Investigation of the Mechanical and Thermal Properties of MWCNT/SiC-Filled Ethylene–Butene–Terpolymer Rubber

Abstract

Rubber is widely used in daily lives, such as in automobile tires, conveyor belts, sealing rings, and gaskets. The performance of rubber determines its service life. Therefore, it is of crucial importance to improve the performance of rubber. Theoretical studies have found that the inherent properties of nanofillers themselves, the interfacial bonding force between fillers and the matrix, and the uniform dispersibility of nanofillers in the polymer matrix are the most significant factors for enhancing the performance of rubber nanocomposites. This study systematically investigated the synergistic enhancement effect of silicon carbide (SiC) and multi-walled carbon nanotubes (MWCNTs) on the mechanical and thermal properties of ethylene–butene–terpolymer (EBT) composites. By optimizing the addition amount of fillers and improving the interfacial bonding between fillers and the matrix, the influence of filler content on the properties of composites was studied. The results demonstrate that the addition of SiC and MWCNTs significantly improved the storage modulus, tensile strength, hardness, and thermal stability of the composites. In terms of mechanical properties, the tensile strength of the composites increased from 6.68 MPa of pure EBT to 8.46 MPa, and the 100% modulus increased from 2.14 MPa to 3.81 MPa. Moreover, hardness was significantly enhanced under the reinforcement of SiC/CNT fillers. In terms of thermal stability, the composites exhibited excellent resistance to deformation at high temperatures. Through the analysis of the mechanical and thermal properties of the composites, the synergistic enhancement mechanism between SiC and MWCNTs was revealed. The research results provide a theoretical basis for the design and engineering applications of high-performance ethylene–butylene rubber composites.

1. Introduction

In the innovative development of rubber materials, ethylene–butene–terpolymer (EBT) has gradually emerged as a promising new synthetic rubber. EBT is synthesized through the polymerization of ethylene, butylene, and a small amount of a third monomer, such as Ethylidene Norbornene (ENB), using a metallocene catalyst. Compared to conventional ethylene–propylene–diene monomer (EPDM) rubber, the replacement of propylene with butylene in the molecular structure of EBT significantly improves the flexibility of the polymer backbone. This structural modification endows EBT with a series of superior properties, opening up broad prospects for its application in various fields [1]. For instance, Japan’s NOK Corporation has adopted EBT as a substitute for EPDM in the manufacturing of sealing products, resulting in enhanced processing performance, improved low-temperature resistance, and reduced production costs. As technology advances and market demand grows, the range of applications for EBT is expected to expand further. However, the shortcomings of EBT in terms of mechanical properties limit its widespread use in fields that require high material strength and durability, presenting a key challenge that needs to be addressed in the ongoing research and development of EBT.

In recent years, the rapid development of nanotechnology has opened new avenues for optimizing the performance of rubber materials. Silicon carbide (SiC) and multi-walled carbon nanotubes (MWCNTs), as two highly promising nanofillers, have attracted widespread attention for their ability to enhance the properties of rubber [2,3,4,5,6]. SiC powder exhibits excellent thermal stability, high strength, high hardness, and outstanding antioxidant properties, making it crucial in applications under extreme conditions such as high temperature, high frequency, and high power, as well as in semiconductor device manufacturing [7,8]. In the field of polymer-based composites, SiC is commonly used as a reinforcement, and it is widely applied in systems such as epoxy resins, polyurethanes, and rubbers, where it significantly improves the mechanical performance and thermal stability of the composite materials [9,10,11,12]. Kumar [13] et al. prepared SiC/carbon fiber (CF)/ethylene–propylene–diene monomer (EPDM) composite materials and investigated the effects of different SiC/CF ratios on the mechanical and thermal properties of the composites. They found that when 10 phr CF and 20 phr SiC were added, the tensile strength of the composite reached 12.09 MPa, with a break elongation of 725%, and its thermal stability improved by 20% due to the addition of SiC. Anancharoenwong [14] et al. studied the effect of SiC content on the vulcanization and mechanical properties of natural rubber/EPDM (NR/EPDM) blends. Their results showed that as the SiC content increased, the Tc90 and Ts2 of the NR/EPDM blend rose. When the SiC content reached 10 phr, the tensile properties of the vulcanized rubber were enhanced, but excessive addition led to a decrease in tensile performance.

