Thursday, February 13, 2020

Fabrication of Carbon Fiber Reinforced SiC Composites by Combining 3D Printing and Liquid Silicon Infiltration Process


A novel method has been developed to fabricate carbon fiber reinforced SiC (Cf/SiC) composites by combining 3D printing and liquid silicon infiltration process. Green parts are firstly fabricated through 3D printing from a starting phenolic resin coated carbon fiber composite powder; then the green parts are subjected to vacuum resin infiltration and pyrolysis successively to generate carbon fiber/carbon (Cf/C) preforms; finally, the Cf/C preforms are infiltrated with liquid silicon to obtain Cf/SiC composites. The 3D printing processing parameters show significant effects on the physical properties of the green parts and also the resultant Cf/C preforms, consequently greatly affecting the microstructures and mechanical performances of the final Cf/SiC composites. The overall linear shrinkage of the Cf/SiC composites is less than 3%, and the maximum density, flexural strength and fracture toughness are 2.83 ± 0.03 g/cm3, 249 ± 17.0 MPa and 3.48 ± 0.24 MPa m1/2, respectively. It demonstrates the capability of making near net-shape Cf/SiC composite parts with complex structures.

1. Introduction

Silicon carbide (SiC) ceramic is the prime candidate for high temperature structural components in aerospace, nuclear and transportation areas due to its refractory nature, superior specific strength, high thermal conductivity and excellent tribology performance at elevated temperature [1]. However, low strain tolerance and fractural resistance limit its applications. Carbon fiber reinforced SiC (Cf/SiC) ceramic matrix composites have attracted great attentions for the improved strength and fracture resistance with the addition of high strength and modulus fibers [2,3].

Liquid silicon infiltration (LSI) technique is one of the major manufacturing processes to fabricate Cf/SiC composites [4,5]. This process basically consists of three steps: i) producing a carbon fiber reinforced plastic (CFRP) green part, ii) pyrolyzing the green part into a porous carbon fiber/carbon (Cf/C) preform, and iii) infiltrating the Cf/C preform with liquid silicon that reacts with carbon to form SiC. LSI process is a near net-shape forming process, owning the advantages of low sintering temperature, cost effective and capability of fabricating large- scale parts [6]. As mentioned above, the desired shape of a Cf/SiC composite part is derived from the geometry of the CFRP green part, which is  conventionally  shaped  by  filament  winding,  autoclave molding, warm pressing and resin transfer molding (RTM), etc [7]. Molds and tools are usually needed for these common processing techniques, which intrinsically limit the ability to produce parts with extremely high geometric complexity, such as helical conformal channels and lattice structures.

3D printing, namely additive manufacturing, is a collection of technologies capable of building complex three-dimensional objects directly from CAD models in an additive mode [8].  In  recent  years, some investigations have been conducted by using 3D printing technology to generate green parts, and then combining with LSI process to fabricate SiC based composite parts. Stierlen et al. [9] first employed selective laser sintering (SLS), a powder bed fusion 3D printing process, in order to fabricate porous preforms using the mixture of SiC powder and reactive polymer binder, and subsequently infiltrated the preforms with liquid silicon to form reaction bonded  SiC  composites  [10–12]. Tian et al. [13,14] used high carbon yield and photo-curable resin for stereolithography (SLA) 3D printing process to fabricate resin precursors, which are then pyrolyzed and converted to porous carbon preforms for LSI process. It has been a challenge for 3D printing to fabricate CFRP composites, especially with the polymeric matrices of high carbon yield, such as phenolic resin or other aromatic polymers [15]. Therefore, limited researches about fabricating Cf/SiC composite through 3D printing have been reported. Lu et al. [16,17] proposed an alternative  route by integrating SLA 3D printing with slip casting pro-cess to prepare  Cf/C preforms. In this method, a negative resin mold   was first built by SLA process, and then the slurry of carbon fibers, SiC particles and phenolic resin solution was casted into the as-fabricated resin mold, after curing of the mixture the resin mold was removed by chemical corrosion. Through this method, they succeed in fabricating hollow turbine blades by using of Cf/SiC composites.

