Wear-resistant sneaker outsole and preparation method thereof

Abstract

The invention relates to the field of polymer composite materials, and discloses a wear-resistant sports shoe outsole and a preparation method thereof.The outsole comprises a polyurethane elastomer matrix, ceramic nanoparticles are generated in situ in the matrix, and a dynamic reversible chemical bond cross-linked network is constructed; a gradient dynamic interface is formed between the ceramic nanoparticles and the polyurethane elastomer matrix through a bifunctional interface anchoring molecule; the polyurethane elastomer matrix is prepared by reaction of a polyalcohol component, a diisocyanate component, a first type of reaction monomer introduced with a dynamic reversible chemical bond functional group, a ceramic nanoparticle precursor, a bifunctional interface anchoring molecule, a chain extender component and the like. Through in-situ generation of a reinforced phase, construction of a dynamic reversible cross-linked network and formation of a unique molecular anchoring gradient dynamic interface between the reinforced phase and a matrix, the wear resistance, toughness, fatigue life resistance and durability of the sports shoe outsole material are synergistically improved.

Description

Technical Field

[0001] The present invention relates to the technical field of polymer composite materials, in particular to a wear-resistant sports shoe outsole and a preparation method thereof. Background Art

[0002] As the part of a sports shoe that comes into direct contact with the ground, its performance is crucial to athletic performance, comfort, and the lifespan of the shoe. Polyurethane (PU) elastomers are widely used in the manufacture of sports shoe outsoles due to their excellent elasticity, moderate hardness, and good processing properties. However, with the increasing intensity of athletic activity and increasingly stringent consumer demands for sole performance, traditional polyurethane sole materials are increasingly lacking in terms of wear resistance, tear strength, fatigue life, and overall mechanical properties under extreme conditions.

[0003] In order to improve the performance of polyurethane soles, researchers have tried a variety of modification methods. Among them, the introduction of inorganic fillers is a commonly used reinforcement method, hoping to improve the wear resistance and mechanical strength of the material through the high hardness and high modulus of the filler. Despite this, this simple physical blending method often faces many challenges. The compatibility between inorganic fillers and organic polyurethane matrices is usually poor, which causes the filler particles to easily agglomerate in the matrix and is difficult to achieve uniform dispersion. This uneven dispersion not only fails to fully exert the reinforcing effect of the filler, but may become a stress concentration point, accelerating the destruction of the material. In addition, the interface formed between the filler and the matrix is ​​usually weakly bonded, and the interface is prone to debonding when subjected to force, making it impossible for stress to be effectively transferred from the matrix to the reinforcing phase, thereby limiting the improvement in the overall mechanical properties of the material.

[0004] Furthermore, the sole pursuit of improvements in hardness and wear resistance sometimes comes at the expense of the material's toughness and fatigue resistance. When the material is subjected to repeated impact or bending, microcracks are likely to form and expand in stress concentration areas, ultimately leading to fatigue fracture of the material. Traditional polyurethane materials lack an effective energy dissipation mechanism to inhibit the expansion of microcracks, and it is difficult to repair minor damage generated during use. This makes the soles prone to performance degradation or even damage after high-intensity and long-term use. Therefore, how to ensure the material's high strength and high wear resistance while giving it excellent toughness, outstanding fatigue life and tolerance to damage is a technical problem that needs to be solved urgently in the current research and development of high-performance sports shoe outsole materials. The development of a new sole material that can achieve synergistic enhancement of multiple properties and perform precise regulation at the microstructural level has important practical significance and application prospects. Summary of the Invention

[0005] The technical problem to be solved by this invention is to overcome the shortcomings of the existing technology and provide a sports shoe outsole material with significantly improved wear resistance and durability, as well as a method for preparing the same. This material achieves excellent overall performance by in-situ generation of a reinforcing phase, construction of a dynamic, reversible cross-linking network, and formation of a unique molecularly anchored gradient dynamic interface between the reinforcing phase and the matrix.

[0006] In order to solve the above technical problems, the present invention adopts the following technical solutions: A first aspect of the present invention provides a wear-resistant sports shoe outsole, wherein the outsole comprises a polyurethane elastomer matrix, and the polyurethane elastomer matrix is ​​prepared by reacting the following components: (1) a polyol component; (II) a diisocyanate component; (III) a first type of reactive monomer capable of introducing a dynamically reversible chemical bond functional group; (IV) a ceramic nanoparticle precursor, which, after being converted into in-situ generated ceramic nanoparticles, has a content of 2-15 wt % in the final outsole based on the total mass of the final outsole; (V) a bifunctional interface anchoring molecule, the amount of which is 1-10 mol% of the molar amount of the ceramic nanoparticle precursor; (VI) a chain extender component; (VII) and an initiator (such as water) for initiating in situ hydrolysis and condensation of the ceramic nanoparticle precursor; Wherein, the polyurethane elastomer matrix comprises: (a) ceramic nanoparticles generated in situ from the ceramic nanoparticle precursor; (b) a dynamically reversible chemical bond cross-linked network formed by the first type of reactive monomers and possibly the second type of reactive monomers or molecules; (c) The bifunctional interface anchoring molecule has one end chemically bonded or strongly physically adsorbed to the surface of the in situ generated ceramic nanoparticles, and the other end chemically bonded or physically embedded in the network structure of the polyurethane elastomer matrix, or directly participates in the cross-linked network formed by the dynamic reversible chemical bond.

[0007] During the polyurethane prepolymer preparation stage, the molar ratio (R value) of the NCO groups in the diisocyanate component to the total hydroxyl groups in the polyol component and the first type of reactive monomers is controlled to be between 1.8 and 2.5. The molar ratio of the polyol component to the first type of reactive monomers can be controlled to be between 70:30 and 95:5.

[0008] During the chain extension stage, the molar ratio of the total active hydrogen in the chain extender component (mainly from the hydroxyl or amino groups of the chain extender) to the remaining NCO groups in the prepolymer is controlled between 0.95 and 1.05.

[0009] If the dynamically reversible chemical bond is formed by the reaction of the furan group and the maleimide group, the molar ratio of the maleimide group to the furan group is controlled to be between 0.9:1 and 1.1:1.

[0010] The molar ratio of the initiator (such as water) used to initiate the in-situ hydrolysis and condensation of the ceramic nanoparticle precursor to the molar ratio of the ceramic nanoparticle precursor is controlled between 2:1 and 4:1.

[0011] In some embodiments of the present invention, the in situ generated ceramic nanoparticles may be selected from at least one of silicon dioxide nanoparticles or titanium dioxide nanoparticles.

[0012] By generating these nano-ceramic particles in situ during the polymerization or curing process of the polyurethane matrix, they can be evenly distributed in the matrix in a highly dispersed state, avoiding the agglomeration problem that is prone to occur with traditional externally added nano-fillers, thereby more effectively exerting their reinforcing effect and forming a tighter initial interface with the matrix.

[0013] In some embodiments of the present invention, the crosslinked network formed by the dynamically reversible chemical bonds is formed via a Diels-Alder reaction. For example, furan groups can be introduced into polyurethane molecular segments (via the first type of reactive monomers, such as furan-containing diols), and then a crosslinker or segment containing a maleimide group (as part or all of the chain extender component, or as the second type of reactive monomers that introduce dynamically reversible chemical bond functional groups) can undergo a Diels-Alder reaction under specific conditions to form thermally reversible covalent crosslinks.

