MATERIALS 兴 HIENGE& ENGIEERING ELSEVIER Materials Science and Engineering A 475(2008)217-223 www.elseviercom/locate/msea and silicon nitride ceramic matrix composites prepard ide Microstructure and properties of particle reinforced silicon c by chemical vapor infiltration Yongsheng liu Laifei Cheng, Litong Zhang, Yunfeng hua, Wenbin Yan National Key Laboratory of Thermostricture Composite Materials, Northwestern Polytechnical University. Xi'an Shaanxi 710072, People's Republic of China Received 28 November 2006; received in revised form 10 April 2007: accepted 10 April 2007 Introduction: Particle reinforced silicon carbide andsilicon nitride ceramic matrix composites were fabricated using designed particle agglomeration and chemical vaporinfiltration(CVi)technique. Scanning electron microscopy(SEM)and Tra electron microscopy (TEM)were employed to observe the microstructures of the preforms and as-infiltrated composites. In the preform, the inter-agglomeration and intra-agglomeration pores had an approximate size of 500-800 um and 5-10 um, respectively. After infiltrated, a large amount of silicon carbide and silicon nitride matrix yere infiltrated in the preform, the sizes of inter-agglomeration and intra -agglomeration pores were 200-400 um and 2-4 um, respectively. The Sic and Si3N4 whiskers were observed in the residual intra-agglomeration pores. The flexural strength of SiC(p/SiC composites changed with first-step and second-step pressure. The maximum of the strength was 284 MPa, and the ratio of the retained strength was 95.4% at 1600C. The fracture toughness of SiC(p)/SiC composites was around 7 Mpa m". The Si3 Na(p/Sia N4 composite attained an acceptable strength, 113.4 MPa and low dielectric constant, about 4.2-4.3 @2007 Elsevier B v. All rights reserved Keywords: Particle agglomeration; Ceramic matrix composites; Chemical vapor infiltration; Silicon carbide; Silicon nitride 1. Introduction particular applications depends on their high temperature per- formance, which is intimately linked to the amount and nature Sic and Si3N4 are two of the most important thermo- of the sintering additives, on the microstructure characteristics, structural ceramic materials [1]. In addition, Si3N4 is considered and eventually, on the second phase distribution and content to be a suitable candidate for high temperature radomes because [10-14 of the following properties:(1)a high mechanical strength; (2) The conventional methods to fabricate particle reinforced Sic a good thermal shock resistance; (3)excellent resistance to rain or Si3N4 composites are sintering methods such as pressureless and sand erosion; (4)an acceptable dielectric constant; (5)a low sintering, hot pressing, and hot isostatic pressing, employing dielectric losses[2]. However, the brittleness of these ceramics sintering additives. The residue of sintering additives existed mits their applications. Therefore, these ceramics have been as continuous intergranular glassy phase, which will soften at reinforced by particle [3], whisker [4] and fiber [5]. Particle elevated temperature. Furthermore, CTE mismatch of matrix, reinforced ceramic matrix posites have received great con- intergranular phase and the reinforcements resulted in the for- cern due to their improved fracture toughness and thermal shock mation of microcracks in the composites, which had a notably resistance and isotropic properties for a variety of high temper- negative effect on high temperature mechanical and chemical ture, high stress and severe erosion applications in aerospace, properties [15-19 hot engine, and energy conversion devices [6-9]. The suitabil- SiC and Si3 N4 matrix composites without sintering additives ity of the isotropic SiC and Si3N4 matrix composites to these can be fabricated by using hot isostatic pressing, which needs ultra-high temperature and pressure and ultra-fine and purity powders. Besides the high costs, it is difficult to fabricate SiC or Corresponding author. Tel +8629 8848 6068 823: fax:+8629 88494620. Si3N4 matrix composites by employing HIP process due to the E-mailaddressliuys99067@163.com(y.Liu). echnical difficulties [13, 20]. Chemical vapor infiltration(CVi)
