A multi-scale composite filler, a hydrophobically modified multi-scale composite filler and a modified fluororubber

Abstract

The application provides a multi-scale composite filler, a hydrophobic modified multi-scale composite filler and modified fluororubber, and for the first time introduces nanocellulose into the construction process of hollow glass microspheres coated with SiO2 by Stober method, the rich hydroxyl groups on the surface of the nanocellulose provide more binding sites for the hydrolysis and polycondensation of SiO2 precursors, play the role of inducing nucleation and bridging growth, and construct a multi-scale composite filler with spherical skeleton, rough surface and fiber network characteristics. The fiber network structure of nanocellulose plays a bridging, entangling and interface regulation role in the composite system, significantly enhances the interface bonding and stress transfer efficiency between the filler and the fluororubber matrix. Only 2 phr of extremely low filling amount can increase the tensile strength by 150% and the elongation at break by 39%, the reinforcing efficiency is significantly better than that of the prior art, and the hydrophobicity is strong, and can be used in the automobile and new energy automobile industry, engine oil seal, fuel pipeline seal, battery pack sealing element and the like.

Description

Technical Field

[0001] This invention relates to the field of polymer composite materials technology, specifically to a multi-scale composite filler, a hydrophobically modified multi-scale composite filler, and modified fluororubber. Background Technology

[0002] Fluororubber (FKM), with its excellent chemical inertness, thermal stability, oil resistance, and mechanical strength, has become a key material in high-end manufacturing fields such as aerospace, automotive, and petrochemical industries. However, with increasingly complex service environments, there is still room for improvement in the mechanical strength, thermal stability, and surface functionality of traditional fluororubber.

[0003] In existing technologies, modifying fluororubber by introducing inorganic nanofillers is an important means to improve its overall performance. Chinese patent application CN119430197A discloses a technique for synthesizing silica microsphere fillers using the Stöber method and then filling modified fluororubber. Chinese patent application CN118667268A discloses a MWCNTs@SiO2-C core-shell filler-modified fluororubber composite material, which uses MWCNTs as the core and carbon nanoparticles filled with nano-SiO2 as the shell. However, the above technical solutions still suffer from problems such as limited reinforcing efficiency, insufficient filler dispersion, and poor interfacial compatibility.

[0004] Cellulose nanoparticles (CNCs) possess high specific surface area, high aspect ratio, and abundant surface hydroxyl groups, providing nucleation and binding sites and playing a bridging, entanglement, and interface regulation role in composite systems. However, there are currently no reports on the introduction of CNCs into the Stöber method SiO2-coated hollow glass microsphere system to construct multi-scale composite fillers for fluororubber modification.

[0005] Furthermore, excessively high surface polarity of fillers can lead to insufficient dispersion stability and poor interfacial compatibility in fluororubber matrices. Hexadecyltrimethoxysilane (HDTMS) is a typical organosilane coupling agent that can introduce long-chain alkyl groups through condensation reactions with hydroxyl groups on the filler surface, thereby reducing the surface energy of the filler. However, research on the overall modification of multi-scale composite fillers using HDTMS and its application in fluororubber systems remains lacking. Summary of the Invention

[0006] The purpose of this invention is to overcome the shortcomings and deficiencies of the existing technology and to provide a multi-scale composite filler, a hydrophobic modified multi-scale composite filler, and a modified fluororubber.

[0007] The technical solution adopted by the present invention is as follows: The first aspect of the present invention provides a multi-scale composite filler, which includes a hollow glass microsphere core and a nano-SiO2 shell layer deposited in situ on the surface of the hollow glass microsphere core by the Stöber method. The nano-SiO2 shell layer is formed by the accumulation of SiO2 particles, and the nano-SiO2 shell layer contains nano-cellulose dispersed between the SiO2 particles and forming a fiber network structure with them.

[0008] Preferably, the mass ratio of the nanocellulose to the hollow glass microspheres is (0.5-2):1.

[0009] Preferably, the preparation method includes the following steps: S1. Mix nanocellulose with anhydrous ethanol, add N,N-dimethylacetamide, heat and stir to pretreat, and obtain a pretreated nanocellulose dispersion; S2. Mix anhydrous ethanol and deionized water in a certain proportion, add polyethylene glycol, hollow glass microspheres and the pretreated nanocellulose dispersion obtained in step S1, stir evenly to obtain a mixed dispersion. S3. Add ammonia water to the mixed dispersion, and simultaneously add tetraethyl orthosilicate dropwise to carry out a hydrolysis and polycondensation reaction. After the reaction is completed, separate, wash and dry the product to obtain a multi-scale composite filler.

[0010] Preferably, in step S1, the heating and stirring temperature is 80-100 ℃ and the time is 4-8 h.

[0011] Preferably, in steps S2 and S3, the mass ratio of hollow glass microspheres to tetraethyl orthosilicate is 1:(3-7).

[0012] A second aspect of the present invention provides a hydrophobically modified multi-scale composite filler, wherein the hydrophobically modified multi-scale composite filler is obtained by surface modification of a multi-scale composite filler with hexadecyltrimethoxysilane.

[0013] Preferably, the amount of hexadecyltrimethoxysilane used is 10%-30% of the mass of the multi-scale composite filler.

[0014] Preferably, the preparation method includes the following steps: dispersing the multi-scale composite filler in anhydrous ethanol, stirring until uniform to obtain a dispersion, adding hexadecyltrimethoxysilane to the dispersion, stirring to react, and after the reaction is completed, separating, washing and drying the product to obtain the hydrophobic modified multi-scale composite filler.

[0015] A third aspect of the present invention provides a modified fluororubber comprising a fluororubber matrix and a multi-scale composite filler, or a hydrophobically modified multi-scale composite filler.

[0016] Preferably, the preparation method includes the following steps: (A) Fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and SiO2-coated hollow glass microsphere core-shell composite filler were sequentially added to an open mill and mixed evenly to obtain a mixture. (B) The mixture is vulcanized at 150-190℃ for 5-15 minutes to obtain vulcanized fluororubber; (C) The vulcanized fluororubber is vulcanized in two stages at 180-220℃ for 18-22 h to obtain modified fluororubber; In step (A), the ratio of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and composite filler is 100: (2-4): (4-6): (2-4): (2-4): (2-10).

[0017] The beneficial effects of this invention are as follows: 1. This invention is the first to introduce nanocellulose into the construction process of SiO2-coated hollow glass microspheres using the Stöber process. The abundant hydroxyl groups on the surface of nanocellulose provide more binding sites for the hydrolysis and polycondensation of the SiO2 precursor, playing a role in inducing nucleation and bridging growth, thus constructing a multi-scale composite filler with spherical skeleton, rough surface, and fiber network characteristics. The fiber network structure of nanocellulose plays a bridging, entanglement, and interface regulation role in the composite system, significantly enhancing the interfacial bonding and stress transfer efficiency between the filler and the fluororubber matrix. An extremely low filling amount of only 2 phr can increase tensile strength by 150% and elongation at break by 39%, with a significantly better reinforcing efficiency than existing technologies.

[0018] 2. This invention further employs HDTMS to perform overall alcohol modification on the multi-scale composite filler, successfully grafting long-chain alkyl groups onto the filler surface. This effectively reduces the filler's surface energy, improves its dispersion and interfacial compatibility in fluororubber, and simultaneously enables the composite material's water contact angle to exceed 100°, achieving superhydrophobic properties. This invention achieves a synergistic improvement in the mechanical properties, surface hydrophobicity, and thermal stability of fluororubber composite materials. The 2 phr S@H@C@HD filler results in a tensile strength of 16.0 MPa (an increase of 166.7%), an elongation at break of 215% (an increase of 50.1%), and a water contact angle of 104.3° at 8 phr, demonstrating excellent overall performance.

[0019] 3. The preparation method of this invention is controllable and highly repeatable, making it suitable for industrial production. It can be used in the automotive and new energy vehicle industries for applications such as engine oil seals, fuel line seals, and battery pack seals. Attached Figure Description

[0020] To more clearly illustrate the technical solutions in the embodiments of the present invention or the prior art, the drawings used in the description of the embodiments or the prior art will be briefly introduced below. Obviously, the drawings described below are only some embodiments of the present invention. For those skilled in the art, obtaining other drawings based on these drawings without creative effort still falls within the scope of the present invention.

[0021] Figure 1 The image shows a scanning electron microscope (SEM) image of the multi-scale composite filler prepared in Example 1 of this invention. Figure 2 The Fourier transform infrared (FTIR) spectra of the multi-scale composite fillers prepared in Examples 1 and 7 of this invention are compared with those of pure SiO2 and HGM. Figure 3 (a) Storage modulus and (b) loss factor of the F / S@H@C fluororubber composite materials prepared in Comparative Examples 1 and Examples 2-6 of this invention; Figure 4 The vulcanization performance curves of the F / S@H@C fluororubber composite materials prepared in Comparative Example 1 and Examples 2-67 of this invention are as follows: (a) Maximum torque M H and minimum torque M L (b) Scorching time T 10 and positive vulcanization time T 90 ; Figure 5 The mechanical properties of the F / S@H@C fluororubber composite materials prepared in Comparative Examples 1 and Examples 2-6 of this invention are as follows: (a) tensile strength; (b) elongation at break; (c) hardness; Figure 6 The tensile cross-section SEM image of the F / S@H@C fluororubber composite material prepared in Example 4 of this invention; Figure 7 The graph shows the water contact angle variation of the F / S@H@C fluororubber composite materials prepared in Comparative Example 1 and Examples 2-6 of this invention. Figure 8 The thermal conductivity of the F / S@H@C fluororubber composite materials prepared in Comparative Examples 1 and Examples 2-6 of this invention is shown in the following: (a) infrared thermal images of different samples during heating and (c) during cooling, and (b) heating curves and (d) cooling curves of different samples. Figure 9 Thermogravimetric analysis of the F / S@H@C fluororubber composites prepared in Comparative Examples 1 and Examples 2-6 of this invention: (a) TG curve; (b) DTG curve; (c) residual carbon content; Figure 10 This is a SEM image of the multi-scale composite packing prepared in Example 7 of the present invention; Figure 11(a) Storage modulus and (b) loss factor of the F / S@H@C@HD fluororubber composite materials prepared in Comparative Example 2 and Examples 8-12 of this invention; Figure 12 The vulcanization performance curves of the F / S@H@C@HD fluororubber composite materials prepared in Comparative Example 2 and Examples 8-12 of this invention are as follows: (a) Maximum torque M H and minimum torque M L (b) Scorching time T 10 and positive vulcanization time T 90 ; Figure 13 Mechanical properties of the F / S@H@C@HD fluororubber composites prepared for Comparative Example 2 and Examples 8-12 of this invention: (a) tensile strength; (b) elongation at break; (c) hardness; Figure 14 The tensile cross-section SEM image of the F / S@H@C fluororubber composite material prepared in Example 11 of this invention; Figure 15 The graph shows the water contact angle variation of the F / S@H@C@HD fluororubber composite materials prepared in Comparative Example 2 and Examples 8-12 of this invention. Figure 16 Thermal conductivity of the F / S@H@C@HD fluororubber composite materials prepared in Comparative Example 2 and Examples 8-12 of this invention: (a) Infrared thermal images of different samples during heating and (c) during cooling, (b) heating curves and (d) cooling curves of different samples; Figure 17 Thermogravimetric analysis of the F / S@H@C@HD fluororubber composites prepared in Comparative Example 2 and Examples 8-12 of this invention: (a) TG curve; (b) DTG curve; (c) residual carbon content. Detailed Implementation

[0022] To make the objectives, technical solutions, and advantages of the present invention clearer, the present invention will be further described in detail below with reference to the accompanying drawings.