[0014] This dynamic network enables the material to dissipate energy through the reversible breaking and recombination of chemical bonds when subjected to external forces, inhibiting the expansion of microcracks, thereby significantly improving the material's toughness and fatigue life, and giving the material a certain damage tolerance.

[0015] In some embodiments of the present invention, the bifunctional interface anchoring molecule comprises a ceramic affinity / reactive end and a dynamic network participating end. The ceramic affinity / reactive end can be selected from at least one of an alkoxysilane group (e.g., -Si(OR)3), a phosphonic acid group, or a carboxylic acid group, each of which is capable of forming a stable chemical bond (e.g., a Si-O-ceramic metal bond) or strong adsorption to the surface of the in situ generated ceramic nanoparticles.

[0016] The dynamic network participating end can be selected from at least one of a furan group, a maleimide group, or other active groups (such as an isocyanate group, a hydroxyl group, and an amino group) that can form a covalent bond with a polyurethane segment or a dynamic cross-linked network.

[0017] The introduction of this dual-functional interface anchoring molecule creates a unique gradient interface layer between the inorganic ceramic nanoparticles and the organic polyurethane matrix. This interface layer not only greatly enhances the interfacial bonding strength through chemical bonding, ensuring effective stress transfer between the matrix and the reinforcement phase, but more importantly, due to the presence of its dynamic network participating end, the characteristics of the dynamic reversible cross-linking network can be extended to the interface region, making the interface itself dynamically responsive, effectively dissipating energy at the interface, and further synergistically improving the overall toughness and wear resistance of the material.

[0018] The uniformly dispersed nano-ceramic particles generated in situ provide a rigid and wear-resistant foundation, the dynamically reversible cross-linked network provides toughness and an energy dissipation mechanism, and the synergistic effect of efficiently coupling the two through a molecularly anchored gradient dynamic interface gives the sole material of the present invention comprehensive performance that surpasses traditional materials.

[0019] A second aspect of the present invention provides a method for preparing the wear-resistant sports shoe outsole, which comprises the following steps: (S1) mixing a polyol component, a diisocyanate component, a first type of reactive monomer capable of introducing a dynamically reversible chemical bond functional group, a ceramic nanoparticle precursor, and a bifunctional interface anchoring molecule, and reacting them under specific conditions, wherein the molar ratio of the NCO group in the diisocyanate component to the total hydroxyl group in the polyol component and the first type of reactive monomer is controlled to be between 1.8 and 2.5, the molar ratio of the polyol component to the first type of reactive monomer is controlled to be between 70:30 and 95:5, and the amount of the bifunctional interface anchoring molecule is 1-10 mol% of the molar amount of the ceramic nanoparticle precursor, to prepare a polyurethane prepolymer containing the ceramic precursor and the bifunctional interface anchoring molecule; (S2) mixing the polyurethane prepolymer with a chain extender component (and a possible second type of reactive monomer or molecule that introduces a dynamic reversible chemical bond functional group), and adding an initiator (such as water, the molar ratio of which to the molar ratio of the ceramic nanoparticle precursor is controlled to be between 2:1 and 4:1), and under initiation conditions, causing the ceramic nanoparticle precursor to undergo in-situ hydrolysis and condensation reactions to generate ceramic nanoparticles (the content of which in the final outsole is 2-15wt% of the total mass of the final outsole), while the polyurethane prepolymer completes a chain extension reaction to form a high molecular weight polyurethane elastomer matrix (wherein the total weight of the chain extender component is 100wt%). The molar ratio of active hydrogen to residual NCO groups in the prepolymer is controlled to be between 0.95 and 1.05), and a dynamically reversible chemical bond cross-linked network is formed (for example, if the furan group and the maleimide group react, the molar ratio of the maleimide group to the furan group is controlled to be between 0.9:1 and 1.1:1); during this process, one end of the bifunctional interface anchoring molecule is bonded to or strongly adsorbed to the surface of the in situ generated ceramic nanoparticles, and the other end is integrated into the network structure of the polyurethane elastomer matrix through a chemical reaction or directly participates in the formation of the dynamically reversible chemical bond cross-linked network; (S3) The mixture obtained in step (S2) is subjected to compression molding and curing to obtain the wear-resistant sports shoe outsole.

[0020] In some embodiments of the present invention, the first type of reactive monomer capable of introducing a dynamic reversible chemical bond functional group in step (S1) may comprise a diol having a furan group. Furthermore, / or, the bifunctional interface anchoring molecule itself may comprise a furan group or a maleimide group as its dynamic network participating end.

[0021] In some embodiments of the present invention, the ceramic nanoparticle precursor in step (S1) can be selected from metal alkoxides, such as tetraethyl silicate (TEOS) or tetrabutyl titanate, etc. These precursors can be converted into corresponding nano-ceramic particles in situ through a controllable hydrolysis-condensation reaction in subsequent steps.

[0022] In some embodiments of the present invention, the chain extender component in step (S2) includes a reactive group that matches the dynamically reversible chemical bond functional group in the first type of reactive monomer or the dynamic network participating end in the bifunctional interface anchoring molecule, thereby forming the dynamically reversible chemical bond crosslinked network via, for example, a Diels-Alder reaction. For example, if a furan group is introduced into the prepolymer, the chain extender component may include a maleimide group.

[0023] The innovation of this preparation method lies in the precise control of the reaction steps and the stoichiometric ratios of the components, enabling the in situ generation of nanoceramic particles, the formation of the polyurethane matrix, the construction of a dynamic reversible cross-linking network, and the simultaneous construction of the critical molecular anchoring gradient dynamic interface. The prepolymer stage uniformly disperses the ceramic precursor and interface anchoring molecules, ensuring the uniformity of the subsequent in situ reaction and the interface anchoring effect. The coordinated execution of multiple reactions during the chain extension and curing stages allows for precise control of the final material structure, resulting in the desired synergistic enhancement effect.

[0024] In summary, the present invention includes at least one of the following beneficial technical effects: 1. This invention achieves high dispersion and strong interfacial bonding of ceramic nanoparticles by in-situ generation within a polyurethane matrix. This is achieved through the strong chemical bonding or physical adsorption of bifunctional interfacial anchoring molecules between the nanoparticles and the polyurethane matrix. This not only effectively transfers stress but also enables the high-hardness ceramic nanoparticles to fully exert their reinforcing effect, thereby imparting excellent wear resistance, higher tensile strength, and tear strength to the sole material.

[0025] 2. This invention introduces a cross-linked network of dynamically reversible chemical bonds within the polyurethane matrix, enabling these dynamic bonds to reversibly break and reform when subjected to impact or repeated stress. This process effectively dissipates external energy, blunts microcrack tips, and prevents rapid crack propagation, thereby imparting the sole material with exceptional toughness, significantly extended fatigue life, and increased tolerance to minor damage.

[0026] 3. The present invention introduces and acts through a dual-functional interface anchoring molecule. One end of the molecule is firmly bound to the surface of the in-situ generated ceramic nanoparticles, and the other end participates in the network structure or dynamic reversible cross-linking network of the polyurethane matrix. This design constructs a molecularly anchored gradient dynamic interface layer, so that the energy dissipation mechanism of the dynamic network not only acts on the matrix, but also extends to the filler-matrix interface area where the stress is most concentrated. This not only strengthens the interface bonding and prevents premature failure of the interface, but also enables the interface itself to have dynamic response and energy absorption capabilities, thereby synergistically improving the overall mechanical properties of the material and its durability in long-term use.

