 4e shows the 2D-WAXS patterns of HPWPUE during stretching, with the stretching direction still in the equatorial direction as in the 2D-SAXS patterns. Initially, a broad halo is observed around q = 11.8 nm-1 (Fig. 4f), corresponding to the average separation between soft-soft and soft-hard segments. As stretching progresses, the intensity of this halo decreases. In the region where λ > 20, a bright (blue) spot emerges in the upper right quadrant of the pattern, along with a new peak at q = 12.2 nm−1 in the 1D meridional profiles (Fig. 4f). This implies that delayed co-crystallization occurs after the soft and hard segments have been rearranged during the stretching process. Figure 4g shows the 1D profiles along φ = 70°. The scattering peaks do not disappear during stretching, but rather gradually shift towards smaller q, reaching 11.37 nm-1 at the end of the stretch. Therefore, changes in the intensity in the φ = 70° direction is attributed to the correlations between the PAA-PHG-PNPG domains, which diminishes as the co-crystallization domains between the hard and soft segments grow with stretching. This further reinforces the interpretation that co-crystallization occurs within this range of λ, with crystal orientation aligning parallel to the stretching direction.

The SAXS and WAXS results demonstrate that the microstructural evolution of HPWPUE corresponds with the DCR process illustrated in Fig. 1a during stretching. For comparison, we present a scatter plot of tensile strength (σe,max) versus elongation at break (εe,max) for various strategies aimed at enhancing polyurethane elastomers23,24,25,50,51,52,53,54,55,56,57,58,59,60 (Fig. 4h). This plot demonstrates that the DCR process yields both higher σe,max and εe,max for our HPWPUE (Fig. 4h and Supplementary Table 4), resulting in the highest toughness among these strategies.

Notably, our HPWPUE exhibits significantly greater toughness and tensile strength compared to other WPUEs reported to have high toughness (Fig. 4i and Supplementary Table 5). Furthermore, when comparing the mechanical properties of HPWPUE with those of recently developed tough gels and rubbers, we found that HPWPUE exhibited superior strength and toughness (Supplementary Fig. 16 and Supplementary Tables 6 and 7). Finally, we validated the scalability and application potential of HPWPUE. The large-scale HPWPUE retains its excellent mechanical performance, and the mechanical protective gloves coatings produced from it significantly outperform commercial products (Supplementary Figs. 17 and 18).

Discussion

Our findings underscore the potential of HPWPUE as a suitable green alternative to conventional TPUE, demonstrating that eco-friendly materials need not compromise mechanical integrity. Specifically, the HPWPUE examined achieves impressive mechanical properties through DCR. The effectiveness of DCR in this HPWPUE stems from the balanced hydrogen bonding within the molecular chains, which dissipates stress energy and promotes chain alignment, alongside a strong tendency to crystallize. By leveraging our insights into mechanical enhancement through DCR, this work establishes a blueprint for designing next-generation sustainable elastomers with superior performance. Future research will further refine the applicability of the DCR strategy, expanding its impact across diverse polymeric systems and accelerating the transition toward high-performance, environmentally responsible materials.

Methods

Materials

Hexamethylene diisocyanate (HDI) (99%), methylene-bis(4-cyclohexylisocyanate) (HMDI) (90%), 2,2-bis(hydroxymethyl)propionic acid (DMPA), and 2-(2-aminoethylamino)ethanol (AEEA) were purchased from Shanghai Meryer Chemical Technology Co., Ltd. Triethylamine (TEA) (99%), dibutyltin dilaurate (DBTDL), ethylenediamine (EDA), diethylenetriamine (DETA), N,N’-di-tert-butylethylenediamine (DBEDA), and isophorone diamine (IPDA) were purchased from Shanghai Macklin Biochemical Co., Ltd. Poly(adipic acid-hexylene glycol-neopentyl glycol) (PAA-PHG-PNPG) (\({M}_{n}\) = 2000 g/mol) were obtained from Taishan Fiberglass Inc. and used as received without further processing.