Materials Science and Engineering A 475 (2008) 217–223 Microstructure and properties of particle reinforced silicon carbide and silicon nitride ceramic matrix composites prepared by chemical vapor infiltration Yongsheng Liu ∗, Laifei Cheng, Litong Zhang, Yunfeng Hua, Wenbin Yang National Key Laboratory of Thermostructure Composite Materials, Northwestern Polytechnical University, Xi’an Shaanxi 710072, People’s Republic of China Received 28 November 2006; received in revised form 10 April 2007; accepted 10 April 2007 Abstract Introduction: Particle reinforced silicon carbide and silicon nitride ceramic matrix composites were fabricated using designed particle agglomeration and chemical vapor infiltration (CVI) technique. Scanning electron microscopy (SEM) and Transmission electron microscopy (TEM) were employed to observe the microstructures of the preforms and as-infiltrated composites. In the preform, the inter-agglomeration and intra-agglomeration pores had an approximate size of 500–800 m and 5–10 m, respectively. After infiltrated, a large amount of silicon carbide and silicon nitride matrix were infiltrated in the preform, the sizes of inter-agglomeration and intra-agglomeration pores were 200–400 m and 2–4m, respectively. The SiC and Si3N4 whiskers were observed in the residual intra-agglomeration pores. The flexural strength of SiC(p)/SiC composites changed with first-step pressure and second-step pressure. The maximum of the strength was 284 MPa, and the ratio of the retained strength was 95.4% at 1600 ◦C. The fracture toughness of SiC(p)/SiC composites was around 7 Mpa m1/2. The Si3N4(p)/Si3N4 composite attained an acceptable strength, 113.4 MPa and low dielectric constant, about 4.2–4.3. © 2007 Elsevier B.V. All rights reserved. Keywords: Particle agglomeration; Ceramic matrix composites; Chemical vapor infiltration; Silicon carbide; Silicon nitride 1. Introduction SiC and Si3N4 are two of the most important thermostructural ceramic materials[1]. In addition, Si3N4 is considered to be a suitable candidate for high temperature radomes because of the following properties: (1) a high mechanical strength; (2) a good thermal shock resistance; (3) excellent resistance to rain and sand erosion; (4) an acceptable dielectric constant; (5) a low dielectric losses [2]. However, the brittleness of these ceramics limits their applications. Therefore, these ceramics have been reinforced by particle [3], whisker [4] and fiber [5]. Particle reinforced ceramic matrix composites have received great concern due to their improved fracture toughness and thermal shock resistance and isotropic properties for a variety of high temperature, high stress and severe erosion applications in aerospace, hot engine, and energy conversion devices [6–9]. The suitability of the isotropic SiC and Si3N4 matrix composites to these ∗ Corresponding author. Tel.: +86 29 8848 6068 823; fax: +86 29 8849 4620. E-mail address: liuys99067@163.com (Y. Liu). particular applications depends on their high temperature performance, which is intimately linked to the amount and nature of the sintering additives, on the microstructure characteristics, and eventually, on the second phase distribution and content [10–14]. The conventional methods to fabricate particle reinforced SiC or Si3N4 composites are sintering methods such as pressureless sintering, hot pressing, and hot isostatic pressing, employing sintering additives. The residue of sintering additives existed as continuous intergranular glassy phase, which will soften at elevated temperature. Furthermore, CTE mismatch of matrix, intergranular phase and the reinforcements resulted in the formation of microcracks in the composites, which had a notably negative effect on high temperature mechanical and chemical properties [15–19]. SiC and Si3N4 matrix composites without sintering additives can be fabricated by using hot isostatic pressing, which needs ultra-high temperature and pressure and ultra-fine and purity powders. Besides the high costs, it is difficult to fabricate SiC or Si3N4 matrix composites by employing HIP process due to the technical difficulties [13,20]. Chemical vapor infiltration (CVI) 0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.04.031