[0027] 4. This invention utilizes an in-situ hydrolysis and condensation process of ceramic nanoparticle precursors, allowing the ceramic nanoparticles to be directly generated and uniformly dispersed within the polyurethane matrix during formation. This fundamentally avoids the problem of traditional prefabricated nanofillers, which tend to agglomerate and struggle to disperse uniformly within the matrix. Combined with the stabilizing effect of dual-functional interfacial anchoring molecules, this ensures the long-term, stable distribution of the reinforcing phase within the matrix, thereby guaranteeing the uniformity and reliability of the sole material's macroscopic properties.

[0028] 5. The preparation method of the present invention provides the possibility of customizing sole materials with specific comprehensive properties such as wear resistance, shock absorption, and rebound according to different sports needs by controlling the various components and regulating the conditions of key steps such as prepolymerization, chain extension, in-situ generation, and curing, showing good engineering application prospects. DETAILED DESCRIPTION

[0029] To make the purpose, technical solution and advantages of the present invention clearer, the technical solution of the present invention will be described in detail below. It should be understood that the specific embodiments described herein are only used to explain the present invention and are not intended to limit the present invention.

[0030] Unless otherwise specified, the conventional methods, experimental procedures and reagents used in the examples of the present invention are conventional techniques in the art or commercially available. Unless otherwise defined, all professional and scientific terms used in the text have the same meanings as those familiar to those skilled in the art.

[0031] Unless otherwise specified, the raw materials used in the following examples and comparative examples are commercially available products and can be obtained from conventional commercial channels.

[0032] Polytetramethylene ether glycol (PTMG): number average molecular weight 1950±150g / mol, brand: BASF 2000. Dehydrate at 105±5℃ and vacuum degree less than 200Pa for 3±1 hours before use.

[0033] 4,4'-diphenylmethane diisocyanate (MDI): polymer grade MDI, brand: Wanhua Chemical PM-200 (polymeric MDI). If using solid MDI, it must be melted at 50±5℃ before use.

[0034] 2,5-Furan dimethanol (FDM): Chemically pure, purity ≥98%, supplier: Aladdin Reagent. Vacuum dry at 60±5°C for 4±1 h before use.

[0035] Tetraethyl silicate (TEOS): analytical grade, purity ≥99.0%, supplier: Sinopharm Chemical Reagent Co., Ltd.

[0036] N-(2-Furylmethyl)-3-aminopropyltrimethoxysilane (F-APTMOS): Custom synthesis, purity ≥95%.

[0037] Maleimidopropyltriethoxysilane (M-PTES): Custom synthesis, purity ≥95%.

[0038] N,N'-(Methylenebis-4,1-phenylene)bismaleimide (BMI): chemically pure, purity ≥97%, supplier: Tokyo Chemical Industry Development Co., Ltd. (Shanghai).

[0039] 1,4-Butanediol (BDO): analytical grade, purity ≥99.5%, supplier: Sinopharm Chemical Reagent Co., Ltd. Dry over 4A molecular sieves for at least 24 hours before use.

[0040] Stannous octoate (Sn(Oct)2): chemically pure, Sn content ≥ 28%, supplier: Alfa Aesar (China) Chemical Co., Ltd.

[0041] Dibutyltin dilaurate (DBTDL): chemically pure, purity ≥95%, supplier: Macklin's reagent.

[0042] Toluene: analytical grade.

[0043] Ethanol: analytical grade.

[0044] N-Methylpyrrolidone (NMP): analytical grade, water content ≤0.05%.

[0045] Dimethylformamide (DMF): analytical grade, water content ≤ 0.05%.

[0046] Deionized water: homemade in the laboratory, conductivity ≤2μS / cm.

[0047] The synthesis methods provided below are for reference only. Those skilled in the art may also prepare the products using other equivalent methods based on existing chemical knowledge.

[0048] Reference preparation method of N-(2-furylmethyl)-3-aminopropyltrimethoxysilane (F-APTMOS): (1) In a reaction flask equipped with a stirrer, a dropping funnel, and a condenser, 17.93 g (0.1 mol) of 3-aminopropyltrimethoxysilane (APTMOS) and 150 ml of anhydrous methanol were added and stirred to dissolve.

[0049] (2) Under ice-water cooling, slowly add dropwise a solution of 9.61 g (0.1 mol) of furfural (freshly distilled) dissolved in 50 ml of anhydrous methanol, maintaining the temperature below 10°C. After the addition is complete, continue stirring at room temperature for 4 hours to form the imine intermediate.

[0050] (3) The reaction mixture was placed in an ice-water bath again, and 7.57 g (0.2 mol) of sodium borohydride (NaBH4) was slowly added in batches. After the addition was complete, the ice bath was removed and the reaction was continued at room temperature with stirring for 12 hours.

[0051] (4) After the reaction is complete, remove most of the methanol by rotary evaporation. Add 100 ml of ethyl acetate and 50 ml of saturated ammonium chloride aqueous solution to the residue and separate the layers. Extract the aqueous phase with ethyl acetate (50 ml x 2). Combine the organic phases and dry them over anhydrous sodium sulfate.

[0052] (5) The desiccant is removed by filtration, and the filtrate is evaporated under reduced pressure to remove the solvent to obtain a light yellow oily liquid, namely N-(2-furylmethyl)-3-aminopropyltrimethoxysilane (F-APTMOS). It can be further purified by column chromatography (silica gel, petroleum ether / ethyl acetate gradient elution).

[0053] Reference preparation method of maleimidopropyltriethoxysilane (M-PTES): (1) In a reaction flask equipped with a stirrer, a dropping funnel, and a thermometer, 22.14 g (0.1 mol) of 3-aminopropyltriethoxysilane (APTES) and 100 ml of anhydrous tetrahydrofuran (THF) were added.

[0054] (2) While stirring and cooling in an ice-water bath, slowly add dropwise a solution of 9.81 g (0.1 mol) of maleic anhydride dissolved in 50 ml of anhydrous THF, maintaining the temperature below 5°C during the addition. After the addition is complete, continue stirring at room temperature for 3 hours to form a white precipitate of N-(3-triethoxysilylpropyl)maleamic acid.

[0055] (3) Add 15.31 g (0.15 mol) of acetic anhydride and 0.82 g (0.01 mol) of anhydrous sodium acetate to the above reaction mixture. Heat the mixture to 60-70°C and stir at this temperature for 4-6 hours to carry out dehydration cyclization.

[0056] (4) After the reaction is completed, cool to room temperature. Filter to remove insoluble matter. The filtrate is rotary evaporated under reduced pressure to remove THF, excess acetic anhydride, and generated acetic acid.

[0057] (5) The residue is dissolved in an appropriate amount of n-hexane. If any insoluble matter is present, filter again. The filtrate is concentrated under reduced pressure to obtain a pale yellow oily liquid, which is maleimidopropyltriethoxysilane (M-PTES). This can be further purified by vacuum distillation or column chromatography (silica gel, n-hexane / ethyl acetate gradient elution).

[0058] Examples 1 to 3: Example 1: Preparation of a polyurethane elastomer containing 10 wt% in-situ generated SiO2, a Diels-Alder dynamic network, and F-APTMOS interface anchoring 1. Preparation of polyurethane prepolymer: In a 500 mL four-necked flask equipped with a mechanical stirrer, a thermometer, a reflux condenser, and a nitrogen protection device, 180.0 g (0.09 mol) of pre-dehydrated polytetramethylene ether glycol (PTMG, Mn = 2000 g / mol) and 1.28 g (0.01 mol) of pre-dried 2,5-furan dimethanol (FDM) were added.

