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. 2026 May 23;47(15):e00967. doi: 10.1002/marc.202500967

Synergistic Enhancement of Environmental Stress Crack Resistance and Mechanical Properties of HDPE via Reactive Blending‐Induced Disulfide‐Linked Covalent Networks

Yue Chen 1, Wei Tian 1, Yongjun Zhu 1, Qiang Xu 1, Ting Li 1, Shibo Wang 1, Jing Huang 1,✉, Weifu Dong 1,✉
PMCID: PMC13434992  PMID: 42177609

ABSTRACT

Environmental stress cracking resistance (ESCR) remains a critical limitation for high‐density polyethylene (HDPE), as conventional strategies to enhance ESCR often compromise mechanical performance or melt processability. Herein, a dynamic covalent network is constructed in HDPE via reactive melt blending using a disulfide‐containing bismaleimide (DTBMI) crosslinker. Radical activation enables DTBMI to capture HDPE macroradicals, forming a dynamically crosslinked network without sacrificing thermoplastic processability. The introduction of 2 phr DTBMI significantly enhances ESCR, with the failure time (f50) increasing by fourfold compared to neat HDPE, while tensile strength, fracture energy, and strain‐hardening behavior are simultaneously improved. Rheological analysis reveals a pronounced increase in low‐frequency modulus and effective network density, indicating strong constraints on long‐time chain relaxation, whereas high‐frequency viscosity remains comparable to that of neat HDPE, preserving melt flow. Fracture morphology after ESCR testing shows pronounced fibrillation, evidencing enhanced stress redistribution and suppressed chain disentanglement. These results demonstrate that time‐scale‐selective dynamic crosslinking offers an effective strategy to reconcile ESCR, mechanical robustness, and processability in HDPE, providing new insights for the design of durable polyolefin materials.

Keywords: disulfide exchange, dynamic covalent network, environmental stress crack resistance, HDPE, reactive melt blending

Reactive melt blending of HDPE with 2phr disulfide‐containing bismaleimide (DTBMI) forms a dynamic covalent network, quadrupling ESCR, enhancing tensile strength and fracture energy while maintaining yield strength, providing a facile, robust strategy to overcome the trade‐off between ESCR, mechanical performance, and processability.

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1. Introduction

Environmental stress cracking (ESC) is a critical failure mode in polymeric materials subjected to sustained mechanical stress in the presence of aggressive environments such as surfactants or solvents [1, 2]. Unlike a catastrophic fracture caused by instantaneous overload, ESC typically proceeds through a slow crack growth (SCG) process and may culminate in sudden brittle failure without obvious macroscopic warning [3]. For semicrystalline polymers, the widely accepted model proposed by Lustiger and Markham attributes ESC to stress‐induced disentanglement and slippage of load‐bearing interlamellar chains that bridge adjacent crystalline lamellae [3, 4]. Accordingly, ESC resistance (ESCR) is highly sensitive to molecular‐level structural features, including molecular weight and its distribution, chain architecture, crystallinity, and the topology and density of tie molecules [5, 6, 7, 8].

High‐density polyethylene (HDPE) is one of the most widely used semicrystalline polyolefins owing to its favorable balance of mechanical strength, chemical resistance, and ease of melt processing [9, 10]. Nevertheless, the predominantly linear molecular architecture and low branching density of HDPE result in relatively weak intermolecular interactions in the amorphous phase, rendering it intrinsically susceptible to ESC. This limitation has become particularly critical in applications such as pressure pipes, containers, and rigid packaging, where HDPE components are exposed to long‐term mechanical loading and chemically aggressive environments [11, 12, 13]. The inferior ESCR performance commonly observed in postconsumer HDPE further highlights this intrinsic vulnerability and underscores the need for effective postmodification strategies applicable to HDPE materials in general [14, 15, 16].

Based on the mechanistic understanding of SCG, various strategies have been developed to enhance the ESCR of HDPE, including copolymerization with α‐olefins, blending with elastomers or other polyethylene grades, and chemical modification [16, 17, 18, 19, 20, 21, 22]. In many cases, ESCR improvement is achieved by increasing chain mobility or entanglement density, which facilitates stress relaxation and delays crack initiation. However, such approaches are often accompanied by a reduction in crystallinity or stiffness, leading to compromised mechanical strength and dimensional stability. Alternatively, permanent covalent crosslinking can effectively suppress chain slippage and improve crack resistance, but typically sacrifices ductility, melt processability, and often requires stringent processing conditions or postcrosslinking steps. Achieving a simultaneous improvement in ESCR and mechanical performance, while maintain processability therefore remains a long‐standing challenge for HDPE‐based materials.

