SiC composites

SiC composites

Materials Chemistry and Physics 74 (2002) 300–305 Effect of sintering additives on mechanical properties of Cf /SiC composites Xin-Bo He Department o...

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Materials Chemistry and Physics 74 (2002) 300–305

Effect of sintering additives on mechanical properties of Cf /SiC composites Xin-Bo He Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China Received 20 April 2001; received in revised form 28 June 2001; accepted 10 July 2001

Abstract The Cf /SiC composites were prepared by precursor pyrolysis-hot pressing with AlN and B as sintering additives, and the effect of sintering additives on the microstructure and mechanical properties of the composites was investigated. In the composites, a SiC–AlN solid solution was formed, which was considered to be mainly due to the reactions between PCS and AlN during the pyrolysis of PCS instead of the inter-diffusion between original SiC and AlN. Although it was beneficial to the microstructure and mechanical properties of the matrix because of the formation of SiC–AlN solid solution that could strengthen the grain boundaries and suppress the coarsening of SiC grains in the matrix, AlN exhibited a small effect on the fiber/matrix interfacial bonding strength. As a result, AlN had little effect on the density and mechanical properties of the composites. In contrast, the density and mechanical properties of the composites were improved significantly by the introduction of a trace of B, and the density and strength increased with the amount of B. However, the maximum of fracture toughness was obtained in the composite with 0.5 wt.% B, which was mainly due to the desirable fiber/matrix interfacial bonding. In the case of composites with 1 wt.% B, the fracture toughness decreased in spite of the enhancement in the density and strength, which was mainly due to the strong fiber/matrix interfacial bonding. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Cf /SiC composites; Precursor pyrolysis-hot pressing; Sintering additives; Mechanical properties

1. Introduction Silicon carbide (SiC) ceramic has attracted considerable interest for applications in wide temperature range due to its excellent properties at elevated temperatures, such as strength and modulus, creep resistance and chemical stability. However, it is very sensitive to defects, such as pores, cracks and large grains, which reduce their reliability as structural materials and limit its use in many structural applications. Therefore, a number of investigations have studied to the improvement in reliability of this ceramic [1–3]. Over the past many years, various kinds of SiC-based composites, including particulate-, whisker- and fiberreinforced SiC composites, have been extensively investigated, which exhibit significant improvement in mechanical properties in comparison with the monolithic SiC ceramic [4–6]. Much attention has been focused on fiber-reinforced ceramic matrix composites, especially continuous fiber-reinforced ceramic matrix composites, which have been demonstrated to be the most promising approach to enhancing the toughness of ceramics. This is mainly due to the assumption that strong fibers can prevent catastrophic brittle-like behavior in ceramics by providing various energy-dissipation processes during propagation of

cracks, and many reports demonstrated the potential of this approach [7–10]. Composites consisting of carbon fibers embedded in SiC matrix are excellent potential candidates for applications in aerospace industry, which requires structural materials intended for high temperatures and aggressive environments, because of their excellent high-temperature mechanical properties, especially toughness, low density and thermal stability. Many reports have shown that sintering additives have great effect on the microstructure and resultant mechanical properties of fiber-reinforced ceramic matrix composites. In the present paper, Cf /SiC composites with AlN and B as sintering additives were prepared by precursor pyrolysis-hot pressing. The effect of the additives on the microstructure and mechanical properties of the composites was investigated.

2. Experimental procedure 2.1. Fabrication of composites The fiber applied in this study was high-modulus carbon fiber (Model M40JB, Toray Industries Inc., Japan),

0254-0584/02/$ – see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 4 - 0 5 8 4 ( 0 1 ) 0 0 4 8 2 - 5

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Table 1 Characteristics of powders Powder

Mean grain size (␮m)

Purity (%)

Major purities (wt.%)

