3.1 Microstructure of coated fiber and different composites
Figure 3 shows the thermogravimetric (TG), differential thermogravimetric (DTG), and differential thermal analysis (DTA) curves of phenolic resin pyrolysis at RT-1000 ℃. Carbon residue rate of the resin is about 38.8%, weight loss mainly occurs at 300-600 ℃, and the weight loss at 350-400℃ is the fastest (see DTG curve). The characteristic peaks of the DTA curve correspond to those of the DTG curve, indicating that obvious chemical reactions occur at specific temperatures. According to the trend of TG curve, the pyrolysis process can be divided into three temperature sections: in low temperature section (RT-300℃), total weight loss rate is about 6.69%. According to the DTG curve, following zero weight loss rate at 0-100℃, the weight loss rate reach peak (14.4·10−4/℃) at 100-150℃ corresponding to the endothermic peak of DTA, due to the endothermic process of moisture evaporation, and then the weight loss rate maintain at about 5·10−4/℃ (150-300℃), due to dehydration of the active hydrogen and hydroxymethyl groups during curing [18]. In medium temperature region (300-600℃), resin has the most weight loss of 38.8 %. At 450-600℃, weight decreases linearly at the rate of about 9·10−4/℃ with increasing temperature, while there is an obvious peak of weight loss rate at 300-450℃ (44·10−4/℃). In this temperature region, due to the violent thermal polycondensation and decomposition reactions, polymer chain gradually transforms into graphite crystallite structure [18] and the peak at 350-400℃ corresponds to decomposition of main chain. In high temperature region (600-1000℃), the weight loss is about 5%. The weight loss rate gradually decreases and tends to zero, with the weight loss reaching saturation (61.15%). The weight loss at this stage is due to the dehydrogenation of aromatic hydrocarbon [18].
Figure 4(a-b) and (c-d) show the surface morphology of degummed carbon fibers and carbon fibers coated with resin derived carbon, respectively. Many longitudinal grooves are exposed on the surface of carbon fibers after degumming. After preparation of fiber coating, a layer of dense and uniform resin-derived carbon coating is formed on the surface of carbon fibers, covering the grooves. The coating is well bonded to carbon fibers, with coating peeling and shedding only in very few areas. Fibers are well dispersed indicating that preparation of coating does not result in severe agglomeration between fibers.
With increasing sintering temperature from 1800℃ to 1900℃, density of the composites increases significantly. The density of T1800 is only 1.92 g/cm3, whereas the density of T1900 and T2000 increases to 2.14 g/cm3 and 2.18g/cm3, respectively, because the initial sintering temperature of SiBCN powder is about 1830℃. Only when the temperature is higher than about 1830℃, can particle rearrangement, high temperature creep and various diffusion mechanisms overcome the activation energy to function, causing obvious sintering between SiBCN particles. The theoretical density of amorphous SiBCN is generally considered as 2.83g/cm3, so relative density of T1800, T1900 and T2000 composites can be calculated as 71%, 81% and 83%, respectively. With the same sintering processes, the relative density of SiBCN monolithic sintered at 1800, 1900 and 2000℃ can reach about 74%, 88% and 93%, respectively [19]. Obviously, the addition of fibers obviously hinders the sintering densification, mainly due to bridging between fibers. The fiber bridging hinders the flow and rearrangement of particles and increases the diffusion distance, causing sintering mechanisms not to fully function. Therefore, defects such as pores and cracks tend to leave in fiber-dense area (see Fig. 5).
Figure 4(a-c) shows surface morphology of T1800, T1900 and T2000 composites parallel to pressing direction (vertical plane). Most of fibers are distributed in plane perpendicular to the pressing direction (horizontal plane), because during the hot pressing process, SiBCN powder drives fibers to flow, causing preferential distribution of fibers in the plane perpendicular to the pressure direction. In fiber dense area, different degrees of matrix shedding occurs, exposing the fiber end, because fiber bridging in the fiber-dense area causes difficulty of sintering densification and decrease of matrix strength. Besides, the axial thermal expansion coefficient of carbon fiber (-0.7·10-6) is much lower than the thermal expansion coefficient of the matrix (3-4·10-6) [19]. Due to the serious thermal expansion mismatch, thermal stress between the fibers and the matrix exceeds the tensile strength of the matrix, resulting in a large number of micro-cracks in the matrix and promoting the shedding of the matrix. With the increase of sintering temperature, the area of matrix shedding decreases gradually and only a small amount of matrix shedding occurs on the surface of T2000 composites (see in Fig. 5(c)), because further sintering densification enhances matrix strength and the fiber/matrix bonding strength. The illustration in the upper right corner of Fig. 5(a-c) is the surface morphology parallel to horizontal plane. Due to less fiber bridging in the horizontal plane, the surface is more flat and has less matrix shedding than that of vertical plane. With increasing sintering temperature, matrix shedding in horizontal plane also further decreases and there is almost no matrix shedding on the surface of T2000 composites, which further confirm the above analysis.