[0059] Start stirring and raise the temperature to 70±2℃.

[0060] 6.62 g (0.02439 mol) of N-(2-furylmethyl)-3-aminopropyltrimethoxysilane (F-APTMOS) was added, and the mixture was stirred at 70±2° C. for 45 minutes.

[0061] Subsequently, 60.58 g of 4,4'-diphenylmethane diisocyanate (MDI) (corresponding to 0.44878 mol of NCO groups, calculated based on the NCO content of the MDI used, ensuring a molar ratio of NCO to total active hydrogen (-OH from PTMG, -OH from FDM, and -NH from F-APTMOS) of 2.0:1) was added. The reaction was stirred and continued at 70 ± 2°C for 2 hours.

[0062] After the reaction system was cooled to 60±2°C, 101.62 g (0.4878 mol) of tetraethyl silicate (TEOS) was added and stirring was continued at 60±2°C for 45 minutes to obtain a polyurethane prepolymer containing TEOS and F-APTMOS. The entire process was carried out under nitrogen protection.

[0063] 2. Chain extension, in-situ ceramic formation, dynamic network construction and molding curing: The polyurethane prepolymer prepared above was maintained at 60±2°C.

[0064] In a separate container, 6.16 g (0.017195 mol) of N,N'-(methylenebis-4,1-phenylene)bismaleimide (BMI) and 10.11 g (0.112195 mol) of 1,4-butanediol (BDO) were mixed thoroughly (ensuring a molar ratio of total active hydrogen to remaining NCO groups in the prepolymer of approximately 1.0:1, and a molar ratio of maleimide groups in BMI to total furan groups in the prepolymer of approximately 1.0:1). 0.132 g of stannous octoate (Sn(Oct)2) was added as a polyurethane reaction catalyst.

[0065] The mixture of the chain extender and the catalyst was quickly added to the prepolymer and stirred vigorously for 3 minutes.

[0066] Then, 26.37 g (1.4634 mol, ensuring a molar ratio of water to TEOS of 3.0:1) of deionized water was added and the mixture was rapidly stirred for 1 minute.

[0067] The mixed material was quickly subjected to vacuum degassing treatment (vacuum degree lower than 500 Pa, degassing for 3 minutes).

[0068] Pour the degassed mixture into a clean metal mold preheated to 100±2℃.

[0069] The mold was immediately placed in a flat-plate vulcanizer and press-cured at a temperature of 100±2°C and a pressure of 10±1 MPa for 2 hours.

[0070] 3. Post-processing: After curing is complete, carefully demould the mold after cooling to room temperature.

[0071] The demoulded sample was placed in an oven at 80±2° C. for post-treatment for 12 hours to obtain the polyurethane elastomer sample of Example 1. The theoretical content of in-situ generated SiO 2 in the sample was approximately 10 wt %.

[0072] Example 2: Preparation of a polyurethane elastomer containing 3 wt% in-situ generated SiO2, a Diels-Alder dynamic network, and F-APTMOS interface anchoring The preparation method of Example 2 is basically the same as that of Example 1, the main difference being that the content of in-situ generated SiO2 is reduced. The specific adjustments are as follows: In the polyurethane prepolymer preparation step, the amount of tetraethyl silicate (TEOS) added was adjusted to 28.60 g (0.1373 mol) so that the theoretical content of in-situ generated SiO2 in the final sample was about 3 wt%.

[0073] Accordingly, the amount of N-(2-furylmethyl)-3-aminopropyltrimethoxysilane (F-APTMOS) added was adjusted to 1.87 g (0.00689 mol, maintaining 5 mol% of the TEOS molar amount). Due to the change in the amount of F-APTMOS used, the exact amount of MDI needed to be recalculated and adjusted based on a molar ratio of NCO to total active hydrogen (from PTMG, FDM, and adjusted F-APTMOS) of 2.0:1.

[0074] In the chain extension step, the amount of catalytic water (deionized water) added was adjusted to 7.42 g (0.4118 mol, maintaining a water to TEOS molar ratio of 3.0:1).

[0075] The amounts of other raw materials (PTMG, FDM, BMI, BDO), reaction conditions, and post-treatment conditions remained the same as in Example 1. The molar ratio of BMI to total furan groups (from FDM and adjusted F-APTMOS) was maintained at approximately 1.0:1. The amount of BDO was adjusted based on the amount of residual NCO groups in the prepolymer after adjustment to maintain a molar ratio of total active hydrogen to residual NCO groups of approximately 1.0:1. The amount of catalyst was also adjusted proportionally based on the adjusted total mass of organic matter.

[0076] Example 3: Preparation of a polyurethane elastomer containing 10 wt% in situ generated SiO2, a Diels-Alder dynamic network, and M-PTES interface anchoring The preparation method of Example 3 is basically the same as that of Example 1, with the main difference being that maleimidopropyltriethoxysilane (M-PTES) is used instead of N-(2-furylmethyl)-3-aminopropyltrimethoxysilane (F-APTMOS) as the bifunctional interface anchoring molecule. The specific adjustments are as follows: In the polyurethane prepolymer preparation step, F-APTMOS is not added. Instead, maleimidopropyltriethoxysilane (M-PTES) is added in an amount such that its molar number is the same as the molar number of F-APTMOS in Example 1 (0.02439 moles), and the specific weight is calculated based on the molecular weight of M-PTES. The maleimide group of M-PTES will directly participate in the subsequent Diels-Alder reaction. Since M-PTES does not contain amino or hydroxyl groups that can pre-react with MDI, the amount of MDI used only needs to consider the hydroxyl groups of PTMG and FDM to ensure that the NCO / (OH) PTMG +OH FDM ) in a molar ratio of 2.0:1.

[0077] During the chain extension step, the amount of N,N'-(methylenebis-4,1-phenylene)bismaleimide (BMI) needs to be adjusted to ensure a molar ratio of total maleimide groups (from M-PTES and BMI) to total furan groups (from FDM) of approximately 1.0:1.

[0078] The amounts of other raw materials (PTMG, FDM, TEOS, BDO, water, catalyst) and the corresponding stoichiometric ratios (except for those involving the interface anchoring molecules and BMI), reaction conditions and post-treatment conditions were consistent with those in Example 1.

[0079] Comparative Examples 1 to 4: Comparative Example 1: Polyurethane elastomer without a dynamically reversible cross-linked network Compared with Example 1, the difference is: During the polyurethane prepolymer preparation step, 2,5-furan dimethanol (FDM) was omitted. The amount of MDI was adjusted accordingly to ensure that the molar ratio of NCO to total active hydrogen (-OH from PTMG and -NH from F-APTMOS) remained at 2.0:1.

[0080] During the chain extension step, N,N'-(methylenebis-4,1-phenylene)bismaleimide (BMI) was omitted, and only 1,4-butanediol (BDO) was used as the chain extender. The amount of BDO was calculated based on the amount of residual NCO groups in the prepolymer, such that the molar ratio of BDO hydroxyl groups to residual NCO groups was approximately 1.0:1. The remaining raw material amounts, reaction conditions, and post-treatment conditions were the same as those in Example 1.

[0081] Comparative Example 2: Polyurethane elastomer without in-situ generated ceramic nanoparticles Compared with Example 1, the difference is: In the polyurethane prepolymer preparation step, tetraethyl silicate (TEOS) is not added.