Supplier

B AlN ␣-SiC

0.3 1.0 0.8

>99.9 99.8 98.6

– O: 0.07; C: 0.02 O: 0.35; free C: 0.77; Al: 0.14; Fe: 0.08

H.C. Starck, Berlin, Germany Institute of Steel and Iron, Beijing, China K.K. Showa Denko, Tokyo, Japan

supplied in the form of a continuous yarn. Its average tensile strength and tensile modulus were 4410 MPa and 377 GPa, respectively. Polycarbosilane (PCS) (National University of Defense Technology, Changsha, China) was selected as the precursor and binder, whose mean molecular weight and melting point were 1417 and 246 ◦ C, respectively. Commercially available ␣-SiC was used as the starting powder. AlN and B were chosen as the sintering additives, and xylene was selected as the dispersant. The characteristics of the powders are summarized in Table 1. The composites were prepared by precursor pyrolysis-hot pressing. First, ␣-SiC was mixed with PCS, AlN and B in xylene. For all mixtures prepared, the PCS content was kept at 40 wt.%. The mixture was milled for 24 h in a jar made of agate using SiC grinding balls, and the resultant mixture was the slurry to be used to impregnate the fibers in preparation of the composites. Then, unidirectional fiber-aligned tapes were obtained by infiltrating the fibers with the slurry aboveprepared and winding slurry-infiltrated strands onto a mandrel. After drying, the prepared tapes were cut, stacked and prepressed into green compacts in a metal die at 180 ◦ C. Finally, the green compacts were hot pressed for 1 h at 1850 ◦ C with a pressure of 25 MPa in a flowing argon atmosphere (1 atm). Seven batches of composites with different systems of sintering additives were prepared, which were designated as M1–M7 and are shown in Table 1. The fiber volume fraction of the composites was estimated to be about 0.5. 2.2. Characterization of composites Bulk densities of the composites were measured by the Archimede’s principle with deionized water as immersion medium. Flexural strength and inter-laminar shear strength were determined using a three-point-bending test on specimens of 3 mm×4 mm×35 mm with a span of 30 and 15 mm, respectively, and the cross-head speed was 0.5 mm/min. For fracture toughness, a single-edge-notched-beam (SENB) test was applied on notched specimens of 2.5 mm × 5 mm × 35 mm in size with a cross-head speed of 0.05 mm/min and a span of 20 mm. The notch, 0.3 mm in width and 2.5 mm in depth, was made in the plane normal to the fiber orientation. Fracture toughness of the composites was obtained from       3 Pmax L a 1/2 KIC = a Y (1) 2 2 b wb where Pmax is the maximum of applied load, b the specimen thickness, w the specimen width, a the notch depth, L the

span of support pins, Y the geometrical factor that depends on the test method and the ratio of notch depth (a) to specimen thickness (b). In this case, Y is given by Y

a 

a   a 2 = 1.96 − 3.07 + 13.66 b b b  a 3  a 4 −23.98 + 25.22 b b

(2)

Characterization of the microstructure was performed by a transmission electron microscope (TEM). The fracture surfaces of the samples were observed by a scanning electron microscopy (SEM). The phases contained in the samples were identified by X-ray difractometry (XRD) with Cu K␣ radiation, and compositional information was obtained by a energy-dispersive X-ray spectroscopy (EDS).

3. Results and discussion After Cutler et al. [11] had found that a solid solution between SiC and AlN, which have the same polymorphs (2H structure) and analogous lattice parameters, can be formed in the temperature range of 1800–2100 ◦ C, many investigations have been conducted on the formation of SiC–AlN solid solution [12,13]. TEM image of the matrix in composite M5, which is representative of all the composites, is given in Fig. 1. EDS analysis reveals that the large grains in Fig. 1 (marked by A) mainly consisted of Si and C with little Al and N, however, in the small grains (marked by B), a lot of Al and N was detected besides Si and C, implying that the large grains should be SiC grains and the small grains should be SiC–AlN solid solution grains. If higher sintering temperatures and prolonged soaking time are employed, the

Fig. 1. TEM image of the matrix in composite M5.

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Table 2 Density and mechanical properties of composites Composite

Sintering additives (wt.%)

Inter-laminar shear strength (MPa)

Flexural strength (MPa)