Carbon-based fillers also exhibit excellent reinforcing properties and thermal stability, along with being lightweight and cost-effective. These include graphite, carbon black, graphene, carbon nanotubes (CNTs), and carbon fibers (CFs) [15,16,17,18,19,20]. Among them, carbon nanotubes are considered outstanding candidates for replacing or complementing traditional reinforcing particles due to their superior mechanical properties, including an extremely high elastic modulus of approximately 1 TPa and a tensile strength of 60 GPa [21,22,23,24]. Liu [25] et al. prepared nitrile rubber/polyamide elastomer/CNT composites using a physical blending method and studied the degradation behavior of the composites in hot oxidative and hot oil environments. The results showed that the incorporation of CNTs significantly improved the composites’ resistance to thermo-oxidative aging and oil resistance. Xu [26] et al. employed density functional theory to investigate the mechanism by which carbon nanotubes scavenge free radicals and enhance the antioxidant properties of polymers. Their results indicated that carbon nanotubes function as a shell system, with the spin density primarily distributed around the aromatic oxygen groups and the nearby sp²-hybridized carbon atoms. Atoms with a spin population ranging from 0 to 0.6 au are inert but can react with highly reactive radicals, including alkyl radicals. The mechanism was further confirmed through electron spin resonance measurements of nanocarbon and heat-treated phenol.

However, the reinforcement process of rubber materials is highly complex, influenced by various factors such as the size, structure, and dispersion state of the reinforcing particles, as well as the interactions between the fillers and between the fillers and the rubber matrix [27,28,29]. Therefore, identifying suitable reinforcing particles that can deliver performance similar to traditional micron-sized carbon black and silica at lower loadings has become a key research focus in the field of rubber reinforcement. This study systematically investigates the mechanical and thermal properties of SiC and MWCNT-reinforced EBT composites. By optimizing the filler content and improving the interfacial bonding between the fillers and the matrix, this study explores the variation in composite material properties and reveals the synergistic enhancement mechanisms, with the aim of providing theoretical guidance for the design and application of high-performance EBT rubber.

2. Materials and Test Procedures

2.1. Materials

Ethylene–butene–terpolymer (EBT, grade K9330M, ethylene mass fraction 50%, third monomer mass fraction 7.1%) was purchased from Mitsui Chemicals (Tokyo, Japan). Multi-walled carbon nanotubes (MWCNTs), with an outer diameter of 5~15 nm, length of 10~30 µm, specific surface area of 220~300 m²/g, and ash content less than 2 wt%, were obtained from Zhongke Times Nanotechnology Co., Ltd. (Chengdu, China). Silicon carbide (SiC, 99.9%, 1~2 µm) and 3-aminopropyltriethoxysilane (KH550, 99%) were purchased from Macklin Co., Ltd. (Shanghai, China). Other reagents, including zinc oxide (ZnO), stearic acid (SA), antioxidant 445, and curing agents (DCP, TAIC), were commercially sourced. In this study, MWCNTs are referred to as CNTs.

2.2. Surface Modification of SiC

To improve the compatibility and dispersion of SiC in EBT, the surface of SiC was modified using the silane coupling agent KH550. In total, 2 phr SiC powder was added to 100 mL of an ethanol–water mixture (volume ratio 19:1), and 0.02 phr of KH550 was slowly added to the solution. The mixture was sonicated for 30 min to ensure adequate contact between SiC and the coupling agent. The solution was then stirred at 60 °C for 5 h to obtain the KH550-modified SiC solution. Then, the modified solution was centrifuged at 3000 rpm for 30 min to separate the solid powder, which was dried in an oven to obtain the modified SiC powder.