In this work, a novel method is proposed to manufacture Cf/SiC composites based on 3D printing technology. The processing scheme includes the following steps: firstly, preparing the carbon fiber powder encapsulated with char-yielding polymer; secondly, producing the carbon fiber green part through SLS process, which serves as the reinforcing framework; then infiltrating the green part with phenolic resin solution and pyrolyzing it to form the Cf/C preform; finally, converting the Cf/C preform to a Cf/SiC composite by LSI process. The purpose of this study is to demonstrate the feasibility of this process to manufacture near net-shape Cf/SiC composite parts with complex geometries. More specifically, the study is focused on establishing correlations between the processing  parameters, resultant microstructures of the Cf/C preform and the performance of the final Cf/SiC composite.

2. Experimental and characterization

2.1. Preparation of samples

The preparation process is schematically shown in Fig. 1. The phenolic resin coated carbon fiber (PF/CF) composite powder was first prepared according to the following procedures: (1) firstly, the phenolic resin powder was completely dissolved in acetone solution with a mass ration of 1:1; (2) then, the carbon fiber powder was added into the solution and a further ball milling at room temperature was performed to obtain the homogenous suspension; (3) after ball-milling process, the acetone was distilled out and the phenolic resin began to crystallized preferentially taking the carbon fibers as nuclei; (4) finally, the obtained precipitate was dried in a vacuum oven at 60 °C and then crushed by universal grinder. In this work, carbon fibers (1.76 g/cm3, T300) with an average fiber diameter of 7 μm, and thermoplastic phenolic resin (1.22 g/cm3, PF2123™) mixed with hardener methenamine (10 wt.%) were used as raw materials. The content of phenolic resin in the starting composite powder was 30 vol. %.

The carbon fiber green parts were manufactured by SLS process. The sintering experiments were conducted on the HK P320™ SLS machine (Wuhan Huake 3D Technology Co.Ltd., China). The SLS system was equipped with a power continuously adjustable CO2 laser with a wavelength and beam diameter of 10.6 μm and 200 μm, respectively. A counter-rotating roller was used to spread the powder across the building area, the moving speed was 200 mm/s. The processing parameters were set as follows: the laser power was 4–10 W, with an interval of 2 W; the scanning speed was 2000 mm/s; the powder layer thickness was 0.1 mm; the scan spacing was 0.15 mm and the part bed temperature was set as 60 °C.

The green parts were firstly carbonized to increase their strengths and meanwhile reduce close porosities. Then the carbonize parts were vacuum infiltrated (∼102 Pa) with low viscosity thermosetting phenolic resin, followed by curing and secondary carbonization, in order to increase the carbon density and form the Cf/C preform. The alcohol soluble B-stage phenolic resin (THC-400™) with a carbon yield over  60 wt.% was supplied by Shaanxi Taihang Impedefire Polymer Co.Ltd., China. The resin infiltrated parts were cured under a multi-step curing schedule suggested by the manufacturer, 120 °C for 1 h then 150 °C for 2 h and finally 180 °C for 3 h, to minimize or avoid internal stress and shrinkage distortion. The carbonization procedures were all conducetd at 900 °C for 1 h with a heating rate of 2 K/min under the protection of flowing argon atmosphere. The resulting porous Cf/C preforms were designated as Cf/C-4 W, Cf/C-6 W, Cf/C-8 W and Cf/C-10 W, according to the laser power used for making carbon fiber green parts in SLS process.

The porous Cf/C preforms were infiltrated with liquid silicon to form the Cf/SiC composites. The Cf/C preforms were placed in a BN- coated graphic crucible, and the pure silicon granules (industrial grade, average particle size of 1∼3 mm) were put on the top and bottom of the Cf/C preforms. The infiltration process was conducted at 1550 °C for 1 h under a pressure of about 102 Pa. The corresponding Cf/SiC composites were labelled as Cf/SiC-4 W, Cf/SiC-6 W, Cf/SiC-8 W and Cf/SiC-10 W, respectively.

