Resolving the Strength–Modulus–Elasticity Tradeoff in Elastomers Using Dual Phase‐Separated Nanodomains

ABSTRACT

Achieving elastomers that simultaneously combine ultrahigh strength, high modulus, and excellent elasticity remains a longstanding challenge because these properties are intrinsically conflicting. Here, we report a dual phase‐separated nanodomain strategy that resolves this trade‐off by transforming reversible cross‐links into spatially confined reinforcing nanodomains. Elastomers are fabricated via copolymerization of rigid aromatic polyurea segments with flexible poly(urethane‐urea) chains containing acylsemicarbazide moieties. The resulting elastomers exhibit an exceptional combination of mechanical properties, including tensile strength of 104.6 MPa, Young's modulus of 43.1 MPa, toughness of 350 MJ m⁻³, and full recovery after 600% strain. Small‐angle x‐ray scattering and electron microscopy reveal two distinct nanodomains originating from self‐assembled aromatic polyurea segments and acylsemicarbazide‐stacked hydrogen‐bond arrays, respectively. Their synergistic reinforcement increases matrix rigidity while preserving entropy elasticity, enabling the simultaneous realization of ultrahigh mechanical robustness and excellent elastic recovery. The elastomers further demonstrate outstanding puncture resistance, environmental stability, healability, and reprocessability. When used as binders for carbon‐fiber fabrics, the composites achieve record‐high fracture energies of up to 2059 kJ m⁻², owing to the exceptional mechanical robustness and energy dissipation of the elastomer, together with strong elastomer‐fiber interfacial adhesion. This dual nanodomain design provides a novel route to high‐performance elastomers that transcend conventional strength‐modulus‐elasticity trade‐offs.

Elastomers integrating ultrahigh strength, high modulus, and excellent elastic recovery are realized via copolymerization of rigid aromatic polyurea segments with flexible poly(urethane–urea) chains bearing acylsemicarbazide moieties. Dual nanodomains from aromatic polyurea assemblies and hydrogen‐bonded acylsemicarbazide stacks synergistically endow the elastomers with a rare combination of 104.6 MPa strength, 43.1 MPa modulus, and full recovery after 600% strain.

1. Introduction

Elastomers are indispensable components of both everyday products and devices used in advanced engineering applications owing to their reversible deformability [1, 2, 3, 4, 5, 6, 7, 8, 9, 10, 11]. Among the key parameters that govern their mechanical performance, tensile strength and Young's modulus play important roles in elastomer design and application [1, 3, 4, 11, 12, 13, 14]. Tensile strength represents the maximum stress that an elastomer can withstand before failure, whereas Young's modulus quantifies the resistance of the elastomer to elastic deformation within the elastic regime [1, 11, 12, 13]. Enhancing these mechanical parameters improves the load‐bearing capability, dimensional stability, and resistance to mechanical damage of elastomers [3, 12]. Therefore, elastomers with an exceptionally high tensile strength and Young's modulus are desired for traditional applications, such as tires and sealing technologies, and are increasingly sought for cutting‐edge applications that require enhanced load‐bearing capacity and structural stability [12, 15, 16, 17]. However, the simultaneous realization of high modulus, high tensile strength, and excellent elasticity in a single elastomer remains a fundamental and longstanding challenge [3, 17]. Increasing the cross‐linking density of polymer networks enhances their strength and modulus; however, excessive cross‐linking restricts segmental mobility and reduces free volume, driving the elastomer toward a rigid, plastic‐like state with severely compromised elasticity [1, 16, 17, 18, 19, 20, 21]. By contrast, high elasticity requires a low cross‐linking density and substantial chain mobility, which inherently conflict with the requirements for high mechanical robustness [15, 17, 20, 21, 22]. This fundamental incompatibility leads to the long‐standing strength–modulus–elasticity tradeoff. The growing emphasis on sustainability further complicates elastomer design. Noncovalent interactions and dynamic covalent bonds offer viable routes to recyclable and healable elastomers [3, 12, 23, 24, 25, 26, 27, 28, 29]; however, their reversible nature often limits the mechanical robustness of the system. Moreover, the dynamic dissociation of the reversible cross‐links in elastomers leads to substantial energy dissipation, reducing their recovery efficiency [17, 19, 20, 21, 26, 27, 28, 29, 30, 31]. These tradeoffs become particularly pronounced when attempting to integrate an ultrahigh modulus and strength with excellent elasticity and sustainability within a single elastomer.

