Investigation on microstructure and mechanical properties of short-carbon-fiber/Ti 3 SiC 2 composites

In this paper, short-carbon-fibers (C sf ) reinforced Ti 3 SiC 2 matrix composites (C sf /Ti 3 SiC 2 , the C sf content was 0, 2, 5 and 10 vol.%) were fabricated by spark-plasma-sintering (SPS) using Ti 3 SiC 2 powders and C sf as starting materials at 1300 o C. The effects of C sf addition on the phase compositions, microstructures and mechanical properties (including hardness, flexural strength and fracture toughness) of C sf /Ti 3 SiC 2 composites were investigated. The C sf , with a bi-layered transition layers, i.e. TiC and SiC layer, were homogeneously distributed in the as-prepared C sf /Ti 3 SiC 2 composites. With the increase of C sf content, the fracture toughness of C sf /Ti 3 SiC 2 composites increased, but the flexural strength decreased, while the Vickers hardness decreased initially then increased steadily when the C sf content was higher than 2 vol.%. These changed performances could be attributed to the introduction of C sf and the formation of much stronger interfacial phases.


Introduction
Ti3SiC2, a typical member of the MAX phase family, is a ternary layered compound. It has attracted extensive attention for its unique combination of metals and ceramics properties, such as low density, easy machinability, high melting point, high tensile strength and damage tolerance, excellent thermal shock resistance, good chemical stability and oxidation resistance (below 900 o C) [1][2][3][4][5][6][7]. Therefore, Ti3SiC2 is a potential structural material for the applications in the environments of high temperature, oxidation, corrosion and wear etc. [1,2,7,8]. However, low fracture toughness and 3 hardness limit its practical applications. The introduction of secondary phase as reinforcement into Ti3SiC2 has been proved to be an effective approach to overcome these disadvantages. For example, numerous reinforcements, including TiB2 [9], SiC [10], c-BN [11], TiC [12], Al2O3 [13] and ZrO2 [14], have been introduced to improve the mechanical properties of Ti3SiC2. Commonly, the additions of ceramic particles can improve Vickers hardness and flexural strength of Ti3SiC2 matrix composites. For example, Górny et al. [9] reported that the addition of TiB2 could effectively improve the Vickers hardness and elastic modulus of Ti3SiC2. Tian et al. [12] prepared Ti3SiC2-TiC composites and found that the Vickers hardness of the composite increased with increasing TiC content up to 90 vol.%, the flexural strength was enhanced by 64% when the content of TiC was 50 vol.%. Wang et al. [13] fabricated Ti3SiC2-xAl2O3 (x= 0-20 vol.%) composites by spark-plasma-sintering (SPS), found that Ti3SiC2-20 vol.% Al2O3 exhibited the highest Vickers hardness, while other mechanical properties deteriorated because of the aggregation of Al2O3 in the composite. On the other hand, the toughness of Ti3SiC2 matrix composites has hardly ever been strengthened due to intrinsic brittleness of these ceramic phases.
Carbon fibers, with numerous attractively comprehensive properties, such as low density, high strength and modulus, excellent chemical inertness and small thermal expansion coefficient, are considered to be the most promising reinforcing phase for many structural ceramics [15][16][17]. Especially, Short-carbon-fibers (Csf) have been applied extensively in the preparation of ceramic matrix composites by using SPS or 4 hot-pressing (HP), due to their low cost, easy to add, and chemical inert at high temperatures. Interestingly, Csf reinforced ceramic matrix composites usually possess relatively higher fracture toughness. For example, Wang et. al. [18] fabricated Csf (up to 1 wt.%) reinforced B4C composites, and found that the composites had higher fracture toughness compared with monolithic B4C due to the occurrence of crack deflection and bridging resulting from interface debonding between fiber and matrix. Similarly, when Csf was introduced in ZrB2-SiC composites, the toughness of the composites was significantly improved, resulted from fiber debonding, pulling-out and bridging as well as crack deflection [19]. Unfortunately, to date, only Lagos et. al. prepared Ti3SiC2-Cf composites by SPS, and processing, microstructure and thermo-mechanical properties (thermal expansion coefficient and thermal conductivity) were investigated [20], while the mechanical properties of Csf reinforced Ti3SiC2 composites are not available.
In this work, Csf/Ti3SiC2 composites (the content of Csf was 0, 2, 5 and 10 vol.%) were fabricated by using SPS technique. The interfacial microstructure between Csf and Ti3SiC2 matrix and effects of Csf introduction on room-temperature mechanical properties of Csf/Ti3SiC2 composites were systematically investigated, and the role of Csf addition played in reinforcement effect was also discussed.