Fig. 1. The schematic diagram of the procedure for fabricating Cf/SiC composite parts.

Fig. 2. The morphologies and EDS patterns of (a), (d) raw carbon fibers and (b), (e) PF/CF composite powder; (c) the length distribution of carbon fibers.

2.2. Characterization

The morphologies of the powders and green parts, microstructures of the polished and fracture surfaces of the Cf/SiC composites were observed by field scanning electron microscope (JSM-7600 F, JEOL, Japan; Quanta 650, FEI, USA) and optical microscope (Axio Lab.A1 MAT, Zeiss, Germany). The bulk densities (ρbulk ) and open porosities (Popen) of the Cf/C preforms and Cf/SiC composites were measured by the Archimedes method. The skeleton densities (ρskeleton) of the Cf/C preforms were measured by means of a gas pycnometer (AccuPyc 1330, Micromeritics Instrument Corp., USA). All measurement of the Cf/C preforms were performed on the milled powders, at least ten data were recorded for each sample. The closed porosities (Pclosed) inside the Cf/C preforms were measured by means of a gas pycnometer (AccuPyc 1330, Micromeritics Instrument Corp., USA). All measurement of the Cf/C preforms were performed on the milled powders, at least ten data were recorded for each sample. The closed porosities (Pclosed) inside the Cf/C preforms were calculated by Eq. (1).
The pore size distributions of the Cf/C preforms were measured with a mercury porosimeter (Autopore IV 9500, Micromeritics Instrument Corp., USA), three samples for each Cf/C preform were tested. The dimension deviation ratio (DDR) was used to quantify the dimensional accuracy of the Cf/C preform using the following equation:
where A is the measured dimension using a digit caliper with an accuracy of 0.01 mm, A0 is the norminal dimension of the designed CAD model. In this study, rectangle bars with dimension of 40 mm (L) ×  10 mm (W) ×4 mm (H) were designed for the dimensional accuracy and the subsequent mechanical tests of the Cf/C preforms. The DDRs in length (L) and width (W) directions were measured, four samples with total twelve replicates were performed for each direction. The compo- sitions of the specimens were analyzed by X-ray diffractometer (XRD-7000, SHIMADZU Corporation, Japan), using Cu Kα radiation in the 2θ range between 10° and 80° using a scan speed of 10°/min. Three point bending tests were performed to determine the flexural strength of the Cf/C preform and the Cf/SiC product using a universal testing machine (Zwick/Roell  Z010,  Ulm,  Germany).  The  Cf/SiC  specimens  with  dimension of 40 mm × 4 mm × 3 mm were prepared by cutting and grinding in accordance with ISO 14704-2008. The crosshead speed was 0.5 mm/min  and  the  supporting  span  was  32 mm.  Four  specimens  for each  measurement  were  tested.  The  fracture  toughness  was  evaluated using  a  single  edge  notched  beam  (SENB)  method  according  to  ASTM C1421-16   standard.   Four   notched   samples   with   a   dimension   of 20 mm × 3 mm × 4 mm were prepared. The edge notch was introduced by a diamond saw with width of 0.2 mm, and ratio of the notch depth to specimen height (a/H) was in the range of 0.40-0.55. The notch depth was  measured  using  a  caliper.  Three-point  bending  tests  were  carried out at a loading speed was 0.05 mm/min and a support span of 16 mm. The  fracture  toughness,  KIC  was  calculated  using  the  following  equations [18].
where P is the maximum load during three-point bending test, L is the support  span,  W  is  the  specimen  width,  H  is  the  specimen  height,  a  is the notch depth, α is the ratio of a and H, and Y is the calibration factor.

