High-anti-cracking silicon elastomer based on multi-scale stress dispersion mechanism and preparation method thereof

Abstract

The application belongs to the field of polymer composites, and provides a high-anti-cracking silicone elastomer based on a multi-scale stress dispersion mechanism and a preparation method thereof, which comprises a polydimethylsiloxane crosslinking network and carboxylated multi-walled carbon nanotubes dispersed in the polydimethylsiloxane crosslinking network; the polydimethylsiloxane crosslinking network forms a long-chain entanglement structure under the catalysis of platinum from Si-H functionalized polydimethylsiloxane and vinyl functionalized polydimethylsiloxane; the carboxylated multi-walled carbon nanotubes are connected to the polydimethylsiloxane crosslinking network through a covalent bond to form a particle cluster structure; the long-chain entanglement structure and the particle cluster structure jointly constitute a multi-scale stress dispersion structure of polymer chain scale and particle cluster scale synergy, so that crack tip stress is gradually dispersed among different scales. The application improves the fatigue threshold, fracture toughness and anti-cracking performance of the material.

Description

Technical Field

[0001] This invention belongs to the field of polymer composite materials, specifically relating to a high crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism and its preparation method. Background Technology

[0002] The statements in this section are merely background information related to the present invention and do not necessarily constitute prior art.

[0003] Silicon elastomers are widely used in flexible electronics, soft robots, and biomedical devices due to their softness, high elongation, and biocompatibility. However, defects and cracks inevitably occur during their long-term service. Traditional silicon elastomers often experience rapid and catastrophic crack propagation due to the high stress concentration at the crack tip, which severely limits the lifespan of devices. Summary of the Invention

[0004] To address the aforementioned problems, this invention proposes a high-crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism and its preparation method. This invention establishes a strong interfacial adhesion between highly entangled long polymer chains and rigid particles, enabling the stress at the crack tip to be dispersed stepwise at both the polymer chain and particle cluster scales, thereby significantly improving the material's crack resistance, fatigue life, and fracture toughness.

[0005] According to some embodiments, the present invention adopts the following technical solution: A highly crack-resistant silicone elastomer based on a multi-scale stress dispersion mechanism includes a polydimethylsiloxane crosslinking network and carboxylated multi-walled carbon nanotubes dispersed in the polydimethylsiloxane crosslinking network. The polydimethylsiloxane crosslinking network is formed by the formation of long-chain entangled structures of Si-H functionalized polydimethylsiloxane and vinyl functionalized polydimethylsiloxane under platinum catalysis. The carboxylated multi-walled carbon nanotubes are covalently linked to the polydimethylsiloxane crosslinking network to form a particle cluster structure. The long-chain entanglement structure and the particle cluster structure together constitute a multi-scale stress dispersion structure with the synergistic effect of polymer chain scale and particle cluster scale, so that the stress at the crack tip is dispersed step by step between different scales.

[0006] In the present invention, by forming a multi-scale stress dispersion structure, the stress at the crack tip is dispersed or dissipated by the synergistic effect of long-chain polymer entanglement, covalent connection between polymer and particles, and rigid particle aggregation, thereby suppressing crack propagation.

[0007] As an alternative implementation, the carboxyl groups on the surface of the carboxylated multi-walled carbon nanotubes react with the Si-H groups in polydimethylsiloxane to form a silane bond connection structure, wherein the covalent bond includes a Si-OC=O structure.

[0008] As an alternative implementation, the functional group molar ratio R1 of the Si-H functionalized polydimethylsiloxane to the vinyl functionalized polydimethylsiloxane is 1.10 to 1.40. Alternatively, the functional group molar ratio R1 of the Si-H functionalized polydimethylsiloxane to the vinyl functionalized polydimethylsiloxane is 1.20.

[0009] As an alternative implementation, the volume fraction F of the carboxylated multi-walled carbon nanotubes in the composite material is 0.03 to 0.10; Alternatively, the volume fraction F of the carboxylated multi-walled carbon nanotubes in the composite material is 0.06.