Y Liu et al. Materials Science and Engineering A 475(2008)217-223 has been demonstrated to be a very effective and matured enough machined, and cut into specimens with the same nat preparation method to fabricate SiC or Si3N4 matrix with ultra- of SiC(p)/SiC composites to further deposit Si3N4 matrix fo pure and controllable grain sizes[21-25. An alternative method 100h to fabricate particle reinforced ceramic matrix composites is to deposit SiC or Si3 N4 matrix on the internal surfaces of the 2.3. Characterization and tests porous particle preforms by CVI, similar to the fabrication of fiber reinforced SiC matrix composites. The SiC or Si3N4 matrix The macro-structure characterizations of the agglomera composites reinforced by particles can be expected to possess tions and preform were performed by a numeral camera. The excellent high temperature mechanical and chemical properties microstructure of the preform and as-infiltrated composite were if CVi is used to fabricate SiC or Si3N4 matrix with virtual examined by scanning electron microscopy (SEM)(Hitachi elimination of sintering additives In previous works [26-29]the $-4700, Japan), and transmission electron microscopy(TEM silicon carbide and silicon nitride agglomerate particle preform JEM3010) using foils prepared by ion-beam thinning. The bend specimens with dimension 3 mm x 4 mm x 40 mm were In the present work, two kinds of ceramic matrix compos- used to evaluate the 3-point flexural strength of two kinds of ites, SiCp/SiC and Si3 N4(p)/Si3 N4 composites were prepared composites with a cross head speed of 0.5 mm min-and a using particle agglomerations and chemical vapor infiltration span length of 20 mm using a SANS CMT4304 instrument rocess.The relationship between the pore structure and gas Five specimens were tested to obtain an average value. The 3- diffusion mechanism in preform were discussed. Finally, the point flexure strength of SiC(p)/SiC composites was also tested microstructure and properties of the two kinds of composites with temperature increasing from room temperature to 1600oC were reported Micro-hardness and fracture toughness were determined at roon temperature by Vickers indentation method with 196N load 2. Experiment procedure for 30s, 10 times on each of five samples per data point. The dielectric constant of the Si3N4(p/Si3 N4 composite was mea- 2.1. Preparation of the particle preform sured using Precision Impedance Analyse Instrument(Agilent 4294A, American), five times on each of three samples per data The raw silicon carbide(SiC wt %>98%)and silicon nitride point. (Si3N4 wt %>98%) particles had a mean grain size of 3. 0 um and were mixed for 5 h in a planetary milling with Al2O3 milling 3. Results and discussion balls using binder, poly-vinyl butyral(PVB), and plasticizer, dibutyl phthalate. The as-treated Sic or Si3N4 particles were 3. 1. Microstructure of particle agglomerations and preform packed in a stainless-steel die, and cold-pressed at room tem- perature and a pressure of 15 MPa or 20 MPa. The cold-press Fig. I shows typical morphologies of the particle agglom- pressure to form particle agglomeration can be named as the erations and preform. Fig. 1(a) shows that the particle first-step pressure. The as received preforms were then crushed agglomerations were grain-like shape, which were prepared and sieved to select SiC or Si3 Na particles agglomerations with through cold-press and crush process. The grain-like agglome size ranging from 0.3 to 0.6 mm. Finally, the selected particle ations would inlay in ceramic matrix when the particle preforms agglomerations were cold-pressed at room temperature and a were infiltrated. The reinforcing mechanisms of agglomerations pressure of 3.0-9.0 MPa to form the particle preform used to include micro-crack deflection, micro-crack divarication and fabricate composites. The cold-press pressure to form parti- agglomerations fracture, which were discussed in [29]. Fig. 1(b) cle preform for infiltration can be named as the second-step is the whole image of agglomerations preform prepared by cold- press process. The particle preform is uniform and porous, with a density of 1.29 g/cm, and porosity 59.6% 2. 