3. Result and discussion

3.1. Powder characteristics

Fig. 2 shows the SEM micrographs of the raw carbon fibers and the PF/CF composite powder. As shown in Fig. 2a, the raw carbon fibers have relatively smooth surfaces with small grooves, which can enhance the surface adhesion with the phenolic resin through mechanical interaction. The fiber length and distribution were analyzed using the digital image processing software Image J. More  than 300 fibers were measured, and the result is shown in Fig. 2c. It shows that the length distribution  of  the  carbon  fiber  ranges  from  20∼220 μm,  and the average fiber length is 59.03 μm. The PF/CF composite powder shows  an inherited morphology from the original carbon fiber (Fig. 2b), and a thin layer of phenolic resin is uniformly coated on their surfaces, which serves as the binder connecting the adjacent carbon fibers to a free- standing compact. The energy dispersive x-ray spectrometry (EDS) patterns in Fig. 2d and e indicate that the content of oxygen on the carbon fiber surface is increased after coating, which comes from the phenol hydroxyl in the phenolic resin molecular chain.

Fig. 3. Top-surface morphologies of (a) the SLS green part; cross-section morphologies of (b) the carbonized part and (c) the Cf/C preform; (d) picture of complex green parts with lattice-truss structures. The insets shows zoom-in SEM images of each sample.

3.2. The microstructural evolution of the Cf/C preforms

Fig. 3 shows the representative microstructures of the SLS green parts, carbonized parts and Cf/C preforms. The green parts show a  porous but self-standing feature, and the phenolic resin is homo- genously coated on the surfaces of the carbon fibers to bond them together, as shown in Fig. 3a. The principle  bonding  mechanism  is  that the linear thermoplastic phenolic resin which is melted by the  laser, flows to wet adjacent carbon fibers and then is cured by hardener to form cross-linked thermosetting products [19,20]. The laser power has influence on the melt flow of the resin, thus changes the microstructure of Cf/C preform, which will be discussed in the following section. After carbonization, the coating of phenolic resin is converted to a compact pyrolytic carbon layer, and the surfaces of the carbon fibers become smoother due to the chemical bond reorganization, see Fig. 3b. The inter-fiber bondings are achieved by these residual carbon ligaments, which provides strength for the carbonized parts. Fig. 3c shows the fracture morphology of the Cf/C preform. Compared to the carbonized parts, the pores between the carbon fibers are first filled with liquid resin, which is then dissociated into pyrolysis carbon after curing and  carbonization, thus decreasing the porosity and meanwhile improving the carbon density in the Cf/C preform. Indistinct layer structures inherited from SLS process are recognized on the cross section of Cf/C preform, which will also influence the microsturcture of the Cf/SiC composites. The inset SEM image also shows some gaps between carbon fibers and pyrolysis carbon, which is caused by the shrinkage of the matrix during curing and pyrolysis procedure. Fig. 3d shows the 3D- printed green parts with lattice-truss structures. Such a small complex structure can be easily fabricated with no distortion. Furthermore, pores with minimum diameters of 1.5 mm can be distinguished.

3.3. The properties of the Cf/C preform

The energy input, especially the laser power applied in the SLS process is critical to the physical properties, including dimensional accuracy, microstructure and mechanical property, of the green parts  and also the resultant Cf/C preforms. Fig. 4 shows the dimensional deviation ratio results of the Cf/C preforms in length and width direction. The results show negative deviations in both length and width direction for the green parts built at the laser power of 4 W. This negative deviation is associated with the powder  densification  effect  in the SLS process. However, when the laser power is set above 6 W, the dimension deviations of the green parts gradually convert into positive deviations and continues to increase with the laser power. The reason  for the increase in dimension is because the growth of part, which occurs when the edges of the parts become less distinct as the neighboring powder partially melts on the parts [21], overweighs the effect of powder densification. Great negative dilatation of the samples occurs after the subsequent carbonization process. Shrinkage during decomposition offsets the parts expansion from SLS process. During the infiltration and pyrolysis process, the dimensions of the Cf/C preforms shows little variations, which may contributed to the  supporting  effect of the carbon fiber framework [22]. Because of the smaller nominal dimension, the DDRs in width direction appear higher than those in length direction. During the successive silicon infiltration process,  the Cf/C preforms show an average dimension decrease of 0.52%. Therefore, the overall linear shrinkage of the Cf/SiC composites during the whole process is less than 3% compared to the CAD model, which is better than the conventional processing methods [23], and can be considered as a near-net-shape forming process [2].