Covalent adaptable networks (CANs), constructed from dynamic covalent bonds, have recently emerged as a promising materials concept that bridges the gap between thermoplastics and thermosets [23, 24, 25]. By enabling network topology rearrangement at elevated temperatures while maintaining covalent connectivity under service conditions, CANs offer a unique opportunity to simultaneously enhance mechanical robustness and retain melt processability [26, 27]. In polyethylene‐based systems, CANs have been realized through the incorporation of dynamic covalent bonds such as dioxaborolane metathesis, β‐hydroxy esters, disulfide bonds, or other exchangeable linkages, typically enabled by functionalized polyethylene copolymers or radical grafting/crosslinking strategies [28, 29, 30, 31, 32, 33, 34]. While these systems have demonstrated improved thermal and mechanical performance, their impact on long‐term failure phenomena, particularly ESCR, remains insufficiently explored.

In our previous work, we demonstrated that incorporating a small fraction of polyethylene‐based CAN into HDPE could significantly improve its ESCR by enhancing intermolecular constraints in the amorphous regions [35]. Nevertheless, the CAN phase in that system was derived from ethylene–methacrylic acid copolymer salts, which are relatively costly and intrinsically flexible, leading to a reduction in the stiffness of the resulting HDPE blends. These limitations highlight the need for a more direct and structurally efficient strategy to introduce dynamic covalent constraints into HDPE without sacrificing its inherent rigidity.

Herein, we report a simple and efficient strategy to construct a disulfide‐containing covalent adaptable network in HDPE via reactive melt blending with a disulfide‐linked bismaleimide (DTBMI) crosslinker. The radical grafting approach employed here provides a practical and scalable route for network construction in polyethylene. Maleimide moieties were selected due to the high reactivity and selectivity of their electron‐deficient C = C bonds under radical conditions, along with their resistance to undesired homopolymerization, enabling controlled network formation during melt processing compared to conventional acrylate‐based crosslinkers [36]. In addition, the incorporation of disulfide bonds as dynamic linkages not only enables catalyst‐free exchange but also introduces relatively hydrophobic and flexible covalent connections, which are advantageous for maintaining ductility and improving ESCR. Moreover, DTBMI can be readily synthesized via a simple two‐step condensation‐cyclization route, further enhancing its practical accessibility. By varying the DTBMI content, HDPE networks with tunable covalent connectivity were obtained. The effects of network formation on crystallization behavior, mechanical properties, melt rheology, and ESCR were systematically investigated. This approach enables a concurrent enhancement of ESCR and mechanical properties while largely preserving melt processability, thereby addressing a key limitation of existing HDPE modification strategies.

2. Experimental Section

Materials and Characterizations are described in Supporting Information.

2.1. Synthesis of Disulfide‐Linked Bismaleimide (DTBMI)

Dithiodianiline (DTDA, 3.92 g, 15.8 mmol) and Maleic anhydride (MAH, 3.46 g, 35.3 mmol) were charged into a three‐necked flask, followed by 50 mL of acetone as solvent. The mixture was stirred at room temperature under a nitrogen atmosphere for 6 h. Subsequently, 16 mL acetic anhydride and 0.29 g sodium acetate were added, and the reaction temperature was raised to 60°C under reflux for an additional 6 h. After completion, the reaction mixture was washed three times with saturated sodium bicarbonate solution and deionized water, followed by one wash with ethanol. The precipitate was filtered and dried in a vacuum oven at 50°C for 12 h to afford the bismaleimide product (DTBMI).

2.2. Preparation of HDPE/DTBMI Blends

HDPE/DTBMIx blends were prepared using a torque rheometer, where x denotes the mass (g) of DTBMI per 100 g HDPE. Specifically, HDPE, dicumyl peroxide (DCP), and DTBMI were premixed and melt‐compounded at 190°C and 40 rpm for 10 min. The resulting samples were compression‐molded using a platen vulcanizer at 200°C and 10 MPa: preheating for 5 min, hot‐pressing for 6 min, and subsequently cold‐pressing at 60°C for 5 min. Neat HDPE and HDPE/DCP0.2 were prepared under identical conditions as controls. The detailed formulations are listed in Table 1.