Incorporating phase‐separated nanodomains formed in situ has emerged as an effective strategy for strengthening elastomers while maintaining satisfactory elasticity (Scheme 1a). Nevertheless, most of the reported high‐strength phase‐separated elastomers exhibit low Young's moduli (typically <10 MPa) [11, 15, 16, 17, 24, 25, 26, 27, 28, 29, 30, 31, 32, 33, 34, 35]. Although their tensile strength can reach 50 MPa or higher, such strength is mainly achieved through the strain‐induced hardening of flexible chains at large deformations. Increasing the cross‐linking density of the networks can further improve both the tensile strength and Young's modulus; however, this enhancement compromises their elasticity (Scheme 1b) [22, 36]. To address these limitations, we propose a dual phase‐separated nanodomain strategy that allows the conversion of dispersed reversible cross‐links in the phase‐separated elastomer matrix into secondary nanodomains (Scheme 1c). This design enables synergistic reinforcement between the primary and secondary nanodomains. By converting dense reversible cross‐links into secondary nanodomains, the effective cross‐linking density is substantially reduced while the rigidity of the matrix is simultaneously enhanced. Thus, elastomers that combine an ultrahigh tensile strength, high Young's modulus, and excellent elasticity can be achieved. Following this design principle, we synthesized reversibly cross‐linked elastomers by copolymerizing aromatic polyurea (AP) with flexible poly(urethane–urea) segments containing acylsemicarbazide (ASC) moieties. In these elastomers, denoted as PUU, the AP segments self‐assemble into rigid phase‐separated nanodomains while the ASC groups form secondary hydrogen‐bonded nanodomains. Consequently, the PUU elastomer simultaneously achieves an ultrahigh tensile strength of 104.6 MPa and a Young's modulus of 43.1 MPa while maintaining excellent elastic recovery, and it fully recovers to its original state after being subjected to 600% strain at room temperature. Moreover, the PUU elastomer exhibits remarkable puncture and tear resistances, satisfactory environmental stability, and exceptional healing and recycling capabilities. Notably, we demonstrate that the PUU elastomer can serve as a high‐performance binder for the fabrication of superstrong and tear‐resistant carbon fiber (CF)‐reinforced polymer composites (CFRCs) that exhibit a record‐high tearing energy of 2059 kJ m⁻².

SCHEME 1.

Three representative elastomers with phase‐separated nanodomains. (a) Elastomer with high strength, low modulus, and high elasticity. (b) Elastomer with ultrahigh strength, high modulus, and limited elasticity. (c) Elastomer with ultrahigh strength, high modulus, and high elasticity.

2. Results and Discussion

2.1. Fabrication of the PUU Elastomer

The PUU elastomer was synthesized as illustrated in Figure 1a. Its synthesis involves the preparation of isocyanate‐terminated AP (I‐AP) and hydrazide‐terminated poly(urethane–urea) containing ASC moieties (H‐P), followed by their copolymerization. I‐AP was prepared by reacting 4,4′‐bis(3‐aminophenoxy)benzophenone (BABP) with 4,4′‐methylenebis(phenyl isocyanate) (MDI) at a 5:6 molar ratio. The reaction afforded I‐AP with a number‐average molecular weight (M _n) of ∼5.8 kDa, as determined by gel permeation chromatography (GPC; Table S1) [37]. H‐P was synthesized using a two‐step process: first, poly(tetramethylene ether)glycol (PTMEG, M _n: ∼2000) and 4,4′‐methylenebis(cyclohexyl isocyanate) (HMDI) were reacted at a 10:20 molar ratio to obtain an isocyanate‐terminated prepolymer. This prepolymer was subsequently reacted with adipic acid dihydrazide (ADH) at a 10:11 molar ratio to obtain H‐P with an M _n of ∼22.6 kDa, as determined by GPC (Table S1). The final PUU sample (M _n of ∼51.3 kDa, as determined by GPC; Table S1) was synthesized by reacting I‐AP with H‐P at a 1:1 molar ratio. The synthesis of PUU was confirmed by Fourier‐transform infrared (FT‐IR) spectroscopy (Figure S1; Tables S2 and S3) [22, 36].

FIGURE 1.

Fabrication of the PUU elastomer. (a) Synthetic routes for I‐AP, H‐P, and PUU. (b) Digital photograph of the PUU elastomer with a thickness of ∼0.2 mm and size of 40 × 30 cm². (c) DMA curves of the PUU elastomer.

PUU elastomer sheets were fabricated by blade coating a solution of PUU in N,N‐dimethylacetamide (DMAc) onto glass substrates, followed by drying at 60°C and peeling off the resulting films. As shown in Figure 1b, the obtained elastomer sheet (∼0.2 mm thickness, 40 × 30 cm²) was highly transparent, had a slight yellow tint, and exhibited an optical transmittance of 91.9% at 550 nm (Figure S2). Figure S3 shows the C═O stretching region of the FTIR spectrum of PUU. The broad C═O band could be deconvoluted into eight subpeaks corresponding to free and hydrogen‐bonded C═O groups from urethane, urea, and amide moieties, as well as benzophenone units (Table S4), which confirmed the presence of multiple hydrogen‐bonded cross‐links in the elastomers. Thermogravimetric analysis (TGA) revealed a 5% weight loss at 308°C, indicating the excellent thermal stability of the PUU elastomer (Figure S4). Further, dynamic mechanical analysis (DMA) revealed the occurrence of two well‐separated relaxation processes in the PUU elastomers (Figure 1c). Upon heating, both the storage (E′) and loss (E″) moduli decreased markedly over the temperature ranges of −90°C to −30°C and 90°C to 125°C, accompanied by two tan δ maxima at −67°C and 120°C. These peaks are attributed to the glass transitions of the soft PTMEG and hard AP segments. The two distinct glass transition temperatures (T _g) indicate the occurrence of phase separation within the PUU elastomer network.