Fig. 4. Dimensional accuracy of the Cf/C preforms in different processing stages.

The microstructures of the Cf/C preforms, including density and porosity, determine the success of formation of silicon carbide by LSI process. The laser power also largely influences the microstructure of  the Cf/C preforms. Singh et al. [24] and Margiotta [25] have reported that in order to fabricate a dense SiC ceramic using LSI method, the    ideal open porosity of the carbon preform is ranging from 31.1% to 57.3%. As shown in Fig. 5a and b, the bulk density of the green  parts first increases with the increase of laser power and then decreases when the laser power reach 10 W, while the open porosity shows the opposite trend. During the SLS process, melting and curing of the phenolic resin binder generates densification and shrinkage in the green parts. Generally, the degree of densification increases with increasing laser power, thus improving the bulk densities of the green parts [26]. However,  when the laser power is excessively high, the localized temperature will dramatically increase above the decomposition temperature of the phenolic resin. As indicated in our previous study [20], the weight loss    of the phenolic resin will significantly increase when the temperature is above 350 °C. For this reason, smoke was observed during the SLS process when a laser power of 10 W was used, which is considered to be a sign of degradation of the polymer material [21]. The carbonized parts show a slightly increase of bulk density and decrease of open porosity in contrast to the green parts. After the resin infiltration and carbonization process, the bulk densities of the Cf/C preforms are distinctly increased to 0.624–0.702 g/cm3, whereas the open porosities are reduced to 64.64%–56.77%. The physical properties of the Cf/C preforms are summarized in Table 1. The skeleton density of the Cf/C preforms varies within a narrow range. The calculated results indicate the existence of closed pores in the Cf/C preforms. The Cf/C-4 W preform shows the lowest closed porosity of 0.37%, while the other preforms have higher closed porosities of about 4%. Therefore, the Cf/C preforms are considered to have a highly connected network of open pore with relatively low closed pore volume. Fig. 5c shows the pore size distributions of the Cf/C preforms varying with the laser power. It is obviously that they all have similar pore size distributions. Pore size mainly distributes in the range from 15 to 30 μm, and there are also a small amount of pores distributing below 10 μm. The median pore diameter varies from 27.1 to 17.85 μm, and the average pore diameter decreases from 13.24 to 6.42 μm with the increase of laser power.

Fig. 6 shows the flexural strength changes of the Cf/C preforms over the whole process. The mechanical property of the green parts is crucial to maintain details of the parts during post-processing. In practice, as the flexural strength of 1.37 ± 0.1 MPa is attained  at laser power of  4 W, cautions should be taken when handle it. With the increase of laser power, the flexural strengths of the green parts increased. A flexural strength of 2.24 ± 0.02 MPa at 6 W has been sufficient to deal with it safely. The strengths of the Cf/C preforms are gradually raised after the carbonization, infiltration and secondary pyrolysis process. A maximum flexural strength of 11.87 ± 1.01 MPa is finally achieved for the Cf/C- 10 W preform.

Fig. 5. The influence of laser power on the (a) bulk density, (b) open porosity and (c) typical pore size distribution of the Cf/C preforms in different processing stages.

Table 1 The physical properties of the Cf/C preforms.

Fig. 6. The flexural strength of the Cf/C preforms in different processing stages.