TABLE 1.

The formula of the samples.

Sample code HDPE (g) DCP (g) DTBMI (g)
HDPE 50 0 0
HDPE/DCP0.2 50 0.1 0
HDPE/DCP0.4 50 0.2 0
HDPE/DTBMI1.0 50 0.1 0.50
HDPE/DTBMI1.5 50 0.1 0.75
HDPE/DTBMI2.0 50 0.1 1.00
HDPE/DTBMI2.5 50 0.1 1.25
HDPE/DTBMI4.0 50 0.2 2.0
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3. Results and Discussions

3.1. Preparation and Network Formation of HDPE/DTBMIX Blends

To construct a dynamic network structure in HDPE, a bismaleimide crosslinker containing dynamic disulfide bonds (DTBMI) was synthesized by end‐group modification of 4,4′‐dithiophenylamine (DTDA) with maleic anhydride, followed by acetylation (Figure 1a). The resulting DTBMI was incorporated into HDPE via a straightforward reactive melt‐blending process. Upon thermal decomposition of the radical initiator dicumyl peroxide (DCP), macroradicals were first generated along the HDPE backbone through hydrogen abstraction [37]. These HDPE macroradicals subsequently attacked the carbon–carbon double bonds of the maleimide moieties in DTBMI, leading to covalent grafting and the formation of a disulfide bond‐linked dynamic covalent network, as illustrated in Figure 1b,c.

FIGURE 1.

FIGURE 1

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(a) Synthetic route of DTBMI (b) Schematic diagram of the radical grafting reactions between HDPE and DTBMI (c) Schematic illustration of the network structure in HDPE/DTBMI blends.

FTIR spectroscopy and 1H NMR analysis were employed to confirm the successful synthesis of DTBMI (Figure 2a,b). In the FTIR spectrum, the characteristic N–H stretching vibrations of the primary amine groups in DTDA at 3413 and 3329 cm− 1 disappear after reaction, while new absorption bands emerge at 3106 cm− 1 (ν = CH), 1701 cm− 1 (νC = O), 1140 and 821 cm− 1, which are characteristic of maleimide ring vibrations and C–N–C stretching, indicating the formation of bismaleimide structures [38]. The 1H NMR spectrum further corroborates this assignment: aromatic proton signals appear between 6.0 and 8.0 ppm, with peaks at 7.3 and 7.5 ppm displaying splitting patterns consistent with coupled aromatic protons, while a singlet at 6.8 ppm is attributed to the olefinic protons of the maleimide moieties. These spectroscopic results confirm the successful formation of DTBMI. Thermogravimetric analysis (TGA) shows that DTBMI exhibits high thermal stability, with the temperature at 5% weight loss (T5%) reaching 396°C and no significant mass loss observed below 370°C (Figure S2a). Furthermore, isothermal TGA at 190°C under an air atmosphere for 1 h reveals a weight loss of less than 1% (Figure S2b). This high thermal stability ensures that DTBMI is thermally compatible with HDPE and suitable for the preparation of HDPE/DTBMIx composites via reactive melt blending.

FIGURE 2.

FIGURE 2

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Synthesis of DTBMI: (a) FTIR spectra (b) 1H NMR spectra of DTBMI Preparation of HDPE/DTBMIx: (c) torque evolution curves (d) gel fractions and (e) FTIR spectra.

Given its confirmed structure and high thermal stability, DTBMI was subsequently incorporated into HDPE via reactive melt blending, and the evolution of torque was recorded to monitor network formation in the HDPE/DTBMIx systems. As shown, neat HDPE exhibits a low and stable torque after melting, whereas HDPE containing DCP and DTBMI shows a torque increase after ∼2 min, reaching a steady state at ∼4 min, indicative of chain branching and network formation (Figure 2c). Compared with HDPE/DCP0.2, all HDPE/DTBMIx samples display significantly higher equilibrium torques, confirming that DTBMI promotes a higher degree of network formation than DCP alone. The equilibrium torque increases with DTBMI content and plateaus at 2.5 phr, suggesting saturation of effective crosslinking.