2.2. Mechanical Properties of the PUU Elastomer

The mechanical properties of the PUU elastomer were characterized by tensile testing with a stretching speed of 50 mm min⁻¹ at 25°C and ∼20% relative humidity (RH). The stress–strain curve in Figure 2a reveals that the elastomer exhibits typical elastomeric behavior, featuring a high tensile strength of 104.6 ± 0.6 MPa, a record‐level Young's modulus of 43.1 ± 0.5 MPa, an elongation at break of 862% ± 10%, and toughness of 350.2 ± 0.7 MJ m⁻³. Remarkably, a PUU sheet weighing 2.0 g (∼0.3 mm thickness, ∼2.0  cm width, and ∼14 cm length) was capable of supporting a heavy object weighing ∼70 kg, approximately 35 000 times its weight, which demonstrates its extraordinary load‐bearing capacity. Furthermore, the elasticity of the PUU elastomer was investigated using cyclic tensile tests at a fixed strain of 100%. As shown in Figure 2b, the elastomer exhibited excellent elastic recovery, with only ∼10% residual strain in the second loading–unloading cycle. To further evaluate elastic restorability, different rest intervals were introduced between test cycles. As shown in Figure 2c, the stress–strain curves gradually recovered their initial form with increasing rest time, and the stress–strain curve fully overlapped with the initial loading curve after 12 h of resting at room temperature. Even after a large deformation of 600%, the PUU elastomer recovered its original length after 36 h of resting, and the corresponding stress–strain curve returned to its initial form (Figure 2d,e). Moreover, the PUU elastomer exhibited relatively rapid elastic recovery when subjected to strains ranging up to 50% (Figure S5). At 100% strain, the PUU elastomer showed a residual strain comparable to or only slightly higher than those exhibited by commercial elastomers (Figure S6). In cyclic tensile tests conducted at 100% strain, the residual strain of the PUU elastomer increased only slightly from 10% to 15% as the number of loading–unloading cycles increased from 1 to 10. By contrast, in the tests conducted at 600% strain, the residual strain increased from 110% in the first cycle to 150% after 10 cycles (Figure S7). Additionally, the combination of ultrahigh mechanical robustness and excellent elastic recovery endowed the PUU elastomer with strong resistance to localized mechanical damage. As shown in Figure 2f,g, a circular PUU film (∼0.75 mm thickness) withstood a puncture force of up to 128.5 N mm⁻¹ without failure. The corresponding puncture energy reached a high value of 1484.9 mJ mm⁻¹. Thus, the PUU elastomer integrates an ultrahigh mechanical strength, high Young's modulus, excellent elastic recovery, and exceptional puncture resistance (Figure 2h; Table S5) [3, 33, 38, 39].

FIGURE 2.

Mechanical performance of the PUU elastomer. (a) Stress–strain curve of the PUU elastomer. (b) Cyclic stress–strain curves of the PUU elastomer subjected to two successive loading–unloading cycles at 100% strain. (c) Cyclic stress–strain curves of the PUU elastomer stretched to 100% strain and then rested at room temperature for different durations. (d) Digital photographs of the original, stretched, and recovered PUU sheet (each small square in the background has the dimensions of 10 × 10 mm²). (e) Cyclic stress–strain curves of the PUU elastomer after two successive loading–unloading cycles at 600% strain and then resting at room temperature for 36 h. (f) Digital photograph demonstrating the puncture resistance of the PUU elastomer. (g) Puncture force–displacement curve of the PUU elastomer. (h) Radar plot comparing the mechanical performance of the PUU elastomer with those of previously reported high‐strength and high‐elasticity elastomers. The parameters being compared include the tensile strength, Young's modulus, recovery ratio at 100% strain, toughness, and puncture force.