[0082] During the chain extension step, no catalytic water (deionized water) was added for TEOS hydrolysis.

[0083] Since TEOS was not added, the silane end of the bifunctional interface anchoring molecule F-APTMOS had no specific reaction partner, but its furan group and amino group still participated in the reaction as in Example 1. The remaining raw material amounts (except TEOS and catalytic water), stoichiometric ratios (for the polyurethane main reaction), reaction conditions, and post-treatment conditions were the same as in Example 1.

[0084] Comparative Example 3: Polyurethane elastomer without bifunctional interface anchoring molecules Compared with Example 1, the difference is: During the polyurethane prepolymer preparation step, N-(2-furylmethyl)-3-aminopropyltrimethoxysilane (F-APTMOS) was omitted. The amount of MDI was adjusted accordingly to ensure that the molar ratio of NCO to total active hydrogen (-OH from PTMG and -OH from FDM) remained at 2.0:1.

[0085] Since F-APTMOS was not added, the amount of BMI was adjusted accordingly to ensure a molar ratio of maleimide groups in BMI to total furan groups in the prepolymer (from FDM alone) of approximately 1.0:1. The amount of BDO was also adjusted based on the amount of residual NCO groups in the prepolymer. The remaining raw material amounts, reaction conditions, and post-processing conditions were the same as in Example 1.

[0086] Comparative Example 4: Conventional thermoplastic polyurethane (TPU) elastomer Conventional thermoplastic polyurethane elastomers are prepared without in-situ generated ceramics, dynamic cross-linked networks, and interfacial anchoring molecules.

[0087] Prepolymer Preparation: To a four-necked flask equipped with a mechanical stirrer, thermometer, reflux condenser, and nitrogen atmosphere, add 180.0 g (0.09 mol) of pre-dehydrated polytetramethylene glycol (PTMG, Mn = 2000 g / mol). Raise the temperature to 70±2°C. Add 45.05 g (corresponding to 0.18 mol of NCO groups, ensuring a molar ratio of NCO to PTMG -OH of 2.0:1) of 4,4'-diphenylmethane diisocyanate (MDI). Continue stirring at 70±2°C for 2 hours to obtain a TPU prepolymer.

[0088] Chain Extension and Molding: Cool the prepolymer to 60±2°C. Add 7.66 g (0.085 mol) of 1,4-butanediol (BDO) (ensuring a molar ratio of BDO hydroxyl groups to remaining NCO groups in the prepolymer of approximately 0.95:1). Add 0.11 g of stannous octoate as a catalyst. After rapid stirring, degas under vacuum. Pour the mixture into a clean metal mold preheated to 100±2°C and press-cure at 100±2°C and 10±1 MPa for 2 hours.

[0089] Post-treatment: After curing is completed, the mold is carefully demoulded after cooling to room temperature. The demoulded sample is placed in an oven at 80±2°C for post-treatment for 12 hours to obtain the TPU elastomer sample of Comparative Example 4.

[0090] Test cases 1 to 4: Test Example 1: Mechanical properties test Experimental description: This test example is intended to evaluate the basic mechanical properties of the polyurethane elastomer materials prepared according to the aforementioned embodiments and comparative examples, including tensile strength, elongation at break, tear strength, and Shore A hardness.

[0091] The experimental steps are as follows: All samples were placed in an environment with a temperature of 23±2°C and a relative humidity of 50±5% for at least 24 hours before mechanical property testing.

[0092] (a) Tensile strength and elongation at break test 1. Refer to GB / T 528-2009 "Vulcanized rubber or thermoplastic rubber - Determination of tensile stress-strain properties" standard.

[0093] 2. Cut the sheet samples prepared in each example and comparative example into standard dumbbell-shaped specimens using a standard dumbbell-shaped cutter (Type 2 or similar suitable size), ensuring that the edges of the specimens are smooth and free of defects. Prepare at least 6 valid specimens for each sample group.

[0094] 3. Use an electronic universal testing machine with a grip suitable for elastomeric materials. Set the tensile rate to 500 ± 50 mm / min.

[0095] 4. Clamp the specimen in the upper and lower fixtures of the testing machine, ensuring that the specimen's force axis coincides with the center line of the fixture, and apply appropriate initial clamping force to prevent slipping.

[0096] 5. Start the test and record the load-displacement curve of the specimen during the stretching process until the specimen breaks.

[0097] 6. Calculate the tensile strength (MPa) and elongation at break (%) of each specimen based on the recorded data.

[0098] 7. For each group of samples, after discarding outliers, take the arithmetic mean of at least 5 valid data as the final result.

[0099] (b) Tear strength test 1. With reference to GB / T 529-2008, Rubber, vulcanized or thermoplastic — Determination of tear strength (trouser-shaped, right-angled and crescent-shaped specimens), this test uses right-angled specimens.

[0100] 2. Cut the sheet samples prepared in each example and comparative example into specimens using a standard right-angle cutter and pre-cut at the right-angle inner corner (length 2.0±0.2 mm). Prepare at least 6 valid specimens for each sample group.

[0101] 3. Use an electronic universal material testing machine and set the tensile rate of the testing machine to 500±50mm / min.

[0102] 4. Clamp the two "trouser legs" of the right-angled specimen in the upper and lower clamps of the testing machine respectively.

[0103] 5. Start the test and record the maximum load or median load during the tearing process.

[0104] 6. Calculate the tear strength (kN / m) based on the recorded load and specimen thickness.

[0105] 7. For each group of samples, after discarding outliers, take the arithmetic mean of at least 5 valid data as the final result.

[0106] (c) Hardness (Shore A) test 1. Refer to GB / T 531.1-2008 "Rubber, vulcanized or thermoplastic — Test method for indentation hardness — Part 1: Shore durometer method (Shore hardness)" standard.

[0107] 2. Stack the sheet samples prepared in each embodiment and comparative example to a thickness of at least 6 mm (if the thickness of a single sheet is insufficient), ensuring that the test surface is flat and smooth.

[0108] 3. Use a Shore A hardness tester. Press the indenter vertically onto the surface of the specimen, with the indenter foot in close contact with the specimen surface.

[0109] 4. Read the hardness value within 3 seconds after the pressure foot is in full contact with the sample.

[0110] 5. Measure the hardness values ​​of at least 5 points at different locations of each sample (each point is at least 6 mm apart).

[0111] 6. Take the arithmetic mean of 5 valid readings as the Shore A hardness of the sample.

[0112] The experimental data are shown in Table 1: Table 1: Basic mechanical properties test results of various examples and comparative examples Sample number Tensile strength (MPa) Elongation at break (%) Tear strength (kN / m) Shore A hardness Example 1 48.2 685 72.5 89 Example 2 36.5 730 58.1 83 Example 3 46.8 670 70.3 88 Comparative Example 1 42.1 490 50.7 91 Comparative Example 2 20.3 750 33.6 76 Comparative Example 3 30.5 610 45.2 85 Comparative Example 4 25.7 580 38.9 80 By comparing the experimental data in Table 1, it can be seen that the comprehensive mechanical properties of the materials represented by Examples 1, 2 and 3, especially tensile strength, elongation at break and tear strength, are significantly different from those of the comparative samples. The sample of Example 1 exhibits higher tensile strength, excellent elongation at break and outstanding tear strength, which is attributed to its unique composite structure: in situ generated ceramic nanoparticles (such as SiO2 in Example 1) serve as reinforcing phases to improve the rigidity and strength benchmark of the material; at the same time, by introducing a dynamic reversible cross-linked network formed by specific reactive monomers (such as FDM and BMI), the polyurethane matrix is ​​endowed with the ability to dissipate energy by reversible breaking and reorganization of bonds when subjected to stress, thereby improving the toughness of the material and its ability to resist crack propagation; and the bifunctional interface anchoring molecules (such as F-APTMOS) construct an effective stress transfer bridge between the inorganic ceramic particles and the organic polyurethane matrix, ensuring the synergistic effect between the two, so that the reinforcing effect of the reinforcing phase and the toughening effect of the dynamic network can be fully utilized.