2.3. Preparation of SiC/CNT/EBT Composite Rubber

In total, 100 phr of EBT was processed in a two-roll mill (XK-160, Zhanjiang Machinery Factory, eZhanjiang, China) to reach a uniform plastic state. The corresponding amounts of CNTs and modified SiC, as specified in

Table 1. SiC/CNT ratio in SiC/CNT/EBT composites.

	Pure EBT	0-SiC/ CNT/EBT	5-SiC/ CNT/EBT	10-SiC/ CNT/EBT	15-SiC/ CNT/EBT	20-SiC/ CNT/EBT	5- SiC/EBT
CNT/phr	0	2	2	2	2	2	0
SiC/phr	0	0	5	10	15	20	5

, were added, with pure EBT samples having no CNTs or modified SiC. The 0-SiC/CNT/EBT to 20-SiC/CNT/EBT samples were slowly blended with 2/0, 2/5, 2/10, 2/15, and 2/20 phr of CNT and modified SiC, respectively, and continuously mixed. In total, 5 phr zinc oxide (ZnO), 1 phr stearic acid (SA), and 1 phr antioxidant 445 were added and thoroughly mixed. Curing agents (3 phr DCP, 2 phr TAIC) were then incorporated and mixed completely to prepare composites with different SiC/CNT contents. The curing process consisted of two stages. In the first stage, the mixture was vulcanized in a hydraulic press at 170 °C, 10 MPa for 13 min. In the second stage, vulcanization was carried out in an electric hot-air oven at 150 °C for 4 h, resulting in rubber sheets with a thickness of 2 mm.

2.4. Characterization

According to the requirements of GB/T 16584-1996, the cure characteristics were tested using a rotorless vulcanizer (MM4310C, Huanfeng Chemical, Beijing, China) at 170 °C for 20 min. Shore hardness tests of the vulcanizates were performed at room temperature using a Shore durometer (Yizhong Precision Instruments Co., Shanghai, China) in accordance with GB/T 531.1-2008 [30]. Five different positions on each sample were measured, and the average value was taken to minimize errors. The tensile strength, 100% modulus, and elongation at break of the vulcanizates were tested using a universal material testing machine (Zwick Roell Z010, Zwick, Germany) at a uniform speed of 500 mm·min⁻¹ following GB/T 528-2009 [31]. Each test was repeated three times (n = 3) with an instrument precision of less than ±2%. The median value of the three replicate samples was finally used to plot the stress–strain curves. Dynamic mechanical analysis was conducted using a dynamic mechanical analyzer (NEXTA DMA200, Hitachi Analytical Instruments Co., Shanghai, China) according to GB/T 40396-2021 [32]. Tests were performed in single-cantilever mode with a specimen size of 25 × 12 × 2 mm³, a constant frequency of 10 Hz, a strain amplitude of 30 μm, a temperature range from −60 °C to 40 °C, and a heating rate of 3 °C/min. The storage modulus, loss modulus, glass transition temperature (T_g), and dynamic loss factor (tanδ) were obtained through analysis. The microstructure and morphology of the composites were characterized using a scanning electron microscope (SEM, JEOL7600F, JEOL Ltd., Tokyo, Japan). Composite samples were cut into pieces of appropriate size, placed in a 170 °C forced-air drying oven using hooks, and removed after 48 h for weighing to calculate thermal mass loss before and after the treatment.

Equilibrium solvent swelling measurements were conducted in n-hexane at 25 °C. Samples were cut into disk-shaped specimens with a diameter of 20 mm and a thickness of 2 mm. Three disks were used for each sample. The samples were dried at 60 °C for 24 h, weighed (m₁), and then immersed in dark capped bottles containing n-hexane [33]. After a soaking time of 72 h to achieve equilibrium, the samples were reweighed (m₂). The specimens were subsequently dried to constant mass, denoted as m₃. The crosslink density (ν_e) was determined using the Flory–Rehner equation [34]:

$$v_e=-{\displaystyle \frac{\ln \left(1-V_r\right)+V_r+\chi {V_r}^2}{V_s(\sqrt[3]{V_r}-\frac{V_r}{2})}}$$

(1)

where χ describes the Flory–Huggins polymer–solvent interaction parameter, V_s is the molar volume of the solvent (130.45 cm³/mol for n-hexane), and V_r is the volume fraction of the sample at swelling equilibrium, which is calculated by Equation (2).

$$V_r={\displaystyle \frac{\frac{m_3}{{\rho }_r}}{\frac{m_3}{{\rho }_r}+\frac{m_2-m_3}{{\rho }_s}}}$$

(2)

where ρ_s is the solvent density and ρ_r is the rubber density.