2.3. Mechanisms Enabling the Ultrahigh Robustness and Excellent Elastic Recovery of the PUU Elastomer

To elucidate the mechanisms underlying the ultrahigh mechanical strength, high Young's modulus, and excellent elastic recovery of the PUU elastomer, its phase‐separated nanodomains were investigated using small‐angle x‐ray scattering (SAXS) and transmission electron microscopy (TEM). The 1D SAXS profile of the PUU elastomer (Figure 3a) shows two distinct scattering peaks at ∼0.20 and ∼0.55 nm⁻¹, suggesting the formation of dual phase‐separated nanodomains. The corresponding long periodicities, calculated using d = 2π/q, are ∼31.4 and ∼11.4 nm, respectively. Figure 3b shows the TEM image of a PUU elastomer sample sectioned using a microtome and then stained with ruthenium tetroxide (RuO₄), demonstrating that two populations of dark spherical nanodomains are uniformly distributed throughout the polymer matrix. Statistical analysis revealed that the larger nanodomains have an average diameter of 22.0 ± 2.4 nm, whereas the smaller ones have a significantly smaller average diameter of 9.3 ± 1.3 nm (Figure 3c). The corresponding average interdomain distances determined from the TEM image are 33.2 ± 1.9 nm for the larger domains and 12.5 ± 1.2 nm for the smaller ones (Figure 3b). The close agreement between the characteristic length scales determined through SAXS and TEM measurements confirmed the formation of well‐defined dual phase‐separated nanodomains in the PUU elastomer. RuO₄ preferentially stains aromatic moieties and urethane and urea groups through redox reactions. Based on this selective staining behavior, the larger nanodomains are attributed to self‐assembled AP segments formed via π–π stacking and hydrogen‐bonding interactions, whereas the smaller nanodomains are attributed to stacked hydrogen‐bonded arrays of ASC groups in conjunction with dicyclohexylmethane units. Considering the relative volume fractions of the constituent segments, the AP‐rich nanodomains are expected to be larger than the ASC‐derived hydrogen‐bonded nanodomains, and the PTMEG chains form a continuous soft matrix. These results support that the PUU elastomer consists of PTMEG chains physically cross‐linked by two chemically distinct phase‐separated nanodomains with different characteristic dimensions.

FIGURE 3.

Structural characterization and mechanistic insights of the PUU elastomer. (a) 1D SAXS profile of the PUU elastomer. The inset shows the corresponding 2D SAXS pattern. (b) TEM image of the PUU elastomer and (c) the corresponding particle‐size distribution. (d) Chemical structures of the PUU‐ASC and PUU‐AP elastomers. (e) Stress–strain curves of the PUU‐ASC and PUU‐AP elastomers. (f,g) Cyclic stress–strain curves of the PUU‐ASC and PUU‐AP elastomers, respectively, recorded during two successive loading–unloading cycles at 100% strain. (h) 1D SAXS profile of the PUU‐ASC elastomer. The inset shows the corresponding 2D SAXS pattern. (i) TEM image of the PUU‐ASC elastomer and (j) the corresponding particle‐size distribution. (k) 1D SAXS profile of the PUU‐AP elastomer. The inset shows the corresponding 2D SAXS pattern. (l) TEM image of the PUU‐AP elastomer and (m) the corresponding particle‐size distribution.

To clarify the role of the dual phase‐separated nanodomains in integrating multiple outstanding mechanical properties, two control elastomers were fabricated following the same procedure used for the PUU elastomer. As shown in Figure 3d; Figures S8 and S9, one control elastomer (denoted as PUU‐ASC) consisted of poly(urethane–urea) chains containing ASC‐stacked hydrogen‐bonded arrays but lacked the rigid AP segments, whereas the other elastomer (denoted as PUU‐AP) contained rigid AP segments and poly(urethane–urea) chains but lacked the ASC‐stacked hydrogen‐bonded arrays. In the synthesis of PUU‐AP, HMDI was replaced with isophorone diisocyanate, which has an asymmetric and sterically hindered structure that suppresses the formation of urea‐derived nanodomains in the PUU‐AP elastomer. The M _n values of the PUU‐ASC and PUU‐AP elastomers were determined to be ∼64.0 and ∼54.5 kDa, respectively (Table S6). The DMA curves in Figure S10 reveal that the T _g values of the PUU‐ASC and PUU‐AP elastomers are −57°C and −65°C, respectively. As shown in Figure 3e and summarized in Table S7, the PUU‐ASC elastomer exhibited a tensile strength of 50.7 ± 0.9 MPa, Young's modulus of 5.1 ± 1.1 MPa, and a strain at break of 1148% ± 21%. In comparison, the PUU‐AP elastomer had a greater tensile strength (63.1 ± 1.4 MPa), Young's modulus (21.2 ± 0.8 MPa), and strain at break (1293% ± 14%). Nevertheless, both the strength and the modulus of the PUU‐ASC and PUU‐AP elastomers were lower than those of the PUU elastomer with dual nanodomains. Cyclic tensile tests performed at a fixed strain of 100% revealed that the PUU‐AP elastomer exhibits a larger residual strain (∼15%) than that of the PUU‐ASC elastomer (∼5%), indicating its inferior elastic recovery (Figure 3f,g). Thus, the PUU‐ASC elastomer is characterized by a relatively low strength and modulus but excellent elasticity, whereas the PUU‐AP elastomer achieves enhanced strength and modulus at the expense of elastic recovery. Remarkably, because PUU‐ASC and PUU‐AP have higher M _n values than PUU, the enhanced tensile strength and modulus of PUU predominantly originate from its dual phase‐separated nanodomains.