[0010] A method for preparing a highly crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism includes the following steps: Si-H functionalized polydimethylsiloxane and vinyl functionalized polydimethylsiloxane were mixed according to a preset functional group molar ratio and hydrosilylated under platinum catalysis to form a long-chain polydimethylsiloxane precursor. Carboxylated multi-walled carbon nanotubes were mixed with Si-H functionalized polydimethylsiloxane to form a covalent bond between the carboxylated multi-walled carbon nanotubes and the polydimethylsiloxane. The two parts are mixed and then water is added, and a cross-linking reaction is carried out under the action of a platinum catalyst. The resulting mixture was poured into a mold, and after degassing and thermosetting, the high crack-resistant silicone elastomer was obtained.

[0011] As an alternative implementation, the functional group molar ratio R1 of the Si-H functionalized polydimethylsiloxane to the vinyl functionalized polydimethylsiloxane is 1.10 to 1.40. Alternatively, the functional group molar ratio R1 of the Si-H functionalized polydimethylsiloxane to the vinyl functionalized polydimethylsiloxane is 1.20.

[0012] As an alternative implementation, the volume fraction F of the carboxylated multi-walled carbon nanotubes in the composite material is 0.03 to 0.10; Alternatively, the volume fraction F of the carboxylated multi-walled carbon nanotubes in the composite material is 0.06.

[0013] As an alternative implementation, the molar ratio R2 of the carboxylated multi-walled carbon nanotubes to Si-H groups is 1.05, and the molar ratio of water to the total Si-H functional groups in the system is 4:1.

[0014] As an alternative implementation, the platinum catalyst is 1,3-divinyl-1,1,3,3-tetramethyldisiloxane platinum (0), and its amount is 1% of the total mass of polydimethylsiloxane.

[0015] Application of a high crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism in flexible electronic devices, flexible strain sensors, soft robots, and biomedical devices.

[0016] Compared with the prior art, the beneficial effects of the present invention are as follows: This invention innovatively constructs a particle-reinforced silicon elastomer with both high crack resistance and high tensile strength based on a multi-scale stress dispersion mechanism. By establishing a strong interfacial adhesion between highly entangled long polymer chains and rigid particles, the stress at the crack tip can be dispersed stepwise at both the polymer chain scale and the particle cluster scale, thereby significantly improving the material's crack resistance, fatigue life, and fracture toughness.

[0017] To make the above-mentioned objects, features and advantages of the present invention more apparent and understandable, preferred embodiments are described below in detail with reference to the accompanying drawings. Attached Figure Description

[0018] The accompanying drawings, which form part of this invention, are used to provide a further understanding of the invention. The illustrative embodiments of the invention and their descriptions are used to explain the invention and do not constitute an improper limitation of the invention.

[0019] Figure 1 This is a schematic diagram of the mechanism by which multi-scale stress dispersion improves the fatigue threshold of composite materials in one embodiment, wherein (a) represents the stress dispersion mechanism of highly entangled long polymer chains, and (b) represents the stress dispersion mechanism of rigid particle aggregation. Figure 2 A schematic diagram illustrating a one-pot reaction for preparing highly entangled long-chain PDMS according to one embodiment; Figure 3 This is a schematic diagram illustrating the covalent bond connection between carboxylated MWCNTs and long-chain PDMS in one embodiment. Figure 4 This is a schematic diagram of Fourier transform infrared spectral difference analysis according to one embodiment; Figure 5 This is a schematic diagram of an optical image of a Teflon mold according to one embodiment; Figure 6 A schematic optical image of a shear geometry stretching component according to one embodiment; Figure 7 This is a schematic diagram illustrating the cyclic stretching of a rectangular composite film using two rigid clamps in one embodiment. Figure 8 This is a schematic diagram of cyclic tension of the composite material when R=1.20 and F=0.06 in one embodiment. (a) is the stress-tension curve after 5000 cycles of tension, (b) is the stress-tension curve under different tension amplitudes, and (c) is the energy release rate under steady state as a function of applied tension. Figure 9 This is a schematic diagram illustrating the cyclic stretching of a rectangular composite film with a crack of length c using two rigid clamps, as one embodiment. Figure 10 This is a schematic diagram of cyclic stretching of different doped particles with the same R and CNTs content in one embodiment, where (a)-(i) are the cyclic stretching parameters of different doped particles, respectively. Figure 11 This is a schematic diagram of the dc / dN-G curves of pure PDMS without rigid particles and CNT / PDMS composite material with and without covalent bonds in one embodiment. Figure 12 This is a schematic diagram of cyclic stretching of different R values ​​under the same CNT content in one embodiment, where (a)-(l) are the cyclic stretching parameters of different R values, respectively. Figure 13 This is a schematic diagram of the dc / dN-G curves of CNT / PDMS composite materials with different R values ​​at the same CNT content in one embodiment. Figure 14 This is a schematic diagram of cyclic stretching with different R values ​​under the same CNT content in one embodiment. (a) is the stress-tension curve after 5000 cycles of stretching, (b) is the stress-tension curve under different stretching amplitudes, and (c) is the energy release rate under steady state as a function of applied stretching. Figure 15 This is a schematic diagram of the dc / dN-G curves of CNT / PDMS composite materials with different R values ​​at the same CNT content in one embodiment. Figure 16 The image shows a scanning electron microscope (SEM) image of the surface of a CNT / PDMS composite material with different F and fixed R according to one embodiment, where the main image scale bar is 50 μm and the inset is a magnified view with a scale bar of 10 μm. Figure 17 This is a schematic diagram illustrating the relationship between the fatigue threshold and tensile strain of different materials in one embodiment. Figure 18This is a schematic diagram illustrating how the doping of rigid particles can increase the fracture toughness of a composite material in one embodiment. In this diagram, (a) is a graph showing the tensile ratio and stress of composite materials with different rigid particle doping, (b) is a graph showing the tensile ratio and strain energy of composite materials with different rigid particle doping, and (c) is a graph showing the crack propagation and energy release rate of composite materials with different rigid particle doping. Figure 19 This is a monotonic stretching process of different materials with pre-cracks in one embodiment, where (a)-(c) are schematic diagrams of different elongation at break. Detailed Implementation

