1. Introduction
During high-speed flight, especially in the re-entry phase, intense aerodynamic heating is induced by shock wave compression and viscous friction between the airframe and the atmosphere. In specific local regions, such as the nose cone and sharp leading edges, temperatures can exceed 2000 ℃ [1], [2], [3]. Consequently, stringent requirements are imposed on thermal protection materials (TPMs). These materials must possess superior mechanical properties, high physicochemical stability, and extremely low ablation rates at elevated temperatures [4], [5]. Currently, candidate materials for high-temperature structural applications primarily include superalloys [6], ultra-high temperature ceramics (UHTCs) [7], and C/C composites [8]. Although superalloys are easily processed and exhibit good thermal shock resistance, their application is restricted by high density and severe creep/oxidation above 1200 ℃ [9]. UHTCs, mainly composed of refractory metal borides or carbides, are characterized by high thermal conductivity, high strength, and excellent oxidation resistance [10], [11], [12]. However, their practical use is severely limited by inherent brittleness, low damage tolerance, and poor thermal shock resistance [8], [14]. Conversely, C/C composites offer advantages such as low density, a low coefficient of thermal expansion, and outstanding thermal shock stability. Nevertheless, their oxidation resistance at moderate temperatures remains a significant drawback [13], [14], [15]. To enhance the high-temperature ablation resistance of C/C composites, matrix modification is currently employed. However, traditional UHTCs (e.g., HfB2, ZrC, HfC, and ZrB2) possess high density (6.1-12.7 g/cm3) and exhibit poor dispersion within fiber preforms or porous C/C matrices [13], [16], [17], [18], [19]. Recently, polymer-derived ceramics (PDCs) have attracted considerable attention due to their designable molecular structures, ease of processing, and excellent high-temperature performance [20], [21], [22], [23], [24]. Various novel ceramics with tailored microstructures and properties can be synthesized via the PDCs route [25], [26], [27], [28]. This approach effectively addresses issues such as large particle size, poor dispersion, and low densification of ceramic phases.
Silicon-based ceramic precursors were initially employed for the modification of C/C composites. C/C-SiC composites were fabricated using polycarbosilane through infiltration and pyrolysis [29]. Although SiO2 provides oxygen blocking effects, a continuous protective layer cannot be formed due to the shear forces during ablation [30]. Furthermore, the high-temperature protection capability is limited by the active oxidation of SiC. With the advancement of high-temperature ceramic precursors, Zr- and Hf-containing polymeric precursors have become preferred modifiers for C/C composites [31], [32], [33], [34], due to their high melting points and excellent ablation resistance. ZrC- and HfC-based matrices are typically prepared via the precursor infiltration and pyrolysis (PIP) method. For instance, C/C-ZrC composites were prepared, where oxygen barrier, thermal insulation, and erosion resistance were provided by a dense ZrO2 layer, the oxidation product of ZrC [31]. It was found that the concentration of the Zr-containing precursor significantly influences the microstructure and properties of the composites [32]. While the ZrC content and ablation resistance of the composite can be improved by increasing the precursor concentrations, particle agglomeration and thus a decrease in flexural strength of the composite are often induced. Additionally, C/C-HfC composites with uniformly dispersed HfC were fabricated [33]. After high-temperature ablation, a tree-coral-like HfO2 thermal barrier and an HfCxOy oxygen-resistant layer were formed. Consequently, the ablation resistance was significantly enhanced compared to pure C/C. For HfB2 precursor-modified C/C composites, a linear ablation rate of 31.4×10-4 mm/s was achieved after ablation for 90 s under a heat flux of 2.38 MW/m² [34]. Despite these advancements, several key challenges remain. First, uniform distribution of the ceramic phase is difficult to achieve during the PIP process, and particle agglomeration frequently occurs, which adversely affects mechanical properties. Second, the continuity, density, and bonding strength of the oxide protective layer are often insufficient, making it prone to spallation under extreme conditions. Finally, the coordinated optimization of the type, content, and microstructure of the ceramic phase is required to maximize mechanical properties and ablation resistance of the composite.