According to the XDR pattern of the composites in Fig. 6(b), the SiBCN matrix of the composites is obviously crystallized, and the precipitated phases include BN(C), α-SiC and β-SiC phase. β-SiC phase is main precipitated phases, and α-SiC is relatively less. With increasing sintering temperature, the diffraction peaks of β-SiC and α-SiC phases gradually become stronger and the half width at half-maximum (FWHM) significantly narrows down, indicating improvement of crystallization degree of SiC phase. The BN(C) phase in T1800 and T1900 composites has only a weak broad peak indicating limited crystallization degree, whereas the diffraction peak of BN(C) phase in T2000 composite is significantly intensified, indicating that BN(C) phase is obviously graphitized at 2000℃. Fig. 6(a) shows the microstructure of SiBCN matrix in T1900 composite observed by TEM. SiC nanocrystals are uniformly distributed in matrix, with mean grain size of about 30-50 nm, but a small number of SiC grains abnormally grown with grain sizes of about 100nm. According to the microstructure (see Fig. 5(a)) and electron diffraction pattern (see Fig. 5(b)) of the SiC grains in matrix, although SiC nanocrystals have high crystallization degree, there are high-density dislocations and twins in SiC nanocrystals. The BN(C) phase is mostly distributed in the edge of the SiC grain (see Fig. 5(a)). Due to low texture, its turbulent layer structure is not clear, which is consistent with the XRD analysis results. The segregation regions of Si and C atom in the amorphous matrix may become the crystal nucleus of SiC phase. With increase of temperature strengthening the thermal vibration of atoms, driven by the chemical potential, the smaller atoms of B and N can overcome activation energy and diffuse out from the segregation regions of Si and C forming the segregation regions of B and N around the SiC crystal nucleus. With the growth of SiC grains, the turbulent layer BN(C) phase is formed around the SiC grains.
3.2 Effect of sintering temperature on mechanical properties of composites
Figure 7(a) shows the bending stress-displacement curve of the composites. For T1900 and T2000 composites, fracture process presents pseudoplastic behavior: after elastic linear deformation, there are jagged stress steps similar to metal yield, due to the gradual pull-out and fracture of fiber bundles. For T1800 composite, fracture process has no obvious linear elastic deformation and yield point, showing typical plastic deformation behavior. The smooth curve without jagged platform indicates that too weak interface bonding leads to little resistance of fiber debonding during fracture. With increasing sintering temperature (1800-2000℃), the bending strength of the composites is significantly improved (30.4-70.5 MPa), but the deformation amount is gradually reduced (0.22-0.075). The main reason is that higher sintering temperature activates more sintering mechanisms to significantly reduce defects and improve matrix strength. Besides, with sintering temperature increasing (1800-2000℃), fiber coating has a certain degree of chemical bonding with matrix, causing higher interfacial bonding strength. Due to higher elastic modulus of carbon fiber (210 GPa) than that of SiBCN matrix (150 GPa) [19], carbon fibers are loaded prior to the matrix, so stronger interfacial bonding is conducive to more loading of the fiber. Compared with SiBCN monolithic prepared by the same hot pressing process, the bending strength of Cf/SiBCN composites is relatively lower, because fiber bridging, poor sintering property and thermal stress release lead to a large number of defects such as micro-cracks and pores in the composites. However, due to the addition of carbon fibers, the composites are not sensitive to the defects and present pseudoplastic behavior during fracture failure.