[0020] The present invention will be further described below with reference to the accompanying drawings and embodiments.

[0021] It should be noted that the following detailed description is illustrative and intended to provide further explanation of the invention. Unless otherwise specified, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this invention pertains.

[0022] It should be noted that the terminology used herein is for the purpose of describing particular embodiments only and is not intended to limit the scope of exemplary embodiments according to the invention. As used herein, the singular form is intended to include the plural form as well, unless the context clearly indicates otherwise. Furthermore, it should be understood that when the terms "comprising" and / or "including" are used in this specification, they indicate the presence of features, steps, operations, devices, components, and / or combinations thereof.

[0023] Where there is no conflict, the embodiments and features described in this application may be combined with each other.

[0024] Example 1 A highly crack-resistant silicone elastomer based on a multi-scale stress dispersion mechanism includes a polydimethylsiloxane crosslinking network and carboxylated multi-walled carbon nanotubes dispersed in the polydimethylsiloxane crosslinking network. The polydimethylsiloxane crosslinking network is formed by the formation of long-chain entangled structures of Si-H functionalized polydimethylsiloxane and vinyl functionalized polydimethylsiloxane under platinum catalysis. The carboxylated multi-walled carbon nanotubes are covalently linked to the polydimethylsiloxane crosslinking network to form a particle cluster structure. The long-chain entanglement structure and the particle cluster structure together constitute a multi-scale stress dispersion structure with the synergistic effect of polymer chain scale and particle cluster scale, so that the stress at the crack tip is dispersed step by step between different scales.

[0025] The mechanism by which stress dispersion leads to the high crack resistance of silicon elastomers can be achieved through the following three aspects: (1) strong covalent bonding between polymer and particles; (2) highly entangled long-chain polymer networks; and (3) particle aggregation or network structure. Specifically, when a crack acts on a polymer chain, the breaking of a single bond can dissipate the energy of the entire chain; when a crack acts on a particle cluster, due to the high rigidity of the particles, the stress is dispersed in multiple particle gaps, and the failure of a single gap can release the energy stored in multiple gaps, thereby achieving cross-scale energy dissipation. Figure 1 Therefore, stress dispersion occurs simultaneously at both the "polymer chain scale" and the "particle cluster scale," thereby significantly improving the material's crack resistance. Figure 1 (b).