In this work, a novel SiHfCB precursor [35], characterized by high Hf content, excellent dispersion, and non-precipitating, was introduced into a porous C/C matrix. This combination was achieved through a high-pressure assisted precursor infiltration and pyrolysis process. Consequently, the precise regulation of a spatially “integrated” three-dimensional multiphase continuous structure was enabled. Subsequently, the microstructure of the obtained C/C-SiHfCB composites was characterized, and the dispersion of the ceramic phase within the C/C matrix was examined. The room-temperature mechanical properties and thermal shock resistance at various temperatures were thoroughly evaluated. The critical thermal shock temperature difference (△Tc) was determined, followed by a detailed analysis of the toughening mechanisms and the failure mechanisms after thermal shock testing. Finally, oxyacetylene flame ablation tests were performed to assess the performance limits of the C/C-SiHfCB composites under extreme conditions. Based on these results, the corresponding anti-ablation mechanisms were elucidated.
2. Experimental
2.1. Materials synthesis and processing
The synthesis of SiHfCB precursor has been reported in previous studies [35]. Here, a SiHfCB precursor with an Hf/Si ratio of 0.2 is selected for the preparation of composites. The molecular weight (Mw) of the SiHfCB precursor (with a Hf/Si ratio of 0.2) is 8890 g/mol, and its viscosity is 1671 mPa·s. The ceramic yield of SiHfCB precursor after pyrolysis at 1500℃ reached 68.6%. The porous C/C composites (72×72×10 mm3, porosity of ~50%) used here was fabricated via a sugar-carbon conversion method [36], with carbon fiber volume fraction of ~20% and a density of ~0.9 g/cm3. The preparation process of C/C-SiHfCB composite is illustrated in Fig. 1. The synthesized liquid SiHfCB precursor was combined with the porous C/C composites through a high-pressure impregnation process. The high-pressure impregnation pressure is set at 100 MPa and holding for 20 min. Subsequently, the impregnated composites were cured stepwise in an oven at 120, 140, 160, 180, and 200℃, holding 4 h at each stage. After curing, it was pyrolyzed at 1500℃ in vacuum for 1 h, with a heating rate of 2 ℃/min. The final C/C-SiHfCB composites was obtained after six cycles of high-pressure precursor impregnation and pyrolysis.
Fig. 1. Schematic illustration of preparation process of C/C-SiHfCB composites. |
2.2. Materials characterization
Phase identification and quantitative analysis were carried out using an X-ray diffractometer (XRD, Empyrean, Panalytical, Netherlands) equipped with monochromatic Cu Kα radiation, operating at a scan speed of 2°/min over a 2θ range of 20-40°. The viscosity of the SiHfCB precursor solution was tested by the digital rotary viscometer(NDJ-1S, Shanghai Youyi Instrument, China). The molecular weights (Mw) of the synthesized SiHfCB precursor were measured using gel permeation chromatography (GPC, Waters E2695, USA). The ceramic yield of the SiHfCB precursor was determined by thermogravimetric analysis (TGA, STA449F3 Jupiter, Netzsch, Germany). The specific method involved heating the precursor from room temperature to 1500℃ at a rate of 10℃/minute in an argon atmosphere (100 ml/min). A scanning electron microscope (SEM, Vega3, Tescan, Czech) with energy dispersive spectroscopy (EDS) was utilized to explore the micromorphology of the sample. The density of porous C/C composites and C/C-SiHfCB composites was determined by Archimedes drainage method. Thermal diffusivity, thermal conductivity, and specific heat capacity of C/C-SiHfCB composites at 25 ℃ to 1000 ℃ was measure by flash heat conduction instrument (LFA 467 HT, Netzsch, Germany).