Synthesis of waterborne polyurethane elastomers

High-performance waterborne polyurethane elastomers (HPWPUE) were made using the prepolymer method. First, 30 g of PAA-PHG-PNPG (0.015 mol) was placed in a dry glass vessel with a mechanical stirrer and a mercury thermometer. It was heated to 120 °C under vacuum for 1 h to remove moisture, then cooled to 70 °C. Next, 2.52 g of HDI (0.015 mol), 7.87 g of HMDI (0.03 mol), and 30 μL of DBTDL were added and stirred for 2.5 h in a nitrogen atmosphere at 90 °C. After that, 0.94 g of DMPA (0.007 mol) and 0.47 g of AEEA (0.0015 mol) were added and mixed for 1 h. Once cooled to 35 °C, TEA was added and left for 30 min to neutralize the DMPA. EDA was then added to react with any remaining -NCO groups. The mixture was emulsified for 30 min, resulting in a waterborne polyurethane emulsion with 40% solid content. Finally, the emulsion was cast into polytetrafluoroethylene molds and dried for 8 h at 60 °C to produce the HPWPUE.

Synthesis of large-scale high-performance waterborne polyurethane elastomer

The raw materials, hexamethylene diisocyanate (HDI), methylene-bis(4-cyclohexylisocyanate) (HMDI), 2,2-bis(hydroxymethyl)propionic acid (DMPA), and 2-(2-aminoethylamino) ethanol (AEEA), were purchased from Wanhua Chemical Group Co., Ltd. All other materials were consistent with those used in laboratory-scale synthesis. Large-scale HPWPUE was synthesized via the prepolymer method. Initially, PAA-PHG-PNPG was dried in an oven at 70 °C for 12 h to remove residual moisture. Next, 4 kg of PAA-PHG-PNPG (2 mol), 336.4 g of HDI (2 mol), 969.4 g of HMDI (4 mol), and 1 g of DBTDL were added and stirred in a 20 L glass reactor at 90 °C for 2.5 h under a nitrogen atmosphere. Then 125.2 g of DMPA (0.93 mol) and 20.8 g of AEEA (0.2 mol) were added and reacted for 1 h. After cooling the prepolymer mixture to 35 °C, triethylamine (TEA) was added and stirred for 30 min to neutralize the DMPA. The prepolymer was then transferred to a water-cooled emulsification bucket and emulsified at 3000 rpm using an industrial emulsifier (Jiangyin Shuangye Machinery Co.,Ltd.). EDA was added to react with any residual -NCO groups. After 30 min of emulsification, a stable waterborne polyurethane emulsion with 40% solid content was obtained. Finally, the emulsion was cast into PTFE molds and dried at 60 °C for 24 h to produce the large-scale HPWPUE.

Process for producing large-scale HPWPUE coated mechanical protection gloves

Pre-treatment: New or returned formers (ceramic hand-shaped molds) are thoroughly washed with an alkaline detergent to remove any old polymer residues, followed by rinsing. The formers are then passed through a hot-air oven to ensure they are completely dry, as any moisture on the molds can lead to non-uniform film formation. Polymer impregnation (Dipping): A waterborne polyurethane thickener (0.1 ~ 0.2 wt%) is added to the large-scale HPWPUE emulsion (~40 wt% solids) to increase its viscosity to 4000–5000 cps. The formers are immersed in the latex tank for 1 min and withdrawn at a rate of 1 m/min to ensure a uniform film thickness. Gelation & Drying: The wet films are first dried at 60 °C for 12 h to transform the liquid polymer into a tacky wet gel. Then, water is expelled at 80 °C for an additional 12 h to complete the coagulation process. Stripping: The film-coated fiberglass-nylon gloves are removed from the formers for mechanical testing.