[0026] Example 2 A method for preparing a highly crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism includes: Si-H functionalized polydimethylsiloxane and vinyl functionalized polydimethylsiloxane were mixed according to a preset functional group molar ratio and hydrosilylated under platinum catalysis to form a long-chain polydimethylsiloxane precursor. Carboxylated multi-walled carbon nanotubes were mixed with Si-H functionalized polydimethylsiloxane to form a covalent bond between the carboxylated multi-walled carbon nanotubes and the polydimethylsiloxane. The two parts are mixed and then water is added, and a cross-linking reaction is carried out under the action of a platinum catalyst. The resulting mixture was poured into a mold, and after degassing and thermosetting, the high crack-resistant silicone elastomer was obtained.

[0027] This invention employs a platinum-catalyzed one-pot reaction to prepare a composite material with a polydimethylsiloxane (PDMS) crosslinked network as the matrix and carboxylated multi-walled carbon nanotubes (COOH-MWCNTs) as the reinforcing phase. The specific characteristics of the raw materials used are as follows: telechelicer type Si-H functionalized PDMS with a number-average molecular weight (Mn) of 17500 and a density of 0.971 g·mL⁻¹. -1 The molecule contains one Si-H functional group at each end; it is a telechelic vinyl-functionalized PDMS with a Mn of 24000 and a density of 0.965 g·mL. -1 The molecule contains a vinyl functional group at each end; the carboxyl content of COOH-MWCNTs is 3.86 wt%, and its apparent bulk density is 0.22 g·cm³. -3 .

[0028] The composite material is constructed based on a multi-reaction synergistic process within a platinum catalytic system, the core of which includes hydrosilylation and subsequent crosslinking reactions involving Si-H groups. According to literature reports, in a one-pot reaction, hydrosilylation between Si-H groups and vinyl groups, as well as water- and oxygen-mediated Si-H crosslinking, occur simultaneously. Since hydrosilylation has a significant kinetic advantage over Si-H crosslinking, it is expected to reach high conversion rates before any significant crosslinking occurs. Assuming these two reactions occur strictly sequentially, the hydrosilylation first generates elongated chains, which are then crosslinked into an elastomer through the reaction of excess Si-H with oxygen and water. In this case, the average molar mass of the network segments can be expressed by the following formula:

[0029] Among them, M DMS-H With M DMS-V Table 1 shows the molecular weights of Si-H functionalized PDMS and vinyl functionalized PDMS, respectively, where R is the molar ratio of Si-H to vinyl functional groups. The molecular weights of PDMS polymerized under different R values ​​calculated by this formula are shown in Table 1.

[0030] Table 1. Molecular weight of PDMS polymerized under different functional group ratios R

[0031] Specifically, the preparation of the composite material consists of two pretreatment parts and a subsequent overall mixing and curing process. First, telechelicerary Si-H functionalized PDMS and telechelicerary vinyl functionalized PDMS are mixed according to a functional group molar ratio R1 (R1 = 1.10, 1.20, 1.30, 1.40). The design of R1 > 1 ensures the presence of excess Si-H groups in the system. These residual Si-H groups not only participate in the subsequent crosslinking reaction but also serve as important reaction sites for interfacial reactions with COOH-MWCNTs. Under the action of a platinum catalyst, hydrosilylation preferentially occurs in the system, i.e., Si-H undergoes an addition reaction with vinyl groups to form Si-CH2-CH2-Si bonds, thereby extending the molecular chain. Due to the significant kinetic advantages of this reaction, it is usually completed rapidly in the early stages of the reaction, thus forming a long-chain or highly entangled network precursor (…). Figure 2 ).

[0032] Secondly, telechelicerable Si-H functionalized PDMS was mixed with COOH-MWCNTs, and the molar ratio R2 of Si-H to carboxyl groups was set to 1.05 to ensure that the carboxyl functional groups could fully participate in the reaction. Under platinum catalysis, COOH-MWCNTs could directly react with Si-H to form a silyl ester structure ( Figure 3Meanwhile, during the reaction, some Si-H undergoes hydrolysis and oxidation reactions with the participation of water and oxygen to generate Si-OH, which further forms a Si-O-Si cross-linked structure through a condensation reaction. Figure 2 Based on this, the generated Si-OH can further undergo a condensation reaction with the carboxyl groups on the CNT surface to form stable silane ester bonds. Figure 3 Therefore, in this system, CNTs are not only introduced as conductive fillers, but also participate in the PDMS network through Si-OC=O covalent bonds, thereby achieving interface strengthening and effective stress transfer.