2.3. Mechanical properties, thermal shock resistance and oxyacetylene torch ablation test
The flexural strength of C/C-SiHfCB composites and porous C/C preform was characterised by a three-point flexural method using an electronic universal material testing machine (Instron-5569, Instron Co., Ltd). The prepared C/C-SiHfCB composites were processed into a 3×4×36 mm3 standard sample, and the sample’s surface was polished with sandpaper to eliminate errors. The effective span is set at 30 mm, and the loading rate of the indenter is set at 3 mm/min. The reported values represent the average of the three tests. The sample flexural strength (σf) calculation Eq is shown in (1).
where Pf is the maximum load imposed by the testing machine (N), L is the span (mm), h is the height of the sample (mm), and W is the width of the sample (mm).
The thermal shock resistance of C/C-SiHfCB composites was measured by water quenching method. The muffle furnace (SX-G03173, Tianjin Zhonghuan Electric Furnace Co., Ltd) has been heated to 220 ℃, 420 ℃, 620 ℃, 820 ℃ and 1020 ℃ respectively, and then the prepared flexural strength test standard samples have been placed in the muffle furnace. The sample was taken out after being kept at this temperature for 15 min, and immediately placed in water at 20℃. The flexural strength of the sample was tested after it was completely cooled. According to ASTM C1525-04 standard [37], the temperature difference reached when the final flexural strength of the sample is reduced to 70% of the flexural strength at room temperature is determined as the critical thermal shock temperature difference (△Tc). Finally, the linear interpolation method is used to calculate the critical thermal shock temperature difference. Three samples were tested in each group, and the final results were averaged.
The ablation resistance of C/C-SiHfCB composites was tested by the oxygen-acetylene test device (Xi'an Zhirui Industrial System Engineering Co., Ltd). The test process strictly complies with GJB323B-2018 standard [38]. The sample size is Φ25 × 5 mm3, placed in the graphite ablation mold. The ablation temperature of the sample surface is controlled by adjusting the oxygen and acetylene flow rate. The nozzle position is 20 mm from the center of the sample surface. An infrared temperature measurement device (SA-2S300A, Wuxi Youtian Environmental Technology Co., Ltd) measured the temperature of the ablated sample surface, with an emissivity of 0.95. An oxyacetylene torch ablated the sample at different temperature. The ablation resistance of samples can be evaluated by mass ablation rate (Rm) and linear ablation rate (Rl). The calculation Eqs are shown in (2) and (3), respectively:
Where m0 is the mass of the material before ablation (g); mt is the mass of the material after ablation (g); S is the surface area of the ablation area (mm2); t is the ablation time (s); l0 is the thickness of the material before ablation (mm); lt is the thickness of the material after ablation (mm).
3. Results and discussion
3.1. Microstructure and thermal properties of C/C-SiHfCB Composites
The surface micromorphology of the C/C-SiHfCB composites subjected to different numbers of PIP cycles was then examined by SEM, as shown in Fig. 2(a-c). It was observed that with an increasing number of PIP cycles, the ceramic content within the composites gradually increased. The pores in the porous C/C preform were progressively filled by the ceramic derived from the pyrolyzed SiHfCB precursor, and no significant pore defects were detected. After six PIP cycles, the density of the composites was calculated to increase from 0.9 g/cm3 to 2.32 g/cm3. Correspondingly, its porosity decreased from 50% to 14%. Subsequently, the XRD pattern of C/C-SiHfCB Composites in Fig. 2(d) reveals that the dominant phase is primarily composed of HfB2 and SiC, along with a minor phase of HfC. Due to the high carbon content in the porous C/C preform, the SiHfCB precursor could partially react with carbon to form HfC during pyrolysis process.