General characterizations

Fourier-transform infrared analysis was conducted using a Nicolet iS50 FT-IR spectrometer with a diamond ATR probe to obtain in-situ FT-IR spectra in the range of 650−4000 cm-1. The samples were mounted onto a homemade tensile testing machine and placed in the test chamber of the spectrophotometer. The spectra were collected by accumulating 32 scans with a resolution of 4 cm-1. Differential scanning calorimetry (DSC) analyses were conducted under a N2 atmosphere. The samples were initially heated from 25 °C to 100 °C at a rate of 10 °C/min to eliminated thermal history, then cooled to -70 °C at 10 °C/min, and finally heated to 180 °C at 5 °C/min. Thermogravimetric analyses (TGA) were conducted using a thermogravimetric analyzer (TGA, STA449F5, Netzsch, Germany) under a nitrogen atmosphere with a heating rate of 10 °C/min. Transparency tests were performed on an UV-Vis spectrophotometer (PE Lambda950), scanning wavelengths from 800−400 nm with air as reference. X-ray diffraction (XRD) patterns were obtained using XRD-6100 (Shimadzu XRD-6100, Japan) with Cu radiation. Data were collected between 10 and 80° at a scanning speed of 10° min-1. The study of the topography and heterogeneity relief from the nano- to micrometer levels was realized by atomic force microscopy (AFM, Bruker-Dimension Icon) of surfaces of HPWPUE. Measurements were performed under ambient conditions using the tapping mode AFM technique. The polarized optical microscopic (POM) images were acquired by Leica DM4 P. Atomic force microscopy (AFM) investigations were carried out in discontinuous contact mode using a SPM-9700HT system (Shimadzu Corp.). For each image, we dynamically adjusted the set point to modulate the tapping force in response to the sample’s varied surface properties. All scans were performed at a frequency of 1 Hz, and topographic phase images were captured at a scan size of 2 × 2 μm2. Specimens were prepared by cutting them into rectangular strips (5 cm × 1.0 mm) and pre-stretched to pre-determined strain levels before being mounted on a custom-made sample holder (diameter: 1 cm; height: 5 mm). A 0.3 mm diameter plastic rod was used to support the sample centrally, providing a stable imaging platform with minimal surface roughness. Molecular weight distribution was analyzed using gel permeation chromatography (GPC) with a 1260 Ⅱ system (Agilent Technologies, Inc., USA) equipped with a refractive index detector. Tetrahydrofuran was used as the eluent at a flow rate of 1.0 mL/min. Calibration of molecular weight was performed using narrow molecular weight polystyrene standards.

In-situ SAXS and WAXS characterizations

In-situ small-angle X-ray scattering (SAXS) and wide-angle X-ray scattering (WAXS) measurements of the HPWPU during tensile deformation were conducted at beamline BL16B1 of the Shanghai Synchrotron Radiation Facility (SSRF) using a custom tensile testing machine. X-ray radiation with a wavelength of 0.124 nm was utilized, and WAXS and SAXS patterns were captured using Pilatus 900 K and Pilatus 2 M detectors, respectively. The detectors had resolutions of 997 × 1080 pixels and 1475 × 1679 pixels, respectively, with a pixel size of 172 × 172 μm². Each WAXS/SAXS image was acquired over a duration of 20 s, with sample-to-detector distances set at 315 nm for WAXS and 2100 nm for SAXS. The patterns were corrected for background scattering under identical measurement conditions. The experiment continued until the sample fractured.

Tensile tests

Tensile tests were conducted in a universal tensile machine (CMT6503, MTS) equipped with a 5000 N load cell. Dumbbell-shaped samples with an active region measuring 12.0 mm × 2.0 mm × 0.4 mm were prepared. The stretching rate was set at 50 mm/min. True stress (σt) and true strain (εt) were calculated from the engineering stress-strain (σ e-ε e) curve using the following equation:

$${\sigma }_{t}={\sigma }_{e}\frac{L}{{L}_{0}}={\sigma }_{e}({\varepsilon }_{e}+1)1$$
$${\varepsilon }_{t}={\int }_{{L}_{0}}^{L}\frac{{dL}}{L}={In}\frac{L}{{L}_{0}}={In}({\varepsilon }_{e}+1)$$
(2)

where σe is the engineering stress, L is the instantaneous length of the deformed specimen, L0 is the original length of the specimen, and εe is the engineering strain. Tensile toughness (τ) of the sample is determined by integrating the area under the engineering σ e-ε e curves using the equation:

$$\tau={\int }_{\varepsilon=0}^{\varepsilon={\varepsilon }_{\max }}{\sigma }_{e}{d}_{\varepsilon }$$
(3)

where εmax is the elongation-at-break of the sample.

During cyclic tensile tests, both loading and unloading processes were performed at a strain rate of 20 mm/min at room temperature. In each cycle, the sample was stretched to a strain of 500% and then allowed to relax at room temperature for a specific waiting time (1 min, 5 min, 10 min, 30 min, 1 h, 3 h, or 6 h) before the subsequent loading process.

Puncture resistance tests

To conduct puncture resistance tests, a universal testing machine (Model CMT1503 by SUST, Zhuhai, China) was used along with a special film puncture fixture (WZDP502F). The thickness of the samples was 0.2 mm. The puncture needle had a diameter of 1 mm and a tip diameter of 0.5 mm. The puncture resistance was assessed by pressing the needle against the sample films at a rate of 50 mm/min until it pierced through, while recording both the force and displacement of the puncture needle.