Gel fraction measurements support this trend (Figure 2d). HDPE/DCP0.2 remains fully soluble (Figure S3), indicating that a low loading of DCP in PE induces radical reactions dominated by chain extension and limited branching, with no crosslinked network formed, which is consistent with previous reports [28, 38]. In contrast, the gel content of HDPE/DTBMIx increases from 23.1% to 37.8% as the DTBMI loading increases from 1.0 to 2.5 phr. FTIR spectra of the gel part (Figure 2e) show characteristic carbonyl absorption at 1718 cm− 1 and C–N bending at 1385 cm− 1, confirming the incorporation of DTBMI into the HDPE network.

3.2. Crystallization, Mechanical and ESCR Properties of HDPE/DTBMIX

Differential scanning calorimetry (DSC) and x‐ray diffraction (XRD) measurements were employed to evaluate the effect of DTBMI incorporation on the crystalline structure of HDPE (Figure 3a,b and Table S1). It reveals that incorporation of DTBMI slightly reduces the melting temperature and broadens the melting endotherm of HDPE, accompanied by a decrease in crystallinity from 73.1% to approximately 60%. This behavior is attributed to the formation of the disulfide crosslinks that restricts chain mobility and disrupts regular crystal packing. A similar reduction in crystallinity is observed for HDPE/DCP0.2, indicating that chain branching induced by DCP also interferes with crystallization. Consistent with the DSC results, XRD patterns (Figure 3c and Table S1) show a decrease in crystallinity upon DTBMI incorporation. Meanwhile, no new diffraction peaks are detected, and the crystalline lamellar thickness estimated from the (110) reflection using the Scherrer equation remains nearly unchanged, indicating that the crystal structure of HDPE is essentially preserved.

FIGURE 3.

FIGURE 3

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(a) The second heating DSC curves (b) crystallinity calculated from DSC curves (c) XRD patterns (d) stress‐strain curves (e) stress at break (f) fracture energy (g) f50: the time it takes for 50% of test specimens to fail under specified stress and environmental conditions, and (h) photographs of the samples during ESC tests of neat HDPE, HDPE/DCP0.2 and HDPE/DTBMIX blends.

Although crystallinity is reduced, the formation of disulfide‐linked network structures effectively compensates for this loss, resulting in even improved mechanical performance of HDPE. The tensile stress‐strain curves, stress at break, and fracture energy of HDPE and modified samples are shown in Figure 3d–f, with detailed mechanical parameters summarized in Table S2. Neat HDPE exhibits typical ductile behavior with a yield strength of 27.7 MPa, a tensile strength of 15.5 MPa, and an elongation at break of 836%. All DTBMI‐crosslinked samples display significantly increased plateau stress and pronounced strain hardening after yielding, indicating enhanced resistance to chain slippage and disentanglement during deformation. The yield strength of the crosslinked samples remains comparable to that of neat HDPE, demonstrating that material stiffness is well preserved. Among them, HDPE/DTBMI1.5 exhibits the highest tensile strength (26.4 MPa) together with an elongation at break of 940%, suggesting an optimal balance between network constraints and chain extensibility.

The tensile fracture energy further reflects the energy dissipation capability of the materials. Neat HDPE exhibits a fracture energy of 130.5 MJ/m3, which increases to 183 MJ/m3 for HDPE/DCP0.2 due to enhanced chain branching and entanglement. For DTBMI‐crosslinked samples, the fracture energy increases from 152 MJ/m3 at 1.0 phr to a maximum of 193 MJ/m3 at 1.5 phr, followed by a slight decrease at higher DTBMI loadings. This trend indicates that a moderate dynamic crosslinking maximizes energy absorption during tensile deformation, whereas excessive crosslinking partially limits chain extensibility. Although HDPE/DCP0.2 also exhibits increased elongation at break and moderate strain hardening, its overall reinforcing effect remains inferior to that achieved by the DTBMI‐based dynamic network.

Building on the enhanced resistance to chain slippage and disentanglement observed during tensile deformation, the ESCR of HDPE/DTBMIx was further evaluated (Figure 3g,h). Neat HDPE exhibits a short failure time (f50) of 12 h, whereas HDPE/DTBMIx samples show a pronounced and progressive increase in cracking lifetime with increasing DTBMI content, reaching a maximum f50 of 60 h for HDPE/DTBMI2.0. Further increasing the DTBMI loading does not lead to additional improvement, which is consistent with the saturation behavior observed in torque curves (Figure 2c) and gel fraction (Figure 2d). In contrast, HDPE/DCP0.2 shows only a modest increase in f50 to 21 h, indicating that chain branching and limited crosslinking contribute to ESCR enhancement but are far less effective than the DTBMI‐based network.