Fig. 3 SEM images and XRD pattern of C/C-SiHfCB composites surface after different high-pressure PIP cycles: (a) 2 PIP cycles, (b) 4 PIP cycles, (c) 6 PIP cycles, (d) XRD pattern. |
Fig. 3. (a-c) SEM images of porous C/C preform section, and (d-f) BSE images of C/C-SiHfCB composites section after 6 high-pressure impregnation-pyrolysis cycles. |
The thermal diffusivity, thermal conductivity, and specific heat capacity of the C/C-SiHfCB composites are listed in Table 1. Both the thermal diffusivity and thermal conductivity decrease with increasing temperature, while the specific heat capacity increases with temperature. However, even at 1000℃, the thermal conductivity remains higher than 40 W/(m·K). This enables the C/C-SiHfCB composites to maintain good thermal conduction performance at high temperatures. Particularly in high-temperature ablation environments, local heat can be rapidly transferred throughout the composite, thereby reducing its surface temperature.
Table 1. Thermal diffusivity, thermal conductivity, and specific heat capacity of C/C-SiHfCB composite. |
| Temperature (℃) | Thermal diffusivity (mm2/s) | Thermal conductivity (W/(m·K)) | Specific heat capacity (J/(g·K)) |
|---|---|---|---|
| 25 | 50.65 | 71.12 | 0.78 |
| 400 | 21.50 | 56.85 | 1.469 |
| 600 | 17.17 | 49.46 | 1.601 |
| 800 | 14.70 | 45.31 | 1.713 |
| 1000 | 13.19 | 43.32 | 1.825 |
3.2. Mechanical properties and thermal shock resistance of C/C-SiHfCB Composites
The flexural strength and fracture surface micromorphology of the prepared C/C-SiHfCB composites was tested via three-point flexural tests, with the corresponding stress-strain curves shown in Fig. 4(a, b). The room-temperature flexural strength of the C/C-SiHfCB composites was 237±42 MPa. The curves in Fig. 4(a, b) display a non-brittle fracture mode. As the applied stress increased gradually, the curve remained linearly elastic up to point A. Beyond point A, surface cracking initiated in the composite. The stress then increased to point B, where the flexural strength reached its maximum value. When the stress exceeded the ultimate strength of the composite, cracking occurred within the bulk material. On the tension side, fiber fracture and fiber debonding were observed. With further increase in strain, the stress exhibited a segmented, non-brittle failure pattern, accompanied by visible surface deformation and fiber pull-out damage in Fig. 4(c, d). The test was stopped after a pronounced drop in flexural strength, indicating flexural failure of the composite. Meanwhile, crack branching and deflection are evident on the composites surface, which is indicative of material toughening. Fig. 4(e, f) reveals that the primary fracture mechanisms include fiber fracture, bridging, debonding, and pull-out. This further confirms that the toughening mechanisms in the C/C-SiHfCB composites play a crucial role in resisting stress-induced damage.
Fig. 5 (a, b) Flexural stress-strain curves of C/C-SiHfCB and porous C/C composite, (c-f) micromorphology of fracture section of C/C-SiHfCB composites after flexural damage. |
Fig. 5. Morphology of surface and fracture section of C/C-SiHfCB composites after thermal shock test at different temperatures: (a, g) RT,(b,h) 200℃, (c,i) 400℃,(d,j) 600℃,(e,k) 800℃,(f,l) 1000℃. |
The load-displacement curves of the C/C-SiHfCB composite, presented in Fig. 6(a), indicate that at lower thermal shock temperature differences, the displacement corresponding to the maximum load was similar. A notable degradation in the flexural strength of the composites after the 1000℃ thermal shock. The primary cause is attributed to the surface oxidation within the C/C-SiHfCB composites in the high-temperature oxidative environment. Consequently, the mechanical properties of the composites were significantly compromised. To calculate the critical thermal shock temperature difference of the C/C-SiHfCB composite, the flexural strength of each sample group was tested after thermal shock. The results are shown in Fig.7(b). The curve of residual strength versus thermal shock temperature reveals a gradual decline in residual strength with increasing temperature. The calculated critical thermal shock temperature difference, corresponding to a strength retention rate of 70%, was determined to be 912℃, demonstrating the excellent thermal shock resistance of the C/C-SiHfCB composite.