The Flory–Huggins polymer–solvent interaction parameter χ, a measure of the interaction energy between solvent molecules and the rubber, was estimated using Equation (3):

$$\chi =\beta +{\displaystyle \frac{V_s}{RT}}{({\sigma }_r-{\sigma }_s)}^2$$

(3)

where β is the lattice constant of the polymer–solvent system (β = 0.34), R is the universal gas constant, T is the absolute temperature, and σ_r and σ_s are the solubility parameters of the rubber sample and the solvent, respectively.

3. Results and Discussion

3.1. Vulcanization Performances

The presence of nanofillers significantly affects the vulcanization process of rubber, which in turn influences the physical and mechanical properties of nanocomposites. To investigate the impact of nanofillers on the vulcanization process, the vulcanization curves and related parameters for unfilled EBT and the corresponding composites were tested and summarized, as shown in Figure 1a and

Table 2. Vulcanization parameters of pure EBT and SiC/CNT/EBT composites.

	ML (dNm)	MH (dNm)	MH-ML (dNm)	Tc10 (min)	Tc90 (min)
Pure EBT	0.23	19.24	19.01	1.15	11.21
0-SiC/CNT/EBT	0.27	20.43	20.16	1.14	11.04
5-SiC/CNT/EBT	0.29	21.40	21.11	1.15	10.94
10-SiC/CNT/EBT	0.30	21.88	21.58	1.11	10.72
15-SiC/CNT/EBT	0.32	22.72	22.40	1.09	10.60
20-SiC/CNT/EBT	0.34	23.12	22.78	1.10	10.86
5-SiC/EBT	0.20	19.11	18.91	1.09	11.15

. The results indicate that the incorporation of CNTs and SiC increases the minimum torque (M_L), maximum torque (M_H), and torque difference (M_H-M_L) of the rubber. The change in M_L is related to the viscosity of the compound. The dispersion of CNTs and SiC in the rubber matrix occupies part of the free volume, reducing the available space for the movement of the rubber molecular chains, which decreases the flowability. Furthermore, the active adsorption sites on the surface of the fillers interact with the rubber molecular chains, restricting the relative slippage and extension of the molecular chains, thereby causing an increase in M_L. This effect is typical in elastomer blends filled with rigid filler particles [35,36].

The values of M_H and M_H-M_L reflect the shear modulus and hardness of the rubber, which result from the combined effects of physical reinforcement from the nanofillers and the crosslinking density of the vulcanized rubber. The results in Table 2 show that the effect of CNTs on the M_H-M_L values of the composites is most significant. After adding 2 phr of CNTs, the M_H-M_L value of the composite increased from 19.01 dNm to 20.16 dNm, whereas the addition of 20 phr SiC resulted in a 4.86 dNm increase in M_H-M_L. This is attributed to the high modulus of carbon nanotubes, which adsorb rubber molecular chains and intertwine during vulcanization, forming a crosslinked network with the rubber. Meanwhile, as the content of SiC gradually increases, the active groups on its surface form covalent bonds with rubber molecules, increasing the crosslinking points at the filler–rubber interface. This further enhances the crosslinking density of the composites (Figure 1b), resulting in a gradual increase in both M_H and M_H-M_L. The scorch time (Tc10) and cure time (Tc90) of the vulcanized rubber are negatively correlated with the amount of SiC, which can be attributed to the excellent thermal conductivity of SiC. During the vulcanization process, heat is transferred quickly and evenly within the material, accelerating the crosslinking reaction.