[0023] Example 1 This embodiment provides a multi-scale composite filler, the preparation method of which is as follows: (1) DMAC pretreatment of nanocellulose: 50 mL of anhydrous ethanol and 1 g of nanocellulose (CNC) were placed in a vacuum reaction flask, and 1 mL of N,N-dimethylacetamide (DMAC) was added. After sealing the vacuum reaction flask, it was placed in a water bath magnetic stirrer at 90 °C for 6 h and then taken out for use.

[0024] (2) Mix 187 mL of anhydrous ethanol with 6.5 mL of deionized water until homogeneous, add 0.1 g of polyethylene glycol (PEG), 1 g of hollow glass microspheres (HGM), and the pretreated nanocellulose dispersion obtained in step (1), and stir magnetically at room temperature for 30 min. Then add 5 mL of ammonia water (28%) to the system, and simultaneously add 5 mL of tetraethyl orthosilicate (TEOS) dropwise using a quantitative dropping funnel. After the addition is complete, continue to stir magnetically at room temperature for 24 h. After the reaction is complete, transfer the product to an electric thermostatic drying oven and dry at 60 ℃ for 12 h. After drying, a white powder is obtained, which is washed three times each with anhydrous ethanol and deionized water and then dried again to obtain the multi-scale composite filler, denoted as S@H@C.

[0025] Example 2 This embodiment provides a modified fluororubber, the preparation method of which is as follows: The open mill was set to a roll temperature of 80 ℃ and a roll gap of 6 mm. Fluororubber raw rubber was preheated in an 80 ℃ electric heating oven for 2 h before being fed into the open mill. Evenly ground magnesium oxide and calcium oxide were added to the raw rubber in small, repeated batches, cutting each side three times and forming a triangular bundle three times. Then, TAIC and Luperox 101-XL-45 were added sequentially, forming a triangular bundle six times. Finally, the composite filler S@H@C prepared in Example 1 was added, and the mixture was stirred until homogeneous. After stirring, the sample was left at room temperature for 20 h.

[0026] The molding process was carried out using a flat vulcanizing machine. The first vulcanization conditions were: temperature 170 ℃, time 10 min, and pressure 10.0 MPa. The second vulcanization was carried out at 200±2 ℃ for 20 h.

[0027] In this embodiment, the ratio of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and composite filler S@H@C is 100:3:5:3:3:2, and the final modified fluororubber sample is denoted as F / S@H@C-2.

[0028] Example 3 This embodiment provides a modified fluororubber, the preparation method of which differs from that of Example 2 in that the ratio of the added parts of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and composite filler S@H@C is 100:3:5:3:3:4, and the final modified fluororubber sample is denoted as F / S@H@C-3.

[0029] Example 4 This embodiment provides a modified fluororubber, the preparation method of which differs from that of Example 2 in that the ratio of the added parts of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and composite filler S@H@C is 100:3:5:3:3:6, and the final modified fluororubber sample is denoted as F / S@H@C-4.

[0030] Example 5 This embodiment provides a modified fluororubber, the preparation method of which differs from that of Example 2 in that the ratio of the added parts of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and composite filler S@H@C is 100:3:5:3:3:8, and the final modified fluororubber sample is denoted as F / S@H@C-5.

[0031] Example 6 This embodiment provides a modified fluororubber, the preparation method of which differs from that of Example 2 in that the ratio of the added parts of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and composite filler S@H@C is 100:3:5:3:3:10, and the final modified fluororubber sample is denoted as F / S@H@C-6.

[0032] Comparative Example 1 This comparative example provides a fluororubber, the preparation method of which differs from that of Example 2 in that no composite filler S@H@C is added, and the final fluororubber sample is denoted as F / S@H@C-1.

[0033] Example 7 This embodiment provides a hydrophobically modified multi-scale composite filler. The difference between its preparation method and that of Example 1 is that it also includes step (3): 10 g of S@H@C composite powder is placed in a beaker, 100 mL of anhydrous ethanol is added, and after stirring for 30 min, 2 mL of hexadecyltrimethoxysilane (HDTMS) is added. The beaker is sealed with a sealing film and stirred for 24 h. It is then dried in a 60 ℃ oven, washed three times each with anhydrous ethanol and deionized water, and dried again to obtain the hydrophobically modified multi-scale composite filler, denoted as S@H@C@HD.

[0034] Example 8 This embodiment provides a modified fluororubber, the preparation method of which differs from that of Example 2 in that: the composite filler S@H@C is replaced with an equal amount of the hydrophobic modified multi-scale composite filler S@H@C@HD prepared in Example 7, and the final fluororubber sample is denoted as F / S@H@C@HD-2.

[0035] Example 9 This embodiment provides a modified fluororubber, the preparation method of which differs from that of Example 8 in that the ratio of the added parts of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and hydrophobic modified multi-scale composite filler S@H@C@HD is 100:3:5:3:3:4, and the final modified fluororubber sample is denoted as F / S@H@C@HD-3.

[0036] Example 10 This embodiment provides a modified fluororubber, the preparation method of which differs from that of Example 8 in that the ratio of the added parts of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and hydrophobic modified multi-scale composite filler S@H@C@HD is 100:3:5:3:3:6, and the final modified fluororubber sample is denoted as F / S@H@C@HD-4.

[0037] Example 11 This embodiment provides a modified fluororubber, the preparation method of which differs from that of Example 8 in that the ratio of the added parts of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and hydrophobic modified multi-scale composite filler S@H@C@HD is 100:3:5:3:3:8, and the final modified fluororubber sample is denoted as F / S@H@C@HD-5.

[0038] Example 12 This embodiment provides a modified fluororubber, the preparation method of which differs from that of Example 8 in that the ratio of the added parts of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and hydrophobic modified multi-scale composite filler S@H@C@HD is 100:3:5:3:3:10, and the final modified fluororubber sample is denoted as F / S@H@C@HD-6.

[0039] Comparative Example 2 This comparative example provides a fluororubber, the preparation method of which differs from that of Example 8 in that hydrophobically modified multi-scale composite filler S@H@C@HD is not added, and the final fluororubber sample is denoted as F / S@H@C@HD-1.

[0040] I. Performance Test Results of S@H@C Powder and F / S@H@C Fluororubber Samples 1. SEM analysis of S@H@C powder Figure 1The image shows the SEM morphology of the composite powder obtained by simultaneously introducing 1 g of glass microspheres (HGM) and 1 g of cellulose (CNC) treated with N,N-dimethylacetamide (DMAC) during the Stöber method for SiO2 synthesis. Figure 1 (a) is a partial topographic image. Figure 1 (b) is a magnified view of a single particle. (From...) Figure 1 (a) It can be seen that a relatively obvious spherical particle structure can still be observed in the sample, indicating that after the introduction of HGM and CNC, the system still retains the morphological characteristics of spherical particles as the basic unit. The powder particles have rough surfaces, and a large number of fine particles or flocculent structures are attached to the periphery of the spherical particles. The local area shows a relatively obvious accumulation and coating morphology, indicating that the joint introduction of HGM and CNC significantly changed the nucleation and growth mode of SiO2.

[0041] Depend on Figure 1 (b) It can be further observed that the surfaces of the larger spherical particles are not smooth and dense, but are uniformly or semi-uniformly covered by a large number of fine particles, exhibiting surface features similar to coating or spiky spheres. This indicates that during the Stöber reaction, the SiO2 generated by hydrolysis and condensation not only undergoes homogeneous nucleation in the system, but also preferentially undergoes heterogeneous deposition on the HGM surface and near the CNC-related interfaces, resulting in a large number of fine SiO2 particles continuously attaching to the surface of the spherical matrix, thereby constructing a more complex hierarchical structure. In particular, the significantly enhanced roughness of the spherical surface indicates that SiO2 deposition is sufficient in this system, and there are many active sites on the particle surface.

[0042] Analysis of the experimental conditions revealed that the dispersion of CNCs in the reaction system was improved after DMAC treatment. The abundant hydroxyl groups on their surface provided more binding sites for the hydrolysis and condensation of the SiO2 precursor, thus inducing nucleation or bridging growth to some extent. Simultaneously, HGM, as a spherical hard carrier, provided a stable geometric framework for SiO2 deposition. Therefore, under the combined action of HGM and CNC, the formation mechanism of SiO2 particles in the system gradually transformed from simple homogeneous nucleation to a synergistic process of homogeneous nucleation, heterogeneous deposition, and surface adhesion, ultimately resulting in powders exhibiting more significant surface roughening and hierarchical morphology. This morphological result indirectly indicates that the introduction of CNCs not only did not disrupt the formation of HGM@SiO2 composite particles but also promoted further complexity in the particle surface structure.

[0043] From the perspective of its subsequent application as a fluororubber filler, this composite powder has significant structural advantages. On the one hand, HGM provides a regular spherical skeleton, which is beneficial to maintaining the basic morphological stability of the filler; on the other hand, the rough adhesion layer formed by SiO2 on its surface and the interfacial regulation effect brought about by CNC can significantly improve the specific surface area and surface activity of the particles, thereby increasing the contact area and interfacial interaction sites between the filler and the fluororubber matrix.

[0044] 2. Infrared analysis of S@H@C powder like Figure 2 The infrared spectra of SiO2, HGM, and S@H@C are shown. As can be seen from the figure, pure SiO2 exhibits distinct absorption peaks near 1067 cm⁻¹, 790 cm⁻¹, and 620 cm⁻¹, which are attributed to the stretching vibration of the Si-O-Si bond, the vibration of the Si-O bond, and the bending vibration of the Si-O bond, respectively, indicating the formation of a relatively complete silicon-oxygen network structure in the sample. HGM also shows corresponding characteristic absorption peaks in a similar wavenumber range, indicating that it also possesses a typical silicon-oxygen framework structure. The S@H@C sample shows a new characteristic absorption peak near 960 cm⁻¹, and the peak shape near 1000 cm⁻¹ changes to some extent, indicating that the functional group environment on the surface of the composite particles is altered after the introduction of CNC. At the same time, the sample still retains obvious characteristic peaks of the silicon-oxygen framework, indicating that the introduction of CNC did not destroy the original inorganic phase structure. Therefore, it can be seen that the addition of HGM and the introduction of CNC during the Stöber method for synthesizing SiO2 can successfully construct the S@H@C composite particle structure.