2. CVI process Fig. 1(c)and(d) show the microstructure of the preform There are a large amount of connectable pores with a size of C matrix was deposited from MTS(content CH3 SiCl3> approximately 500-800 um inter-agglomeration in Fig. 1(c) 980wt %). Hydrogen(content H2299.99%)was a carrier gas However, the size of intra-agglomeration pore changes from 5 of MTS Argon(content H2>99.9%)was used as dilution gas. to 10 um in Fig. 1(d). The small intra-agglomeration pores are The deposition conditions were as follows: (MTS)/H2= 1/10 for also connectable basically, but they are more small and complex 400h at P=3 kPa, Ar=350 ml/min, and T=1000C. The spec- The pore size and distribution are in accordance with [29] imens were machined, and cut into specimens with dimension In CVI process, the transfer of reaction gases is very imp 3mm x4 mm x 40 mm to further deposit SiC matrix for 100h. tant to infiltration. However, the gas m In porous solids, the gas Si3N4 matrix was deposited from silicon tetrachloride size and shape of pores in the preform (SiCl4>99.99 wt %o and iron99.99%). Hydrogen(content H2299999%) was free path of the gas molecule (1). The diffusion mechanisms a carrier gas of Sicl4. The process conditions are as fol- include Fick diffusion, Knudsen diffusion and transition diffu- lows: SiClA/NH3=1/3 for 200 h at P=2 kPa, H2=100 ml/min, sion, when d> 100A, d s01A, and d=0. 1-100A respectively. Ar=200 ml/min, and T=900C. The specimens were also In a multi-phase gas system, i can be calculated from formula
218 Y. Liu et al. / Materials Science and Engineering A 475 (2008) 217–223 has been demonstrated to be a very effective and matured enough preparation method to fabricate SiC or Si3N4 matrix with ultrapure and controllable grain sizes[21–25]. An alternative method to fabricate particle reinforced ceramic matrix composites is to deposit SiC or Si3N4 matrix on the internal surfaces of the porous particle preforms by CVI, similar to the fabrication of fiber reinforced SiC matrix composites. The SiC or Si3N4 matrix composites reinforced by particles can be expected to possess excellent high temperature mechanical and chemical properties if CVI is used to fabricate SiC or Si3N4 matrix with virtual elimination of sintering additives. In previous works[26–29] the silicon carbide and silicon nitride agglomerate particle preform was designed. In the present work, two kinds of ceramic matrix composites, SiC(p)/SiC and Si3N4(p)/Si3N4 composites were prepared using particle agglomerations and chemical vapor infiltration process. The relationship between the pore structure and gas diffusion mechanism in preform were discussed. Finally, the microstructure and properties of the two kinds of composites were reported. 2. Experiment procedure 2.1. Preparation of the particle preform The raw silicon carbide (SiC wt.% > 98%) and silicon nitride (Si3N4 wt.% > 98%) particles had a mean grain size of 3.0m, and were mixed for 5 h in a planetary milling with Al2O3 milling balls using binder, poly-vinyl butyral (PVB), and plasticizer, dibutyl phthalate. The as-treated SiC or Si3N4 particles were packed in a stainless-steel die, and cold-pressed at room temperature and a pressure of 15 MPa or 20 MPa. The cold-press pressure to form particle agglomeration can be named as the first-step pressure. The as received preforms were then crushed and sieved to select SiC or Si3N4 particles agglomerations with size ranging from 0.3 to 0.6 mm. Finally, the selected particle agglomerations were cold-pressed at room temperature and a pressure of 3.0–9.0 MPa to form the particle preform used to fabricate composites. The cold-press pressure to form particle preform for infiltration can be named as the second-step pressure. 2.2. CVI process SiC matrix was deposited from MTS (content CH3SiCl3 ≥ 98.0 wt.%). Hydrogen (content H2 ≥ 99.99%) was a carrier gas of MTS. Argon (content H2 ≥ 99.9%) was used as dilution gas. The deposition conditions were as follows: (MTS)/H2 = 1/10 for 400 h at P = 3 kPa, Ar = 350 ml/min, and T = 1000 ◦C. The specimens were machined, and cut into specimens with dimension 3 mm × 4 mm × 40 mm to further deposit SiC matrix for 100 h. Si3N4 matrix was deposited from silicon tetrachloride (SiCl4 ≥ 99.99 wt.