Fig. 6. Load-displacement curves and critical thermal shock temperature difference curve of C/C-SiHfCB composites after thermal shock resistance test at different thermal shock temperature differences: (a) load-displacement curves, (b) critical thermal shock temperature difference curve. |
Fig. 7. (a) Time-temperature curves of C/C-SiHfCB composites, and (b) optical photos of the samples before and after undergoing different temperatures of oxyacetylene torch ablation test. |
3.3. Ablation resistance and mechanism of C/C-SiHfCB Composites
3.3.1. Ablation resistance of C/C-SiHfCB composites
To further explore the temperature limit of the C/C-SiHfCB composite, oxyacetylene torch ablation tests aiming to study the ablation resistance and behavior of the composites at different temperatures were conducted. The temperature-time curves during the ablation process and the macromorphology before and after testing were compared, as shown in Fig.7. The ablation time was fixed at 60 s, with oxyacetylene torch temperatures set at 2000℃, 2150℃, and 2300℃. Fig. 7(b) presents the overall surface macromorphology of the C/C-SiHfCB composites before and after ablation. After ablation at 2000℃, the surface remained intact with no localized oxide layer spallation or structural damage. The HfB2, HfC and SiC phases inside the C/C-SiHfCB composites at the ablation center are usually covered by HfO2 and SiO2 formed by high-temperature oxidation, which probably provide excellent ablation resistance. Under the 2150℃/60 s condition, more HfO2 and SiO2 phases were generated from the oxidation of the C/C-SiHfCB composites surface. However, no large pores were observed in the matrix fiber-oriented perpendicular to the oxyacetylene torch direction. This is primarily attributed to the filling and covering of some pores by HfO2 and a small amount of SiO2. After exposure at 2300℃ for 60 s, partial oxide layer spallation appeared at the ablation center. This is mainly due to the volatilization of SiO2, which significantly reduced their bonding strength with the composite matrix. Consequently, partial matrix ablation and minor oxide layer spallation occurred under the high-temperature oxyacetylene torch impact. These results indicate that after the 2150℃/60 s test, the overall contour of the C/C-SiHfCB composites remained intact. The matrix structure perpendicular to the oxyacetylene torch direction was preserved without significant structural damage, demonstrating relatively excellent ablation resistance.
Table 2 presents the linear and mass ablation rates of the C/C-SiHfCB composites after exposure to different oxyacetylene torch temperatures. The composite exhibited a linear ablation rate of 5.7×10⁻⁴ mm/s at 2000℃, which rose to 15.6×10⁻⁴ mm/s at 2150℃ and further increased to 38.4×10⁻⁴ mm/s at 2300℃. This significant rise indicates that the protective oxide layer on the composites surface was compromised at 2300℃, leading to partial ablation of the internal material. Concurrently, the mass ablation rate increased with rising oxyacetylene torch temperature. This trend is primarily caused by the volatilization of the B2O3 and SiO2 oxide layers and the oxidation of the C/C matrix, which is also the reason for the increase in linear ablation rate. Based on the ablation test results, the C/C-SiHfCB composites exhibited a relatively low ablation rate at 2150℃, which could be considered as a novel near-zero-ablative TPM.