3.2. Mechanical Properties

Tensile tests were conducted to evaluate the mechanical properties of pure EBT and SiC/CNT/EBT composites. The stress–strain curves are shown in Figure 2a, while Young’s modulus, tensile strength, 100% modulus, and elongation at break are presented in Figure 2b–e, respectively. As indicated in Figure 2b, Young’s modulus increases gradually with increasing filler content, and the increment in the hybrid filler system is significantly higher than that in single-filler systems, indicating that the SiC/CNT hybrid filler system can effectively enhance the material’s resistance to elastic deformation.

The 100% modulus and tensile strength of pure EBT rubber are 2.1 MPa and 6.7 MPa, respectively. After adding 2 phr CNT, these values rise to 2.5 MPa and 7.0 MPa, while adding 20 phr SiC increases them to 3.8 MPa and 8.5 MPa. The performance improvement in the CNT-only system is attributed to the unique one-dimensional nanostructure and high aspect ratio of carbon nanotubes, which enable effective interactions with rubber molecular chains via van der Waals forces. When the composite is subjected to external forces, CNTs can transfer stress through the tube walls and distribute the load, allowing the material to withstand more external loading before complete failure. The incorporation of SiC further enhances the mechanical properties, with tensile strength showing a positive correlation with filler content. Although elongation at break decreases slightly with increasing SiC content, it still meets engineering requirements.

As a reinforcing filler, SiC particles disperse in the rubber matrix, and under the action of coupling agent KH550 and CNTs, good interfacial adhesion is established between SiC and rubber. With increasing SiC loading, more crosslinking points form in the rubber matrix, contributing to a higher overall crosslink density. This further restricts the mobility of rubber molecular chains and reduces chain deformation, leading to a decrease in elongation at break. However, surface modification of SiC using only coupling agent KH550 is limited by single-chemical interfacial interactions and insufficient dispersibility, failing to completely avoid interfacial defects caused by filler agglomeration. Many studies have shown that CNTs can improve filler dispersion [37,38,39,40,41]. Their unique one-dimensional nanostructure forms strong interactions with filler particles and rubber chains via van der Waals forces, effectively preventing SiC agglomeration during mixing and optimizing the uniformity and continuity of the filler network.

Considering the strong anisotropic characteristics of CNTs and their positive implementation in composite production, the mechanical properties of SiC/CNT/EBT composites in different directions were tested, as shown in Figure S2. Tensile tests perpendicular to the sheet extrusion direction indicate that the 100% modulus is higher than that in the parallel direction, while the tensile strength and elongation at break are significantly reduced. This is attributed to the anisotropic network structure formed by the orientational alignment of CNTs along the sheet extrusion direction during processing, which reduces the stress transfer efficiency at the filler–matrix interface and restricts molecular chain mobility in the perpendicular direction. Therefore, in practical applications, attention should be paid to matching the load direction with the processing orientation of the material, ensuring that the principal stress direction is consistent with the filler orientation to avoid performance degradation caused by anisotropy.

The variation in hardness with filler content is shown in Figure 2f, where the hardness of the composites gradually increases with the increase in SiC mass fraction. This increase in hardness may reduce the deformation energy absorption capacity of the composites, while the uniformly dispersed SiC particles help rapidly transfer external impacts throughout the material matrix, avoiding concentrated impact stress.

Figure 3 illustrates the volume wear (V_s) of composites with different contents of SiC/CNT. The addition of CNT and SiC significantly reduces the wear of the composites. When 5 phr SiC is added, the volume wear of the composite decreases by 2.1%; after adding 2 phr CNT, the reduction in volume wear reaches 5.5% and further addition of 20 phr SiC leads to a 13.4% decrease in volume wear. The reasons for this phenomenon can be divided into two aspects. On the one hand, carbon nanotubes can form a nanoscale lubrication film at the friction interface, which forms a low shear resistance layer and effectively reduces the direct contact between abrasive particles and the rubber surface. On the other hand, SiC particles rely on their high hardness properties to effectively resist the abrasive plowing effect, thereby reducing the occurrence of material shedding. In addition, under the condition of high content filler and relatively uniform dispersion, a continuous lubrication network will be formed.