Materials preparation
Ti3SiC2 powders (Forsman Technology Co., Beijing, China) and Csf (ZLXC, Co., 5 Cangzhou, China) were used as raw materials. The purity and mean particle size of Ti3SiC2 powders was 98% and 3 ~ 5 m, respectively, and the main impurity was Al2O3.
The preparation process of Csf/Ti3SiC2 composites is schematically illustrated in Fig. 1. Firstly, Csf was dispersed into deionized water with carboxymethyl cellulose sodium (CMC-Na) as the dispersant [21]. The mass ratio of CMC-Na and Csf was 1:3.
Then, Ti3SiC2 powders were added into the as-prepared solution. After stirring with a magnetic stirrer, the slurry of Ti3SiC2 powders and Csf was obtained. Dried in an oven, the mixed material was packed into a cylindrical graphite die. Followed that, a green body with a diameter of Φ 40 mm was obtained by cold pressing under a load of 20 MPa. Subsequently, the sintering was conducted in SPS facility. After the vacuum of 10 Pa in the sintering chamber was acquired, the green body was heated to 1300 °C, under a constant pressure of 40 MPa soaked for 8 min, and the heating rate of 50 o C/min. Then cooled to room temperature in the sintering chamber. All composites were fabricated under such identical SPS condition.

Characterizations of composition and microstructure of the composites
Determinations of actual densities of Csf/Ti3SiC2 composites by Archimedes' method. The phase compositions were performed on a D/max-2400 X-ray diffractometer (XRD) (Rigaku, Tokyo, Japan) with a Cu Kα radiation (λ = 0.1542 nm).
The tube voltage was 50 kV, and the current was 100 mA. The 2θ range was (8°-80°) with count time of 1 s per 0.02°/step. The microstructure, cross-section morphologies and fracture surfaces of the composites were performed by scanning electron microscope (SEM, Oberkochen, Germany, EHT=20.00 kV) with energy dispersive spectrometer (EDS, Oxford Instruments, UK) in vacuum. 7

Determinations of mechanical properties
For mechanical tests, the samples were cut into a series of bars from the as-prepared composites by electrical discharge machining. Before testing, the surfaces of all samples were sanded to 2000 # SiC sandpaper, polished to a mirror surface with 1.0 mm diamond paste, then ultrasonically cleaned in ethanol and distilled water, finally dried.
Vickers hardness test was carried out using Vickers indenter (432SVD, WOLPERT, USA) at a load of 9.8 N for 15 s. The parallel operations were conducted for nine times, finally the mean value was obtained for each sample.
All mechanical performance tests were conducted on a universal testing machine Where, S is span distance, C is notch length, B and W are the width and height of the sample, respectively. which will be discussed in the following section.