% and iron ≤ 10−5) and ammonia gas (content NH3 ≥ 99.99%). Hydrogen (content H2 ≥ 99.999%) was a carrier gas of SiCl4. The process conditions are as follows: SiCl4/NH3 = 1/3 for 200 h at P = 2 kPa, H2 = 100 ml/min, Ar = 200 ml/min, and T = 900 ◦C. The specimens were also machined, and cut into specimens with the same size as that of SiC(p)/SiC composites to further deposit Si3N4 matrix for 100 h. 2.3. Characterization and tests The macro-structure characterizations of the agglomerations and preform were performed by a numeral camera. The microstructure of the preform and as-infiltrated composite were examined by scanning electron microscopy (SEM) (Hitachi S-4700, Japan), and transmission electron microscopy (TEM, JEM3010) using foils prepared by ion-beam thinning. The bend specimens with dimension 3 mm × 4 mm × 40 mm were used to evaluate the 3-point flexural strength of two kinds of composites with a cross head speed of 0.5 mm min−1 and a span length of 20 mm using a SANS CMT4304 instrument. Five specimens were tested to obtain an average value. The 3- point flexure strength of SiC(p)/SiC composites was also tested with temperature increasing from room temperature to 1600 ◦C. Micro-hardness and fracture toughness were determined at room temperature by Vickers indentation method with 196 N load for 30 s, 10 times on each of five samples per data point. The dielectric constant of the Si3N4(p)/Si3N4 composite was measured using Precision Impedance Analyse Instrument (Agilent 4294A, American), five times on each of three samples per data point. 3. Results and discussion 3.1. Microstructure of particle agglomerations and preform Fig. 1 shows typical morphologies of the particle agglomerations and preform. Fig. 1(a) shows that the particle agglomerations were grain-like shape, which were prepared through cold-press and crush process. The grain-like agglomerations would inlay in ceramic matrix when the particle preforms were infiltrated. The reinforcing mechanisms of agglomerations include micro-crack deflection, micro-crack divarication and agglomerationsfracture, which were discussed in [29]. Fig. 1(b) is the whole image of agglomerations preform prepared by coldpress process. The particle preform is uniform and porous, with a density of 1.29 g/cm3, and porosity 59.6%. Fig. 1(c) and (d) show the microstructure of the preform. There are a large amount of connectable pores with a size of approximately 500–800m inter-agglomeration in Fig. 1(c). However, the size of intra-agglomeration pore changes from 5 to 10m in Fig. 1(d). The small intra-agglomeration pores are also connectable basically, but they are more small and complex. The pore size and distribution are in accordance with [29]. In CVI process, the transfer of reaction gases is very important to infiltration. However, the gas transfer mainly lies on the size and shape of pores in the preform. In porous solids, the gas diffusion mechanism depends on pore diameters (d), and mean free path of the gas molecule (λ). The diffusion mechanisms include Fick diffusion, Knudsen diffusion and transition diffusion, when d ≥ 100λ, d ≤ 0.1λ, and d = 0.1–100λ respectively. In a multi-phase gas system, λ can be calculated from formula
Y Liu et al. /Materials Science and Engineering A 475(2008)217-22 Agglomerati Particle eration pores Intra-particle pores 10 and preform prepared by silicon nitride particles(a) particle agglomerations; (b) particle preform; (c)the glomeration pores; and (d)the intra-agglomeration pores (1)[30 ent that a large amount of Sic is infiltrated in the inter- and intra-agglomerations pores. The reinforcing SiC particles P and agglomerations are enveloped by CVI SiC. The resid (1) ual pore inter- and intra-agglomeration is unavoidable,but the amount of pore decreases obviously. The residual pore where k is Boltzmann's constant. k=1. 10-23 J/K: Tis the size inter-agglomeration is about 200-400 um, which is rather smaller than that before infiltration. The residual pore size intra- temperature, K; oi or oj is the effective diameter of molecule i agglomeration is about 2-4um, which is also smaller than or j, m: P is partial pressure of gas, Pa; m; or m; is the molecule that before infiltration. The size changes of inter-and intra- nass of i or j specimen gas. As an example, the a values of all gas agglomeration pores before and after infiltration shows that the in SiCl4-NH3-H2-Ar system are listed in Table 1. Therefore, amount of deposit in inter-agglomeration pore is more than that e can conclude that the gas diffusion mechanism is transition of intra-agglomeration pore diffusion in the inter-agglomeration