3.3. Morphology

The dispersion of nanofillers in rubber determines the effectiveness of reinforcement. Therefore, it is essential to systematically investigate the dispersion behavior of fillers in SiC/CNT/EBT composites. The cross-sectional morphology of pure EBT rubber and composites with varying SiC contents was observed using scanning electron microscopy (SEM) (Figure 4), combined with SEM characterization of the modified SiC and CNT (Figure S3), revealing the mechanism of filler dispersion and interface interactions.

Figure S3 shows that the modified SiC exhibits an irregular block-like structure with an average particle size of approximately 1 μm. The smaller particle size provides a higher specific surface area, which is beneficial for the formation of a thermal conductivity network and enhancement of mechanical properties. The SiC surface is coated with the KH550 coupling agent, which increases the chemical and physical adsorption between SiC and the rubber molecular chains, providing conditions for the formation of additional crosslinking points within the matrix. On the other hand, CNTs exhibit a typical one-dimensional tubular structure, and their high aspect ratio makes it easy to form a continuous network framework in the rubber matrix. Notably, the unique structure of CNTs can act as a “dispersant” through steric hindrance, effectively suppressing the agglomeration of SiC particles. As shown in Figure 4c–f, SiC particles in the 5-SiC/CNT/EBT sample exhibit uniform dispersion in the matrix. As the SiC content increases, SiC particles and CNTs start to contact each other, eventually forming a three-dimensional network structure. The formation of this hierarchical network significantly enhances the mechanical properties of the composites: tensile strength, 100% modulus, and hardness increase by 26.7%, 78.3%, and 12.3%, respectively, compared to pure EBT (Figure 2). In samples without CNTs (5-SiC/EBT), partial agglomeration of SiC particles occurs, as observed in Figure 4g. EDS mapping was used for the quantitative visualization analysis of Si elements (Figure 4h–l) to evaluate the dispersion uniformity of SiC in the rubber. The analysis results show that the distribution pattern of Si elements is consistent with the microstructural characteristics observed by SEM, further confirming the reliability of CNTs in improving the dispersion state of SiC.

The above-mentioned phenomena are closely related to the synergistic reinforcement mechanism of the fillers. Alam et al. [42] observed that when a binary filler system consisting of graphite powder and silica was used in natural rubber, the graphite powder, acting as a solid lubricant, facilitated the dispersion of silica fillers within the rubber matrix. Based on this study, graphene and carbon black were also found to influence the dispersion of SiC in the EPDM matrix, helping to achieve better mechanical properties in samples with hybrid filler systems. Similarly, in our work, CNTs, through physical entanglement and steric hindrance effects, confined SiC particles to a nanoscale dispersion state, thereby maximizing their specific surface area advantages and enhancing the thermal conductivity network and mechanical properties.

3.4. Payne Effect

Strain sweep tests were performed on pure EBT and various SiC/CNT filler systems using a rubber process analyzer (RPA) to calculate the Payne effect index (ΔG′ = G′₀ − G′_∞, where G′₀ and G′_∞ represent the storage modulus (G′) of the composites at low and high strains, respectively). The results (Figure 5) show that the G′ of the 5-SiC/EBT sample at low strain increased only slightly compared to pure EBT rubber. This is because an effective filler network was not established in this system, and the weak interfacial adhesion between SiC and EBT limited the efficiency of stress transfer.

In the SiC/CNT dual-filler system, ΔG′ increased progressively as the SiC content increased from 0 phr to 20 phr. This is attributed to CNTs bridging SiC particles via their high aspect ratio, forming an interpenetrating CNT-SiC network structure. At low strains, the network provides elastic support, significantly enhancing G′; at high strains, the network fragments stepwise to dissipate energy, causing a sharp drop in G′, which is characteristic of dynamic filler network reconstruction. Compared with the single-filler system, the Payne effect index of the double-filler composite system is significantly improved, confirming that CNTs, through nanoscale dispersion regulation and multi-scale interface coordination, transforms SiC from inefficient single-point enhancement to efficient network collaborative enhancement, achieving an essential breakthrough in dynamic mechanical properties.