[0033] To verify the occurrence of the above interfacial reaction, Si-H PDMS and COOH-MWCNTs were mixed separately at R² = 1.05, and Pt catalyst was added and reacted at 100 °C for 24 h. Subsequently, Fourier transform infrared (FTIR) spectroscopy was performed on the samples before and after the reaction. After baseline correction and normalization, the spectra were analyzed using OMNIC software using difference spectral analysis (the post-reaction spectrum minus the pre-reaction spectrum). The results showed that at 2171 cm⁻¹... -1 and 2988 cm -1 A distinct negative peak appears at 1092 cm⁻¹, corresponding to the consumption of Si-H bonds and carboxyl OH bonds, respectively; Meanwhile, at 1092 cm⁻¹... -1 A significant positive peak appears at this location, which is attributed to the formation of Si-OC bonds. Figure 4 The above results directly demonstrate that a stable covalent interface is formed between COOH-MWCNTs and PDMS through the aforementioned reaction pathway.

[0034] Subsequently, the two systems were mixed, and deionized water was added to maintain a water-to-total Si-H functional group molar ratio of 4:1. The introduction of water promotes partial Si-H hydrolysis to generate Si-OH, providing an additional reaction pathway for the CNT interface reaction; it also participates in the subsequent construction of the cross-linked network. The mixed system was stirred at 500 rpm for 12 h on a magnetic stirrer to ensure sufficient reaction and uniform dispersion of all components. After mixing, the total volume of PDMS (excluding CNTs) in the system was fixed at 1500 mm². 3 .

[0035] By adjusting the amount of COOH-MWCNTs added, its volume fraction F in the composite material was controlled. This study set four levels for F: 0, 0.03, 0.06, and 0.10. The calculation formula is as follows:

[0036] The required mass of COOH-MWCNTs is calculated based on their bulk density.

[0037] Subsequently, a platinum catalyst (1,3-divinyl-1,1,3,3-tetramethyldisiloxane platinum (0)) was added to the system at 1% of the total mass of PDMS, and the mixture was stirred for 1 min to ensure uniform catalyst distribution. Figure 5 The uniformly mixed slurry was poured into a Teflon mold (cavity size: 50 mm × 60 mm × 1 mm), and after vacuum degassing, it was placed in an oven at 100 ℃ for 24 h to cure. During the curing process, the excess Si-H groups in the system that did not participate in the hydrosilylation reaction gradually crosslinked through slower reactions such as oxidation, hydrolysis, and condensation under the combined action of platinum catalysis, water, and oxygen, ultimately forming a stable three-dimensional network structure, in which the surface-functionalized CNTs were firmly embedded.

[0038] After curing, the composite film was removed from the mold and cut in half lengthwise. Each sample measured 50 mm × 30 mm × 1 mm, and tensile test specimens were prepared. An adhesive primer (Loctite 7701) was applied to both sides of the short side of the sample (the clamping area) and the surfaces of the four glass slides. The sample was allowed to stand for 3 minutes under ventilation to promote solvent evaporation. Then, cyanoacrylate adhesive (Krazy Glue) was used to attach both ends of the sample to the glass slides, ensuring that the effective test area (unbonded area in the middle) was 10 mm in length (gauge length) and 50 mm in width. The tensile test specimen preparation was thus completed. Figure 6 As shown.

[0039] The specific amounts of raw materials used in each formulation, including the volume of telechelic Si-H functionalized PDMS, telechelic vinyl functionalized PDMS, and deionized water, as well as the mass of COOH-MWCNTs and catalyst, are all precisely calculated based on the above molar ratio, volume fraction, and total volume requirements. Detailed data are listed in Table 2.

[0040] Table 2. Distribution ratio of each group in the preparation of composite materials

[0041] Fatigue resistance under cyclic tension To systematically evaluate the fatigue resistance of composite materials under long-term cyclic loading, cyclic tensile tests were first conducted using a pure shear geometry. Specifically, a rectangular thin film sample was held along its two long sides using rigid clamps, and subjected to a constant tensile ratio (λ). amp Applying periodic loading under () Figure 7 Experimental results show that the stress-strain response of the material undergoes a significant evolution in the initial cyclic stage, but gradually stabilizes with increasing cycle number, reaching a steady state after approximately 1000 cycles. Figure 8 (a).