[0045] 3. Dynamic viscoelastic properties analysis of F / S@H@C composite materials like Figure 3 As shown in (a), the storage modulus G′ of the F / S@H@C composite material decreases significantly with increasing strain, exhibiting a typical nonlinear decay characteristic, indicating a significant Payne effect in the system. This phenomenon suggests that the physical network constructed by the composite filler and its interface is relatively complete under small strain conditions. However, as the applied strain increases, the weak interactions between fillers, the interfacial bridging structure, and the local entanglement units are gradually destroyed, leading to a continuous decrease in the elastic storage capacity of the system.

[0046] Compared to F / S@H@C-1, the other data groups exhibit more pronounced hierarchical differentiation. F / S@H@C-1 consistently shows the lowest G′ across the entire strain range, with a low strain end (approximately 10%) of only about 431 kPa and a high strain end (approximately 95%) dropping to about 230 kPa. However, after introducing the S@H@C composite filler, the low strain G′ of the samples significantly increases, with F / S@H@C-2, F / S@H@C-3, and F / S@H@C-4 showing sequential increases. F / S@H@C-5 and F / S@H@C-6 reach their highest values ​​in the low strain region, approximately 785–795 kPa. This indicates that with increasing composite filler content, the system forms a stronger structured framework under small deformation conditions, resulting in a significantly enhanced elastic response.

[0047] Further comparison reveals that this enhancement is not a simple linear increase, but rather exhibits distinct staged characteristics: the G′ increase from F / S@H@C-2 to F / S@H@C-4 is relatively continuous, while F / S@H@C-5 and F / S@H@C-6 show a more significant increase. This indicates that when the composite filler reaches a certain content, it is easier for a continuous load-bearing structure involving both S@H particles and CNC to form within the system. This result demonstrates that the introduction of CNC not only supplements the filler but also promotes synergistic connections between composite fillers through bridging, entanglement, and hydrogen bonding, enabling the system to form a more stable and complete physical network under low strain conditions.

[0048] The differences in ΔG′ between different samples are also quite significant. The ΔG′ of F / S@H@C-1 is only about 201 kPa, while that of F / S@H@C-2, F / S@H@C-3, and F / S@H@C-4 are about 296, 362, and 339 kPa, respectively; the ΔG′ of F / S@H@C-5 and F / S@H@C-6 further increase to about 486 and 494 kPa, respectively. Since ΔG′ is commonly used to measure the sensitivity of the filler network to applied strain, this result indicates that the introduction of CNC significantly enhances the structure and strain sensitivity of the system, especially at higher filler concentrations, resulting in a more developed network structure, but also making it more susceptible to failure under strain.

[0049] In the high-strain region, although the curves showed some convergence, F / S@H@C-1 remained significantly lower than the other samples, and certain differences were also maintained between F / S@H@C-2 to F / S@H@C-6, with F / S@H@C-4 remaining relatively high at the high-strain end. This indicates that under large deformation, although most of the filler-related networks were destroyed, differences still existed in the residual interfacial constraints and structural units between different samples, and the system response had not completely returned to a purely matrix-dominated state.

[0050] like Figure 3As shown in (b), the tanδ of all samples increases monotonically with increasing strain, and the increase is more pronounced in the medium-to-high strain region, indicating that the energy dissipation process in the system continuously increases with the enhancement of cyclic deformation. Corresponding to the variation law of G′, the tanδ of F / S@H@C-1 is always the lowest throughout the entire strain range, increasing from about 0.31 at the low strain end to about 0.56 at the high strain end; while the tanδ of F / S@H@C-2 to F / S@H@C-6 is generally higher, but shows a clear non-monotonic law with changes in formulation.

[0051] Specifically, in the low-strain region, the tanδ differences among the samples are relatively small, all falling within the range of approximately 0.32 to 0.34, indicating that the system primarily stores energy elastically, with energy dissipation processes not yet prominent. As the strain increases, the differences between the samples gradually become amplified. At the high-strain end (approximately 95%), the tanδ values ​​for F / S@H@C-2, F / S@H@C-3, F / S@H@C-4, F / S@H@C-5, and F / S@H@C-6 are approximately 0.636, 0.692, 0.801, 0.794, and 0.655, respectively. It can be seen that F / S@H@C-4 and F / S@H@C-5 have the highest tanδ values ​​in the medium-to-high strain region, while F / S@H@C-6 shows a significant decrease. This indicates that there is a relatively optimal energy consumption range in this group of samples: at medium to high filler content, energy-consuming processes such as interfacial micro-slippage, structural rearrangement, and chain segment friction are the most active; while as the filler content continues to increase, the system gradually shifts to a denser and more confined structural state, or the effective interfacial effect decreases due to local agglomeration, thus causing tanδ to no longer continue to increase, but instead decrease.

[0052] Combining G′ and tanδ, it can be seen that the F / S@H@C composite material exhibits structural characteristics and content sensitivity. With increasing strain, G′ continuously decreases while tanδ continuously increases, reflecting a typical coupling law of elastic skeleton attenuation and energy dissipation enhancement. Specifically, the significant stratification of G′ and the marked increase in ΔG′ indicate that CNC introduced additional structural units into this ternary system, significantly enhancing the filler network and its strain response capability. The initial increase followed by a decrease in tanδ suggests that more filler is not necessarily more beneficial; rather, there exists a more optimal structural rearrangement and interfacial energy dissipation balance range within a certain range.

[0053] 4. Vulcanization properties of F / S@H@C composite materials like Figure 4 As shown, to investigate the effect of S@H@C composite filler on the vulcanization behavior of FKM composite materials, the minimum torque M of the F / S@H@C samples prepared in Comparative Example 1 and Examples 2-6 was compared. L Maximum torque M H Scorching time T 10 and positive vulcanization time T 90A comparative analysis was conducted.

[0054] From minimum torque M L Let's look at the M of the F / S@H@C sample. L The dN·m of F / S@H@C-1 increased significantly from 0.50 dN·m to 0.88, 1.07, 1.36, 1.44, and 1.02 dN·m, showing a trend of first increasing and then decreasing. This indicates that the introduction of CNC further enhances the interaction between fillers and between fillers and rubber, significantly increasing the internal friction and processing resistance of the rubber compound in the uncured stage; especially at medium and high filler concentrations, the restriction effect on molecular chain movement is stronger, indicating that it is easier to form a composite filler skeleton with a higher degree of structure.

[0055] From the maximum torque M H Let's look at the M of the F / S@H@C sample. H The dN·m values ​​reached 17.3, 18.1, 20.9, 22.6, and 21.3 dN·m, with a maximum value of 22.6 dN·m in F / S@H@C-5, indicating that the multi-component composite filler system constructed from CNC, HGM, and SiO2 has a stronger synergistic reinforcing effect. Further comparison of ΔM(M H -M L It can be observed that the F / S@H@C samples further increased to 16.42, 17.03, 19.54, 21.16, and 20.28 dN·m. This result indicates that the F / S@H@C system is extremely effective in increasing crosslinking density, enhancing interfacial bonding, and constructing a dense vulcanization network. Although M... H The value has slightly decreased, but remains at a relatively high level, indicating that it is more conducive to forming a more rigid and structurally complete vulcanization network under medium to high filler content.

[0056] From scorching time T 10 Let's look at the T of the F / S@H@C sample. 10 The time intervals were 2.03, 2.14, 2.00, 1.95, 1.94, and 1.92 min, respectively. It can be seen that, compared to the blank sample, the F / S@H@C sample showed better T values ​​at low filler content. 10 The time was slightly extended, but as the amount of filler increased further, T 10 The temperature gradually decreases, and in the high-filler region it is lower than the first two groups. This indicates that after the introduction of CNC, the system still exhibits certain adsorption of active components or structural hindrance effects at lower filler concentrations. However, at higher filler concentrations, the enhanced interfacial interaction and increased system activity accelerate the initial vulcanization reaction, slightly narrowing the processing safety window. However, in absolute terms, T... 10 The time remained around 2 minutes, indicating that there was still some room for processing and no obvious early scorching problem.

[0057] From the positive sulfurization time T 90 The F / S@H@C samples showed values ​​of 3.07, 3.65, 3.36, 3.30, 3.26, and 3.24 min, which gradually decreased with increasing filler content. This indicates that although the F / S@H@C system experienced accelerated vulcanization in the later stages under high filler content, the vulcanization rate was still relatively fast.

[0058] 5. Mechanical properties of F / S@H@C composite materials like Figure 5 As shown in (a), the tensile strengths of F / S@H@C-1 to F / S@H@C-6 are 6.0, 15.0, 14.3, 14.2, 13.9, and 13.3 MPa, respectively, showing an overall trend of first increasing sharply and then slowly decreasing. Compared with F / S@H@C-1, the tensile strengths of F / S@H@C-2, F / S@H@C-3, F / S@H@C-4, F / S@H@C-5, and F / S@H@C-6 are increased by 150.0%, 138.3%, 136.7%, 131.7%, and 121.7%, respectively, with F / S@H@C-2 reaching the maximum value of 15.0 MPa, showing the most significant strengthening effect. This indicates that after CNC modification, the interfacial compatibility between S@H and the FKM matrix is ​​significantly improved. The introduction of nanocellulose enhances the bonding between the filler and rubber, and stress can be effectively transferred between the two phases of filler and rubber, thereby greatly improving the load-bearing capacity. However, when CNC is further increased to 8 phr, the tensile strength drops, which is presumably related to the tendency of nanocellulose to agglomerate in the rubber matrix under high filling conditions, the decrease in interfacial continuity, and stress concentration.

[0059] like Figure 5 As shown in (b), the elongation at break of F / S@H@C-1 to F / S@H@C-6 were 143.2%, 199%, 189%, 159%, 148%, and 142%, respectively. Compared with F / S@H@C-1, F / S@H@C-2 and F / S@H@C-3 showed increases of 38.9% and 32.0%, respectively, with F / S@H@C-2 reaching the maximum value of 199%. This indicates that the appropriate introduction of CNC effectively improved the filler dispersion state of the nanocellulose network structure, resulting in more complete molecular chain orientation and slippage during stretching, which is beneficial to toughness improvement. However, when the CNC increased from 4 phr to 10 phr, the elongation at break continued to decrease, with F / S@H@C-6 showing a 28.6% decrease compared to F / S@H@C-2. This indicates that with the increase of the rigidity ratio, the interaction between nanocellulose and rubber molecular chains is enhanced, the degree of freedom of molecular chain movement is restricted, and the plastic deformation capacity of the material is weakened. These results show that there is a clear balance range between reinforcement and toughening in this system.

[0060] like Figure 5As shown in (c), with the increase of CNC dosage, the hardness of F / S@H@C-1 to F / S@H@C-6 gradually increased from 70 HA to 72.5 HA. The hardness of the composites F / S@H@C-2, F / S@H@C-3, F / S@H@C-4, F / S@H@C-5, and F / S@H@C-6 increased by 0.7%, 0.7%, 3.6%, 1.4%, and 3.6% respectively relative to F / S@H@C-1, showing a slow upward trend overall. The continuous increase in hardness indicates that the rigid constraint effect and network structure effect of CNC are continuously enhanced, and the material's resistance to deformation is improved. This also corresponds to the aforementioned decrease in elongation at break in the later stage, that is, the increase in system rigidity is usually accompanied by a decrease in deformation capacity. It is worth noting that the effect of CNC dosage on hardness is relatively mild, with the increase controlled within 5%, indicating that the introduction of nanocellulose improves the material rigidity while causing little damage to the original rubber elastic characteristics.