3.5. Dynamic Mechanical Analysis (DMA)

Dynamic mechanical analysis (DMA) was performed on pure EBT and its composites, and the results are shown in Figure 6. As shown in Figure 6a, it can be observed that with increasing temperature, the G′ of the composite materials gradually decreases. This is primarily due to the increased movement of rubber chains during the glass–rubber transition, which leads to greater energy dissipation. At the same time, the mobility of the polymer chains increases with rising temperature.

The DMA test shows that the G′ of pure EBT under dynamic small strain load is relatively low, indicating that its ability to resist elastic deformation is weak and it is more prone to reversible elastic deformation. When SiC/CNT fillers are added, the G′ of the composite is significantly increased, indicating that SiC/CNT effectively enhances the elastic stiffness of the material under dynamic loads through rigid support and strong interfacial interaction. Additionally, strong interface interactions between the SiC/CNT fillers and the matrix effectively improve the efficiency of stress transfer from the matrix to the fillers. This enhancement effect is further reflected in the reduction in the maximum tanδ value and the improvement in the tensile properties of the composites. As the filler content increases from 0 to 20 phr, the storage modulus further improves. Even within the glass transition temperature (T_g) range, the modulus reduction in the composites is smaller than that in pure EBT, indicating that the filler network effectively restrains the temperature-induced intensification of molecular chain mobility. This enhances the composite’s dimensional stability and dynamic mechanical properties across different temperatures.

Tanδ is used to measure the efficiency of energy loss in materials due to the molecular rearrangement and internal friction of molecular chains. As shown in Figure 6b, the addition of only 5 phr SiC introduces an interface between SiC and the rubber matrix in the composite. In the interfacial region, the mobility of molecular chains is restricted, increasing internal friction loss and leading to a slight increase in tanδ. By contrast, in the SiC/CNT/EBT composites, the height of the tanδ peak decreases significantly. This is because the network structure formed by the rubber matrix and fillers, together with enhanced interfacial bonding between fillers and the matrix, reduces energy loss caused by interfacial slippage. Such strong interfacial interactions also improve the efficiency of stress transfer from the matrix to the fillers, thereby enhancing the tensile strength of the composites.

During dynamic mechanical analysis, as the temperature increases, the mobility of the rubber molecular chains gradually enhances. Within a certain temperature range, the friction between rubber chains significantly intensifies, leading to an increase in internal energy dissipation, and thus a rise in internal loss. On the tanδ–temperature curve, this results in a peak, with the temperature corresponding to the tanδ peak being the material’s glass transition temperature (T_g). After adding SiC/CNT fillers, the position of the tanδ peak remains relatively stable, indicating that the fillers do not significantly alter the T_g of the EBT matrix. This suggests that the composite material maintained its original glass transition characteristics while optimizing other properties.

3.6. Thermal Stability

In high-temperature service environments, the thermal stability of composites directly affects their long-term reliability and functional integrity. Therefore, studying their thermal degradation behavior is of significant guiding importance for material design. The thermal degradation results shown in Figure 7a indicate a pronounced stepwise decrease in the thermal loss rate as the SiC/CNT filler content increases from 0 phr to 20 phr, which corresponds well with the trend of enhanced storage modulus observed in dynamic mechanical analysis.