Table 2. Linear and mass ablation rates of C/C-SiHfCB composites after oxyacetylene torch ablation. |
| Composite | Ablation condition | Linear ablation rate (mm/s) | Mass ablation rate (g/(mm2·s)) |
|---|---|---|---|
| C/C-SiHfCB | 2000℃/60 s | 5.7×10-4 | 2.9×10-6 |
| 2150℃/60 s | 15.6×10-4 | 17.7×10-6 | |
| 2300℃/60 s | 38.4×10-4 | 32.3×10-6 |
Fig. 8(a, b) presents the micromorphology of the ablation center of the C/C-SiHfCB composites after testing at 2000℃ for 60 s. No distinct ablation pits are observed on the surface. However, slight oxidation of the substrate occurs, resulting in the formation of a small number of pores. The pore formation is mainly attributed to the diffusion of a limited amount of oxygen into the interior of the composites during ablation. The diffused oxygen reacts with the constituents of the composite, generating gaseous products such as CO2 and CO, accompanied by the volatilization of B2O3. The escape of these gaseous products from the matrix leads to the formation of pore structures. Because the exposure time at this temperature is relatively short, the oxidation of the matrix remains shallow. Fig. 8(c, d) presents the micromorphology of the ablation center after testing at 2150℃ for 60 s. Under this condition, the ablation temperature is higher, resulting in more pronounced oxidation of the C/C matrix. A significant amount of SiO2 and HfO2 is exposed on the composite surface, although no large bubbles are observed. At this elevated temperature, the SiO2 layer formed during ablation becomes thicker and exhibits improved fluidity, enabling it to effectively fill matrix cracks and surface pores. However, partial volatilization of SiO2 at this temperature reduces the viscosity of the oxide melt and facilitates oxygen diffusion. This promotes internal oxidation of the matrix, leading to increased gas generation and higher internal pressure. Consequently, gases more easily rupture the oxide layer and escape, resulting in the formation of numerous smallpores at the ablation center. Fig. 8(e, f) shows the micromorphology of the ablation center after testing at 2300℃ for 60 s. A larger portion of the C/C-SiHfCB composites is ablated, consistent with the highest mass ablation rate reported in Table 2. The HfO2 particles in the central ablation region begin to grow and exhibit a compact, sintered micromorphology. This sintering behavior promotes crack healing within the ablation zone, which partially impedes further oxygen ingress into the composite interior and suppresses dynamic ablation. As a result, a substantial amount of HfO2 and SiO2 remains at the ablation center. During the rapid cooling after ablation, Hf-containing phases preferentially precipitate in the ablation center, forming compact blocks distributed within the remaining surface layer.
Fig. 8. Micromorphology of the ablation center of C/C-SiHfCB composites surface after oxyacetylene torch ablation testing at different temperature: (a, b) 2000℃/60 s, (c, d) 2150℃/60 s, (e, f) 2300℃/60 s. |
Fig. 9 presents the XRD patterns of the ablation center on the C/C-SiHfCB composites after oxyacetylene torch at different ablation temperatures. After ablation test, the surface phases mainly consist of m-HfO2, t-HfO2, and a small amount of exposed carbon matrix. This is because a continuous oxide protective layer cannot be fully established during the short-duration test at 2000℃, resulting in the partial exposure of the graphitized carbon matrix on the composite surface. The t-HfO2 phase generally forms at elevated temperatures. During the ablation process, the surface oxide layer experiences rapid cooling once the heat source is removed. Such rapid cooling suppresses the complete transformation of t-HfO2 into the thermodynamically stable m-HfO2 during cooling. In addition, the oxidation-derived HfO2 tends to possess a relatively fine grain size, which can further stabilize the tetragonal phase because t-HfO2 is more readily retained in fine grains.
Fig. 9. XRD patterns of C/C-SiHfCB composites surface ablation center after different oxyacetylene torch ablation testing temperatures. |
As the ablation temperature increases, the formation of high-temperature t-HfO2 becomes more significant, and its retention after cooling is further promoted by the combined effects of rapid quenching and grain-size stabilization. Upon exposure to an oxyacetylene torch at 2000℃, oxidation of the SiHfCB-derived ceramic phases (containing HfB2, HfC, and SiC) generated HfO2 and SiO2. Any B2O3 formed from the oxidation of HfB2 rapidly volatilized due to its high vapor pressure at this temperature. The volatilization, coupled with flame scouring, led to the formation of pores within the oxide layer. These pores subsequently acted as rapid diffusion channels for oxygen, further accelerating the oxidation process. As the surface ablation temperature increased to 2150℃ the SiC component underwent active oxidation to form SiO(g), while the B2O3(l) began to vaporize and escape from the surface. Meanwhile, the escaping gases helped to impede the inward diffusion of oxygen, thereby providing effective protection for the underlying substrate. Upon reaching 2300℃, the previously formed SiO2(l) also rapidly vaporized, ultimately leaving behind a porous oxide layer composed predominantly of HfO2. The remaining porous HfO2 skeleton, while not fully dense, provided structural support and further contributed to the composite's robust ablation resistance. The chemical reactions occurring during the oxyacetylene torch ablation of the C/C-SiHfCB composites between 2000℃ and 2300℃ are consistent with (4), (5), (6), (7), (8), (9), (10), (11).