[0042] Under steady-state conditions, further measurements were performed on the same thin film at different λ values. amp Stress-strain curves under ( Figure 8 (b) It can be observed that the loading and unloading paths do not overlap significantly, exhibiting a significant lag behavior; at the same time, as λ amp As the stress increases, the overall stress level shifts upward, and residual strain still exists even at zero stress. This indicates that the material undergoes some irreversible structural evolution during cycling (such as interfacial debonding, chain segment breakage, or microvoid formation), but this structural change tends to stabilize in the steady-state stage, thus ensuring the repeatability of subsequent energy analysis and crack propagation tests. Composite materials with different polymer chain lengths (controlled by the functional group molar ratio R1) and different COOH-MWCNTs volume fractions (F) also exhibit systematic differences in their steady-state mechanical responses. Figure 10 , 12 and 14).

[0043] Based on the above steady-state stress-strain curves, the corresponding values ​​for each λ can be calculated. amp The energy release rate G is given. For pure shear geometry, G equals the product of the effective height of the specimen and the area of ​​the integral of the stress-strain curve of the crack-free specimen. Since the experimental data are discrete points, the quadratic polynomial least squares method is used to calculate G-λ. amp The relationship is fitted to obtain a continuous function expression. Figure 8 (c). This establishes a quantitative mapping relationship between applied strain and steady-state energy release rate G, providing a basis for subsequent crack propagation behavior analysis.

[0044] In the crack propagation experiment, an initial edge crack (c=10 mm) was introduced on the side of the sample, and different λ values ​​were applied. amp The crack propagation length was measured after N=50,000 cyclic loading cycles were applied. c, thus calculating the crack propagation rate dc / dN ( Figure 9 Furthermore, by correlating dc / dN with the corresponding steady-state energy release rate G, the fatigue crack propagation rate curve (dc / dN-G curve) is obtained. Figure 11 The G value corresponding to dc / dN approaching zero in this curve is defined as the fatigue threshold G of the material. th .

[0045] Experimental results show that composite materials without covalent bonds (such as unmodified CNTs systems) have a higher fatigue threshold (G). th The concentrations are similar to those of pure PDMS matrix, at 189.6 mJ·m⁻¹. -2 With 191.9 mJ·m -2 ; and after introducing COOH-MWCNTs and forming covalent interconnects with PDMS, G th Significantly increased to 228.5 mJ·m-2 This result indicates that the presence of rigid particles alone cannot improve the fatigue threshold; effective stress transfer must be achieved through strong polymer-particle interfacial adhesion, allowing the particles to participate in load-bearing and thus inhibiting crack propagation.

[0046] Based on this, to further elucidate the mechanism of action of long polymer chains, the dc / dN-G relationship of materials with different R1 values ​​(1.40, 1.30, 1.20, 1.10) was systematically studied under the condition of fixed F=0.03. Figure 13 According to Formula 1, a decrease in R1 corresponds to a decrease in crosslinking density and an increase in chain length. Experimental results show that as R1 decreases (chain length increases), the fatigue threshold G of the material decreases. th Significantly improved (Table 3). This trend indicates that long-chain polymers can achieve greater stress dispersion at the crack tip, allowing the energy stored throughout the chain to be dissipated simply by breaking a single bond, thereby effectively improving crack resistance.

[0047] Table 3 Fatigue thresholds of different composite materials

[0048] Furthermore, the effect of particle content on fatigue performance was investigated under the condition of R1=1.20. As the volume fraction F of COOH-MWCNTs increased from 0.03 to 0.10, the fatigue threshold of the composite material continuously increased, reaching a maximum of 381.1 mJ·m⁻¹. -2 ( Figure 15 This phenomenon indicates that as the particle content increases, the particles gradually transform from isolated and dispersed to aggregated and even form a continuous network, causing the stress at the crack tip to be dispersed in multiple particle-particle gaps, thereby further improving the energy dissipation capacity. Figure 16 ).

[0049] Finally, through analysis Figure 8 G-λ of the c-type graph amp Relationship, comparing G in different material systems th The corresponding critical strain shows that the highest critical tensile strain, reaching 46.44%, is achieved when R1 = 1.20 and F = 0.06. Figure 17 This indicates that the formulation maintains excellent stretchability while ensuring a high fatigue threshold. Therefore, this parameter combination was selected as the optimal material system in the subsequent monotonic tensile characterization.