Pure EBT, lacking rigid fillers for physical restraint, exhibits increased molecular chain mobility at high temperatures, accelerating oxidation and degradation reactions and demonstrating the highest thermal loss rate (1.189%). When 2 phr CNT is introduced, the thermal loss rate dramatically drops to 1.004%, which is attributed to the physical barrier effect of the fillers that delays oxygen diffusion. With the further addition of 5 phr SiC, the thermal loss rate decreases to 0.923%, as the chemical grafting layer (e.g., Si-O-C bonds) formed between the modified fillers and the matrix significantly enhances the interfacial bonding. Meanwhile, the one-dimensional thermal conduction pathways of CNTs and the high thermal conductivity of SiC work synergistically to construct an efficient heat dissipation path, effectively suppressing chain degradation reactions caused by local hotspots. Notably, the thermal loss rate of the 20 phr system (0.808%) is significantly lower than that of the 15 phr system (0.868%), indicating that this filler combination maintains good dispersion and synergistic effects even at high filler contents.

After being aged at 170 °C for 48 h, the composites exhibited excellent stability and anti-deformation ability in terms of the compression set under high-temperature compressive loads. As shown in Figure 7b, the compression set of the composites with different filler ratios was stable within the range of 20% to 25%. The specific values were 22.74%, 25.41%, 24.05%, 21.69%, 23.74% and 21.33%, respectively. This indicates that the materials possess outstanding anti-creep performance under high-temperature conditions. The addition of fillers had a relatively small impact on the compression set performance of the composites.

We systematically investigated the regulatory effect of the SiC/CNT hybrid filler system on the thermal oxidation aging behavior of vulcanizates. The composites were placed in an environment at 170 °C for 48 h to conduct an accelerated aging experiment, and then the changes in their tensile properties were tested (Figure 8). The results show that the thermal oxidative environment can lead to significant decreases in the tensile strength and elongation at break of vulcanizates. In the pure EBT system, due to the lack of protection for molecular chains, severe thermal oxidative degradation occurs under high temperatures, demonstrating the highest performance attenuation—the tensile strength reduction rate reaches 10.03%, and the elongation at break decreases by 32.07%. This is attributed to the synergistic effect of the breakage of the crosslinking network and the disentanglement of molecular chains.

It is worth noting that the introduction of the SiC/CNT hybrid filler system significantly delays the thermal oxidative degradation process of the composites, which is manifested as an improvement in the property retention rate. After adding 2 phr CNT, the variation rate of the tensile strength slows down to 5.80%, and the variation rate of the elongation at break drops to 26.94%. Further addition of 20 phr SiC enables the system to achieve a better level of property retention (the variation rate of the tensile strength is 1.50%, and the variation rate of the elongation at break is 17.51%). This optimization effect stems from the electron delocalization effect of the sp² structure of CNTs. It reduces the activity of endogenous free radicals and slows down the oxygen diffusion rate, weakening the free radical crosslinking reaction under high temperatures. Meanwhile, the high thermal conductivity of SiC promotes the uniform distribution of heat and suppresses the degradation caused by local overheating. CNTs and SiC form a “thermal-conductive-barrier” composite network through the interfacial synergy effect, effectively maintaining the integrity of the crosslinking structure.

4. Conclusions

The present study demonstrates that the SiC/CNT composite filler system exhibits both high efficiency and economic viability. It averts the issues of cost escalation and agglomeration that would arise from excessive utilization of CNTs. Meanwhile, through the dispersion optimization mechanism, it ameliorates the agglomeration defect of single SiC, significantly strengthening the interfacial bonding force and facilitating the formation of the thermal conductive network, thereby enhancing the mechanical properties and thermal stability of EBT composites. In optimizing the addition amount of SiC and improving the interfacial bonding between the filler and the matrix, remarkable enhancements have been achieved in the tensile strength, modulus at a given elongation, hardness, and thermal stability of the composite materials. Specifically, the tensile strength of the composites increased from 6.7 MPa of pure EBT to 8.5 MPa, the modulus at 100% elongation rose from 2.1 MPa to 3.8 MPa, and the hardness witnessed an increment of 10.5%. Moreover, the compression set remains at a relatively low level even under high filler contents. The design of EBT composites provides an effective material optimization strategy, offering valuable theoretical references for the design and development of high-performance and low-cost rubber composites. It also contributes to the expansion of their application scope in fields with stringent requirements for mechanical and thermal properties, such as the automotive, aerospace, and industrial manufacturing sectors.