Effects of heat treatment temperatures on microstructures and mechanical properties of the chopped carbon fibres SiC composites

Raw materials and fabrication

The chopped carbon fibre paper (C_sf-paper) was fabricated by wet papermaking, which involves several steps: solution preparation, fibre dispersion, vacuum filtration. Afterwards, the C_sf/SiC composites were obtained after laminating and CVI. More specific experimental procedures are described as follows.

Commercially available carbon fibres (T-300^™) were used as the reinforcement of ceramics, with a diameter of 7 µm and an average length of 4–7 mm. By the aid of 10-minute ultrasonic dispersion and nonionic surfactant triton (Triton X-100, 1 wt-%), the chopped carbon fibres (C_sf) were dispersed homogeneously in carboxymethyl cellulose (CMC) at a concentration of 0.4 g/400 ml [13]. The suspension was vacuum-filtrated through a polytetrafluoroethylene filter membrane with a 0.22 µm pore size under the pressure of −0.1 MPa. After filtering, the wet C_sf-paper with filter membrane was blown out from vacuum filter to a piece of filter paper. Afterwards, the filter membrane was peeled to obtain a single wet C_sf-paper with a diameter of 100 mm and thickness of ca. 500 μm, and then dried it in an oven at 70°C for 30 min.

Afterwards, six layers of C_sf-paper were laminated to produce a multilayer template. Graphite fixture with a thickness of 3 mm was applied to put pressure on the pile of single layers to get a multilayer template. After that, fibre mat preforms have been accomplished successfully. Finally, pyrolytic carbon (PyC) was introduced onto the fibres as interface; followingly, SiC matrix was infiltrated into preforms via CVI. At 960°C, C₃H₆ was selected as source materials and introduced into a reaction chamber. PyC interface was fabricated on the fibres in Ar. PyC interface has a good effect in protecting the carbon fibres, alleviating the interface stress and reducing the interface bonding strength. SiC matrix was introduced into the preforms under the conditions of 1000°C. The CVI of the matrix was carried out for eight times (72 × 8 h). The specific CVI process is as follows: First of all, the CVI method was used to infiltrate SiC matrix for two times (72 × 2 h), and re-infiltrate for two times after demolding, which is followed by roughing to polish off SiC coating and to open the closed pores. Afterwards, CVI was carried out for two times to infiltrate SiC matrix to obtain C_sf/SiC composites with higher density. After standard three-point bending specimens were prepared, final three-point bending specimens were prepared by CVI for two times for further densification. Here, methyltrichlorosilane (MTS, CH₃SiCl₃), argon and hydrogen were selected as source material, diluent and carrier gas, respectively. The molar ratio between H₂ and MTS was 10.

Characterisation

Archimedes principle was employed in the measurement of bulk densities and open porosities of the composites. All the samples were cut into the dimension of 36 mm (length) × 3 mm (width) × 2 mm (thickness) and separated into five groups. The five groups were heat treated at different temperatures in room temperature (RT, i.e. without heat treat), 1200°C, 1400°C, 1600°C and 1800°C. They were denoted as S0, S12, S14, S16 and S18, respectively.

A three-point bending test (SANS CMT 4304, Sans Materials Testing Co., Shenzhen, China) was performed at room temperature to measure the flexural strength with a span of 30 mm. The loading rate was 0.5 mm min⁻¹. The data for each specimen were averaged over five tests to ensure the accuracy of the data. Microstructures of the specimens were observed through scanning electron microscopy (SEM, Hitachi S-2700, Tokyo, Japan) while the phase composition was identified by X-ray diffraction (XRD, Rigaku-D/Max2400, Rigaku Corp., Tokyo, Japan).