Notably, HDPE/DCP0.2 and HDPE/DTBMI2.0 exhibit comparable crystallinity (∼60%), yet the ESCR of HDPE/DTBMI2.0 is nearly three times higher. This comparison clearly indicates that crystallinity reduction alone cannot account for the observed ESCR improvement. Instead, the dominant factor is the disulfide‐crosslinked network constructed by DTBMI, which imposes strong constraints on chain slippage and disentanglement in the amorphous phase. Such constraints effectively retard interlamellar chain relaxation and slow down slow crack growth under environmental stress, thereby leading to a significant enhancement in ESCR performance.

3.3. Intermolecular Entanglement and Chain Mobility Restriction Induced by DTBMI

To further elucidate the molecular origin of the simultaneously enhanced mechanical performance and ESCR, the thermo‐mechanical behavior and melt rheology of HDPE/DTBMIx were investigated. Dynamic mechanical analysis (DMA) reveals that all samples exhibit a sharp drop in storage modulus (E′) at around 150°C due to crystal melting (Figure 4a). In contrast to neat HDPE, DTBMI‐crosslinked samples display a pronounced rubbery plateau above the melting temperature. The plateau modulus increases systematically with DTBMI content, rising from 0.27 MPa for HDPE/DTBMI1.0 to 0.35 MPa for HDPE/DTBMI2.0, accompanied by a continuous decrease in tan δ (Figure 4b). HDPE/DCP0.2 shows only a moderate increase in E′, confirming that DTBMI introduces much stronger intermolecular constraints than peroxide‐induced branching alone.

FIGURE 4.

FIGURE 4

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Thermo‐mechanical response and melt rheology of HDPE‐based systems. (a) Storage modulus (E′) and (b) loss factor (tan δ) as a function of temperature from DMA. Frequency‐dependent rheological behavior at 190°C: (c) storage modulus (G′), (d) effective elastic chain density (νe) derived from the low‐frequency plateau modulus, (e) tan δ, and (f) complex viscosity (η*).

To exclude the influence of crystallinity and enable quantitative analysis, oscillatory shear rheology was performed in the melt state at 190°C (Figure 4c–f). All samples exhibit a low‐frequency plateau in storage modulus (G′), reflecting long relaxation times associated with chain entanglements or network structures (Figure 4c). Notably, the plateau modulus of HDPE/DTBMIx is more than two orders of magnitude higher than that of neat HDPE, indicating the formation of an effective elastic network rather than purely physical entanglements. Based on the rubber elasticity relationship (νe = G′/RT), the effective elastic chain density (νe) was estimated using the low‐frequency plateau modulus. As summarized in Figure 4d, the νe values increase from 0.03 mol/m3 for HDPE to 0.56 mol/m3 for HDPE/DCP0.2, and further to 3.19–8.62 mol/m3 for HDPE/DTBMIx, providing direct quantitative evidence for a dynamically crosslinked network with strongly constrained chain mobility.

The frequency‐dependent tan δ curves further support this conclusion (Figure 4e). Neat HDPE exhibits a clear relaxation peak within the tested frequency range, whereas no relaxation peak is observed for HDPE/DTBMIx with a high DTBMI content, indicating that long‐time chain relaxation is effectively suppressed by the DTBMI‐induced network.

Despite the strong network constraints at long time scales, all DTBMI‐crosslinked samples exhibit pronounced shear‐thinning behavior (Figure 4f). At high frequencies, the complex viscosities of HDPE/DTBMIx converge with those of neat HDPE, suggesting that mechanical force accelerates disulfide exchange and enables rapid network rearrangement under high shear [31]. In contrast, the significantly elevated low‐frequency viscosity and the absence of a viscosity plateau reflect inhibited chain disentanglement and prolonged relaxation times. This time‐scale‐dependent mobility, restricted at long times yet accessible at short times, provides a molecular basis for the enhanced resistance to slow crack growth while retaining good melt processability.