3.4.  Microstructure  and  phase  composition  of  the  Cf/SiC  composite  parts

Microstructures of the Cf/SiC composites are shown in Fig. 7. The  dark gray, light gray and black regions in the back scattered electron SEM images are identified as silicon carbide, silicon and carbon, respectively. It obviously shows that the microstructures of Cf/SiC composites correlates to the bulk density or porosity in the Cf/C preforms. For Cf/SiC-4 W (Fig. 7a), it shows a high content of residual Si and little unconverted carbon. Generally, sufficient porosity and accessible  porous carbon preform favors the complete reaction of liquid  Si  with  the carbon matrix [27]. The low bulk density, high open porosity and  also large pore diameters of the Cf/C-4 W preform are the reasons for a resulting Cf/SiC material with  high  residual Si  and  low  SiC  content.  Fig. 7b, c show the microstructures of Cf/SiC-6 W and Cf/SiC-8 W prepared from Cf/C-6 W and Cf/C-8 W, respectively. The contents of residual Si are reduced, and meanwhile the contents of SiC are increased, which results from the lower bulk density, higher open porosity and smaller pore diameters compared to Cf/SiC-4 W. Micrograph of Cf/SiC- 10 W (Fig. 7d) shows a relatively higher volume fraction of SiC and less and finer residual Si phase. Besides, an amount of carbon residues consisting of unconverted carbon fibers and pyrolysis carbon was distinctly observed. The preform Cf/C-10 W has slightly increased bulk density, decreased open porosity as well as the smallest average pore size, which can induce a certain extent of blockage of the channels  during Si infiltration. As a result, some carbon fibers were remained in the Cf/SiC-10 W parts. Fig. 7e shows the high magnification  of Cf/SiC-   10 W, most of the carbon fibers appear serrated edges, suggesting that these carbon fibers are partially siliconized during the reaction. The representative cross section of the Cf/SiC composite is also evaluated.  As shown in Fig. 7f, the layer structure inherited from  the Cf/C  preform is still recognized, the residual Si distributes mainly in the interlayer between two adjacent building layers.

The XRD patterns of the Cf/SiC composite parts fabricated under different conditions are shown in Fig. 8. It confirms the presence of residual silicon and silicon carbide, and also reveals the disappearance  of the unreacted carbon. Their corresponding diffraction peaks are labelled. The XRD patterns suggests that the face centered cubic β-SiC is the only crystalline phase of silicon carbide. The lattice parameter of β- SiC cubic cell calculated from the XRD data is 4.36 Å, which is in good agreement with the known value (JCPDS Card No. 65-0360).

3.5.  Phase  formation  and  reaction  mechanism

The polished surfaces were treated with the acid mixture (1 HF: 1 HNO3) for 1 h at room temperature to remove the residual silicon, so as to elucidate the mechanism governing the siliconization of the Cf/C preform. The microstructure formation is strongly affected by the open porosity as well as the different carbon sources in the Cf/C preform.     Fig. 9a shows the morphology of the silicon carbide in the silicon-rich area, which consists of two distinctly different β-SiC grains, with grain sizes of about 0.5 μm and 20 μm. Obviously, the fine grained β-SiC crystals with grain sizes of 0.5–2 μm are generated from the siliconization of carbon fibers, which has been totally transformed into SiC. The larger faceted β-SiC (>20 μm) crystals are formed due to the sili- conization of the polymer derived glassy carbon. Carbon fibers which have experienced high temperature carbonization (∼1500 °C) are relatively unreactive compared with the glassy carbon [28]. Therefore, more nucleation sites exist on the carbon fibers, thus leading to more independently nucleated fine β-SiC cryatals. On the contrary, the glassy carbon with higher reactivity are more readily dissolved  by  liquid  Si  [27], and the local temperature rising due to the violent exothermic reaction causes the growth of coarse SiC grains into the Si melt [29,30]. Fig. 9b shows the groove formed by the etching of siliconized fiber, around which are the fine grained β-SiC particles. In the area with moderate silicon, the carbon fiber first partially dissipates in liquid silicon to produce a verge of fine crystallized SiC grain; and then the remaining carbon continues  to  react with the silicon, diffusing across the formed SiC layer, to form a non-dense granular structure [31]. As indicated by the previous study, the dissolved carbon is likely existed in the form of C, C-Si, C-Si4 pairs or even  SiC4  units  [30,32],  which  is  highly unstable to acid etching. In the carbon-rich region, as shown in Fig. 9c, the carbon fibers are partially siliconized,  only  after  the  polymer derived carbon coating has been totally transformed. The EDS result in Fig. 9d indicates that  the center of the fiber mainly consists of    C atoms.  It confirms that  there exist  carbon  fibers that  have  not  been converted into SiC, which are beneficial to improve the fracture toughness of the composites.