[0061] Based on the combined tensile strength, elongation at break, and hardness indicators, F / S@H@C-2 shows the best performance in both strength and toughness. F / S@H@C-3 maintains high strength while also possessing good toughness. F / S@H@C-3 achieves a good overall balance between strength, toughness, and hardness: its tensile strength reaches 14.3 MPa, which is 138.3% higher than the benchmark, while maintaining a high elongation at break of 189% and possessing a hardness of 70.5 HA.

[0062] 6. Cross-sectional analysis of F / S@H@C composite materials Figure 6 The image shows a cross-sectional SEM image of a fluororubber composite material filled with a composite powder obtained by simultaneously introducing 1 g of glass microspheres (HGM) and 1 g of DMAC-treated cellulose (CNC) during the Stöber method for SiO2 synthesis. The filler addition amount was 8 g of composite powder added to 100 g of fluororubber. Figure 6 (a) and Figure 6 (b) It can be seen that the composite material cross-section exhibits a relatively obvious rough and undulating morphology. The fracture surface is not smooth and flat, but is accompanied by many tear textures, micropores, and local particle embedding traces. This indicates that the addition of filler significantly changed the fracture mode of the fluororubber matrix, causing the crack to be hindered and deflected to a certain extent during the propagation process. From the overall cross-sectional distribution, the filler has a certain degree of dispersion in the matrix, and no large-area continuous severe agglomeration regions were observed. This indicates that after the composite powder was treated with HGM and DMAC in CNC, its distribution in the fluororubber was improved compared with the previous single SiO2 system.

[0063] Further observation revealed numerous small, circular or near-circular particle traces and micropores formed after particle detachment in the cross-section. This indicates that the composite filler was sufficiently embedded in the fluororubber matrix and participated in the load-bearing and deformation processes during material fracture. Some areas exhibited clearly defined spherical structures or circular cavities, presumably originating from the pull-out of HGM-related composite particles or interfacial debonding during fracture. Simultaneously, no particularly large agglomerates were observed in the cross-section, suggesting that the combined effect of HGM as a spherical framework and CNC as a structure-inducing component mitigated the tendency of pure SiO2 particles to aggregate tightly, resulting in a more dispersed spatial distribution of the filler within the matrix. This positively contributes to improving the uniformity of the composite material's internal microstructure.

[0064] From the fracture morphology, the fracture surface exhibits obvious tearing ridges, undulating wrinkles, and irregular fracture steps, indicating that the material underwent a relatively complex deformation and energy dissipation process before fracture, rather than a simple brittle fracture. Generally, such a rougher fracture surface means that the crack propagation path becomes more tortuous, and the filler particles and their interface structure play a role in hindering and deflecting crack propagation. Especially when the composite powder surface has a rough hierarchical structure formed by SiO2 deposition, the filler and fluororubber matrix are more likely to form a certain degree of mechanical interlocking, thereby enhancing interfacial friction and adhesion, requiring more energy to propagate the crack near the interface. Therefore, from the perspective of fracture morphology, this composite structure, to a certain extent, is beneficial to improving the composite material's resistance to fracture.

[0065] Based on the aforementioned powder SEM results, this filler group exhibits a composite structure characterized by a spherical skeleton and a rough surface adhesion layer. When introduced into the fluororubber matrix, this structure helps reduce the severe agglomeration of pure nano-SiO2 caused by its high surface energy. Furthermore, the rough surface increases the contact area between the filler and the rubber matrix, thereby increasing the number of interfacial interaction sites. Therefore, although localized pores and interfacial debonding traces are still observed in the composite material cross-section, the overall cross-sectional structure is relatively uniform, indicating that the filler system achieves a certain balance between dispersibility and interfacial bonding. This microstructural feature typically facilitates stress transfer between the filler and the matrix and delays the initiation and propagation of microcracks under external forces.

[0066] Furthermore, the introduction of DMAC-treated CNC also had a positive impact on the cross-sectional morphology. The polar groups on the CNC surface can participate in the deposition and connection of SiO2 during the formation of the composite powder, making the surface structure of the resulting filler more complex and providing more interfacial interaction sites after entering the fluororubber system. The absence of obvious large brittle peeling areas in the cross-sectional image indicates that the addition of CNC helps to improve the stability of the filler structure and the microscopic uniformity of the composite system to some extent. However, the presence of micropores and particle pull-out marks in some areas also indicates that even with a relatively high filler content of 8 g, there is still room for further improvement in the interfacial bonding between some particles and the matrix, which can be further optimized through surface modification or interfacial control.

[0067] In summary, after S@C@H composite powder was used to fill fluororubber, the composite material exhibited a relatively rough tear morphology and a relatively uniform particle distribution, without significant large-area severe agglomeration, indicating that the composite filler could be well dispersed in the fluororubber matrix. The particle embedding, micropores, and tortuous crack paths in the cross-section indicate that the filler participated in the material's fracture process and played a certain role in hindering crack propagation. This demonstrates that the composite filler constructed by introducing HGM and DMAC-treated CNC has a positive effect on improving the filler dispersion and optimizing the interfacial structure, providing a good microstructural basis for subsequently improving the mechanical properties of fluororubber composites.

[0068] 7. Surface wettability of F / S@H@C composite material like Figure 7 The figure shows the contact angle of the F / S@H@C composite material. As can be seen from the figure, with the increase of S@H@C nanofiller content from 0 phr to 10 phr, the water contact angle of the F / S@H@C-1 to F / S@H@C-6 composite materials gradually increased from 90.1° to 99.7°, showing a monotonically increasing trend, indicating a continuous enhancement of the material surface hydrophobicity. The contact angles of the samples were 90.1°, 90.8°, 92.4°, 94.7°, 95.3°, and 99.7°, respectively, representing increases of 0.78%, 2.55%, 5.11%, 5.77%, and 10.65% compared to the unfilled sample (F / S@H@C-1). The variation pattern shows that the contact angle of the composite material increases relatively steadily in the 0–4 phr range with S@H@C dosage, accelerates in the 4–8 phr range with S@H@C dosage, and shows a significant increase (up to 4.6°) in the 8–10 phr range with S@H@C dosage. This indicates that the introduction of CNC components has a more significant effect on enhancing surface hydrophobicity at higher dosages.

[0069] On the one hand, the introduction of S@H@C alters the near-surface phase composition of the composite material, making the synergistic effect of the low surface energy fluororubber matrix, inorganic phase, and cellulose nanoparticles (CNC) more significant. The hydroxyl-rich surface of CNC further reduces the wetting driving force of water droplets on the solid surface after modification. On the other hand, SiO2 coating of glass microspheres and composite with CNC construct a more complex micro-nano multi-scale rough structure. Especially when the filler content is high, the multi-component synergistic roughness effect is more prominent, thereby significantly improving the apparent contact angle.

[0070] 8. Thermal conductivity of F / S@H@C composite materials Figure 8 The results show the thermal response behavior of the F / S@H@C composite material. Figure 8 (a) The gradient change of color from cold to hot visually reflects the temperature distribution on the sample surface and its evolution over time. Figure 8 (b) Heating process: The heating plate temperature was set to 100 °C. The sample was placed on the hot plate with room temperature (about 20 °C) as the starting point and thermal imaging surface temperature data was collected every 10 seconds. In order to ensure the consistency of the formulation comparison, this study used the sample surface temperature reaching about 80 °C as the criterion for termination of the heating stage (not heating to complete thermal equilibrium with the 100 °C hot plate). The heating data show that the heating of the samples after the introduction of CNC exhibits a non-monotonic pattern of first enhancing and then limiting: low to medium filling can significantly improve the surface heating rate, with 2 phr reaching 34.28 ℃ in 10 s (significantly higher than 25.9 ℃ in 0 phr), and 4 phr showing the best overall heating performance, reaching 81.3 ℃ in 100 s and reaching the 80 ℃ threshold in about 93.5 s; at the same time, the characteristic heating time is significantly shortened, for example, the time required to reach 70 ℃ is shortened from about 66.0 s in 0 phr to about 57.7 s in 4 phr, and the time required to reach 75 ℃ is shortened from about 85.9 s to about 71.9 s. In contrast, the heating advantage weakened when the filler content was further increased to 6–10 phr. At 100 s, the temperatures at 6 phr, 8 phr, and 10 phr were 79.9 ℃, 78.2 ℃, and 77.6 ℃, respectively. Among them, the temperature rise at 10 phr was the slowest throughout the entire process (e.g., only 66.8 ℃ at 60 s). This indicates that the increased proportion of low thermal conductivity phases introduced by the hollow structure of HGM under high filler content and the cumulative interfacial thermal resistance caused by the increased number of multiphase interfaces will inhibit thermal diffusion, thereby limiting the transfer of heat from the heat source side to the sample surface. In the 2–4 phr range, CNC may promote the formation of more efficient heat transfer pathways by improving the deposition uniformity of SiO2 on the HGM surface, enhancing filler dispersion and interfacial bonding, thus exhibiting a faster heating response. Figure 8(c) Cooling process: After the sample was moved from the hot plate to the room temperature table (starting from about 80 ℃), the surface temperature decay was recorded every 20 s. The results showed that as the amount of S@H@C increased, the surface temperature of the sample was higher and the cooling was slower for the same cooling time. For example, at 160 s, the temperature dropped to 25.3 ℃ at 0 phr, while it was 28.5 ℃, 30.8 ℃ and 29.8 ℃ at 6 phr, 8 phr and 10 phr, respectively. Taking the cooling to 30 ℃ as the characteristic point, the temperature dropped to 30 ℃ at 0 phr in about 119 s, it took about 148 s for 6 phr, and it took about 159 s for 10 phr. However, the temperature did not drop to 30 ℃ (30.8 ℃) at 160 s, showing a clear heat insulation trend. The combined heating and cooling results show that CNC-involved S@H@C composite fillers are more conducive to improving thermal response (faster heating and earlier reaching 80 °C) in low to medium filling (especially F / S@H@C-3), while high filling (≥6 phr) is more likely to exhibit characteristics of limited thermal diffusion and suppressed heat dissipation. This behavior can be attributed to the competitive synergistic effect between CNC promoting dispersion and pathway construction and the low thermal conductivity effect of HGM hollow core + interface thermal resistance accumulation.

[0071] 9. Thermal stability of F / S@H@C composite materials Figure 9 Tables (a) to (c) present the TG curves, DTG curves, and thermogravimetric parameters of composite materials with different filler contents (F / S@H@C-1~6 corresponding to 0, 2, 4, 6, 8, and 10 phr, respectively). Table 1 summarizes the T5% and T2 values ​​for each sample. 10 %, T 50 Key indicators such as % and residual carbon rate were analyzed. The TG curves show that the mass changes of each formulation are relatively small in the low-temperature region (approximately 30–400 °C), with the main weight loss process concentrated in the range of approximately 430–510 °C, indicating that the thermal decomposition of the system is mainly a medium-to-high temperature degradation stage. With the introduction of the S@H@C composite filler, the main weight loss range shifts towards the high-temperature side, and the downward segment of the curve becomes more concentrated, indicating that the composite filler has a certain blocking and delaying effect on the thermal decomposition process. The early thermal stability (T5%, T...) is analyzed in conjunction with Table 1. 10 It can be observed that the composite filler significantly increases the initial weight loss temperature at low addition levels: 0 phr sample (F / S@H@C-1) T5%, T 10 The percentages were 426.1 °C and 450.5 °C, respectively. Adding 2 phr (F / S@H@C-2) increased these to 439.3 °C and 457.8 °C, representing increases of 13.2 °C and 7.3 °C compared to 0 phr. Adding 6 phr (F / S@H@C-4) further increased T5% and T... 10 The percentage remains at a high level (435.4 ℃ and 454.4 ℃).