SAXS and TEM analyses revealed that the PUU‐ASC and PUU‐AP elastomers contained only a single type of phase‐separated nanodomain (Figure 3h–m). The 1D SAXS profile of the PUU‐ASC elastomer (Figure 3h) exhibits a single scattering peak at ∼0.57 nm⁻¹, corresponding to a long periodicity of ∼11.0 nm (d = 2π/q). Further, the TEM image and statistical analysis results of the RuO₄–stained PUU‐ASC elastomer (Figure 3i,j) revealed the presence of uniformly dispersed nanodomains with an average diameter of 8.5 ± 0.9 nm and interdomain distance of 11.0 ± 0.7 nm, in good agreement with the SAXS results. By contrast, the 1D SAXS profile of the PUU‐AP elastomer (Figure 3k) revealed a longer period, reaching ∼28.6 nm, for its phase‐separated nanodomains. Consistently, the TEM image and statistical analysis results confirmed that PUU‐AP contained larger nanodomains with an average diameter of 24.2 ± 1.9 nm and interdomain spacing of 28.8 ± 2.1 nm (Figure 3l,m). In the PUU‐ASC elastomer, the nanodomains originate exclusively from the aggregates of ASC‐derived hydrogen‐bonded arrays and adjacent dicyclohexylmethane units because of the absence of rigid AP segments. By contrast, in the PUU‐AP elastomer, only the rigid AP chains self‐assemble into homogeneously dispersed nanodomains because the isolated and weakly cooperative hydrogen bonds within the poly(urethane–urea) chains are inadequate for inducing phase separation. Notably, the dimensions (sizes and long periodicities) of the nanodomains in the PUU‐ASC and PUU‐AP elastomers matched those of the small and large nanodomains of the PUU elastomer, respectively. This correspondence unambiguously confirms the compositional origin of the dual phase‐separated nanodomains in the PUU elastomer.

In the PUU‐AP elastomer, the rigid AP nanodomains serve as stiff nanofillers to endow it with a high strength and modulus. However, the weak, non‐aggregated hydrogen bonds within the soft poly(urethane–urea) chains readily dissociate during deformation to dissipate energy, resulting in a relatively poor elastic recovery. By contrast, in addition to the primary AP nanodomains, the PUU elastomer contains secondary ASC‐derived hydrogen‐bonded nanodomains. The formation of these secondary nanodomains effectively decreases the cross‐linking density by converting the dispersed hydrogen bonds into spatially confined hydrogen‐bonded nanodomains; this phenomenon suppresses excessive energy dissipation and thereby endows the PUU elastomer with markedly improved elastic recovery compared with that of the PUU‐AP elastomer. Notably, the ASC‐derived hydrogen‐bonded nanodomains also serve as additional nanofillers, further strengthening the PUU elastomer. Consequently, the synergistic reinforcement from the AP and ASC‐derived hydrogen‐bonded nanodomains enables the PUU elastomer to achieve a higher strength and modulus than those of the PUU‐AP elastomer. In the case of the PUU‐ASC elastomer, although the ASC‐derived hydrogen‐bonded nanodomains impart excellent elasticity, its strength and modulus remain substantially lower than those of the PUU elastomer owing to the absence of rigid AP nanodomain reinforcement. These results demonstrate that the combination of an ultrahigh strength, high modulus, and excellent elastic recovery in the PUU elastomer originates from the synergistic action of the dual phase‐separated nanodomains. When the PUU elastomer is stretched, the rigid AP nanodomains undergo deformation and disruption earlier than the softer ASC‐derived hydrogen‐bonded nanodomains (Figure S11).

2.4. Water Resistance, Healability, and Reprocessability of the PUU Elastomer

Hydrogen bonds are generally susceptible to disruption by water, which often leads to the poor stability of hydrogen‐bonded polymers in aqueous environments. Given its importance in practical applications, the water resistance of the PUU elastomer was evaluated by immersing it in water, 1 mol L⁻¹ HCl and 1 mol L⁻¹ NaOH aqueous solutions for 72 h. As shown in Figure 4a, the PUU elastomer retained its original shape and transparency after immersion in these media. Compared with the pristine sample, the PUU elastomers retrieved from the aqueous media (not dried) exhibited only a marginal decrease in their tensile strength, accompanied by a slight increase in the strain at break (Figure S12). Upon drying at 60°C for 12 h, the stress–strain curve of the elastomer nearly overlapped with that of the pristine elastomer, indicating the complete recovery of the mechanical properties (Figure 4b). These results demonstrate that the PUU elastomer has satisfactory water resistance, even under acidic and alkaline conditions. This satisfactory water stability can be attributed to the rigid and hydrophobic AP nanodomains formed in situ and the ASC‐derived hydrogen‐bonded nanodomains with hydrophobic dicyclohexylmethane units, which confine the hydrogen bonds within densely packed and hydrophobic microenvironments. Moreover, the intrinsically hydrophobic PTMEG chains in the soft domains further enhance the water repellency of the polymer matrix. The synergistic effect of the hydrophobic phase‐separated nanodomains and polymer matrix not only prevents the water from penetrating the elastomer network but also suppresses the dissociation of the hydrogen bonds in aqueous environments, thereby endowing the PUU elastomer with excellent water resistance.