Microstructures of the composites
SEM micrographs of the polished surfaces of the as-prepared Ti3SiC2 and Csf/Ti3SiC2 composites are presented in Fig. 3. Minor micro-pores existed in the Ti3SiC2 sample ( Fig. 3(a)), the black spots belong to the Al2O3 phase. For the Csf/Ti3SiC2 composites ( Fig. 3(b, c, d)), the black circle-like and stripe-like Csf were uniformly dispersed in the gray Ti3SiC2 matrix, and exhibited various orientations in a three-dimensional space. Additionally, the length of Csf in the Ti3SiC2 matrix was apparently shorter than their original length (3-5 mm). The reduction in Csf length is caused by the applied pressure during sintering process or Csf might hide in the inner of the samples [23]. And, no pores were found in these composites, indicating that the well dispersed Csf could promote the densification of Ti3SiC2 matrix. The most conceivable 10 reason is that the incorporation of Csf resulted in the existence of more phase boundaries, which is beneficial to the elimination of pores. In addition, Csf well dispersed in the Ti3SiC2 matrix also confirms the desirability of preparing such composites by this method. To get a better understanding of the interface reaction and phase evolution, the polished cross-section morphology of 10 Csf/Ti3SiC2 was observed by SEM, and element distribution was identified by EDS, the results are presented in Fig. 4.
According to Fig. 4 (a) and (b), there were two carbon fibers, which orientation was perpendicular to the viewing plane. No pores and defects appeared at the interface zone, and a chemical reaction between Csf and Ti3SiC2 matrix was observed. The chemical 11 compositions of the selected zones 1-5 in Fig. 4 (b) were determined by EDS, as listed in Table 1. It was found that the interface phase with duplex structure was formed between Csf and Ti3SiC2 matrix. Based on EDS analysis results, the inner interface layer adjacent to Csf was identified as TiC, while the outer interface layer was SiC. Such result was quite different from the results of Ti3SiC2-Cf composites prepared by Lagos et al [20]. The main reason for the formation of the interface microstructure is that our experiment actively reduces the sintering temperature of the samples to slow down the interface reaction activity, and increases the holding time from 5 min to 8 min to increase the relative density of the Csf/Ti3SiC2 composites.  Fig. 4 (a); element distributions of (c) Ti, (d) Si, (e) Al, and (f) C. Fig. 4(b). The interfacial reaction mechanism can be proposed tentatively based on the above SEM observations and EDS results. From Fig 2 and Fig 3, we deemed that Ti3SiC2 can maintain thermal stability at about 1300 o C. The above-mentioned phenomenon could be attributed to the environment-dependent decomposition behavior of Ti3SiC2, i.e., a carbon-rich environment will promote the decomposition of Ti3SiC2 and transform them into TiC0.67 and Si [24][25][26], as described by the reaction (4).

Table 1 EDS results of the marked spots 1-5 in
Ti3SiC2 → 3 TiC0.67 + Si (g) (ΔG (1600 K)=181 kJ/mol) Furthermore, carbon fibers constitute a C source whereas reaction between carbon fibers and Ti3SiC2 (Eq. 4) leads to the formation of a Si source. Thus, SiC formation is promoted (Eq. 5) [26], and results from out diffusion of C atoms through the formed TiCx layer.
The formation of a nearly stoichiometric TiC phase is due to the diffusion of C from the Csf into the carbon vacancies of TiC0.67, as shown in the reaction (6).
C + TiC0.67 → TiC As a result, a double-layered interface phases are formed. Such a double layer 13 interface can serve as a C diffusion barrier and it prevent carbon fibers to provide a C source. Therefore, the decomposition of Ti3SC2 is inhibited.
From the EDS result in Fig. 4(e), the bright agglomerates were enriched in Al.
Meanwhile, the EDS analysis also confirmed the existence of minor well-dispersed Al2O3 particles with the size of 2-3 μm, which corresponds to the XRD results.
The actual densities of the as-prepared composites are listed in Table 2. With increasing the content of Csf, the density of the composite decreased. This is because that the density of Csf (1.78 g/cm 3 ) was much lower than that of Ti3SiC2 (4.53 g/cm 3 ).
Based on the nominal ratio of the initial contents of Ti3SiC2 and Csf, the theoretical densities of the as-prepared Ti3SiC2 and Csf/Ti3SiC2 composites were calculated using the rule of mixtures, also listed in Table 2. Obviously, the relative densities of all materials prepared by SPS were higher than 98% in this work. In addition, under the condition of the same SPS processing, the relative densities of the as-prepared Csf/Ti3SiC2 composite increased with increasing Csf content. It should be noted that, during the calculation of the theoretical densities of the composites, only the nominal compositions of the composite were considered, the formation of high density TiC (4.93 g/cm 3 [1]) during the sintering of the bulk material was neglected. Therefore, correspondingly, the higher contents of Csf as well as TiC led to the greater deviation of calculated density.