3.4. Fracture Morphology and ESCR Mechanism of HDPE/DTBMIx

To further elucidate the mechanism underlying ESCR enhancement, the fracture morphologies of HDPE‐based samples after ESCR testing were examined by scanning electron microscopy (SEM). As shown in Figure 5, neat HDPE and HDPE/DCP0.2 exhibits a relatively smooth fracture surface with several large protrusions, characteristic of brittle failure dominated by rapid crack propagation. By contrast, all DTBMI‐crosslinked samples exhibit distinct fibrillar structures on the fracture surfaces, indicative of extensive chain stretching, orientation, and plastic deformation during crack propagation. These features demonstrate that disulfide‐linked dynamic covalent network effectively enhances intermolecular interactions in the amorphous phase and promotes stress redistribution at the microscale. With increasing DTBMI content, the fibrillar structures become more developed; however, no significant morphological differences are observed beyond 2.0 phr, consistent with the saturation behavior observed in ESCR, gel fraction, and rheological measurements.

FIGURE 5.

FIGURE 5

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SEM images of fracture cross‐sections after ESCR testing (Scale bar = 50 µm, Original magnification: ×1000): (a) HDPE, (b) HDPE/DCP0.2, (c) HDPE/DTBMI1.0, (d) HDPE/DTBMI1.5, (e) HDPE/DTBMI2.0, and (f) HDPE/DTBMI2.5. (g) Schematic illustration of the environmental stress cracking mechanism in neat HDPE and DTBMI‐crosslinked HDPE.

Based on these observations, the ESCR enhancement mechanism is schematically illustrated in Figure 5g. In neat HDPE, weak physical entanglements permit interlamellar chain slippage and lamellar separation under environmental stress, leading to premature crack initiation and rapid slow‐crack propagation. In contrast, the DTBMI‐induced dynamic covalent network, preferentially distributed in the amorphous regions, constrains chain disentanglement and bridges adjacent lamellae. Upon loading, these covalent junctions resist stress localization and promote cooperative chain orientation rather than interfacial debonding, thereby retarding crack growth and significantly improving ESCR.

3.5. Dynamic Characteristics of the DTBMI‐Crosslinked HDPE Network

Although the HDPE/DTBMIx samples exhibit clear improvements in mechanical performance and ESCR, their relatively low gel fractions indicate that stress relaxation may arise from a combination of dynamic bond exchange and chain disentanglement. Under such conditions, it is difficult to unambiguously isolate the contribution of dynamic bond exchange from that of uncrosslinked or weakly entangled chains. Notably, under a fixed DCP content of 0.2 phr, increasing the DTBMI content from 1.5 to 2.5 phr leads to only a limited increase in gel fraction (34.1%–37.8%), suggesting that the extent of network formation is constrained by the availability of radicals. To demonstrate the intrinsic dynamic nature of the DTBMI‐based network more clearly, a highly crosslinked sample (HDPE/DTBMI4.0) was therefore prepared by increasing both DTBMI and DCP contents. The torque evolution is shown in Figure S4a. For a consistent comparison, a corresponding control sample with the same DCP content (HDPE/DCP0.4) was used. Gel fraction analysis reveals that HDPE/DCP0.4 alone remains fully soluble (Figure S4b,c), whereas HDPE/DTBMI4.0 reaches a gel content of 48.6%, confirming that a higher degree of network dynamic crosslinking was achieved.

Stress relaxation experiments of HDPE/DTBMI4.0 were conducted by DMA to probe the dynamic behavior of the crosslinked network (Figure 6a). Despite its relatively high gel fraction, HDPE/DTBMI4.0 is capable of fully relaxing stress at elevated temperatures (170°C–200°C), demonstrating that the network undergoes topological rearrangement via dynamic bond exchange. In contrast, a control sample containing only DCP (HDPE/DCP0.4) shows much faster stress relaxation (Figure S5), which is characteristic of viscoelastic flow of uncrosslinked or weakly branched HDPE rather than network rearrangement. The significantly different relaxation behavior between the DCP‐only sample and the DTBMI‐crosslinked network indicates that the observed stress relaxation in HDPE/DTBMI4.0 is not dominated by peroxide‐induced branching, but is closely associated with dynamic bond exchange within the covalent network. It should be noted, however, that the observed stress relaxation arises from the combined contributions of dynamic network rearrangement and the relaxation of uncrosslinked chain segments, owing to the overall moderate crosslinking density. This is further evidenced by the rapid stress decay observed at short relaxation times. Meanwhile, the surface scratch on HDPE/DTBMI4.0 self‐healed after thermal treatment at 190°C for 2 h, further supporting the dynamics of the network (Figure S6).