Fig. 7. Back scattered electron SEM images of the top surface of Cf/SiC composite parts. (a) Cf/SiC-4 W, (b) Cf/SiC-6 W, (c) Cf/SiC-8 W, (d) Cf/SiC-10 W, (e) high magnification optical micrograph and (f) representative cross section of Cf/SiC-10 W.

Fig. 8. X-ray diffractogram of the Cf/SiC composite parts.

3.6.  Mechanical  properties  of  the  Cf/SiC  composite  parts

During the infiltrating process,  the molten Si penetrates into the Cf/ C preforms and reacts with the polymer derived carbon and carbon  fibers to form SiC matrix. There is an approximately 58% volume increase when one mole of amorphous carbon reacts to form one mole of SiC [33]. Therefore, the formed SiC and the unreacted molten Si fill the pores in the preforms, leading to the increase of density. From the result showing in Table 2, bulk  density of  2.80–2.83 g/cm3, flexural  strength of 179–249 M Pa and fracture toughness of 3.02–3.48 MPa m1/2 are obtained. The density, flexural strength and fracture toughness of the Cf/SiC composite parts are relevant to their microstructures, which are strongly affected by the properties of the Cf/C preforms, such as bulk density, open porosity and pore size. The maximum density, flexural strength and fracture toughness are simultaneously attained for the Cf/ SiC-10 W parts. The improved mechanical properties should be attributed to the microstructure that contains the most quantity of SiC phase and the increased amount of unreacted carbon fibers. As compared to a commercialized short carbon fiber reinforced SiC composite—Cesic® (type MF), which is fabricated by conventional powder mixing, compression molding, and infiltration techniques, the Cf/SiC composite fabricated by this method exhibits higher strength but lower fracture toughness.

Fig. 10 shows the representative stress-strain curves of the Cf/SiC composites recorded during the SENB tests. It can be seen that the curves for Cf/SiC-4 W, Cf/SiC-6 W and Cf/SiC-8 W samples are almost linear up to the maximum load and then followed by sharp load de- creases. In the case of Cf/SiC-10 W, however, a gradual load decrease takes place, indicating a sequential failure process, which can be attributed to the crack deflection and other toughening mechanisms caused by the increased amount of unreacted carbon fibers. Moreover, the strains to failure of Cf/SiC-4 W, Cf/SiC-6 W and Cf/SiC-8 W samples are less than 0.08%, while Cf/SiC-10 W sample shows  an  increased strain to failure up to 0.16%. Both observations indicate that with the increaseing amount of unreacted carbon fibers, the fracture behavior of the Cf/SiC composite changes from a brittle to a more damage-tolerant failure mode [35,36].