FIGURE 4.

Water resistance, healability, and recyclability of the PUU elastomer. (a) Digital images of the PUU elastomer soaked in different aqueous media for 72 h. (b) Stress–strain curves of the PUU elastomer soaked in different aqueous media for 72 h and subsequently dried at 60°C for 12 h. (c) Digital images of the PUU elastomer cut into two pieces (i) and healed at 80°C for 18 h (ii) (each small square in the background is 1 × 1 cm² in size). (d) Digital image of a healed PUU elastomer sheet lifting a 5 kg weight without failure. (e) Stress–strain curves of the PUU elastomer sheet healed at 80°C for different durations. (f) Digital images showing the reprocessing of the PUU elastomer via hot pressing. (g) Stress–strain curves of the PUU elastomer over three cycles of cutting and recycling via hot pressing.

The reversibility of hydrogen bonds imparts excellent healing and reprocessing properties to the PUU elastomer. Its healing behavior was examined by cutting a rectangular PUU sheet into two separate pieces using a razor blade and subsequently bringing the fractured surfaces into contact (Figure 4c(i)). To facilitate chain mobility across the interface, a trace amount of DMAc was applied to the fractured surfaces before contact. The samples were then held in contact and heated at 80°C for different durations. After 18 h of healing, the cut became indistinguishable (Figure 4c(ii); Figure S13), and the healed elastomer could be used to lift a 5.0 kg weight without fracture (Figure 4d). Accordingly, the stress–strain curves revealed that the elastomer fully recovered its original mechanical properties after 18 h of healing (Figure 4e). The healed elastomer also re‐established its dual phase‐separated nanodomains (Figure S14). The combined effects of the solvent and heat promoted the dynamic dissociation of the hydrogen bonds at the fractured interfaces, thereby enhancing the mobility of polymer chains. Upon contact, the polymer chains diffused across the interface and re‐established hydrogen bonds. Subsequent solvent evaporation and cooling to room temperature reconstructed the cross‐linked network, enabling the complete healing and full recovery of the mechanical performance of the PUU elastomer. The dynamic rupture and reformation of hydrogen bonds could be exploited to efficiently reprocess the PUU elastomer through hot pressing. As shown in Figure 4f, fragmented PUU elastomer pieces could be reprocessed into an intact sheet by hot pressing at 120°C under a pressure of 4 MPa for 15 min. The stress–strain curves in Figure 4g demonstrate that the mechanical properties of the reprocessed elastomer remained essentially unchanged after three cutting–reprocessing cycles.

2.5. Fabrication of Tear‐Resistant CF/PUU Composites

CFRCs with a high strength and exceptional resistance to tearing are highly attractive for personal protection and damage‐tolerant structural applications that critically require resistance to tear‐induced crack propagation for safety and reliability [40, 41, 42, 43, 44, 45]. Elastomeric binders that combine a high strength, high modulus, and excellent energy dissipation are required for fabricating such CFRCs. Accordingly, CF fabrics were combined with the PUU elastomer binder at a PUU:CF fabric mass ratio of 35:65 to fabricate CF/PUU composites via a solution impregnation method, followed by hot pressing, as illustrated in Figure 5a. The resulting composites are denoted as n‐CF/PUU, where n represents the number of CF fabric layers. Figure 5b displays a digital photograph of the 3‐CF/PUU composite with three CF fabric layers. Top‐view scanning electron microscopy (SEM) images of the pristine CF fabric and 3‐CF/PUU composite confirmed the successful impregnation and integration of the PUU binder with the CF fabrics (Figure S15). Cross sectional SEM images revealed thicknesses of approximately 0.25, 0.50, and 0.75 mm for the 1‐, 2‐, and 3‐CF/PUU composites, respectively (Figure S16). Notably, no delamination of the CF fabric layers was observed, indicating strong interfacial adhesion between the fiber bundles and the PUU binder. As shown in Figure 5c, the lap‐shear adhesion strength between two CF fabrics bonded with the PUU elastomer binder reached 1.5 MPa (see Lap‐Shear Test for details in Supporting Information). This strong adhesion originates from the abundant hydrogen‐bonding interactions between the urea, urethane, and amide groups in the PUU binder and the carboxyl/carbonyl groups located on the CF fabric surfaces.