Figure 8(b-d) shows the fracture morphology of T1800, T1900 and T2000 composites, respectively. The fracture surface of composites is uneven, indicating that cracks have obvious deflection near fiber bundles. The large number of pulled out fibers and the residual pits indicate that debonding and pulling out are main toughening mechanisms after fiber fracture. As sintering temperature increases, fiber pull-out length gradually decreases due to increase of interfacial bonding strength. Specifically, stress concentration at interface cannot be alleviated by fiber debonding and at fiber ends, stress quickly reaches the fiber strength, resulting in a closer fracture position to the fiber end and shorter fiber pull-out length. Figure 8(a) shows fracture morphology of Cf/SiBCN composite without fiber coating prepared by hot pressing at 1800℃. Carbon fiber morphology becomes blurred with obvious distortion, indicating that even at lower sintering temperature (1800℃) fibers without coating have severe physical damage and chemical corrosion. Compared with T1800 composite, pull-out length of carbon fibers is shorter and many matrix particles are adhered on the surface, showing the characteristics of brittle fracture, because the chemical bonding between carbon fibers and matrix causes too strong interfacial bonding. The fractured fibers of T1800, T1900 and T2000 composites retains the original morphology without signs of corrosion and deformation damage indicating that fiber coating has good physical and chemical compatibility with matrix and protects carbon fibers from matrix corrosion and damage of thermal stress and external stress.
Figure 7(b) shows changes of bending strength, vickers hardness, elastic modulus and fracture toughness of Cf/SiBCN composites with sintering temperature. With sintering temperature increasing (1800-2000℃), comprehensive mechanical properties of composites are significantly improved: bending strength increases from 30.4 MPa to 70.5MPa, vickers hardness increases from 0.91 GPa to 2.25 GPa, and modulus increases from 20.3 GPa to 41.6 GPa, while the change of fracture toughness is not obvious (2.24-2.38 MPa·m1/2). The improvement of vickers hardness and elastic modulus is mainly due to reduction of defects in the matrix, and the little effect of sintering temperature on the fracture toughness is caused by offset effect of two factors. On the one hand, increase of sintering temperature leads to higher matrix strength and stronger interface bonding, so greater external stress is needed for debonding and pull-out of fibers and fracture of matrix, resulting in the improvement of fracture toughness. On the other hand, with sintering temperature increasing, stronger interfacial bonding decreases the number and length of fiber pull-out, resulting in decrease of fracture toughness. Compared with sintering temperature, fiber coating has a more significant effect on fracture toughness of composites. Fracture toughness of the Cf/SiBCN composites without fiber coating is only about 1.47 MPa·m1/2, indicating that fiber coating can effectively improve the toughening effect of fibers. According to its fracture morphology in Fig. 8(a), due to the serious damage of fibers and less pull-out of fibers, energy consumption of fiber fracture, debonding and pull-out is less, causing less fracture toughness.
3.3 The oxidation and thermal expansion properties of the composites
The thermal shock resistance of composites is determined by the oxidation resistance, mechanical properties and thermo-physical properties of composites. Therefore, the oxidation and thermal expansion properties were investigated to lay a foundation for the study on their thermal shock properties. Fig. 9(a) shows thermogravimetric curves of the composites during non-isothermal oxidation with temperature increasing from RT to 1400℃. The curves of all composite have the similar variation trend and according to the trend of the curves, oxidation process can be divided into three stages. In low temperature section (RT-620℃), water volatilization and gas desorption cause obvious weight loss (5%) at 180-260℃ and apart from that, weight remain constant, because carbon fibers and matrix remain chemical inertia to oxygen due to the low temperature.
In medium temperature region (620-1100℃), the weight of composites reduces rapidly, reaching the minimum at about 800-900℃, and then rapidly increases. In this temperature region, the weight variation of composites is due to the weight loss from the oxidation of carbon fibers with coating and the B2O3 volatilization by the reactions (1) and (3), and the weight gain from the oxidation of amorphous/nanocrystalline matrix by chemical reactions (2) and (4). However, due to the high oxidation activation energy and the protective effect of the oxide film, oxidation rate of h-BN and SiC in the matrix is quite slower than that of the carbon component at relatively low temperatures, so at 620-900℃, weight loss of carbon fibers is the dominant factor (see Fig. 11(a-d)). However, above 1000 ℃, h-BN is rapidly oxidized and SiC begins to undergo obvious oxidation [20–22]. Therefore, with the temperature increasing from 900 to 1100℃, the accelerated oxidation of SiC and h-BN leads to the weight gain (see Fig. 11(e-f)). Besides, with the outer carbon fibers consumed, a porous layer, containing amorphous/nanocrystalline matrix and its oxidation products is left outside, covering the unreacted core containing carbon component and the porous layer is gradually healed by molten oxidation product (see Fig. 12). Therefore, as temperature increases with oxidation time, the porous layer gradually becomes thicker and denser, causing more difficult internal diffusion of oxygen in the porous layer, so oxidation rate of carbon fibers decreases with oxidation time and temperature. The above factors lead to the obvious weight gain of composites at 900-1100℃. Compared with T1900 (906℃) and T2000 (891℃) composites, weight of T1800 composite starts to increase at lower temperature (860℃), with less weight loss (17%) due to the looser matrix. The looser matrix favors oxygen diffusion and oxidation faster in the matrix, causing more weight gain. In addition, more oxidation products can more quickly heal porous layer and hinder the further oxidation of carbon fibers, causing less weight loss.