pores due to the d=5-89A Interface microstructure observed by tEM shows a thin SiC The diffusion mechanism belongs to Knudsen diffusion in the layer can be deposited on the near surface of SiC reinforcing intra- agglomeration pores due to d≤0.1λ particle during CVI(see Fig. 2(c). The incomplete burnout of residual carbon detaches from SiC reinforcing particle and exists 3.2. Microstructure and properties of SiCpysic composites between the CVI column SiC and the common SiC matrix, which will exist as a defect layer, resulting in weak interfacial bonding Fig. 2 shows typical microstructures of SiC(Py/SiC compos- strength for SiC(P)/SiC composites CvI SiC whiskers are found ites fabricated by CVI. Fig. 2(a)and(b)show the cross-section in some intra-agglomeration pores as shown in Fig. 2(d). The morphologies of the composites infiltrated for 500 h. It is appar- diameters of the whiskers are almost the same, about several Table Mean free paths of reactant gases and other coefficients in the SiCL-NH3-H2-Ar system Gas species Flux(ml min-) a(nm) M(gmol-) Partial pressure(Pa) 0.2915 0.5484 99 0.2900 17
Y. Liu et al. / Materials Science and Engineering A 475 (2008) 217–223 219 Fig. 1. Typical morphologies of particle agglomerations and preform prepared by silicon nitride particles (a) particle agglomerations; (b) particle preform; (c) the inter-agglomeration pores; and (d) the intra-agglomeration pores. (1) [30]. λi = ⎡ ⎣ n j=1 1 + mi mj π σi + σj 2 2 P kT ⎤ ⎦ (1) where k is Boltzmann’s constant, k = 1.38 × 10−23 J/K; T is the temperature, K; σi or σj is the effective diameter of molecule i or j, m; P is partial pressure of gas, Pa; mi or mj is the molecule mass ofi orjspecimen gas. As an example, the λ values of all gas in SiCl4–NH3–H2–Ar system are listed in Table 1. Therefore, we can conclude that the gas diffusion mechanism is transition diffusion in the inter-agglomeration pores due to the d = 5–89λ. The diffusion mechanism belongs to Knudsen diffusion in the intra-agglomeration pores due to d ≤ 0.1λ. 3.2. Microstructure and properties of SiC(p)/SiC composites Fig. 2 shows typical microstructures of SiC(P)/SiC composites fabricated by CVI. Fig. 2(a) and (b) show the cross-section morphologies of the composites infiltrated for 500 h. It is apparent that a large amount of SiC is infiltrated in the interand intra-agglomerations pores. The reinforcing SiC particles and agglomerations are enveloped by CVI SiC. The residual pore inter- and intra-agglomeration is unavoidable, but the amount of pore decreases obviously. The residual pore size inter-agglomeration is about 200–400m, which is rather smaller than that before infiltration. The residual pore size intraagglomeration is about 2–4 m, which is also smaller than that before infiltration. The size changes of inter- and intraagglomeration pores before and after infiltration shows that the amount of deposit in inter-agglomeration pore is more than that of intra-agglomeration pores. Interface microstructure observed by TEM shows a thin SiC layer can be deposited on the near surface of SiC reinforcing particle during CVI (see Fig. 2(c)). The incomplete burnout of residual carbon detaches from SiC reinforcing particle and exists between the CVI column SiC and the common SiC matrix, which will exist as a defect layer, resulting in weak interfacial bonding strength for SiC(P)/SiC composites. CVI SiC whiskers are found in some intra-agglomeration pores as shown in Fig. 2(d). The diameters of the whiskers are almost the same, about several Table 1 Mean free paths of reactant gases and other coefficients in the SiCl4–NH3–H2–Ar system Gas species Flux (ml min−1) σ (nm) M (g mol−1) Partial pressure (Pa) λ (m) H2 100 0.2915 2 270 37.7 SiCl4 30 0.5484 169.9 81 50.6 NH3 40 0.2900 17 108 90.1 Ar 200 0.3432 40 540 90.1