FIGURE 6.

FIGURE 6

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(a) Stress relaxation curves of HDPE/DTBMI4.0 and (b) stress‐strain curves and (c) photographs of HDPE/DTBMI4.0 after multiple processing.

To further assess reprocessability, HDPE/DTBMI4.0 was subjected to repeated hot‐pressing cycles (Figure 6b,c). It should be noted that the first‐cycle sample is already a reprocessed material, since crosslinking occurs during the initial reactive blending, indicating that the dynamically crosslinked HDPE retains good melt processability upon reprocessing. Mechanical properties of the reprocessed samples are summarized in Table S2. The yield strength keeps almost unchanged at ∼22 MPa and even increases slightly over multiple cycles, indicating that the stiffness is well preserved. In contrast, the elongation at break decreases progressively, from ∼600% after the first cycle to ∼178% after two cycles and ∼138% after three cycles. This progressive decrease in elongation at break can be attributed to a combined effect of possible irreversible C–C crosslink accumulation and thermo‐oxidative side reactions during processing. FTIR analysis reveals the emergence of a new absorption band at 1762 cm− 1 after the third reprocessing cycle (Figure S7a), indicating the occurrence of thermo‐oxidative side reactions. These reactions, together with unavoidable chain scission under repeated thermal and shear stress, preferentially reduce large‐strain deformability. Notably, the gel fraction remains essentially unchanged upon reprocessing (Figure S7b), indicating that the overall network connectivity is largely preserved.

4. Conclusions

In this work, a dynamic covalent network was successfully introduced into HDPE through reactive melt blending using a disulfide‐containing bismaleimide (DTBMI). During melt blending, DTBMI captures HDPE macroradicals to form a dynamically crosslinked network that imposes strong constraints on long‐time chain motion while retaining thermoplastic processability. With the incorporation of 2 phr DTBMI, the ESCR failure time (f50) increases from 12 to 60 h. Simultaneously, tensile strength is enhanced from 15.5 to 25.9 MPa, and tensile fracture energy increases from 130.5 to 191.3 MJ·m− 3, accompanied by pronounced strain‐hardening behavior, while the yield strength remains essentially unchanged. Rheological and DMA analyses reveal a pronounced enhancement in rubbery and melt‐state modulus and a substantial increase in effective network density, confirming restricted chain disentanglement and prolonged relaxation times. Importantly, the network exhibits time‐scale‐selective behavior: chain mobility is strongly suppressed under low‐frequency or long‐term stress, whereas short‐time molecular motion remains active due to dynamic disulfide exchange, enabling efficient melt processing. Fracture morphology after ESCR testing further supports this mechanism, showing extensive fibrillation and improved stress redistribution. This study demonstrates a facile and robust strategy to overcome the long‐standing trade‐off between ESCR, mechanical robustness, and processability in HDPE. The proposed design principle may be extended to other polyolefin systems where long‐term durability under environmental stress is required.

Author Contributions

Y. Chen performed most of the experiments and wrote the original draft. W Tian, Y. J. Zhu, and Q. Xu helped with some of the experiments. J Huang conceived the presented idea, supervised the project, and revised the manuscript. T. Li and S. Wang helped with partial data analysis. W. F. Dong supervised the project and revised the manuscript. All authors discussed the results and contributed to the final manuscript.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: marc70324‐sup‐0001‐SuppMat.docx.

MARC-47-e00967-s001.docx (1.2MB, docx)

Acknowledgements

This work was supported by the National Natural Science Foundation of China (52303127, 52273089) and Jiangsu Young Sci‐Tech Talent Support Program (JSTJ‐2025‐762).

Contributor Information

Jing Huang, Email: jinghuang@jiangnan.edu.cn.

Weifu Dong, Email: wfdong@jiangnan.edu.cn.

Data Availability Statement

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

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Associated Data

This section collects any data citations, data availability statements, or supplementary materials included in this article.

Supplementary Materials

Supporting File: marc70324‐sup‐0001‐SuppMat.docx.

MARC-47-e00967-s001.docx (1.2MB, docx)

Data Availability Statement

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

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