In order to investigate the toughening mechanism of the short  carbon fibers in the Cf/SiC composite, the fracture surface was characterized by SEM. Fig. 11 shows the typical fractured surfaces of the Cf/ SiC-10 W composite after the bending tests. Fig. 11a displays the fracture surface of the silicon-rich region. Obvious crack deflections can be observed due to the cracks pass through the carbon fiber domain. Some fiber breakages and fiber debondings are also detected, which indicate the interfacial shear resistance is caused by the fiberdebondings [36]. The broken fibers, along with the SiC matrix (indicated in the upper left corner), show typical transgranular fractures. A close view of the fiber fracture surface shows lots of  fine SiC grains (Fig. 11b),  implying that the carbon fibers in this area have been siliconized. It is also seen that the siliconized carbon fiber are strongly bonded to the matrix, which should be avoid for the purpose of improving toughness [37]. Fig. 11c shows the fracture surface of the carbon-rich region, which appears a rough fracture surface and ragged cracks propagation path in this area. Two different kinds of carbon fiber are observed. One shows the same smooth surface as the received fiber, the other one shows some pits and grooves on the surface, indicating the occurrence of partially siliconization. Fiber debondings are considered to be the main toughening mechanism in this area, which evidently consumes the propagating energy of the cracks and leads to the improved toughness.

Fig. 9. SEM micrographs of etched Cf/SiC specimens revealing the siliconization mechanism of the carbon fibers.

Table 2 Mechanical properties of the Cf/SiC composite parts.

Fig. 10. Stress-strain diagrams in the SENB tests for the Cf/SiC composites.

4. Conclusions

It is demonstrated that the Cf/SiC composite can be fabricated by combining 3D printing and liquid silicon infiltration process. The main challenge of this novel method is the fabrication of Cf/C preform with desirable microstructure, mechanical property and dimensional accuracy. The laser power shows critical impacts on the microstructure formation, mechanical strength and dimension stability of the green  part, and also the resultant Cf/C preform. The derived Cf/C preforms are characterized by bulk densities of 0.624–0.702 g/cm3, open porosities of 56.77–64.64% and average pore diameters of 6.42–13.24 μm. The green parts possess sufficient strength for handling during the post- processing. The linear shrinkage of the Cf/C preforms compared to the CAD model is less than 2%, and the overall linear shrinkage of the resulting Cf/SiC composite is less than 3%. The microstructure of the Cf/ SiC composite mainly consists of β-SiC, residual silicon, and a small fraction of unreacted carbon fibers. The microstructure formation of the Cf/SiC composite is strongly affected by the properties, such as bulk density, open porosity as well as pore size of the Cf/C preform. The maximum density, flexural strength and fracture toughness of the Cf/ SiC composite are 2.83  ± 0.03 g/cm3, 249  ± 17.0 MPa and 3.48 ± 0.24 MPa m1/2, respectively. It is shown that the residual carbon fibers can improve the fracture toughness and also change the fracture behavior of the Cf/SiC composite from a brittle to a non-catastrophic mode. The predominant toughening mechanisms deduced from the fractographic analysis are concluded to be crack deflection, fiber breakage and fiber debonding.

The major advantages of this method over other ceramic processing techniques are the enhanced capability for making near net-shape  Cf/ SiC composite parts with complex structures without the need  of  tooling. However, there still exist challenges to fabricate large-scale parts. Besides, there are also spaces for the further improvement of mechanical performance.

Fig. 11. Toughening mechanisms of short carbon fibers in the Cf/SiC composite parts.


The study was supported by the Special Fund of Technology Innovation from Hubei Province (No. 2016AAA021), the Guangdong Innovative and  Entrepreneurial  Research  Team  Program  (No. 2013C071), the Graduates’ Innovation Fund, Huazhong University of Science and Technology (No. 003100008) and the Fundamental Research Funds for the Central Universities (No. 2016YXZD069). The authors would also like to thank the Analytical and Testing Center of Huazhong University of Science and Technology (HUST) for the SEM observation tests, and the State Key Laboratory of Material Processing and Die & Mould Technology of HUST for mechanical and XRD tests. Special thanks are to Ningbo Vulcan Mechanical seals Manufacturing Co. Ltd for the LSI experiments.

Source: Wei Zhun - Nanyang Technological University, Rongzhen Liu - Huazhong University of Science and Technology, Yusheng Shi - Huazhong University of Science and Technology, Chunze Yan - Huazhong University of Science and Technology


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