[0113] Specific analysis of the comparative data can further support the above mechanism. As a conventional polyurethane elastomer, the mechanical properties of Comparative Example 4 are all at a low level. Although Comparative Example 2 (without in-situ generated ceramic nanoparticles) may maintain a higher elongation at break due to the presence of a dynamic network, its tensile strength, tear strength and hardness all decrease significantly, indicating the key contribution of in-situ generated ceramic nanoparticles to the overall strength and wear-resistant foundation of the material. The sample of Comparative Example 1 (without dynamic reversible cross-linked network) may show a higher initial modulus or hardness in some cases due to the lack of an energy dissipation mechanism of a dynamic bond, its elongation at break and tear strength are significantly lower than those of the example sample containing a dynamic network, which highlights the importance of a dynamic reversible cross-linked network for improving material toughness and resistance to damage.

[0114] The performance of Comparative Example 3 (without bifunctional interface anchoring molecules) is between Comparative Example 2 and Example 1. Although its tensile strength and tear strength are better than Comparative Example 2 without ceramic particles, it is still inferior to Example 1 with interface anchoring, which clearly shows the key role played by bifunctional interface anchoring molecules in strengthening the interface bonding between the inorganic reinforcing phase and the organic matrix and promoting the effective transfer of stress. In the absence of such effective interface anchoring, even if there are in-situ generated ceramic particles and dynamic networks, the synergistic reinforcement effect will be greatly reduced. The data of Example 2 (lower ceramic content) and Example 3 (different interface anchoring molecules) further illustrate that by regulating the content of ceramic nanoparticles or the type of interface anchoring molecules, the specific properties of the material can be adjusted on the basis of maintaining the core technical characteristics of the present invention to adapt to different application requirements, but the core reinforcement and toughening mechanism is consistent.

[0115] Test Example 2: Wear resistance test Experimental description: This test example is intended to evaluate the wear resistance of the polyurethane elastomer materials prepared according to the aforementioned examples and comparative examples. Wear resistance is a key indicator of sports shoe outsole materials and is directly related to the service life of the outsole.

[0116] The experimental steps are as follows: 1. Refer to GB / T 9867-2008 “Rubber, vulcanized or thermoplastic — Determination of wear resistance — Rotating roller abrader method”.

[0117] 2. From the sheet samples prepared in each Example and Comparative Example, use a dedicated sampler to prepare cylindrical specimens with a diameter of 16.0 ± 0.2 mm and a thickness of 6.0-10.0 mm. Ensure that the upper and lower surfaces of the specimens are parallel and the edges are smooth. Prepare at least four valid specimens for each sample group.

[0118] 3. Before testing, place all samples in an environment with a temperature of 23±2℃ and a relative humidity of 50±5% for at least 24 hours.

[0119] 4. Use a rotating roller abrasion tester. Use No. 60 standard abrasive paper and calibrate or replace it as required.

[0120] 5. Accurately weigh the initial mass (m0) of each specimen before wear to the nearest 0.1 mg.

[0121] 6. Install the specimen in the specimen fixture of the abrasion machine and apply a normal load of 10.0 ± 0.2N.

[0122] 7. Set the wear stroke to 40.0 ± 0.4 m (or the equivalent number of wear roller revolutions).

[0123] 8. Start the abrasion testing machine and allow the specimen to wear on the rotating abrasive sandpaper.

[0124] 9. After the abrasion is completed, remove the specimen from the specimen holder and use a soft brush to remove the grinding debris on the specimen surface.

[0125] 10. Accurately weigh the mass of the abraded specimen (m1) to the nearest 0.1 mg.

[0126] 11. Calculate the mass wear loss of each specimen: Δm = m0 - m1 (mg).

[0127] 12. For each group of samples, after discarding outliers, take the arithmetic mean of at least 3 valid data as the final mass wear value.

[0128] The experimental data are shown in Table 2: Table 2: DIN abrasion test results of various examples and comparative examples Sample number Mass wear loss (mg) Example 1 45.3 Example 2 68.7 Example 3 48.1 Comparative Example 1 85.9 Comparative Example 2 152.4 Comparative Example 3 95.6 Comparative Example 4 135.1 The wear test results shown in Table 2 clearly reveal the differences in wear resistance of different materials. The samples of Examples 1, 2 and 3 all showed wear resistance significantly better than that of the comparative examples, among which the mass wear loss of Examples 1 and 3 was the lowest, showing the best wear resistance. The acquisition of this excellent performance is the result of the synergistic effect of multiple key technical features in the material structure of the present invention. First, by uniformly generating high-hardness ceramic nanoparticles (such as SiO2) in situ in the polyurethane matrix, a solid anti-wear foundation is provided for the material; these nanoparticles can effectively resist the cutting and scratching of abrasives, reducing material loss. Secondly, the bifunctional interface anchoring molecules (such as F-APTMOS or M-PTES) construct a strong interface bond between the ceramic nanoparticles and the polyurethane matrix, ensuring that the nanoparticles are not easily detached from the matrix during the wear process, thereby being able to continuously exert their wear-resistant effect and enabling stress to be effectively transferred from the matrix to the high-strength nanoparticles, avoiding premature damage to the matrix.

[0129] Further analysis shows that the mass wear loss of Comparative Example 2 (without in-situ generated ceramic nanoparticles) is the highest, which directly proves the decisive role of in-situ generated ceramic nanoparticles in improving the wear resistance of the material. Without these hard fillers, the simple polyurethane matrix (even if it contains a dynamic network) will be rapidly lost under the action of wear. The wear resistance of Comparative Example 4 (conventional TPU) is also poor, far less than that of the embodiments of the present invention. Although the wear resistance of Comparative Example 3 (without bifunctional interface anchoring molecules) is better than that of Comparative Example 2 without ceramic particles, it is worse than that of Example 1. This shows that even if there are in-situ generated ceramic particles, if there is a lack of effective interface anchoring, the binding force between the nanoparticles and the matrix is ​​insufficient, and the interface is easily damaged during the wear process, resulting in premature failure or shedding of the nanoparticles, thereby weakening their wear resistance contribution. This once again highlights the importance of molecular-level interface design for achieving excellent wear resistance.

[0130] The wear resistance of Comparative Example 1 (without a dynamically reversible cross-linked network) is also inferior to that of Example 1. Although the main contribution of the dynamic network is to improve toughness and fatigue resistance, it also has a positive effect on wear resistance to a certain extent. The presence of a dynamic network can enable the material to absorb and dissipate part of the energy through the dynamic reorganization of bonds when it is subjected to microscopic impact and cutting by abrasive particles, reducing the brittle fracture and peeling of the material, thereby indirectly improving the wear resistance and durability. Although Example 2 has better wear resistance than all comparative examples due to its low content of ceramic nanoparticles, it is slightly inferior to Example 1 and Example 3 with a higher ceramic content. This is also in line with the general rule that the reinforcing phase content affects the wear resistance of composite materials. Taken together, it is the organic combination of in-situ generated ceramic nanoparticles, strong interface anchoring and dynamically reversible cross-linked networks that together give the material of the present invention excellent wear resistance.