3.3.2. Ablation resistance mechanism of C/C-SiHfCB composites
The cross-sectional microstructure of the C/C-SiHfCB composites after exposure to an oxygenacetylene torch at 2000℃ for 60 s is depicted in Fig. 10. A glassy oxide protective layer with a thickness of approximately 50 μm was observed on the ablated composites surface. The EDS elemental maps and spectrum in Fig. 10 (b-f) indicate that the oxide layer primarily consists of SiO2, with a minor amount of HfO2. Apart from the SiO2, HfO2, and CO generated from the oxidation of ceramic phase within C/C-SiHfCB composites and a small portion of the C/C matrix directly exposed on the surface, the internal matrix of the composites remained largely unoxidized. As shown in Fig. 10, the outermost SiO2-porous HfO2 layer plays a critical role in regulating the ablation behavior of C/C-SiHfCB composites under a 2000℃ oxyacetylene torch. Acting as an effective oxygen diffusion barrier, this layer significantly restricts oxygen ingress into the composite interior.
Fig. 10. Cross-sectional SEM image, EDS elemental distribution and content of C/C-SiHfCB composites ablation center after 2000℃/60 s oxyacetylene torch ablation test: (a) cross-sectional SEM image, (b-e) the corresponding EDS elemental mapping of (b) Si, (c) C, (d) O, (e) Hf, and (f) EDS elemental content of (a). |
Consequently, the oxygen partial pressure within the interior is substantially reduced, effectively suppressing internal oxidation. Moreover, the SiO2 phase can form a relatively stable protective structure at elevated temperatures, which further enhances the barrier against oxygen transport. Therefore, SiO2 and HfO2 play a dominant role in determining the ablation resistance of C/C-SiHfCB composites under an oxyacetylene torch at 2000 ℃.
The cross-sectional microstructure of the C/C-SiHfCB composites after exposure to an oxygenacetylene torch at 2150℃ for 60 s is depicted in Fig. 11. A denser oxide protective layer, approximately 260 μm in thickness, was formed on the material surface. The EDS elemental maps and corresponding spectrum (Fig. 11 b-f) show that the ablation layers mainly consist of an HfO2 layer and an HfO2-SiO2 layer. The HfO2 layer contains only a small amount of Si, whereas the surface HfO2-SiO2 layer contains minor Hf. This difference is attributed to the ablation of molten SiO2 from the outermost oxide layer by the high-temperature oxyacetylene torch at 2150℃. Consequently, under these ablation conditions, Most of the SiO2 and HfO2 are generated by the oxidation of HfB2, HfC, and SiC within the C/C-SiHfCB composites. Instead, HfO2 forms a continuous oxide skeleton, with SiO2 partially filling the pores within this framework. This composite oxide structure effectively impedes oxygen ingress, thereby protecting the internal matrix of the composite.