Vickers hardness
The dependence of Vickers hardness of Ti3SiC2 and Csf/Ti3SiC2 composites on the theoretical content of Csf is shown in Fig. 5. The Vickers hardness of the pristine Ti3SiC2 sintered by SPS in the present work was 5.40 ± 0.06 GPa (measured at the indentation load of 9.8 N). It was reported previously that the hardness of monolithic Ti3SiC2 was about 4 GPa [3], lower than that of Ti3SiC2 synthesized by SPS in this work. The reasons may be due to our Ti3SiC2 sample contains a small amount of Al2O3 phase, it may also be due to the impact of indentation size. El Raghy et al. [3] found that indentation size had a great influence on the hardness. From Csf/Ti3SiC2, similar situation took place. Although the increased contents of TiC and SiC gave rise to the increased hardness of the composite, which was still lower than that of the matrix Ti3SiC2. The hardness of the composite was degraded due to the chemical reaction between Csf and Ti3SiC2 matrix, but further evaluate the content of the interface reaction products to compensate for the decrease in the hardness of the composite is beyond the scope of this paper. When the content of Csf was 10 vol.%, the hardness of the as-prepared composite was 6.69 GPa, which was larger than that of the Ti3SiC2. For the 10 Csf/Ti3SiC2, the contents of formed TiC and SiC interfacial phases with higher hardness was sufficient to compensate the hardness loss by the introduction of Csf.
Zhang et al. [29] presented that the maximum Vickers hardness of Ti3SiC2-40 vol.% TiC composite was about 13 GPa. Accordingly, it is reasonable to believe that the contents of added Csf and TiC derived from the interfacial reaction between Csf and Ti3SiC2 simultaneously determine the hardness of Csf/Ti3SiC2 composite. It also can be clearly seen from Fig. 6 that as the increase of Csf content, the fracture toughness of Csf/Ti3SiC2 composites increased, reaching a maximum value of 17 6.48 ± 0.29 MPa•m 1/2 for 10 Csf/Ti3SiC2. After adding Csf, the fracture toughness of Csf/Ti3SiC2 composites were enhanced, which was related to the activation of certain kinds of toughening mechanisms [30,31], as described as follows.

Fig. 6
The dependence of flexural strength and fracture toughness of the as-prepared Csf/Ti3SiC2 composites on the content of Csf.
The mechanical properties of fiber-reinforced composites rely on not only the intrinsic properties of fiber and matrix, but also the characteristics of fiber/matrix interface [32,33]. Interface delamination and the formation of a "weak" interphase are significantly vital to the comprehensive properties of fiber-reinforced ceramic matrix composites [34]. During the failure process of composites, their fracture toughness could be improved effectively by interface debonding, fiber bridging and fiber pulling-out caused by such characteristics mentioned above. To further understand the possible toughening mechanisms of the as-prepared Csf reinforced Ti3SiC2 composites, the fractured surfaces of the composites after SENB test were observed by SEM, as shown in Fig. 7. Obviously, these composites exhibited a fully brittle fracture. The breakage of Csf and interface debonding could be found, but fiber pull-out did not 18 appear, corresponding to a strong interface bonding between Csf and Ti3SiC2. During the preparation of Csf/Ti3SiC2 composites, original Csf did not undergo any surface treatment. As mentioned above, the formation of interfacial phases of TiC and SiC during heat-pressing sintering caused an enhanced interface bonding between Csf and Ti3SiC2 matrix, and such chemical bonding was much stronger than van der Waals' force in the matrix [35,36]. Therefore, the fiber bridging or fiber pulling-out became very difficult to occur due to the presence of strong interface between fiber reinforcement and matrix [33,37]. On the other hand, under this condition, as the Csf content increases, more energy is consumed during fracture process due to interface debonding and Csf breakage, thereby resulting in the increased fracture toughness of the as-prepared composite.

Conclusions
Csf reinforced Ti3SiC2 composites were prepared by SPS process. The behavior of Csf in the as-prepared Csf/Ti3SiC2 composites, microstructures and mechanical properties of the as-prepared Csf/Ti3SiC2 composites were studied. The main results are as follows: (1) Dense Csf/Ti3SiC2 composites with 2, 5, 10 vol.% Csf were prepared by SPS at 1300 o C. The interfacial reaction layer with duplex structure of TiC inner layer and SiC outer layer was formed between Csf and Ti3SiC2 matrix.
(3) The contents of added Csf and TiC produced by the interfacial reaction between Csf and Ti3SiC2 matrix played a critical part in the mechanical properties of Csf/Ti3SiC2 composites.