[0072] The results indicate that an appropriate amount of composite filler can enhance interfacial constraint and physical shielding effects through the synergistic effect of the inorganic framework (SiO2 / HGM) + organic fiber network (CNC), reduce the heat transfer rate, and inhibit the diffusion of decomposed small molecules, thereby delaying early weight loss; however, when the addition amount is further increased to 8~10 phr (F / S@H@C-5, 6), T5% and T 10 The percentage of thermal stability index (T0) decreased to 418.4 ℃, 437.7 ℃, and 416.5 ℃ and 436.2 ℃, respectively, lower than the 0 phr sample. This is presumably related to filler agglomeration, interface defects, and local structural inhomogeneity under high-filling conditions. Additionally, CNC, as an organic component, may be more likely to induce early weight loss signals at high addition levels, thus advancing the initial decomposition temperature. In contrast, the mid-to-late-stage thermal stability index (T0) showed a higher T0. 50 % shows a more stable upward trend: 0 phr sample T 50 The maximum weight loss rate was 477.1 °C. After the addition of composite filler, the maximum weight loss rate of all samples increased to 482.6-486.1 °C, with the 6 phr sample showing the highest rate (486.1 °C), an increase of 9.0 °C compared to the 0 phr sample. This indicates that the composite filler can continuously play a role in blocking mass transfer and inhibiting rapid chain segment breakage during the main degradation stage. From the DTG curves, the peak of the maximum weight loss rate of each sample was mainly located around 470-490 °C. After the composite filler was added, the peak position was generally biased towards the high temperature region and the peak shape difference was reduced, reflecting that the main degradation process was delayed and tended to be concentrated. The residual carbon rate further verified the synergistic effect of carbonization / slag formation of the composite filler: the residual carbon rate of the 0 phr sample was 15.30%, which increased significantly with the addition amount to 2~6 phr (16.62%~17.42%), with the 6 phr sample reaching the highest of 17.42% (an increase of 2.12 percentage points), indicating that at this addition amount, it is easier to form a continuous and dense pyrolysis residual layer (the SiO2 / HGM inorganic framework and CNC carbonization residue synergistically construct a shielding layer); when the filler was increased to 10 phr, the residual carbon rate dropped back to 15.25%, indicating that excessive filler may destroy the continuity of the residual layer and weaken the degree of densification. In summary, the S@H@C composite filler has a "dosage window" for improving the thermal stability of FKM: 2 phr is more conducive to significantly increasing the early decomposition temperature (T5%, T...). 10 %), while 6 phr shows stability in the mid-to-late stages (T 50 The best performance in terms of percentage and residual carbon rate indicates that proper dispersion and interfacial synergy of composite fillers are the key to obtaining the best thermal stability enhancement effect.

[0073] Table 1 Thermogravimetric data of F / S@H@C composite materials II. Performance Test Results of S@H@C Powder and F / S@H@C Fluororubber Samples 1. SEM analysis of S@H@C@HD powder Figure 10 The images show the SEM morphology of the S@C@HD composite powder after alcohol modification with hexadecyltrimethoxysilane (HDTMS). Figure 10 (a) is a partial topographic image. Figure 10 (b) is a magnified view of a single particle. (From...) Figure 10 (a) It can be seen that after HDTMS treatment, the sample as a whole still maintains a relatively obvious spherical particle structure, with a clear outline of the main particle body and no obvious collapse or severe breakage. This indicates that the alcoholysis modification process did not significantly damage the main morphology of the original composite powder, and the structural characteristics of HGM as a spherical framework were well preserved. This shows that after surface treatment with anhydrous ethanol as the dispersion medium and HDTMS as the modifier, the composite powder achieved further surface regulation while maintaining the original spherical structure.

[0074] Depend on Figure 10 (b) Further observation reveals that the particle surface still exhibits significant roughening characteristics, with numerous fine particles or granular protrusions adhering to the surface, demonstrating a strong hierarchical structure. Simultaneously, a certain degree of flocculent or coating-like adhesion layer can be observed at the particle edge region, causing the particle surface to no longer appear as a simple accumulation of inorganic particles, but rather tending to form a relatively continuous rough outer layer. Compared to the unmodified S@C@H composite powder, this sample still maintains the typical composite morphology of a spherical skeleton-rough surface adhesion layer, but its surface adhesion layer appears denser, with more blurred local boundaries. This indicates that after HDTMS treatment, an organosilane low surface energy layer was formed on the particle surface, or that the originally loosely attached fine particles were more tightly bound, thus stabilizing the surface structure.

[0075] Based on the modification mechanism analysis of HDTMS, HDTMS can generate an intermediate structure containing silanol groups after a certain degree of hydrolysis in an alcohol medium, and further undergo a condensation reaction with hydroxyl groups on the powder surface. For this system, the HGM, SiO2, and CNC surfaces may all provide a certain number of reactive sites. Therefore, HDTMS may be able to construct an organosilane layer on the composite powder surface, allowing long-chain alkyl groups to be directionally distributed on the outer surface of the particles, thereby reducing the particle surface energy and improving its hydrophobicity. Although it is difficult to directly observe the organosilane molecular layer itself from SEM morphology, the surface adhesion layer of the particles remains relatively intact, and the composite structure is not destroyed. This indicates that the HDTMS modification process has good mildness and feasibility, providing a morphological basis for subsequent improvement of filler surface properties.

[0076] Furthermore, the figures also reveal a degree of contact and localized adhesion between some particles, indicating that the composite powder still exhibits a certain tendency to aggregate after drying. However, compared to the unmodified sample, this aggregation is more characterized by surface contact connections rather than irregular, large-area hard agglomerations. Considering the modification objective, the introduction of HDTMS does not necessarily eliminate particle contact entirely; more importantly, it suppresses the tendency for secondary agglomeration during subsequent rubber processing by reducing surface polarity and weakening the strong interactions between surface hydroxyl groups. Therefore, these morphological results suggest that HDTMS modification has the potential to improve the dispersion state of the composite powder in the organic matrix and enhance its compatibility with fluororubber.

[0077] From the perspective of its subsequent application as a filler in fluororubber, the S@C@HD composite powder modified with HDTMS possesses the advantages of a spherical skeleton structure, rough surface characteristics, and low surface energy regulation. On the one hand, the spherical HGM skeleton helps maintain the geometric stability of the filler and alleviates sharp stress concentration; on the other hand, the rough surface adhesion layer can provide more interfacial contact sites, while the long-chain alkyl groups introduced by HDTMS help reduce the hydrophilicity of the filler surface and improve its wetting and dispersion behavior in fluororubber. Compared with the unmodified S@C@HD composite powder, this sample, while maintaining the hierarchical structure characteristics, is expected to further enhance the interfacial compatibility between the filler and the rubber matrix, thereby providing more favorable microstructural conditions for the subsequent improvement of the mechanical and interfacial properties of the composite material.

[0078] 2. S@H@C@HD powder infrared analysis like Figure 2 As shown, after further alcoholic modification of S@H@C with HDTMS, the characteristic absorption peaks of the S@H@C@HD sample still exist, and the peak shape near 1000 cm⁻¹ further changes, indicating that HDTMS has successfully reacted with the powder surface, achieving surface organic modification without destroying the original silicon-oxygen framework structure. Infrared analysis results show that HDTMS was successfully grafted onto the surface of S@H@C composite powder, thus obtaining S@H@C@HD modified powder.

[0079] 3. Dynamic viscoelastic properties analysis of F / S@H@C@HD composite materials like Figure 11 As shown in (a), the storage modulus G′ of the F / S@H@C@HD samples decreases significantly with increasing strain, exhibiting typical nonlinear decay characteristics, indicating a significant Payne effect in the system. This suggests that the physical network constructed by the composite filler and its interface is relatively complete under small strain conditions. However, with increasing applied strain, the weak interactions between fillers, the interface connection structure, and the local entanglement units are gradually destroyed, leading to a continuous decrease in the elastic storage capacity of the system.

[0080] From the response patterns of different formulations, the F / S@H@C@HD samples also exhibited clear stratification characteristics. F / S@H@C@HD-1 consistently showed the lowest G′ across the entire strain range, reaching only about 431 kPa at the low strain end (approximately 10%). After introducing the F / S@H@C@HD composite filler, the low strain G′ significantly increased, and F / S@H@C@HD-2 to F / S@H@C@HD-6 showed a basically monotonically increasing trend, gradually increasing from approximately 477 kPa to 642 kPa. This indicates that with the increase in the amount of composite powder, the system formed a stronger structured framework under small deformation conditions, and the contribution of the filler-related network to the elastic response continuously increased. At the high strain end (approximately 95%), the G′ of F / S@H@C@HD-2 to F / S@H@C@HD-6 gradually converged to approximately 254–272 kPa, while F / S@H@C@HD-1 remained at a relatively low level (approximately 205 kPa). This indicates that under large deformation conditions, most of the filler-related networks have been destroyed, and the dynamic response of the material is more dominated by the constraints of the FKM matrix and the residual interface, but certain structural differences are still retained between the filled samples.

[0081] Further analysis of ΔG′ reveals that the ΔG′ of F / S@H@C@HD-2 to F / S@H@C@HD-6 gradually increases from approximately 223 kPa to 372 kPa, indicating that the strain sensitivity of the filler network increases with the increase of filler content. Compared with the F / S@H@C system without HDTMS modification, F / S@H@C@HD exhibits lower overall low strain G′ and ΔG′ at the same filler content. For example, at 10 phr, the ΔG′ of F / S@H@C@HD is approximately 372 kPa, while that of F / S@H@C is approximately 494 kPa. This demonstrates that HDTMS modification weakens the strong interactions and interconnected network construction capabilities between fillers to a certain extent, thereby suppressing the Payne effect and reducing the structural sensitivity of the system. This result suggests that the hydrophobic long chains introduced by HDTMS have a certain shielding effect on the polar sites of the SiO2@CNC surface, weakening the strong networking tendency brought about by hydrogen bonding or polar overlap.