$$\text{C}\text{(s)}+{\text{O}}_{2}\text{(g) }\text{→ }{\text{C}\text{O}}_{2} \text{(g)}$$
1
$$4\text{B}\text{N}\text{(s)}+{3\text{O}}_{2}\text{(g) }\text{→ 2}{{\text{B}}_{3}\text{O}}_{2} \text{(l)}+{2\text{N}}_{2}\text{(g)}$$
2
$${{\text{B}}_{3}\text{O}}_{2} \text{(l) }\text{→ }{{\text{B}}_{3}\text{O}}_{2}\text{(g) }$$
3
$$\text{S}\text{i}\text{C}\text{(s)}+{2\text{O}}_{2}\text{(g) }\text{→ 2}{\text{S}\text{i}\text{O}}_{2} \text{(s)}+{\text{C}\text{O}}_{2}\text{(g)}$$
4
In high temperature region (1100-1400℃), the weight of the composites increases linearly at a lower rate. After rapid oxidation at 900-1100℃, matrix is coated by oxide film, so oxidation rate of matrix is reduced significantly and more controlled by internal diffusion of oxygen in oxide film. Diffusion rate is less sensitive to temperature, so in this temperature region, weight increases linearly with temperature. Compared with T1900 and T2000 composites, T1800 composite still maintains a relatively higher weight gain rate, because its matrix is not dense with many micro-pores and is difficult to form a dense oxide film, aggravating oxidation weight gain of matrix.
Except for oxidation properties, coefficient of thermal expansion (TCE) is one of the important indicators that determine its thermal shock resistance. Fig. 9(b) shows the TCE of the Cf/SiBCN composites perpendicular to hot pressing direction at different temperatures (RT-1200℃). The TCE of the T1800, T1900 and T2000 composites all increases with temperature (0.43-1.6·10−6 K−1, 1.6-3.18·10−6 K−1, 1.33-3.74·10−6 K−1) and has the similar increasing trend. In lower temperature region (RT-600℃), the TCE increases rapidly, and in higher temperature region (600-1200℃), the TCE increases slowly and tends to constant. The average TCE of the composites is smaller than that of SiC, SiC-BN and SiBCN monolithic, due to the limitation of fibers. Fibers are mainly distributed in horizontal plane and the axial TCE of fibers (-0.14-1.7·10−6/K) is much lower than matrix (4-5·10−6/K). During the heating process, TCE mismatch of matrix and fibers in the horizontal plane causes compressive stress in matrix to limits the thermal expansion of the matrix. With increasing sintering temperature, the thermal expansion coefficient of the composites increases: the TCE of T1900 and T2000 composites is close to that of C/SiC composites, while the TCE of T1800 composite is much lower, due to more amorphous phase and micro-pore in the matrix. Many atom gaps in amorphous phase, due to the disorder arrangement of atoms, provide space for the thermal vibration of atoms, and more micro-pores, due to lower relative density, provide space for the thermal expansion of grains, resulting in a lower thermal expansion coefficient.
3.4 Effect of sintering temperature on thermal shock properties of composites
The thermal shock resistance of the composites is characterized by residual bending strength after thermal shock. Fig. 10 shows the residual bending strength of the composites after thermal shock at different temperatures (700-1100℃). After thermal shock at 700 ℃, the bending strength of T1900 and T2000 composites is reduced by 40.9% and 43.1%, respectively. For T1900 and T2000 composites, due to stronger interface bonding, the thermal stress at the interface is hard to be alleviated by fiber debonding and the tensile stress in the matrix is more likely to accumulate to the yield strength of the matrix causing more cracks in the matrix and decrease of mechanical properties. In addition, at 700℃ although fibers are less damaged (see Fig. 11(a)), resin-derived carbon coating as interfacial phase is damaged seriously due to oxidation and thermal stress, so the interface with weakened bonding strength cannot effectively transfer the load to fibers (see Fig. 11(b)), resulting in a significant decline in mechanical properties. However, after thermal shock at 700 ℃, the strength of T1800 reduces by only 20.4%. For one thing, due to the lower CET, the internal stress caused by thermal shock is less and due to the weak interface bonding, the thermal stress at the interface can be alleviated by fiber debonding, causing less damage to matrix. For another, due to the weak interface bonding, interface in T1800 composite originally cannot transfer the load to fibers effectively, so after thermal shock, the further damage on interface has less effect on mechanical properties.