Technologic predominance

Using a vacuum filtration model as shown in Figure 1, the C_sf was found to show a homogeneous dispersion in the typical C_sf-paper fabricated by vacuum filtration technique. The model of Figure 1 is a process in which the dispersed fibres in the suspension gradually fall on the 2D plane into a single-layer wet paper during vacuum filtration. Meanwhile, the amount of chopped fibres in a single-layer wet paper can be controlled by adjusting the amount of chopped fibres in the suspension, as shown in Figure 1. The fibres distributing randomly in 2D plane are shown in Figure 2. Therefore, vacuum filtration could effectively avoid single fibre orientation which is inevitable in tape casting, realising the contractive and densified isotropy. Compared with the ball-milling method (where fibre length reduction is inevitable [15]), the solution dispersion and further vacuum infiltration treatment in this paper can obtain homogeneous fibre dispersion and avoid fibre length reduction simultaneously, which ensured the integrity and reinforcing effects of the fibres. Overlapping each other, the fibre sheet with layered structure was obtained. The excellent structural integrity of preform was maintained after SiC infiltration for eight times. There was no significant damage or breakage in the fibres during the CVI process, implying that compared with the sintering method in which fibres would be damaged anyhow, the introduction of matrix through CVI process could be a useful way for connecting matrix and fibres. Figure 3 shows the fracture surface of samples under different heat treatment temperatures. Infiltrated SiC effectively filled the space in the template and fibres were connected through SiC matrix. The matrix was devised by pores and the fibres’ pull-out were more apparent with the temperature increases. This indicates that the bonding strength between fibres and matrix decreased with the treatment temperature increases, which does harm to the strengthening effect.

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Figure 1

The model to simulate vacuum filtration of C_sf-paper.

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Figure 2

Typical SEM image of (a) surface morphology and (b) microstructure of fibre sheet without CVI.

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Figure 3

The fracture morphologies of the composites at different heat treatment temperatures (a) 1200°C, (b) 1400°C, (c) 1600°C and (d) 1800°C.

The volume fraction of the fibres in the composites was controlled by the lamination of C_sf-paper. Compression was applied to define the thickness of preforms after lamination. Meanwhile, it was evident that the number of the effective bonding points between inter-layer (fibres between different layers) and intra-layer (fibres in the single wet paper layer) fibres increased and inter-layer gaps decreased after compression, indicating the practical pre-layering. In this work, SiC matrix was introduced into fibre sheets by CVI for eight times. Relatively, a weak interface between matrix and fibres formed in chronological CVI process. Fibres’ pull-out could be observed in the fracture surface as shown in Figure 3, which has been documented beneficial to toughen.

Composition of composites

Figure 4 reveals the XRD patterns of specimens treated at different temperatures. Only two types of peaks were discovered, which is assigned to β-SiC as the main phase and α-SiC as a secondary phase, respectively. In addition, a significant development of β-SiC could be observed as the heat treatment temperature rose while there was not any fluctuation in the intensity of α-SiC for all processing temperatures. As β-SiC could be stable at low temperature and translate to high-temperature stable α-SiC only when the temperature above 1800°C, this variation of the relative content between the two implied the presence of residual carbon in the composite. That residual carbon brings harmful effects to both the densification of SiC matrix composite and shrinkage of the whole composite, producing a large amount of porosity which plays an important role in the reduction of flexural strength [18-20].

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Figure 4

XRD patterns of the specimens treated at different temperatures.

Mechanical properties

Table 1 indicates that the C_sf/SiC composites fabricated by vacuum infiltration have advantages in comprehensive mechanical properties over those obtained through other methods, such as tape casting, reaction bonding and spark plasma sintering. Compared with the samples obtained by tape casting, samples in this work performed slightly high flexural strength and low elastic modulus. The slightly low modulus might be due to the random fibres orientation in wet papermaking technology, while fibres tend to align along the tape casting direction in tape casting process. Apparently, the fracture toughness of S0 was nearly twice that of reaction bonded samples when the content of fibres was equivalent, which was in agreement with the efficiency of fibre pull-out. S0 had not only the lowest bulk density among all these samples but also excellent mechanical properties, indirectly indicating a potential improvement of the fracture toughness of S0.