M1 M2 M3 M4 M5 M6 M7

1AlN 1AlN–0.25B 1AlN–0.5B 1AlN–1B 5AlN–0.5B 5AlN–1B 10AlN

18.6 30.4 34.9 45.8 36.3 43.1 17.1

315.1 418.5 492.3 569.7 526.6 578.0 302.7

amount of Al and N in the large grains will be increased. Because SiC and AlN are highly covalently bonded compounds, the inter-diffusion rate between SiC and AlN, which is the predominant formation mechanism of SiC–AlN solid solution, is very slow in conventional processing conditions [14]. So, it is reasonable that a uniform solid solution from original SiC and AlN was not obtained in the present processing conditions, and higher temperatures and longer time were required to achieve a uniform solid solution. EDS results also showed that no AlN grains were found in the composites, which was consistent with XRD results of the composites. Since the pyrolysis of PCS at high temperatures contains the formation of amorphous SiC, nucleation of SiC and notable coarsening of SiC crystallite [15], it can be proposed that the SiC–AlN solid solution in the composites was formed mainly through the reactions between PCS and AlN grains during the pyrolysis of PCS, instead of the inter-diffusion between original SiC and AlN. The density and mechanical properties of the composites are summarized in Table 2. It can be seen that the composites with 1 wt.% AlN (M1) and 10 wt.% AlN (M7) showed rather low densities less than 78% of theoretical density and poor mechanical properties, indicating that AlN had a small effect on the density and mechanical properties of the composites. From the low inter-laminar shear strength, it can be concluded that, although it could enhance the density and mechanical properties of the matrix by strengthening the

± ± ± ± ± ± ±

33.8 27.4 45.7 38.5 31.3 40.1 24.7

Fracture toughness (MPa m1/2 ) 9.4 13.7 16.8 14.1 17.2 14.6 10.2

± ± ± ± ± ± ±

2.7 1.7 3.0 2.3 2.1 3.4 2.8

Relative density (%) 77.9 89.6 92.8 95.6 92.8 96.4 75.5

grain boundaries and suppress the coarsening of SiC grains [4], SiC–AlN solid solution had a small effect on the interfacial bonding strength between the fibers and matrix that plays a key role in determining the mechanical properties of fiber-reinforced ceramic composites [16,17]. As a result, AlN showed a small effect on the density and mechanical properties of the composites. On the contrary, B showed a great effect on the density and mechanical properties of the composites. It was observed that the density and mechanical properties of the composites were improved significantly by the introduction of a trace amount of B, and the density and strength increased with the amount of B. The composites with 1 wt.% B (M4 and M6) were well densified with high densities over 95% of theoretical density and high flexural strength over 569 MPa. It is apparent that the great difference in the mechanical properties of the composites was caused mainly by the difference in the inter-laminar shear strength, which was an indicator of the fiber/matrix interfacial bonding strength. Therefore, it was believed that B exhibited a great effect on the density and mechanical properties of the composites primarily by affecting the interfacial bonding strength between the fibers and matrix. Fig. 2a depicts the microstructure of composite M5. As shown, the composite is well densified, suggesting that B can drastically promote the densification of the composites. This is similar to the previous work [18], which has reported that B can promote the densification of SiC ceramic

Fig. 2. TEM images of composite M5.

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by improving the grain-boundary diffusivity as a result of the segregation of B to the grain boundaries. Fig. 2b is a TEM image of the fiber/matrix interface in Fig. 2a. It was observed that there existed a layer of interphase (∼70 nm in thickness) between the fibers and matrix, whose characteristics were responsible for the interfacial bonding strength. EDS results showed that the interphase was mainly composed of C, which should be primarily attributed to the preferential deposition of PCS-derived C on the surfaces of the fibers. Due to the lack of B, the composites with l wt.% AlN (M1) and 1 wt.% AlN–0.25 wt.% B (M2), especially the composite M1, could not be well densified, leading to very low densities and very weak fiber/matrix interfacial bonding. So, the composites exhibited lower mechanical properties. In

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contrast, better mechanical properties were achieved in the composites with 0.5 wt.% B (M3 and M5) and 1 wt.% B (M4 and M6), which was believed to be due to the high densities resulting from the addition of more B. However, the maximum of fracture toughness was obtained in the composite with 0.5 wt.% B (M5), suggesting that the fiber/matrix interfacial bonding was the most desirable for fracture toughness. As compared with composites with 0.5 wt.% B (M3 and M5), the composites with 1 wt.% B (M4 and M6) exhibited decreased fracture toughness in spite of the enhancement in the density and strength (including inter-laminar shear strength), indicating a relative strong fiber/matrix interfacial bonding. Based on the facts described above, it can be proposed that B strengthened the fiber/matrix interfacial bonding by strengthening the interphase and fiber/interphase

Fig. 3. Fracture surfaces of composite: M5 (a); M6 (b); and M7 (c).