With further increasing thermal shock temperature (T=900℃, 1100℃), the strength of the composites decreases further, but the decrease is limited. Although the oxidation of carbon fibers is significantly intensified (see Fig. 11(c-f)), its effect on mechanical properties is less because strength of the composites is more determined by interface bonding strength rather than fiber strength. After the thermal shock at 1100℃, the color of the composite surface turns white. Based on the EDS analysis (Table 1) and SEM observation, the white material is oxidation products (SiO2), which heals pores, cracks, inhibiting further decrease of composite strength. In summary, due to the lower coefficient of thermal expansion, lower fiber loading ratio, less stress concentration at the interface, and better defect healing effect, T1800 composite has the best thermal shock resistance, which is consistent with the analysis of the oxidation and thermal expansion properties of the composites.
Figure 11 shows the surface and fracture morphology of T2000 composite after thermal shock at different temperatures. As shown in Figure 11(a-b), after thermal shock at 700℃, there is no obvious oxidation of the matrix. Carbon fibers are only slightly oxidized: the active points of the carbon fibers are oxidized to form pits. However, obvious oxidation of the resin-derived carbon interface forms obvious gaps between fibers and matrix, causing the failure of the fibers to load through the interface, and significant decrease in the mechanical properties, as is analyzed above. Compared with the fracture morphology before thermal shock (see Fig. 8(d)), the fiber pull-out length increases considerably, due to the failure of the interface.
Figure 11(c-d) shows the surface morphology and fracture morphology of T2000 composites after thermal shock at 900℃. The carbon fibers on surface are oxidized seriously to shape of bamboo shoots and even are completely consumed to leave many holes as channel of oxygen internal diffusion. At the fracture surface, the carbon fibers near the surface are completely consumed, and the carbon fibers away from the surface are not oxidized obviously due to the excessive oxygen diffusion resistance. Figure 11(e-f) shows surface morphology and fracture morphology of T2000 composite after thermal shock at 1100℃. The carbon fibers on surface are completely consumed to leave holes. According to EDS analysis results, the matrix is obviously oxidized to form molten SiO2 (see in Table 1) that tends to cover the surface and heal pores and cracks. According to the fracture morphology, deeper carbon fibers are also obviously oxidized, indicating that the oxygen diffusion depth further increases.
Figure 12 (a-d) shows enlarged view of the fracture morphology from the surface to the inside successively, and obviously the oxidation degree gradually decreases with depth. The fibers of the composite near the surface are almost completely oxidized, leaving large holes in the matrix and the matrix is oxidized to form the oxidized layer with the thickness of 3-4µm (see Figure 12(a)). At depth of around 100µm, fibers are oxidized seriously, with bamboo-shoot shape and the fiber surface is covered by molten oxide that plays a protective role of coating (see Figure 12(b)). Based on the EDS analysis, the main component of the molten oxide is SiO2 (see Table 1). Similarly, the spherical molten oxide melt is also formed on the surface of hole left by fiber oxidation, indicating that the oxidation product can heal pores and cracks, hindering oxygen diffusion and further oxidation of composites. At depth of about 200µm, fiber front becomes significantly thinner, the interface phase is completely oxidized, and there is clear gap between fibers and matrix (see Figure 12(c)). Due to weak bonding or even complete debonding of fibers and matrix, the fiber pull-out length is significantly greater than that before the thermal shock. However, at depth of 300µm, the fibers still maintain their initial morphology and have good bonding with the matrix (see Figure 12(d)). The fiber pull-out length is close to that before the thermal shock, indicating that the degradation of interface bonding strength is less.
Table 1
Element analysis result at different spot by EDS
EDS
|
Atomic content of element %
|
Si K
|
C K
|
O K
|
Spot 1
|
0
|
100
|
0
|
Spot2
|
31.53
|
0
|
78.47
|
Spot3
|
37.21
|
0
|
62.79
|
In summary, damage to the interface makes fibers fail to carry load effectively causing the decrease in the strength of T2000 composite. Due to the high activity of the resin-derived carbon, the strength of the composite is significantly degraded after thermal shock at a lower temperature (700℃). With the temperature difference further increasing, although carbon fibers are more severely oxidized, further degradation of strength is relatively less, due to the damaged interface and the self-healing effect of oxidation product.