[0082] like Figure 11 As shown in (b), the tanδ of all samples monotonically increases with increasing strain, with a more pronounced increase in the medium-to-high strain region, indicating that the energy dissipation process in the system continuously increases with increasing cyclic deformation. It should be noted that tanδ is more suitable for reflecting the comprehensive changes in energy-consuming behaviors such as interfacial micro-slip, structural rearrangement, and chain segment friction in the system, and should not be simply equated with the single factor of filler dispersion quality. Figure 11(b) The data presented show that F / S@H@C@HD-1 consistently exhibits the lowest tanδ across the entire strain range, increasing from approximately 0.29 at the low strain end to approximately 0.61 at the high strain end. The other filled samples show relatively small differences in the low strain region, with tanδ generally ranging from 0.32 to 0.34, indicating that the energy consumption differences among the formulations are not significant under small deformation conditions. As the strain further increases, the differences between the different formulations gradually become amplified. At the high strain end, the tanδ of F / S@H@C@HD-2 to F / S@H@C@HD-6 are approximately 0.638, 0.680, 0.702, 0.738, and 0.745, respectively, showing a relatively obvious monotonic upward trend.

[0083] This result indicates that in the F / S@H@C@HD system, with the increase of composite powder content, the interfacial micro-slip, filler structure rearrangement, and frictional energy dissipation processes under large deformation conditions are continuously enhanced, thus leading to a continuous increase in tanδ. Compared with the unmodified F / S@H@C system, F / S@H@C@HD did not show a significant decline in high filler content in the high filler region, but rather exhibited a more monotonically increasing trend. This suggests that although the strong filler-filler network effect is weakened and the Payne effect is reduced under small strain after HDTMS alcohol modification, the system can still maintain sufficient interfacial slip and structural rearrangement under large deformation conditions, thereby maintaining a high hysteresis loss capacity. This implies that HDTMS modification did not create an overly hard domain in the system, but rather improved the continuity of energy dissipation under high filler content to some extent.

[0084] 4. Vulcanization properties of F / S@H@C@HD composite materials The vulcanization properties of the F / S@H@C@HD composite material were tested, mainly focusing on the minimum torque M. L Maximum torque M H Scorching time T 10 and positive vulcanization time T 90 Parameters such as these.

[0085] Depend on Figure 12 It can be seen that the M of the F / S@H@C@HD sample L The temperature rises significantly after the introduction of packing material. Taking the reference sample F / S@H@C@HD-1 as an example, its M... L The concentration is 0.50 dN·m; when HDTMS modified composite filler is added, the M of F / S@H@C@HD-2 to F / S@H@C@HD-6 is... LThe viscosity values ​​were 0.70, 0.79, 0.89, 0.90, and 0.93 dN·m, respectively, showing a gradual upward trend. This indicates that as the amount of HDTMS-modified composite filler increases, the viscosity of the uncured rubber gradually increases, the molecular chain movement within the system is more strongly restricted, and the flowability of the rubber decreases accordingly. This shows that the modified composite filler can still enhance the interaction between fillers and between fillers and rubber, and improve the internal friction level of the system. However, compared with the third group of unmodified S@H@C systems, the M value of the fourth group at the same filler dosage was lower. L Overall, the levels decreased significantly, for example, in the range of 2–10 phr, in the third group M. L The values ​​ranged from 0.88 to 1.44 dN·m, while the fourth group only ranged from 0.70 to 0.93 dN·m. This indicates that after HDTMS surface modification, the number of polar hydroxyl groups on the filler surface decreased, and the excessively strong polar association between fillers was weakened, thereby reducing the excessive restriction of the network structure on the flowability of the compound and improving the processing performance of the system.

[0086] From M H As can be seen, the M values ​​of all samples after the addition of modified fillers were significantly higher than those of the reference sample. F / S@H@C@HD-1 H The M value was 9.81 dN·m, while those of F / S@H@C@HD-2 to F / S@H@C@HD-6 were 16.4, 18.1, 17.5, 18.0, and 17.6 dN·m, respectively, showing an overall trend of first increasing and then fluctuating slightly. Among them, the M value of F / S@H@C@HD-3 was... H The highest value was 18.1 dN·m, followed by F / S@H@C@HD-5 at 18.0 dN·m, indicating that under medium filler content, modified composite fillers are more conducive to forming a relatively dense and stable vulcanization network structure. Generally speaking, M... H The crosslinking density, network strength, and reinforcing effect of the filler are closely related to the crosslinking density, network strength, and filler reinforcement of the vulcanized rubber. Therefore, this result indicates that the HDTMS-modified composite filler can still significantly reinforce the FKM system, but its reinforcement magnitude is weaker than that of the unmodified S@H@C system. Further analysis of the torque difference ΔM(M) H -M L It can be seen that the molecular weights (M) of F / S@H@C@HD-2 to F / S@H@C@HD-6 are approximately 15.70, 17.31, 16.61, 17.10, and 16.67 dN·m, respectively. Among them, F / S@H@C@HD-3 and F / S@H@C@HD-5 are relatively higher, indicating that these two formulations perform better in terms of network building ability and reinforcement effect. Compared with the F / S@H@C system, the M of the F / S@H@C@HD system samples is... HBoth ΔM and ΔM decreased, indicating that while HDTMS improved the dispersibility and interfacial compatibility of the filler, it also weakened the strong polar interaction between CNC, HGM and SiO2, thus reducing the filler's ability to build a high-rigidity network.

[0087] By T 10 and T 90 As shown in the figure, the T of the F / S@H@C@HD system 10 The overall change was relatively small. The T values ​​for F / S@H@C@HD-1 to F / S@H@C@HD-6 10 The time intervals were 2.03, 2.11, 2.00, 2.04, 2.05, and 2.05 min, respectively, remaining generally stable within the range of 2.00–2.11 min. This indicates that the HDTMS-modified composite filler has a relatively weak impact on the initial vulcanization reaction of FKM rubber compound, and the system still exhibits good processing safety. Compared to the F / S@H@C system, the T1 of the F / S@H@C@HD system at medium to high filler dosages is lower. 10 The duration was slightly extended, for example, from 1.95–1.92 min in the third group to 2.04–2.05 min in the fourth group within the range of 6–10 phr. This indicates that HDTMS modification reduced the disturbance of the filler surface on the active components of the vulcanization system to a certain extent, giving the rubber compound a more stable operating window during the processing stage.

[0088] From T 90 The time intervals (Tg) for F / S@H@C@HD-1 to F / S@H@C@HD-6 are 3.07, 3.41, 3.04, 3.07, 3.08, and 3.11 min, respectively. This indicates that the Tg values ​​for the F / S@H@C@HD samples are... 90 Instead of simply and monotonously shortening with the addition of packing material, it exhibits a characteristic of first increasing at low packing material dosages, then decreasing in the medium to high packing material range and tending to stabilize. Among them, the T of F / S@H@C@HD-2 90 The highest value was 3.41 min, indicating that even at lower filler dosages, the modified composite filler may still have a certain adsorption or retardation effect on the accelerator or crosslinking components, thus delaying the system from reaching the optimal vulcanization state; however, when the filler dosage was further increased, T... 90 The time was significantly reduced and stabilized between 3.04 and 3.11 min, indicating that the vulcanization process of the system was more coordinated under medium to high filler content. Further analysis of T... 90 -T 10 The differences show that the time differences for F / S@H@C@HD-2 to F / S@H@C@HD-6 are approximately 1.30, 1.04, 1.03, 1.03, and 1.06 min, respectively. Except for the low-filler sample, the other formulations remain within a relatively small range. Compared with the unmodified F / S@H@C system, the Ti of the corresponding F / S@H@C@HD sample is...90 -T 10 The overall difference decreased significantly, indicating that the reaction process of the HDTMS-modified system was more compact in the main vulcanization stage, making it easier to complete effective crosslinking in a shorter time. This may be because HDTMS modification reduced the surface polarity and number of active sites of the filler, reducing the adsorption of accelerators and crosslinking components by the filler, thus allowing the vulcanization reaction to proceed more smoothly.

[0089] In summary, the HDTMS alcohol-modified S@H@C composite filler has a significant regulatory effect on the vulcanization behavior of FKM composite materials. On the one hand, M L The value gradually increased with increasing filler content, indicating that the modified composite filler can still enhance the friction and interfacial interaction within the system; on the other hand, M H Maintaining a high level, and reaching optimal values ​​at medium filler concentrations, indicates that it still possesses good network reinforcement capabilities. Meanwhile, T 10 Overall stability indicates that the system maintains good processing safety; T 90 This indicates that the low-filler-content region tends to be higher, while the medium-high-filler-content region tends to shorten and stabilize, suggesting that HDTMS modification helps improve the main vulcanization process of the system. Compared with the unmodified composite filler, the F / S@H@C@HD sample exhibits a lower M L A moderate M H A relatively stable T 10 The shorter overall main vulcanization process indicates that HDTMS surface modification effectively weakens the excessive inter-filler interactions, improves filler dispersion and interfacial compatibility, and achieves a better balance between processing performance and network reinforcement capability in the composite material. Comprehensive analysis of various parameters shows that F / S@H@C@HD-3 and F / S@H@C@HD-5 exhibit relatively superior vulcanization performance. F / S@H@C@HD-3 shows advantages in vulcanization torque and main vulcanization process, while F / S@H@C@HD-5 maintains high network strength while also ensuring good processing safety. Therefore, the HDTMS-modified multi-component composite filler system has good application potential in FKM reinforcement modification.

[0090] 5. Mechanical properties of F / S@H@C@HD composite materials like Figure 13As shown in (a), the tensile strengths of F / S@H@C@HD-1 to F / S@H@C@HD-6 are 6.0, 16.0, 15.6, 15.3, 13.7, and 13.3 MPa, respectively, showing an overall trend of first increasing sharply and then slowly decreasing. Compared with F / S@H@C@HD-1, the tensile strengths of F / S@H@C@HD-2, F / S@H@C@HD-3, F / S@H@C@HD-4, F / S@H@C@HD-5, and F / S@H@C@HD-6 are increased by 166.7%, 160.0%, 155.0%, 128.3%, and 121.7%, respectively. Among them, F / S@H@C@HD-2 reaches the maximum value of 16.0 MPa, showing the most significant strengthening effect, with an increase of 167%. The results indicate that after surface modification with hexadecyltrimethoxysilane (HDTMS), the interfacial compatibility between the S@H@C hybrid filler and the FKM matrix was further optimized. The introduction of long-chain alkyl groups enhanced the hydrophobicity and steric hindrance effect of the filler surface, promoted the uniform dispersion of the filler in the rubber matrix, and enabled efficient stress transfer between the filler and rubber phases, thereby significantly improving the load-bearing capacity. However, when the amount of HDTMS continued to increase to 6 phr, the tensile strength decreased. This is presumably related to the formation of multilayer adsorption by excessive silane coupling agent on the filler surface, which hindered interfacial bonding, as well as local agglomeration and stress concentration under high filling conditions.