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Fig. 4. Surface morphologies of pullout fibers in composite: M6 (a) and M7 (b).

interfacial bonding, which may be caused by the reactions of B with the fibers (C) and interphase (C) through 4B + C → B4 C

(3)

The fracture surfaces of the composites are given in Fig. 3. Fig. 3a illustrated an extensive fiber debonding and pullout, whereas in Fig. 3b, a typical brittle fracture behavior was observed. It is well known that interfacial bonding strength between fibers and matrix can be evaluated by the morphology of fracture surfaces. An extensive fiber pullout indicates a relatively weak fiber/matrix interfacial bonding, while a flat fracture surface indicates a strong fiber/matrix interfacial bonding. Therefore, the extensive debonding and pullout of fibers showed that the fiber/matrix interfacial bonding in the composites with 0.5 wt.% B (M3 and M5) were more desirable than those in the other composites, which was consistent with the above proposal. As a result, the composites achieved the highest values in fracture toughness. In contrast, the fibers in the composites with 1 wt.% B (M4 and M6) were unfavorably strongly bonded to the matrix, which invalidated the predominant energy-consuming mechanisms of fiber debonding and pullout. So, the composites displayed decreased fracture toughness and increased strength. It is evident that the difference in mechanical properties between the composite with 0.5 and 1 wt.% B was primarily attributed to the discrepancy in the interfacial bonding strength, which was associated with the interfacial reactions of B with the fibers and interphase. In composites with 0.5 wt.% B (M3 and M5), moderate interfacial reactions occurred, leading to a desirable fiber/matrix interfacial bonding, which was beneficial to fracture toughness, whereas severe interfacial reactions took place in the composites with 1 wt.% B (M4 and M6), resulting in a strong fiber/matrix interfacial bonding, which was beneficial to strength. In the case of composite with 1.0 wt.% AlN (M7), a step-like fracture surface was observed (Fig. 3c). This, in combination with the smooth surfaces of the pullout fibers, indicated a very weak fiber/matrix interfacial bonding, which should be attributed to the lack of B. Consequently, the applied load could not be effectively

transferred from the matrix to the fibers, resulting in poor mechanical properties. Fig. 4 presents the representative pullout fibers in the composite with 5 wt.% AlN–l wt.% B (M6) and 1.0 wt.% AlN (M7). As shown, some regions with residue from the matrix were observed on the surface of the fiber in Fig. 4b, implying the presence of fiber/matrix interphase in the composites. Moreover, it was noted that the surface of the pullout fiber in Fig. 4b was mostly clean and smooth except for some interphase, indicating that AlN exhibited little effect on the fiber/matrix interfacial bonding, whereas the surface of the pullout fiber in Fig. 4a was rather rough, which was believed to be mainly due to the severe interfacial reactions. These results also confirmed that B exhibited a great effect on the density and mechanical properties of the composites primarily by affecting the interfacial bonding strength between the fibers and matrix.

4. Conclusions In this paper, Cf /SiC composites were prepared by precursor pyrolysis-hot pressing with AlN and B as sintering additives, and their microstructure and mechanical properties were evaluated. The following conclusions can be derived from the present studies. 1. SiC–AlN solid solution between original SiC and AlN were not formed because of the very low inter-diffusion rate between original SiC and AlN. However, SiC–AlN solid solution between PCS-derived SiC and AlN were formed, which was considered to be mainly due to the reactions between AlN and PCS during the pyrolysis of PCS. 2. Because of the formation of SiC–AlN solid solution that could strengthen the grain boundaries and suppress the coarsening of SiC grains in the matrix, AlN was beneficial to the microstructure and mechanical properties of the matrix. However, AlN exhibited a small effect on the fiber/matrix interfacial bonding strength. As a result, AlN

X.-B. He / Materials Chemistry and Physics 74 (2002) 300–305

had little effect on the density and mechanical properties of the composites. 3. The density and mechanical properties of the composites were improved significantly by the introduction of a trace of B, and the density and strength increased with the amount of B. However, the maximum of fracture toughness was obtained in the composites with 0.5 wt.% B, which was mainly due to the desirable fiber/matrix interfacial bonding. In the case of composites with 1 wt.% B, the fracture toughness decreased in spite of the enhancement in the density and strength, which was mainly due to the strong fiber/matrix interfacial bonding. 4. The composite with 5 wt.% AlN–0.5 wt.% B, sintered at 1850 ◦ C with a pressure of 25 MPa, achieved better mechanical properties of 526.6 MPa in flexural strength and 17.2 MPa m1/2 in fracture toughness.

Acknowledgements The author is grateful to Dr. X.M. Zhang for his assistance in this investigation.

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