FIGURE 5.

Fabrication and characterization of CF/PUU composites. (a) Schematic of the preparation of the CF/PUU composites. (b) Digital image of the 3‐CF/PUU composite. (c) Lap‐shear force–displacement curves for two pieces of CF fabrics adhered using the PUU elastomer binder. (d) Typical stress–strain curves of the 1‐, 2‐, and 3‐CF/PUU composites. (e) Digital image of a torn 1‐CF/PUU composite. (f) Tearing force–displacement curves of the 1‐, 2‐, and 3‐CF/PUU and 1‐CF/epoxy composites. (g) Comparison of the stress and puncture energies of the 1‐CF/PUU composite and other recently reported single‐layer CFRCs. (h–j) SEM images showing a magnified view of the torn region of the 1‐CF/PUU composite shown in (e): (h) notch tip region, (i) region near the notch tip, and (j) region away from the notch tip. (k, l) Cross sectional SEM images of the torn 1‐CF/PUU composite. (m) Schematic of notch propagation in the CF/PUU composite.

The mechanical properties of the n‐CF/PUU composites were evaluated using uniaxial tensile tests at a stretching speed of 50 mm min⁻¹. As shown in Figure 5d and summarized in Table S8, the n‐CF/PUU composites exhibited a nearly identical mechanical performance, regardless of the number of CF fabric layers, with the tensile strength, Young's modulus, and toughness reaching 710.2 ± 9.4 MPa, 15.3 ± 0.5 GPa, and 20.7 ± 0.2 MJ m⁻³, respectively. The tear resistance of the n‐CF/PUU composites was further investigated using trouser tear tests, following a previously reported protocol [42, 46, 47]. For this test, rectangular specimens (80 mm × 50 mm) with a central notch of 30 mm were prepared from the 1‐, 2‐, and 3‐CF/PUU composites. In the test, one end of the specimen was clamped while the other was pulled at a constant stretching speed of 10 mm min⁻¹ (Figure 5e). No fiber pull‐out was observed during this tearing process (Figure S17). The representative tearing force–displacement curves reveal a pronounced increase in the tearing force with an increase in the number of CF fabric layers, indicating exceptionally high fracture energies of 2059, 2167, and 2493 kJ m⁻² for the 1‐, 2‐, and 3‐CF/PUU composites, respectively (Figure 5f). As the fracture energy of the pure PUU elastomer sheet is 261.9 kJ m^{− 2}, the tear resistance of PUU contributed only marginally to the exceptional tear resistance of the CF/PUU composites (Figure S18). Notably, the fracture energy of 2059 kJ m⁻² represents the highest value reported to date for a single‐layer CFRC (Figure 5g; Table S9) [46, 47, 48, 49, 50].

To elucidate the origin of the exceptional tear resistance of the CF/PUU composites, the surface morphologies of the 1‐CF/PUU specimen near and far from the notch tip were examined using SEM (Figure 5h–l). A pronounced deformation of the PUU binder was observed near the notch tip, whereas regions remote from the notch remained essentially undeformed (Figure 5h–j). This spatially localized deformation indicates that the applied tearing force is primarily dissipated through the substantial deformation of the PUU binder, accompanied by hydrogen‐bond dissociation. The load is subsequently transferred to the CF bundles aligned perpendicular to the crack path, leading to collective bundle fracture, instead of fiber pull‐out or individual fiber failure (Figure 5k). Notably, the PUU binder remained firmly adhered to the CF fabrics even after the CF bundle fractured, confirming that the interfacial adhesion was sufficiently strong to sustain efficient energy dissipation (Figure 5l). For comparison, a bisphenol A–based epoxy thermoset was used as a binder to fabricate a 1‐CF/epoxy composite. The stress–strain curve of the 1‐CF/epoxy composite (Figure S19) revealed a significantly lower fracture energy of 411 kJ m⁻² (Figure 5f) owing to the absence of a highly deformable, energy‐dissipating binder network. These results demonstrate that the exceptional tear resistance of the CF/PUU composites arises from the synergistic interplay of the mechanically robust, yet highly dissipative PUU binder and its strong interfacial adhesion to the CF fabric (Figure 5m). The high modulus and strength of PUU suppressed stress concentration at the crack tip, whereas its reversible network and phase‐separated nanodomains enabled substantial energy dissipation during deformation and localized fracture (Figure 5m(i, ii)). Moreover, the strong interfacial adhesion between the two components facilitated collective bundle fracture, rather than single fiber pull‐out or individual fiber breakage, thereby significantly increasing the force threshold for fracture (Figure 5m(iii)).