[0050] Fracture resistance under monotonic tension To systematically evaluate the crack resistance of composite materials under quasi-static loading, this study employed a pure shear configuration to conduct monotonic tensile fracture tests on different samples. The test samples were rectangular films with a height H = 10 mm and a pre-fabricated edge crack length c = 10 mm. During the test, the crack tip propagation behavior was observed by gradually increasing the tensile ratio, and whether the sample ultimately experienced unstable fracture was recorded.

[0051] To establish a quantitative relationship between crack propagation and fracture energy, we first used 0.01 s-1... -1 The tensile rate was used to monotonically stretch a crack-free specimen, and the stress-strain curve was obtained, such as... Figure 18 As shown in Figure a. Subsequently, integrating this curve yields the relationship between strain energy density W and stretch ratio λ, as shown in Figure a. Figure 18 As shown in b. An initial crack was then prepared on the sample, and the initial position of the crack tip was recorded using a microscope; subsequently, the sample was loaded to a predetermined value below the critical tensile ratio, i.e., λ < λ. c The load was briefly applied to allow the crack to stabilize before unloading. After unloading, another image of the crack tip was taken, and the crack propagation amount under the corresponding tensile condition was obtained by comparing the changes in crack position before and after loading. c. Then, combining this with the relations under the pure shear configuration:

[0052] This tensile state was converted into the corresponding energy release rate. To ensure the accuracy and comparability of the results, each data point was tested using a new pre-cracked specimen, and crack propagation was precisely observed using a microscope with a resolution of approximately 10 μm. Following the above steps, pre-cracks were created in the specimens, and the tensile strain was gradually increased in increments of 10%. The critical tensile ratio λ for unstable crack propagation was then determined. c The fracture toughness of the material is calculated according to Formula 3.

[0053] Based on the above experiments, the crack propagation rate can be further plotted. c-energy release rate G relationship curve ( Figure 18 (c), and based on this, identify two key parameters of crack propagation behavior under monotonic tension: crack initiation and propagation threshold G. i and instability fracture toughness G c .from From the relationship between cG and γ, all samples exhibited typical three-stage fracture behavior. When γ is below the crack initiation and propagation threshold γc... i When G is between G and G, the crack remains stationary; when G is between G and G, the crack remains stationary. i With G c Between these two points, the crack propagates steadily at a rate comparable to the loading rate and stops after unloading; when G reaches or exceeds G... c At that moment, the crack rapidly became unstable and penetrated the sample. This indicates that Gi It describes a material's ability to resist crack initiation and propagation, while G c This characterizes the material's ultimate ability to resist catastrophic fracture.

[0054] Experimental results show that materials with different network structures exhibit significant differences in macroscopic fracture behavior. When COOH-MWCNTs are introduced and covalent bonds are established between the polymer and particles, the composite material does not fracture even at a λ as high as 5.56. Figure 19 The network with the highest crack resistance (a) exhibited the best crack resistance. In contrast, the long-chain PDMS network with a highly entangled structure (F=0, R1=1.20) did not experience crack instability propagation when stretched to λ=1.96. Figure 19 (b) Its elongation at break is λ=4.28. In contrast, conventional PDMS networks experience rapid instability fracture at a tensile ratio of λ≈1.32, such as... Figure 19 As shown in Figure c, this indicates a high stress concentration at the crack tip, making the material extremely sensitive to pre-existing cracks. Thus, it can be seen that from conventional short-chain networks to long-chain entangled networks, and then to multi-scale composite networks with rigid particle reinforcement and interfacial connections, the crack resistance of materials under monotonic tension exhibits a clear trend of progressively increasing strength.

[0055] Further comparison of the fracture toughness of samples with different rigid particle contents reveals that the fracture toughness of the material significantly improves with increasing COOH-MWCNTs volume fraction F: when F=0, G c =96.2 mJ·m -2 When F=0.03, G c =161.4 mJ·m -2 When F=0.06, G c =257.8 mJ·m -2 When F=0.10, G c =456.4 mJ·m -2 ( Figure 18 (c). This result shows that the introduction of rigid particles can significantly improve the fracture toughness of the material under monotonic tensile conditions.