[0131] Test Example 3: Dynamic Mechanical Analysis Experimental description: This test example aims to evaluate the changes in the viscoelastic properties of the polyurethane elastomer materials prepared according to the aforementioned embodiments and comparative examples as a function of temperature through dynamic mechanical analysis, particularly the changes in the storage modulus (E'), loss modulus (E"), and loss factor (tan δ). These parameters help to understand the glass transition behavior, dynamic network properties, and the effects of different components on the dynamic response of the material.

[0132] The experimental steps are as follows: 1. Use a dynamic mechanical analyzer.

[0133] 2. Cut the sheet samples prepared in each embodiment and comparative example into long strips with typical dimensions of 30±2 mm in length, 5±0.5 mm in width, and 2±0.2 mm in thickness. Ensure that the sample dimensions are regular and the surface is smooth.

[0134] 3. Clamp the specimen in the DMA's tension or three-point bending mode fixture (tensile mode is used in this test example) and apply appropriate preload to ensure that the specimen is always in tension during the test.

[0135] 4. Set test parameters: Temperature scanning range: -80℃ to 180℃.

[0136] Heating rate: 3℃ / min.

[0137] Test frequency: 1Hz.

[0138] Strain amplitude: 0.1% (ensure within the linear viscoelastic region).

[0139] 5. Test under nitrogen atmosphere to prevent high temperature oxidation.

[0140] 6. Record the storage modulus (E'), loss modulus (E") and loss factor (tan δ = E" / E') as a function of temperature.

[0141] 7. Determine the primary glass transition temperature (Tg) from the tanδ-temperature curve, typically taking the temperature corresponding to the tanδ peak. Analyze changes in tanδ at higher temperatures (100-150°C) related to dynamic reversible chemical bond behavior.

[0142] The experimental data are shown in Table 3: Table 3: Main parameters of dynamic mechanical analysis of various examples and comparative examples Sample number E'at 25℃(MPa) Tg (tanδ peak temperature, °C) tanδat 120℃ Example 1 850 -33.5 0.185 Example 2 580 -35.2 0.17 Example 3 825 -34.1 0.192 Comparative Example 1 910 -31.8 0.065 Comparative Example 2 250 -38.5 0.165 Comparative Example 3 650 -32.9 0.155 Comparative Example 4 310 -40.1 0.042 Note: Tg refers primarily to the glass transition of the polyurethane soft segment. Tanδat 120°C is used to characterize the energy dissipation capacity associated with dynamic bonds at high temperatures.

[0143] The dynamic mechanical analysis data shown in Table 3 provide important information for understanding the relationship between the internal structure of the material and the dynamic response. The samples of Examples 1, 2 and 3 all exhibited relatively high storage modulus (E') at room temperature (25°C), especially Examples 1 and 3, which is mainly due to the effective reinforcement effect of the in situ generated ceramic nanoparticles. The uniform dispersion of these hard nanoparticles in the polyurethane matrix limits the movement of the polymer chain segments, thereby increasing the stiffness of the material. The E' value of Comparative Example 2 (without in situ generated ceramic nanoparticles) is significantly lower than that of all the examples, clearly demonstrating the contribution of ceramic nanoparticles to the storage modulus. The presence of bifunctional interface anchoring molecules ensures effective stress transfer between the nanoparticles and the matrix, allowing the reinforcement effect of the nanoparticles to be fully exerted, which can be confirmed by the higher E' values ​​of Examples 1 and 3 compared to Comparative Example 3 (without bifunctional interface anchoring molecules).

[0144] The glass transition temperatures (Tg) of all samples, characterized by the tan δ peak temperature, are all in the range of -30°C to -40°C, which mainly corresponds to the segmental motion of the polyurethane soft segment. The specific value of Tg is affected by a combination of factors such as the crosslinking density of the material, the hard segment content, and the interaction between the filler and the matrix. The slight difference in Tg between the samples reflects the difference in their microstructures. More critical is the tan δ value in the higher temperature region (such as 120°C). Examples 1, 2, 3 and Comparative Examples 2 and 3 containing dynamic bond structures all show relatively high tan δ values ​​at 120°C. This is attributed to the dynamic reversible chemical bonds (such as bonds formed by Diels-Alder reactions) present in their structures. At higher temperatures, these dynamic bonds begin to partially reversibly break and reorganize, a process accompanied by significant energy dissipation, resulting in an increase in the tan δ value. This shows that the material has stronger damping performance and energy absorption capacity in this temperature range, which is crucial for improving the toughness and fatigue resistance of the material.

[0145] In contrast, Comparative Examples 1 and 4, which do not contain a dynamically reversible cross-linked network, have significantly lower tanδ values ​​at 120°C, indicating that they mainly behave like conventional cross-linked elastomers or thermoplastic elastomers at this temperature, lacking an efficient energy dissipation mechanism due to dynamic bond activity. Although Comparative Example 1 also has a certain storage modulus due to its chemically cross-linked network, its energy dissipation capacity at high temperatures is far inferior to that of the embodiment containing a dynamic network. These results strongly support that in the material described in the present invention, the dynamic network constructed by introducing dynamic reversible chemical bonds can dissipate energy through reversible bond exchange at higher temperatures or under stress conditions, thereby giving the material unique dynamic response characteristics and potential damage self-healing or energy absorption capabilities. In situ generated nanoparticles and interface anchoring technology provide a solid structural foundation for the stable existence and efficient operation of these dynamic networks.

[0146] Test Example 4: Macro damage and performance recovery evaluation after heat treatment Experimental description: This test case aims to preliminarily evaluate the ability of some polyurethane elastomer materials containing dynamically reversible chemical bond networks to achieve a certain degree of performance recovery through heat treatment after suffering macroscopic damage.

[0147] The experimental steps are as follows: 1. Samples of Example 1, Comparative Example 1 and Comparative Example 4 were selected for testing.

[0148] 2. Cut the sample into standard dumbbell-shaped specimens. The size and cutting method are the same as those in Test Example 1 (Tensile Properties Test).

[0149] 3. Initial performance test: For three undamaged specimens in each group of samples, test their initial tensile strength according to the method of Test Example 1(a), and record it as σ0.

[0150] 4. Man-made damage introduced: Take at least 3 additional dumbbell-shaped specimens from each set of samples.

[0151] In the center of the narrow parallel portion of each specimen, use a sharp, standardized blade (e.g., a surgical blade) and carefully create a non-penetrating scratch / incision perpendicular to the specimen surface with a length of approximately 2.0 ± 0.2 mm and a depth of approximately 0.5 ± 0.1 mm (approximately 1 / 4 to 1 / 3 of the specimen thickness). Ensure that the degree of damage is as consistent as possible for all specimens.

[0152] 5. Post-damage performance test: For the specimens after the introduction of damage (3 in each group), the tensile strength is immediately tested according to the method of Test Example 1(a), and is recorded as σ1.

[0153] 6. Heat treatment and recovery: The remaining specimens with induced damage (at least 3 per group) were carefully placed in the oven.