Furthermore, because of the intrinsic water resistance of the PUU binder, the 3‐CF/PUU composites maintained their original mechanical properties after exposure to various aqueous environments (Figure S20). Moreover, the dynamic hydrogen bonds within the PUU binder and at the CF/PUU interfaces enabled the dissociation of the CF/PUU composites in DMAc, allowing the recovery of both the PUU binder and the structurally intact CF fabric (Figures S21–S23). Moreover, the recovered components could be reused to fabricate CF/PUU composites that fully maintained their original mechanical properties (Figure S24).

3. Conclusion

We developed a PUU elastomer that simultaneously exhibits ultrahigh tensile strength, extremely high Young's modulus, and excellent elastic recovery through the copolymerization of rigid AP segments and flexible poly(urethane–urea) segments containing high‐density hydrogen‐bonded arrays. In the PUU elastomer, the AP segments self‐assemble into primary rigid nanodomains, while the high‐density hydrogen‐bonded arrays within the poly(urethane–urea) segments stack into secondary nanodomains. The synergistic reinforcement of the two phase‐separated nanodomains enabled a rare combination of a high tensile strength of 104.6 MPa, Young's modulus of 43.1 MPa, and full recovery after 600% strain, which surpass those of the elastomers reinforced by a single type of nanodomain. In addition, the PUU elastomer exhibited excellent puncture and water resistances, healing capability, and reprocessability. The CF/PUU composites prepared using the PUU elastomer as the binder achieved a record‐high fracture energy of 2059 kJ m⁻² owing to the strong interfacial adhesion between the CF fabric and the PUU binder, along with the exceptional mechanical robustness and efficient energy dissipation of the composite system. Owing to the unique combination of strength, resilience, and durability, the PUU elastomer developed in this study is well suited for protective equipment, sealing technologies, and structural composites, which require a high load‐bearing capacity and long‐term reliability. This dual nanodomain strategy combined with reversible cross‐linking establishes a new paradigm for elastomer design by demonstrating how spatially confined reversible interactions can overcome intrinsic mechanical tradeoffs, enabling sustainable high‐performance elastomers.

4. Experimental Section

4.1. Synthesis of the PUU Elastomer

The PUU elastomer was synthesized using a three‐step polymerization process. (i) Synthesis of I‐AP: MDI (1.5 g, 6.0 mmol) was completely dissolved in anhydrous DMAc (30 mL) at room temperature under a nitrogen atmosphere. Subsequently, an anhydrous DMAc solution (40 mL) of BABP (2.0 g, 5.0 mmol) was added dropwise to the MDI solution. The reaction mixture was maintained at 60°C under nitrogen with continuous stirring for 12 h to obtain I‐AP. The resulting I‐AP solution was used directly in subsequent steps, without further purification. (ii) Synthesis of H‐P: PTMEG (M _n: ∼2000, 20 g, 10 mmol) was dried under vacuum at 120°C for 2 h to remove residual moisture and then cooled to 80°C. HMDI (5.3 g, 20 mmol) and ten drops of dibutyltin dilaurate were sequentially added to the molten PTMEG sample. The resulting mixture was continuously stirred at 80°C under a nitrogen atmosphere for 2 h to obtain an isocyanate‐terminated prepolymer (OCN‐prepolymer‐NCO). The reaction mixture was then cooled to 60°C, and an anhydrous DMAc solution (250 mL) of ADH (1.9 g, 11 mmol) was added, and the reaction was continued at 60°C for an additional 2 h under nitrogen to obtain H‐P. The obtained H‐P solution was used directly in the next step, without further purification. (iii) Copolymerization of I‐AP with H‐P: The I‐AP solution was added to the H‐P solution at a 1:1 molar ratio with continuous stirring at 60°C under nitrogen. Upon stirring for 24 h, a viscous, pale yellow PUU solution was obtained. This solution was cast on glass substrates and cured at 60°C to obtain defect‐free PUU elastomer sheets.

4.2. Preparation of the n‐CF/PUU Composites

The preparation of the 1‐CF/PUU composite is described as an example. A piece of pre‐cut CF fabric (2.4 g) was placed in a square silicone mold. A 100 mg mL^{− 1} PUU solution was prepared by dissolving the PUU elastomer (1.3 g) in DMAc (13 mL). The mass ratio of the CF fabric to PUU elastomer was maintained at 65:35. The PUU solution was poured into the silicone mold with the CF fabric and allowed to uniformly impregnate the CF fabric. Thereafter, the solvent was evaporated at 60°C to obtain a 1‐CF/PUU composite sheet. Finally, the resulting 1‐CF/PUU composite was hot‐pressed at 120°C under a pressure of 4 MPa for 15 min. The 2‐CF/PUU and 3‐CF/PUU composites were prepared following a similar procedure, except that the number of CF fabric pieces placed in the silicone mold was increased to two and three, respectively.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: adma74001‐sup‐0001‐SuppMat.docx.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.