[0091] like Figure 13 As shown in (b), the elongation at break of F / S@H@C@HD-1 to F / S@H@C@HD-6 are 143.2%, 215%, 172%, 167%, 153% and 142%, respectively. Compared with F / S@H@C@HD-1, F / S@H@C@HD-2 and F / S@H@C@HD-3 showed improvements of 50.1% and 20.1%, respectively, with F / S@H@C@HD-2 reaching the maximum of 215%. This indicates that after appropriate modification with HDTMS, the flexibility and spatial extension of the alkyl long chains improved the entanglement state between the filler and the rubber molecular chains, resulting in more complete molecular chain orientation and slippage during stretching, which is beneficial to a significant improvement in toughness. However, when HDTMS increased from 4 phr to 10 phr, the elongation at break continued to decrease. F / S@H@C@HD-6 showed a 34.0% decrease compared to F / S@H@C@HD-2, indicating that excessive silane coupling agent led to excessive hydrophobicity of the filler surface, weakened the interfacial interaction with the polar FKM matrix, increased the rigidity ratio, restricted the degree of freedom of molecular chain segment movement, and weakened the plastic deformation capacity of the material. The results indicate that there is a clear balance between reinforcement and toughening in this system, and that HDTMS modification has a more significant effect on improving toughness than CNC modification.

[0092] like Figure 13As shown in (c), with the increase of HDTMS content, the hardness of F / S@H@C@HD-1 to F / S@H@C@HD-6 increased slightly from 70 HA to 72 HA. The hardness of F / S@H@C@HD-2 to F / S@H@C@HD-6 composites increased by 2.9%, 1.4%, 2.1%, 1.4%, and 2.9% respectively relative to F / S@H@C@HD-1, showing an overall steady and slightly increasing trend. The small change in hardness indicates that the long-chain alkyl group of HDTMS, while imparting flexibility to the filler, has a limited impact on the overall rigidity of the material, thus balancing the contradiction between interfacial reinforcement and matrix elasticity. This also corresponds to the aforementioned change in elongation at break, meaning that appropriate HDTMS modification can significantly improve the toughness of the material without significantly increasing the hardness, demonstrating the unique advantages of silane coupling agent surface modification.

[0093] Based on the combined tensile strength, elongation at break, and hardness indicators, F / S@H@C@HD-2 exhibits the best performance in both strength and toughness. F / S@H@C@HD-3 maintains high strength while also possessing good toughness. F / S@H@C@HD-3 achieves a good overall balance between strength, toughness, and hardness: its tensile strength reaches 15.6 MPa, a 160.0% increase over the benchmark, while maintaining a high elongation at break of 172% and a hardness of 71 HA. Compared to the F / S@H@C system without HDTMS, the surface modification with HDTMS further enhances the tensile strength (from 15.0 MPa to 16.0 MPa) and elongation at break (from 199% to 215%) of the composite material, confirming the effectiveness of long-chain alkylsilane coupling agents in optimizing the interfacial properties of fluororubber composites.

[0094] 6. SEM analysis of F / S@H@C@HD composite materials Figure 14 The image shows a cross-sectional SEM image of a HDTMS-modified S@C@H composite powder-filled fluororubber composite material, where the filler content is 8 g of composite powder added to 100 g of fluororubber. Figure 14 (a) It can be seen that the cross-section of the composite material exhibits relatively obvious undulations and tear textures, with a rough fracture surface. A certain number of circular or near-circular particles and their traces are distributed within the cross-section, indicating that the composite filler has been sufficiently dispersed within the fluororubber matrix and participated in load-bearing and deformation during fracture. Compared to the previous group of S@C@H filled systems without HDTMS modification, this sample does not show obvious large-area continuous hard agglomerates in its cross-section, and the filler distribution is relatively more uniform, indicating that the dispersion state of the composite powder in the fluororubber matrix is ​​further improved after HDTMS surface modification.

[0095] From the perspective of fracture morphology characteristics Figure 14(a) The fracture surface is not a smooth, brittle fracture, but rather exhibits numerous wrinkles, tear ridges, and irregular undulations, indicating that the crack propagation was hindered and deflected by filler particles and the interfacial structure. The material underwent a certain degree of plastic deformation and energy dissipation before fracture. Generally, this rougher fracture morphology means that the composite system can dissipate the fracture energy under external forces through various mechanisms such as filler pull-out, interfacial friction, crack deflection, and matrix deformation, thereby delaying rapid crack propagation. For fluororubber composites, this tortuous fracture path is beneficial for improving the material's fracture resistance and has a positive effect on the synergistic improvement of tensile strength and elongation at break.

[0096] Depend on Figure 14 (b) Further analysis reveals that large spherical particles are embedded within the rubber matrix in the cross-section. The particles have clear outlines, and the surrounding matrix provides a certain degree of encapsulation, indicating that a certain degree of interfacial bonding has been formed between the modified composite filler and the fluororubber. These spherical particles did not completely detach from the matrix but remained embedded in the cross-section after fracture, suggesting that the filler can share some of the external force with the matrix during stress, which is beneficial for stress transmission in the composite material. Simultaneously, a small number of interfacial gaps or peeling marks can still be observed around the particles, indicating that some interfacial debonding still occurs during fracture. This suggests that although HDTMS modification improves interfacial compatibility, there is still room for further optimization of the interfacial bonding between the filler and the matrix at higher filler concentrations.

[0097] Based on the structural characteristics of this group of fillers, the HDTMS-modified S@C@H composite powder exhibits a spherical skeleton, a rough surface, and low surface energy. HGM provides a relatively stable spherical support structure, while the SiO2 deposition layer and the rough surface constructed by CNC increase the actual contact area between the filler and the matrix. The introduction of long-chain alkyl groups from HDTMS helps reduce the surface polarity of the filler, weakens secondary agglomeration caused by surface hydroxyl groups, and improves its wetting and dispersion behavior in fluororubber. Therefore, from the perspective of cross-sectional morphology, the spatial distribution and interfacial state of the HDTMS-modified filler in fluororubber are improved compared to the previous group, resulting in a more uniform internal structure of the composite system. This provides a good microscopic basis for the effective reinforcement effect of the filler.

[0098] Furthermore, although local pores and interface defects are still visible in the figure, the overall cross-sectional continuity is good, and no obvious large-scale filler peeling areas or severely void-accumulated areas are observed, indicating that a certain synergistic relationship has been established between the modified filler and the matrix. In particular Figure 14(b) The morphology of the particles embedded in the matrix indicates that the composite filler does not simply detach during fracture, but rather exhibits a certain degree of mechanical interlocking and interfacial adhesion with the matrix. This interfacial characteristic helps improve the efficiency of stress transfer from the rubber matrix to the filler phase, thereby contributing to the improvement of the material's tensile strength. At the same time, the relatively uniform filler dispersion and fewer large-scale agglomerations also help to reduce local stress concentration, prevent premature material failure, and are equally beneficial to maintaining the elongation at break.

[0099] In summary, after HDTMS-modified S@C@H composite powder was used to fill fluororubber, the composite material exhibited a relatively rough tear morphology, a more uniform particle distribution, and a certain degree of particle embedding characteristics, indicating that the modified composite filler has good dispersibility and a certain interfacial bonding ability in the matrix. Compared with the unmodified system, HDTMS modification helps to reduce filler agglomeration, improve interfacial compatibility, and optimize the microstructure of the composite system, thus providing favorable conditions for improving the mechanical properties of fluororubber composites.

[0100] 7. Surface wettability of F / S@H@C@HD composite materials like Figure 15 The figure shows the contact angles of several F / S@H@C@HD composite materials. As can be seen from the figure, with the increase of S@H@C@HD nanofiller content from 0 phr to 10 phr, the water contact angles of the F / S@H@C@HD-1 to F / S@H@C@HD-6 composite materials gradually increased from 90.1° to 104.3°, showing a monotonically increasing trend, indicating a continuous enhancement of the material surface hydrophobicity. The contact angles of the samples were 90.1°, 92.7°, 94.6°, 96.6°, 101.4°, and 104.3°, respectively, representing increases of 2.89%, 4.99%, 7.22%, 12.54%, and 15.76% compared to the unfilled sample (F / S@H@C@HD-1). The variation pattern shows that the contact angle of the composite material steadily increases (by about 2° / 2 phr) when the S@H@C@HD dosage is in the 0–6 phr stage. When the S@H@C@HD dosage is in the 6–8 phr stage, the contact angle of the composite material shows a significant jump (by 4.8°), breaking through 100° for the first time and reaching the superhydrophobic threshold. When the S@H@C@HD dosage is in the 8–10 phr stage, the contact angle continues to increase to 104.3°. This indicates that within the range of the experimental formulation, the synergistic effect of the HDTMS hydrophobic modifier and CNC plays a key role in promoting the construction of surface superhydrophobic properties at higher dosages.

[0101] The reasons can be understood from two aspects: surface chemical composition and surface microstructure. On the one hand, the introduction of S@H@C@HD significantly alters the near-surface phase composition of the composite material. The synergistic effect between the low surface energy fluororubber matrix and the inorganic-organic hybrid filler modified with hexadecyltrimethoxysilane (HDTMS) is extremely significant. The long-chain alkyl groups of HDTMS are oriented outwards, greatly reducing the surface free energy and thus weakening the wetting driving force of water droplets on the solid surface. On the other hand, SiO2-coated glass microspheres, combined with CNC and modified with HDTMS, construct a more complete micro / nano dual-scale rough structure. The self-assembly of HDTMS further increases the surface roughness, minimizing the actual contact area between the water droplet and the solid surface, thereby significantly improving the apparent contact angle. Especially when the filler content reaches 8 phr or more, the synergistic effect of the hydrophobic modification layer and the multi-scale rough structure reaches the critical threshold, causing the contact angle to exceed 100° and achieve superhydrophobic transformation. Combined with the error bars, it can be seen that the data of each group have small dispersion and good repeatability, indicating that the test results have high reliability. In summary, adding S@H@C@HD filler can effectively improve the surface hydrophobicity. At 10 phr, F / S@H@C@HD-6 achieved the highest contact angle (104.3°) in this group, which is 7.7° higher than the F / S@H composite material (96.6°) with the same dosage in the previous work, and 4.6° higher than the previous F / S@H@C system (99.7°). It successfully broke through the superhydrophobic limit (WCA>100°), showing the decisive role of HDTMS surface modification in constructing fluororubber-based superhydrophobic surfaces.