Fig. 11. Cross-sectional SEM image, EDS elemental distribution and content of C/C-SiHfCB composites ablation center after 2150℃/60 s oxyacetylene torch ablation test: (a) cross-sectional SEM image, (b-e) the corresponding EDS elemental mapping of (b) Si, (c) C, (d) O, (e) Hf, and (f) EDS elemental content of (a). |
Fig. 12 presents a schematic of the ablation behaviors and mechanisms specific to the C/C-SiHfCB composites at different ablation temperatures. In Fig. 12 (a), at an ablation temperature around 2000℃, the ceramic phase within C/C-SiHfCB composites is oxidized, forming a protective layer composed of porous HfO2and SiO2. In this process, the porous HfO2 acts as a skeleton structure, minimizing the volatilization and loss of the surface oxide layer under high-temperature gas flow scouring. Simultaneously, surface pores are filled by the B2O3 and SiO2 in a viscous flow state, which reduces the oxygen infiltration rate and protects the internal matrix from further erosion. However, the liquid B2O3 exists only briefly before being vaporized into B2O3(g) and escaping from the surface. Concurrently, gases such as CO(g) and SiO(g) generated during ablation also dissipate from the composites surface. This leads to the formation of pores within the surface oxide layer. A substantial amount of liquid SiO2, which possesses higher temperature resistance, can rapidly fill these pores, thereby minimizing oxygen erosion of the material interior. As shown in Fig. 12 (b), with increasing oxyacetylene torch temperature and gas-flow scouring rate, surface oxidation becomes more severe. Under these extreme ablation conditions, part of the molten SiO2 in the outer oxide layer is volatilized or eroded by the high-temperature gas flow, resulting in the formation of pores and defects within the oxide layer. These pores subsequently serve as rapid diffusion pathways for oxygen, facilitating penetration into the oxide layer and thereby accelerating the oxidation reaction. Consequently, the oxidation process is intensified, leading to the growth and thickening of the surface oxide layer. Meanwhile, the oxidation of HfB2, HfC, and SiC within the C/C-SiHfCB composites generates SiO2 and HfO2. Due to their limited mutual solubility under these conditions, most of the SiO2 and HfO2 do not form a solid solution. Instead, HfO2 tends to form a continuous oxide skeleton, while SiO2 partially fills the pores within this framework. This composite oxide structure effectively impedes further oxygen ingress, thereby protecting the internal matrix of the composites.
Fig. 12. Schematic diagram of mechanism of ablation resistance of C/C-SiHfCB composites at different oxyacetylene torch ablation test temperatures: (a) 2000℃, (b) 2150℃. |
4. Conclusions
In this study, a relatively dense C/C-SiHfCB composite with a residual porosity of only 14% was successfully fabricated via six cycles of high-pressure PIP using a SiHfCB precursor to infiltrate a porous C/C matrix. The resulting HfB2, HfC, and SiC within the C/C-SiHfCB composites was continuously distributed within and tightly bonded to the matrix. Consequently, the composites exhibited a significantly improved average flexural strength of 237±42 MPa at room temperature, characterized by a desirable non-brittle fracture mode. Furthermore, the material demonstrated exceptional thermal shock resistance, with a critical temperature difference (△Tc) of 912℃. During oxyacetylene flame ablation tests, the linear ablation rates remained as low as 5.7×10-4 mm/s at 2000℃ and 15.6×10-4 mm/s at 2150℃, indicating outstanding ablation resistance. This performance is attributed to the synergistic effect of HfO2 and SiO2 formed during ablation: the SiO2 effectively sealed cracks and pores within the HfO2 skeleton, creating a robust barrier that hindered oxygen penetration and protected the internal matrix.
CRediT authorship contribution statement
Yang Lyu: Writing - original draft, Methodology, Conceptualization. Fei Li: Writing - review & editing. Chunlin Wang: Investigation, Formal analysis. Yuhao Fang: Investigation, Formal analysis. Wenzheng Zhang: Investigation, Funding acquisition, Formal analysis. Mingyi Tan: Funding acquisition, Formal analysis. Ping Hu: Investigation, Funding acquisition. Yuan Cheng: Funding acquisition, Formal analysis. Wenbo Han: Writing - review & editing, Resources. Xinghong Zhang: Supervision, Resources, Funding acquisition.
Declaration of competing interest
We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence our work, there is no professional or other personal interest of any nature or kind in any product, service and/or company that could be construed as influencing the position presented in, or the review of, the manuscript entitled “Efficient fabrication of light C/C-SiHfCB composites with excellent thermal shock resistance and high temperature ablation resistance above 2000℃”.