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Table 1

Mechanical properties of C_sf/SiC composites fabricated by vacuum infiltration, tape casting, reaction bonding and spark plasma sintering.

	Fibre content (vol.-%)	Bulk density (g cm−3)	Flexural strength (MPa)	Fracture toughness (MPa m1/2)	Elastic modulus (GPa)
Vacuum filtration	10	2.53	269	5.8	164.7
Tape casting [5]	10	/	206	/	244
Reaction sintering [13]	10	2.83	213	3.2	/
Spark plasma sintering [21]	20	2.65	118	/	141

Fluctuation of flexural strength and Young's modulus as the temperature increases is shown in Figure 5. The flexural strength reached a maximum value of 306 MPa at 1200°C, and decreased continuously with the increase of temperature. Maximum modulus, 196.5 GPa, was also obtained at 1200°C. Variation of Young's modulus showed a similar trend with the flexural strength ranging from room temperature to 1600°C. However, the modulus increased at the temperature of 1800°C while flexural strength decreased. From Figure 6, the flexural stress–displacement curves were observed. The slope of the curves changed with the change of temperature, following the rule of S12 > S14 > S18 > S0 > S16, which corresponded to the variation of Young's modulus in Figure 5.

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Figure 5

Variation of flexural strength and Young's modulus with the change of heat treatment temperature from room temperature to 1800°C.

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Figure 6

The flexural stress–displacement curves of the composites at different heat treatment temperatures.

Heat treatment plays a significant role in changing the morphology and distribution of the microcracks and pores, and accordingly leads to the adjustment of the microstructure of the composites. Moreover, heat treatment influences not only the surface activity and chemical injury of fibres, but also the interfacial bonding strength. The heat treatment at proper temperature will eliminate the residual thermal stress produced during the preparation process, which means that appropriate heat treatment temperature contributes to the enhancement of the fracture toughness. This process corresponds to the improvement of mechanical properties when the temperature reached 1200°C. However, excessive processing temperature will hinder relevant energy dissipation such as fibre debonding and crack deflection, which brings a negative impact on fracture toughness. Excessive processing temperature will hurt fibres chemically and weaken the interfacial bonding strength. And microcracks appear because the expansion of the matrix and fibre under high temperature does not match; thus, mechanical properties worsen [22-23].

Toughening mechanism of the composite

As for the fibres reinforced ceramic matrix composites, the fracture toughness is mainly determined by the interface strength [14]. Proper bonding strength can stimulate crack deflection and hold back regional crack propagation in the meantime [24-25]. Fracture morphologies of S12, S14, S16 and S18 were detected by SEM, and the results were exhibited in Figure 3(a–d). It could be observed that fibre pull-out became more obvious as the processing temperature went up from 1200°C to 1600°C, especially for S16 which fibre layered pull-out became very evident when fracture occurred. By comparison, the fracture of S18 was smoother with longer fibre pull-out length.

Appropriate heat treatment temperature contributes to the enhancement of the mechanical properties because the residual thermal stress from composite fabrication process can be released. Meanwhile, this also reduces the interfacial bonding strength to ensure that the fibre and matrix can be de-bonded. Before pulling out, the fibre has to debond with the matrix in the first place. The fibre debonding presents new interface and then gives rise to the increase of interfacial energy. Fibre pull-out consumed considerable fracture energy, which was regarded as a helpful toughening mechanism. Longer fibre pull-out length implies better toughening results. So it needs a certain amount of systematic energy to expend. However, an excessively high-temperature treatment does cause the mismatch of thermal expansion between matrix and fibres, which results in cracks. Although the cracks do not reduce the strength of the material, the thermal expansion does. The stress from thermal–physical incompatibility causes a great damage to the fibre, which weaken the stress transfer between matrix and fibres [24,26].

However, an increase in modulus of 1800°C might be the consequence of the increase in bonding force originated from the sintering of SiC at 1800°C. This may explain why there is a variation of flexural strength and Young's modulus according to Figure 5 with heat treatment temperature ranging from room temperature to 1800°C.