Y Liu et al. Materials Science and Engineering A 475(2008)217-223 (b) CVI SiC matrix CVI SiC matrix SiC particle Residual pore CvI SiC matrix (d) CVI SiC whisker CVI SiC al microstructure of SiC(P/SiC composites(a) polished cross-section of SiCp/SiC composites; (b)SEM image of intra-agglomeration; (c)TEM image between two SiC reinforcing particles; (d) SEM image of SiC whisker formed intra-agglor hundreds nanometer, and the lengths of the whiskers were about sen diffusion. There is Sic globularity at the whiskers terminal, which shows the growth mechanism of the CVI SiC whiskers is According to the whole CVI process, microstructure differ- vapor-liquid-solid process ence of Sic deposits could be explained based on the deposition The flexural strength of specimen as a function of the first eaction and above gas diffusion process. The whole deposition step pressure and second-step pressure is shown in Fig 3. The reaction of SiC is shown in reaction(2). strengths of Sic(P/SiC composites with 15 MPa first-step pres- sure are all higher than those with 20 MPa. Under the 15 MPa CH3 SiCl3 -> SiC+ 3HCI (2) first-step pressure, the strength of the composites decreases when The reaction gases transfer to the agglomeration surfaces nrough transition diffusion. Some of the reaction gases diffuse into the inner of the agglomerations through Knudsen diffusion then the deposition reaction occurs, and the Sic is deposited in 22s0 the inter- and intra-agglomeration pores as shown in Fig. 2(a) and(b). The CH3SiCl3 will also traverse the residual porous carbon layer through Knudsen diffusion, which leads to the t formation of CVI SiC on the surface of Sic reinforcing par- icle as shown in Fig. 2(c). As deposition time increased, more nd more Sic deposits were deposited on the agglomeration surface and intra-agglomeration. The intra-agglomeration pores would first be blocked out due to the smaller pore size. There- after, the inter-agglomeration pores would also be blocked out There existed inter-and intra-agglomerations pores because of Second-step pressure(MPa) the release of by-product gases. In some intra-agglomeration Fig 3. Flexural strength of composites as a function of the first-step pressure pores,the CvI SiC whiskers grow at the low super-saturation, and second-step pressure. N ) First-step pressure 15MPa; (aa) Second-step because the small amount reaction gas transfers through Knud- pressure 20MPa
220 Y. Liu et al. / Materials Science and Engineering A 475 (2008) 217–223 Fig. 2. Typical microstructure of SiC(P)/SiC composites (a) polished cross-section of SiCp/SiC composites; (b) SEM image of intra-agglomeration; (c) TEM image of interfaces between two SiC reinforcing particles; (d) SEM image of SiC whisker formed intra-agglomeration pores. hundreds nanometer, and the lengths of the whiskers were about 20m. According to the whole CVI process, microstructure difference of SiC deposits could be explained based on the deposition reaction and above gas diffusion process. The whole deposition reaction of SiC is shown in reaction (2). CH3SiCl3 1000 ◦C −→H2 SiC + 3HCl (2) The reaction gases transfer to the agglomeration surfaces through transition diffusion. Some of the reaction gases diffuse into the inner of the agglomerations through Knudsen diffusion, then the deposition reaction occurs, and the SiC is deposited in the inter- and intra-agglomeration pores as shown in Fig. 2(a) and (b). The CH3SiCl3 will also traverse the residual porous carbon layer through Knudsen diffusion, which leads to the formation of CVI SiC on the surface of SiC reinforcing particle as shown in Fig. 2(c). As deposition time increased, more and more SiC deposits were deposited on the agglomeration surface and intra-agglomeration. The intra-agglomeration pores would first be blocked out due to the smaller pore size. Thereafter, the inter-agglomeration pores would also be blocked out. There existed inter- and intra-agglomerations pores because of the release of by-product gases. In some intra-agglomeration pores, the CVI SiC whiskers grow at the low super-saturation, because the small amount reaction gas transfers through Knudsen diffusion. There is SiC globularity at the whiskers terminal, which shows the growth mechanism of the CVI SiC whiskers is vapor–liquid–solid process. The flexural strength of specimen as a function of the firststep pressure and second-step pressure is shown in Fig. 3. The strengths of SiC(P)/SiC composites with 15 MPa first-step pressure are all higher than those with 20 MPa. Under the 15 MPa first-step pressure, the strength of the composites decreases when Fig. 3. Flexural strength of composites as a function of the first-step pressure and second-step pressure. ( ) First-step pressure 15 MPa; ( ) Second-step pressure 20 MPa.