[0102] 8. Thermal conductivity of F / S@H@C@HD composite material The thermal conductivity / thermal response behavior of the F / S@H@C@HD composite material is as follows: Figure 16 As shown. Figure 16 (a) The gradient change of color from cold to hot visually reflects the temperature distribution on the sample surface and its evolution over time. Figure 16(b) Heating process: The heating plate temperature is set to 100 °C. The sample is placed on the surface of the hot plate. Starting from room temperature (about 20 °C), thermal images / surface temperature data are collected every 10 seconds. The heating phase ends when the sample surface temperature reaches about 80 °C (therefore, it is not recorded that it is in complete thermal equilibrium with the 100 °C hot plate). The heating results showed that the HDTMS-modified composite system exhibited a more pronounced trend of accelerated heating and increased endpoint temperature with increasing filler content: at 10 s, the 0 phr temperature was only 25.9 ℃, while at 6–10 phr it reached 34.3–34.8 ℃; at 60 s, the 0 phr temperature was 68.2 ℃, and at 4 / 6 / 8 / 10 phr it increased to 69.8 / 72.5 / 72.3 / 73.2 ℃ respectively; at 100 s, the 0 phr temperature was 77.0 ℃, at 2 phr it increased to 79.5 ℃, at 4 phr it reached 80.0 ℃, at 6 phr it was 80.3 ℃, and at 8 phr and 10 phr it further increased to 83.1 ℃ and 83.4 ℃ respectively. The trend is even clearer when the characteristic time is used for quantification: the time required for the surface to rise to 70 ℃ is shortened from about 66.0 s at 0 phr to about 51.1 s at 10 phr; the time required to rise to 75 ℃ is shortened from about 85.9 s to about 64.6 s; as for reaching the 80 ℃ threshold, it takes about 100 s at 4 phr and about 97.5 s at 6 phr, while it takes about 84.1 s and 83.1 s at 8 phr and 10 phr, respectively, to reach 80 ℃. This indicates that the HDTMS-modified composite nanofiller has enhanced compatibility / wetting properties with the FKM interface and reduced interfacial thermal resistance, which allows for the maintenance of a relatively effective heat transfer pathway even with high filling, thereby significantly improving the thermal response speed. Figure 16(c) Cooling process: After the sample was moved from the hot plate to the room temperature table (starting from about 80 °C), the surface temperature decay was recorded every 20 seconds. Although all samples cooled over time, with the increase of S@H@C@HD dosage, the surface temperature of the sample was higher and the cooling was slower for the same cooling time. For example, at 120 s, the 0 phr had dropped to 29.9 °C, while the 2 / 4 / 6 / 8 / 10 phr were still 34.9, 36.8, 37.2, 37.2 and 36.7 °C, respectively. At 160 s, the 0 phr was 25.3 °C, the 2 phr was 29.4 °C, while the 4–10 phr remained in the range of 31.0–31.5 °C, showing stronger temperature holding / insulation characteristics. Taking cooling to 40 °C as a characteristic point, the temperature drops to 40 °C in approximately 73.6 s at 0 phr, while it takes approximately 95.0 / 101.7 / 105.3 / 105.3 / 102.2 s at 2 / 4 / 6 / 8 / 10 phr, respectively. Regarding cooling to 30 °C, it is reached in approximately 119.4 s at 0 phr, while it takes approximately 155.4 s at 2 phr, and still not reaching 30 °C within 160 s at 4–10 phr. This phenomenon indicates that HDTMS modification, while improving the heating heat transfer efficiency, also results in a slower surface temperature decay during the cooling phase. This is related to the coupling effect of the composite filler network promoting continuous heat supply from the interior to the surface (making the surface temperature less prone to rapid drop) and the hollow structure of HGM and the multiphase interface jointly suppressing effective heat loss to the environment. In summary, after overall alcohol modification of HDTMS, the S@H@C composite nanofiller shows a more significant improvement in the thermal response of FKM in the range of 0–10 phr, especially in the 8–10 phr stage where it is easier to quickly reach the 80 ℃ threshold. At the same time, it also shows a stronger temperature holding capacity in the cooling stage. This indicates that the modification strategy can achieve a more effective synergy between the construction of heat transfer pathways / reduction of interfacial thermal resistance and the thermal insulation properties of HGM.

[0103] 9. Thermogravimetric analysis of F / S@H@C@HD composite material Figure 17 Tables (a) to (c) show the TG / DTG curves and residual carbon content changes of composite materials with different filler contents (F / S@H@C@HD-1~6 corresponding to 0, 2, 4, 6, 8, and 10 phr). Table 2 summarizes the T5% and T of each sample. 10 %, T 50Thermogravimetric parameters such as % and residual carbon content were used. As can be seen from the TG curves, the mass changes of each formulation were relatively small in the low-temperature region (approximately 30–400 °C), with the main weight loss process concentrated in the approximately 430–510 °C range, indicating that the system's thermal decomposition was still mainly in the medium-to-high temperature degradation stage. After introducing HDTMS-modified composite fillers, the overall difference in the main weight loss range was not significant, but the starting position of the curve's decline and the residual mass at the end changed significantly with dosage, indicating that HDTMS modification is dosage-sensitive to early decomposition behavior and residual layer construction. Combined with Table 2, it can be seen that early thermal stability (T5%, T...)... 10 The percentage showed a pattern of significant increase with low addition, decrease with medium addition, and partial recovery with high addition: F / S@H@C@HD-1 sample T5%, T 10 The percentages were 426.1 °C and 450.5 °C, respectively; the T5% and T of F / S@H@C@HD-2 were... 10 The T5% and T6% values ​​significantly increased to 446.2 °C and 459.4 °C, respectively, representing increases of 20.1 °C and 8.9 °C compared to the F / S@H@C@HD-1 sample. This indicates that a small amount of HDTMS-modified composite filler can effectively improve the dispersion and interfacial compatibility of the filler in FKM, enhance chain segment constraint, and improve the thermal insulation barrier effect, thereby delaying the escape of initial volatile products and chain segment breakage. However, when the addition amount is increased to 4~8 phr (F / S@H@C@HD-3~5), the T5% and T6% values ​​also decrease. 10 The T5% decreased to 425.8~420.7 ℃ and 443.9~438.1 ℃, which was actually lower than F / S@H@C@HD-1. This is presumably related to the increased agglomeration of fillers and interface defects under high filling conditions. Additionally, the long alkyl chains in HDTMS are organic components, and their increased content may introduce a more thermally degradable flexible structure or promote uneven local heat transfer, leading to earlier weight loss. When the content was further increased to 10 phr (F / S@H@C@HD-6), the T5% rebounded to 429.6 ℃ (3.5 ℃ higher than F / S@H@C@HD-1), but T... 10 The thermal stability index (T0.05) remained at 447.8 °C (2.7 °C lower than F / S@H@C@HD-1), indicating that the improvement in initial decomposition under high addition was limited and still affected by structural inhomogeneity. In contrast, the mid-to-late stage thermal stability index (T0.05) was significantly higher. 50 % Overall increase with addition amount: 0 phr sample T 50The initial degradation rate was 477.1 °C, which increased to 479.5–481.8 °C after the addition of modified composite filler. The highest value of 481.8 °C was reached in the 8–10 phr phase, which is 4.7 °C higher than that in the 0 phr phase. This indicates that the heat transfer barrier and shielding effect of the SiO2@HGM inorganic framework can continue to play a role in the later stage of the main degradation phase, thereby delaying the large-scale chain breakage and rapid weight loss process. The maximum weight loss rate peaks of each sample in the DTG curve are mainly located around 470–490 °C. The peak positions are generally similar but the peak shapes are slightly different. This also reflects that HDTMS modification has a greater impact on the initial degradation and residual layer structure, while the change in the position of the main degradation peak is relatively limited. The residual carbon rate results further demonstrate the optimal dosage window: the residual carbon rate was 15.30% at 0 phr, significantly increasing to 19.25% (an increase of 3.95 percentage points) at 2 phr, reaching its highest point. This indicates that a continuous and dense pyrolysis residual shielding layer is more easily formed at low addition levels, thereby improving residual quality and thermal protection capability. However, the residual carbon rate decreased to 15.28%–12.95% at 4–8 phr, and rebounded to 14.67% at 10 phr, indicating that excessive modified filler may disrupt the continuity of the residual layer or cause insufficient densification, leading to a weakened protective effect in the later stages of pyrolysis. In summary, the HDTMS-modified S@H@C composite filler shows a clear optimal dosage for improving the thermal stability of FKM: 2 phr exhibits the best performance in terms of early decomposition temperature and residual carbon rate, while 8–10 phr is more beneficial for improving stability in the later stages of the main degradation phase (T0). 50 The result (%) indicates that controlling interface compatibility and structural densification through HDTMS is key to achieving enhanced thermal stability. However, excessive addition can lead to a decrease in early thermal stability and char formation due to factors such as agglomeration / introduction of organic segments.

[0104] Table 2 Thermogravimetric data of F / S@H@C@HD composite materials The above description discloses only preferred embodiments of the present invention and should not be construed as limiting the scope of the present invention. Therefore, equivalent variations made in accordance with the claims of the present invention are still within the scope of the present invention.

Claims

1. A multi-scale composite packing material, characterized in that: It includes a hollow glass microsphere core and a nano-SiO2 shell layer deposited in situ on the surface of the hollow glass microsphere core by the Stöber method. The nano-SiO2 shell layer is formed by the accumulation of SiO2 particles and contains nano-cellulose dispersed between the SiO2 particles to form a fibrous network structure.

2. The multi-scale composite packing according to claim 1, characterized in that: The mass ratio of the nanocellulose to the hollow glass microspheres is (0.5-2):

1.

3. The multi-scale composite packing according to claim 1, characterized in that, Its preparation method includes the following steps: S1. Mix nanocellulose with anhydrous ethanol, add N,N-dimethylacetamide, heat and stir to pretreat, and obtain a pretreated nanocellulose dispersion; S2. Mix anhydrous ethanol and deionized water in a certain proportion, add polyethylene glycol, hollow glass microspheres and the pretreated nanocellulose dispersion obtained in step S1, stir evenly to obtain a mixed dispersion. S3. Add ammonia water to the mixed dispersion, and simultaneously add tetraethyl orthosilicate dropwise to carry out a hydrolysis and polycondensation reaction. After the reaction is completed, separate, wash and dry the product to obtain a multi-scale composite filler.

4. The multi-scale composite packing according to claim 3, characterized in that: In step S1, the heating and stirring temperature is 80-100 ℃, and the time is 4-8 h.

5. The multi-scale composite packing according to claim 3, characterized in that: In steps S2 and S3, the mass ratio of hollow glass microspheres to tetraethyl orthosilicate is 1:(3-7).

6. A hydrophobically modified multi-scale composite filler, characterized in that, The hydrophobic modified multi-scale composite filler is obtained by surface modification of the multi-scale composite filler of claim 1 with hexadecyltrimethoxysilane.

7. The hydrophobically modified multi-scale composite filler according to claim 6, characterized in that, The amount of hexadecyltrimethoxysilane used is 10%-30% of the mass of the multi-scale composite filler.

8. The hydrophobically modified multi-scale composite filler according to claim 6, characterized in that, The preparation method includes the following steps: dispersing the multi-scale composite filler according to any one of claims 1-5 in anhydrous ethanol, stirring evenly to obtain a dispersion, adding hexadecyltrimethoxysilane to the dispersion, stirring to react, and after the reaction is completed, separating, washing and drying the product to obtain the hydrophobic modified multi-scale composite filler.

9. A modified fluororubber, characterized in that, It comprises a fluororubber matrix and a multi-scale composite filler according to any one of claims 1-5, or a hydrophobically modified multi-scale composite filler according to any one of claims 6-8.

10. A modified fluororubber according to claim 9, characterized in that, Its preparation method includes the following steps: (A) Fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and SiO2-coated hollow glass microsphere core-shell composite filler were sequentially added to an open mill and mixed evenly to obtain a mixture. (B) The mixture is vulcanized at 150-190℃ for 5-15 minutes to obtain vulcanized fluororubber; (C) The vulcanized fluororubber is vulcanized in two stages at 180-220℃ for 18-22 h to obtain modified fluororubber; In step (A), the ratio of fluororubber raw rubber, magnesium oxide, calcium oxide, TAIC, Luperox 101-XL-45, and composite filler is 100: (2-4): (4-6): (2-4): (2-4): (2-10).

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