[0154] For the sample of Example 1, the oven temperature was set at 120±2° C. and the heat treatment was performed for 2 hours.

[0155] For the samples of Comparative Example 1 and Comparative Example 4, since they do not have the expected thermally induced dynamic bond repair mechanism, they are also heat-treated at 120±2° C. for 2 hours for comparison.

[0156] 7. Performance test after recovery: After the heat treatment, the sample was taken out of the oven and allowed to cool naturally at room temperature (23 ± 2 °C) for at least 4 hours.

[0157] For the heat-treated sample, the tensile strength is measured according to the method of Test Example 1(a), and is recorded as σ2.

[0158] 8. Performance recovery efficiency calculation: Calculate the performance recovery efficiency (η): η = [(σ2 - σ1) / (σ0 - σ1)] × 100%. If σ2 ≥ σ0, η is calculated as 100%. If σ1 ≥ σ0, this evaluation method is not applicable.

[0159] The experimental data are shown in Table 4: Table 4: Tensile strength recovery of some samples after damage and heat treatment The experimental results in Table 4 preliminarily demonstrate the potential of different materials to achieve performance recovery through heat treatment after suffering macroscopic damage. After the introduction of standard scratches, the tensile strength of the sample in Example 1 was significantly reduced. After heat treatment at 120°C, its tensile strength was significantly restored, and the recovery efficiency reached about 71.2%. This phenomenon strongly demonstrates the unique behavior of the dynamic reversible chemical bond network constructed inside the material (in this case, a system based on the Diels-Alder reaction) under thermal stimulation. When heat is applied, the dynamic reversible bonds (such as the adduct of furan and maleimide) undergo partial reverse reaction dissociation, so that the polymer chain segments gain greater mobility in the damaged area, promote the re-contact of the crack interface and the mutual diffusion of the segments; subsequently, during the heat treatment process or during cooling, the positive Diels-Alder reaction occurs again, and the covalent bond cross-linking network is reformed at the original crack interface, thereby restoring the mechanical properties of the damaged area.

[0160] In contrast, after similar damage and the same heat treatment, Comparative Example 1 (conventional chemically cross-linked polyurethane without a dynamically reversible cross-linking network) and Comparative Example 4 (conventional thermoplastic polyurethane) showed minimal recovery in tensile strength, with performance recovery efficiencies far below 5%. Comparative Example 1, with its internal permanent chemically cross-linked network, cannot be reorganized or repaired by simple heat treatment once covalent bonds break, forming macrocracks. As a conventional thermoplastic material, Comparative Example 4 may soften and flow to a certain degree during heat treatment, but lacks a specific reversible covalent bonding mechanism capable of repairing fractured interfaces at the molecular level. Therefore, even if the material undergoes morphological changes under heat, it is difficult to effectively restore the mechanical properties lost due to the cracks.

[0161] Therefore, the results of this test example clearly show that the special network structure formed by introducing dynamic reversible chemical bonds in the material system described in the present invention is the key to achieving performance recovery after macroscopic damage. This ability does not come from the simple thermoplastic flow of the material, but is based on the reversible breaking and recombination of specific chemical bonds under heat-driven conditions, which gives the material a certain degree of "self-repair" or "damage repair" potential. Combined with the excellent basic mechanical properties brought by the in-situ generation of ceramic nanoparticles and effective interface anchoring of the material itself, the introduction of this dynamic network further expands the functionality of the material, providing it with additional advantages in applications requiring high durability and long life.

[0162] While embodiments of the present invention have been shown and described, it will be appreciated by those skilled in the art that various changes, modifications, substitutions, and variations may be made to these embodiments without departing from the principles and spirit of the invention, and that the scope of the invention is defined by the appended claims and their equivalents.

Claims

1. A wear-resistant sports shoe outsole, characterized in that: The outsole comprises a polyurethane elastomer matrix, and the polyurethane elastomer matrix comprises: In-situ generated ceramic nanoparticles, wherein the content of the ceramic nanoparticles is 2 to 15 wt% of the total mass of the final outsole; cross-linked networks formed by dynamically reversible chemical bonds; a bifunctional interface anchoring molecule, one end of which interacts with the surface of the in situ generated ceramic nanoparticles and the other end of which participates in the network structure of the polyurethane elastomer matrix or the cross-linked network formed by the dynamic reversible chemical bonds; The polyurethane elastomer is prepared by reacting raw materials including polyol, diisocyanate, and monomers introduced with dynamic reversible chemical bond functional groups.

2. The wear-resistant sports shoe outsole according to claim 1, characterized in that: The in-situ generated ceramic nanoparticles are selected from at least one of silicon dioxide nanoparticles or titanium dioxide nanoparticles.

3. The wear-resistant sports shoe outsole according to claim 1, characterized in that: The cross-linked network formed by the dynamic reversible chemical bonds is formed through a Diels-Alder reaction.

4. The wear-resistant sports shoe outsole according to claim 1, characterized in that: The bifunctional interface anchoring molecule comprises a ceramic affinity / reaction end and a dynamic network participation end, wherein the ceramic affinity / reaction end is selected from at least one of an alkoxysilane group, a phosphonic acid group or a carboxylic acid group, and the dynamic network participation end is selected from at least one of a furan group, a maleimide group or an active group capable of forming a covalent bond with a polyurethane segment or a dynamic cross-linked network.

5. A method for preparing the wear-resistant sports shoe outsole according to any one of claims 1 to 4, characterized in that: The following steps are involved: (S1) mixing and reacting a polyol, a diisocyanate, a monomer introducing a dynamically reversible chemical bond functional group, a ceramic nanoparticle precursor, and a bifunctional interface anchoring molecule to prepare a polyurethane prepolymer containing the ceramic precursor and the bifunctional interface anchoring molecule; (S2) mixing the polyurethane prepolymer with a chain extender, and initiating in-situ hydrolysis and condensation of the ceramic nanoparticle precursor to generate ceramic nanoparticles, while simultaneously forming a polyurethane elastomer matrix and a dynamic reversible chemical bond cross-linked network, wherein one end of the bifunctional interface anchoring molecule interacts with the surface of the in-situ generated ceramic nanoparticles, and the other end participates in the network structure of the polyurethane elastomer matrix or the cross-linked network formed by the dynamic reversible chemical bond; (S3) The mixture obtained in step (S2) is molded and solidified.

6. The preparation method according to claim 5, characterized in that In step (S1), the monomer for introducing the dynamic reversible chemical bond functional group comprises a diol with a furan group, and / or the bifunctional interface anchoring molecule comprises a furan group or a maleimide group.

7. The preparation method according to claim 5, characterized in that In step (S1), the ceramic nanoparticle precursor is selected from metal alkoxides.

8. The preparation method according to claim 5, characterized in that The chain extender in step (S2) contains a reactive group that matches the dynamic network participating end of the monomer of the dynamic reversible chemical bond functional group or the bifunctional interface anchoring molecule to form a dynamic reversible chemical bond cross-linked network through a Diels-Alder reaction.

9. The preparation method according to claim 5, characterized in that In step (S1), the amount of the bifunctional interface anchoring molecule is 1 to 10 mol% of the molar amount of the ceramic nanoparticle precursor.

10. The preparation method according to claim 5, characterized in that In step (S2), in-situ hydrolysis and condensation of the ceramic nanoparticle precursor is initiated by adding a catalytic amount of water, and the molar ratio of water to the molar ratio of the ceramic nanoparticle precursor is 2:1 to 4:1.

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