Dense monolithic (Ti,Zr,Hf)C/SiC ceramic nanocomposites with four different molar ratios of metallic elements in the (Ti,Zr,Hf)C phase (i.e., Ti:Zr:Hf=1:1:1, 2:3:5, 2:3:3, and 1:2:1) were prepared upon pyrolysis of novel (Ti,Zr,Hf)-containing single-source-precursors (SSPs), followed by spark plasma sintering. A thorough characterization was conducted to elucidate the synthesis of the SSPs, polymer-to-ceramic transformation, chemical/phase compositions and microstructure of the SiTiZrHfC-based ceramics. The results revealed the feasibility of synthesizing the nanocomposites with high (Ti,Zr,Hf)C content using SSP method. These nanocomposites were characterized by a unique microstructure with in situ generated (Ti,Zr,Hf)C@C core-shell nanoparticles homogeneously mixed with β-SiC. The ablation behavior of the nanocomposites was evaluated on an air-plasma device for 60 s. Impressively, the nanocomposites exhibited excellent ablation resistance, and the lowest linear ablation rate reached -0.58 μm/s at 2200 °C. Notably, the ablation resistance can be dramatically improved by precisely tailoring the atomic ratios of metal elements within the (Ti,Zr,Hf)C phase via molecular design of the SSPs. The formation of a multiple-oxides layer with both high-melting-point phase ((Ti,Zr,Hf)O₂) and low-melting-point phases ((Zr,Hf)TiO₄) and glassy SiO₂ as well as their structure played a critical role in the enhanced ablation resistance. The uniform distribution of the high-melting-point (Ti,Zr,Hf)O₂ nano-/micro- particles throughout the glassy SiO₂ matrix significantly enhanced the viscosity and stability of the oxide layer by pinning effect, offering superior protection against the ingress of oxygen atoms and excellent resistance to mechanical erosion.

C/C-SiC-(Zr_xHf_1−x)C composites were prepared by the reactive melt infiltration method. The microstructure of the porous C/C skeleton and the C/C-SiC-(Zr_xHf_1−x)C composites, as well as the structural evolution and ablation behavior of the C/C-SiC-(Zr_xHf_1−x)C composites, were systematically investigated. The results show that the C/C-SiC-(Zr_xHf_1−x)C composites were mainly composed of carbon fiber, carbon matrix, SiC ceramic, (Zr_xHf_1−x)C and (Zr_xHf_1−x)Si₂ solid solutions. The refinement of the pore structure is beneficial to promote the formation of (Zr_xHf_1−x)C ceramic. The C/C-SiC-(Zr_xHf_1−x)C composites exhibited outstanding ablation resistance under an air–plasma environment at around 2000 °C. After ablation for 60 s, CMC-1 appeared to possess the minimum mass and linear ablation rates of only 2.696 mg/s and −0.814 µm/s, respectively, which are lower than those of CMC-2 and CMC-3. During the ablation process, a Bi-liquid phase and a liquid–solid two-phase structure were formed on the ablation surface which could act as an oxygen diffusion barrier to retard further ablation, which is responsible for the excellent ablation resistance of the C/C-SiC-(Zr_xHf_1−x)C composites.

Carbon-fiber-reinforced carbon and silicon carbide (C/C-SiC) composites were prepared using chemical vapor infiltration (CVI) combined with reactive melt infiltration (RMI). The microstructure and flexural properties of C/C-SiC composites after oxidation in different temperature water vapor environments were studied. The results indicate that the difficulty of oxidation in water vapor can be ranked from easy to difficult in the following order: carbon fiber (CF), pyrolytic carbon (PyC), and ceramic phase. The surface CFs become cone-shaped under corrosion. PyC has a slower oxidation rate and lower degree of oxidation compared to CF. The SiO₂ layer formed by the oxidation of SiC and residual Si was insufficient to fully cover the surface of CFs and PyC. As the temperature increased, the oxide film thickened, but the corrosion degree of CF and PyC intensified, and the flexural performance continuously deteriorated. The flexural strength of C/C-SiC composites was 271.86 MPa at room temperature. Their strength retention rates were all higher than 92.19% after water vapor corrosion at 1000 °C, still maintaining the “pseudoplastic” fracture characteristics. After water vapor corrosion at 1200 °C, the CFs inside the composites sustained more severe damage, with a strength retention rate as low as 48.75%. The fracture mode was also more inclined towards brittle fracture.

In this study, two types of porous C_f/C_m composites were obtained by introducing pyrolytic carbon (PyC) and pyrolytic carbon/furan resin carbon (PyC/FRC). Subsequently, C_f/C_m-SiC-(Zr_xHf_1−x)C composites with different carbon matrices were prepared by introducing SiC and (Zr_xHf_1−x)C matrices into the porous C_f/C_m composites via the reactive melt infiltration method, specifically termed as C_f/PyC-SiC-(Zr_xHf_1−x)C and C_f/PyC/FRC-SiC-(Zr_xHf_1−x)C composites. The microstructures of the porous C_f/C_m and C_f/C_m-SiC-(Zr_xHf_1−x)C composites with different carbon matrices were examined, and a comprehensive analysis was conducted on microstructural evolution and mechanical behaviors of the C_f/C_m-SiC-(Zr_xHf_1−x)C composites. The results indicate that both C_f/C_m-SiC-(Zr_xHf_1−x)C composites underwent similar microstructural evolution processes, differing only in terms of evolution kinetics and final microstructure. Differences in the pore structures of porous C_f/C_m composites, as well as in the reactivities of carbon matrices, were identified as primary influencing factors. Additionally, both C_f/C_m-SiC-(Zr_xHf_1−x)C composites exhibited “pseudo-ductile” fracture characteristics, with flexural strengths of 214.1 ± 8.8 MPa and 149.6 ± 12.2 MPa, respectively. In the C_f/PyC-SiC-(Zr_xHf_1−x)C composite, crack initiation during loading primarily originated from the ceramic matrix, while in the C_f/PyC/FRC-SiC-(Zr_xHf_1−x)C composite, failure initially arose from the residual FRC matrix. Excessive fiber corrosion and the presence of residual low-modulus FRC matrix resulted in lower mechanical performance.