[0056] The above description is merely a preferred embodiment of the present invention and is not intended to limit the invention. Various modifications and variations can be made to the present invention by those skilled in the art. Any modifications, equivalent substitutions, improvements, etc., made by those skilled in the art without creative effort within the spirit and principles of the present invention should be included within the scope of protection of the present invention.

Claims

1. A highly crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism, characterized in that, Includes a polydimethylsiloxane crosslinked network and carboxylated multi-walled carbon nanotubes dispersed in the polydimethylsiloxane crosslinked network; The polydimethylsiloxane crosslinking network is formed by the formation of long-chain entangled structures of Si-H functionalized polydimethylsiloxane and vinyl functionalized polydimethylsiloxane under platinum catalysis. The carboxylated multi-walled carbon nanotubes are covalently linked to the polydimethylsiloxane crosslinking network to form a particle cluster structure. The long-chain entanglement structure and the particle cluster structure together constitute a multi-scale stress dispersion structure with the synergistic effect of polymer chain scale and particle cluster scale, so that the stress at the crack tip is dispersed step by step between different scales.

2. The high crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism as described in claim 1, characterized in that, The carboxyl groups on the surface of the carboxylated multi-walled carbon nanotubes react with the Si-H groups in polydimethylsiloxane to form a silane bond connection structure, and the covalent bond includes a Si-OC=O structure.

3. The high crack resistance silicon elastomer based on a multi-scale stress dispersion mechanism as described in claim 1, characterized in that, The functional group molar ratio R1 of the Si-H functionalized polydimethylsiloxane to the vinyl functionalized polydimethylsiloxane is 1.10 to 1.

40. Alternatively, the functional group molar ratio R1 of the Si-H functionalized polydimethylsiloxane to the vinyl functionalized polydimethylsiloxane is 1.

20.

4. The high crack resistance silicon elastomer based on a multi-scale stress dispersion mechanism as described in claim 1, characterized in that, The volume fraction F of the carboxylated multi-walled carbon nanotubes in the composite material is 0.03–0.

10. Alternatively, the volume fraction F of the carboxylated multi-walled carbon nanotubes in the composite material is 0.

06.

5. A method for preparing a highly crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism, characterized in that, Includes the following steps: Si-H functionalized polydimethylsiloxane and vinyl functionalized polydimethylsiloxane were mixed according to a preset functional group molar ratio and hydrosilylated under platinum catalysis to form a long-chain polydimethylsiloxane precursor. Carboxylated multi-walled carbon nanotubes were mixed with Si-H functionalized polydimethylsiloxane to form a covalent bond between the carboxylated multi-walled carbon nanotubes and the polydimethylsiloxane. The two parts are mixed and then water is added, and a cross-linking reaction is carried out under the action of a platinum catalyst. The resulting mixture was poured into a mold, and after degassing and thermosetting, the high crack-resistant silicone elastomer was obtained.

6. The method for preparing a highly crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism as described in claim 5, characterized in that, The functional group molar ratio R1 of the Si-H functionalized polydimethylsiloxane to the vinyl functionalized polydimethylsiloxane is 1.10 to 1.

40. Alternatively, the functional group molar ratio R1 of the Si-H functionalized polydimethylsiloxane to the vinyl functionalized polydimethylsiloxane is 1.

20.

7. The method for preparing a highly crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism as described in claim 5, characterized in that, The volume fraction F of the carboxylated multi-walled carbon nanotubes in the composite material is 0.03–0.

10. Alternatively, the volume fraction F of the carboxylated multi-walled carbon nanotubes in the composite material is 0.

06.

8. The method for preparing a highly crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism as described in claim 5, characterized in that, The molar ratio R2 of the carboxylated multi-walled carbon nanotubes to Si-H groups is 1.05, and the molar ratio of water to the total Si-H functional groups in the system is 4:

1.

9. The method for preparing a highly crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism as described in claim 5, characterized in that, The platinum catalyst is 1,3-divinyl-1,1,3,3-tetramethyldisiloxane platinum (0), and its amount is 1% of the total mass of polydimethylsiloxane.

10. The application of a high crack-resistant silicon elastomer based on a multi-scale stress dispersion mechanism as described in any one of claims 1-4 in flexible electronic devices, flexible strain sensors, soft robots, and biomedical devices.

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