Y Liu et al. /Materials Science and Engineering A 475(2008)217-22 Table 2 shown in Table 2. fracture toughness and micro-hardness val Fracture toughness and micro-hardness value as a function of the second-step ues changes with the second-step pressure. The highest fracture roughness is 7. 11 MPamin2 The second-step pressure(MPa) Fig 4 shows strength measurements of SiC(py/Sic Micro-hardness(GPa) Fracture roughness(Mpam"2) 25.83 25.72 26 17 24.39 ites as a function of temperature. The room temperature strength is around 284 MPa and the value is almost retained up to 1600C This is due to SiC(PSiC composites can be fabricated by CV with virtual elimination of sintering additives, this is consistent with the study of Sic ceramics and its composites with intergral ular glassy phase free grain boundaries, whose high temperature 20 mechanical properties are much superior to that of Sic ceramic and particle reinforced SiC matrix composites with intergranular glassy phase [7] 3.3. Microstructure and properties of Si3 N4(p/Si3 N4 Fig 5 shows the typical microstructure of the Si3N4(p)/Si3N4 Fig. 4. Flexural strength with temperature as a function of SiCp/SiC composites used 15 MPa the first step pressure and 3 MPa second step pressur composites fabricated by CVI. There are large amounts of silicon nitride infiltrated inter-agglomerations and intra-agglomerations as shown in Fig. 5(a). Known from the cross-section morphol the second-step pressure changes from 3 to 5 MPa, then the ogy of composites, the composites is dense, but the residual strength increases when the second-step pressure changes from pores with an average size of 200-300 um inter-agglomerations 5 to 7 MPa, finally the strength decreases when the second- exist as shown in Fig. 5(a). The composite edge is denser than step pressure is above 7 MPa. The highest flexural strength is the inner because the edge is more favorable for gas transport 284 MPa, when the first-step pressure is 15 MPa and the second- and deposition reaction than the inner( Fig. 5(b). There is also step pressure is 7 MPa. The same function of strength changing residual pore intra-agglomeration as shown in Fig. 5(c). In some exists when the first-step pressure is 20 MPa and the highest intra-agglomeration pores, the CVi Si3 N4 whiskers also grow at strength is 265 Mpa. the low super-saturation(Fig. 5(d)), which is around 5 um. Fracture toughness and micro-hardness values, determined The flexural strengths and dielectric constants of sev by means of Vickers indentation method with a load of 196N, are eral Si3N4 materials are shown in Table 3. Compared with Coating Agglomeration (c) CvISin4 whisker Intra-agglomerations pore 5.0 um Fig. 5. Typical microstructure of Si3 N4(P/Si3N4 composites(a)cross-section; (b)composites edge coating; (c)intra-agglomerations: (d)CVI Si3N4 whisker
Y. Liu et al. / Materials Science and Engineering A 475 (2008) 217–223 221 Table 2 Fracture toughness and micro-hardness value as a function of the second-step pressure The second-step pressure (MPa) 3579 Micro-hardness (GPa) 25.83 25.72 26.17 24.39 Fracture roughness (Mpa m1/2) 6.93 6.87 7.11 6.74 Fig. 4. Flexural strength with temperature as a function of SiCP/SiC composites used 15 MPa the first step pressure and 3 MPa second step pressure. the second-step pressure changes from 3 to 5 MPa, then the strength increases when the second-step pressure changes from 5 to 7 MPa, finally the strength decreases when the secondstep pressure is above 7 MPa. The highest flexural strength is 284 MPa, when the first-step pressure is 15 MPa and the secondstep pressure is 7 MPa. The same function of strength changing exists when the first-step pressure is 20 MPa and the highest strength is 265 Mpa. Fracture toughness and micro-hardness values, determined by means of Vickers indentation method with a load of 196 N, are shown in Table 2. Fracture toughness and micro-hardness values changes with the second-step pressure. The highest fracture roughness is 7.11 MPa m1/2. Fig. 4 shows strength measurements of SiC(p)/SiC composites as a function of temperature. The room temperature strength is around 284 MPa and the value is almost retained up to 1600 ◦C. This is due to SiC(P)/SiC composites can be fabricated by CVI with virtual elimination of sintering additives, this is consistent with the study of SiC ceramics and its composites with intergranular glassy phase free grain boundaries, whose high temperature mechanical properties are much superior to that of SiC ceramic and particle reinforced SiC matrix composites with intergranular glassy phase [7]. 3.3. Microstructure and properties of Si3N4(p)/Si3N4 composites Fig. 5 shows the typical microstructure of the Si3N4(p)/Si3N4 composites fabricated by CVI. There are large amounts of silicon nitride infiltrated inter-agglomerations and intra-agglomerations as shown in Fig. 5(a). Known from the cross-section morphology of composites, the composites is dense, but the residual pores with an average size of 200–300m inter-agglomerations exist as shown in Fig. 5(a). The composite edge is denser than the inner because the edge is more favorable for gas transport and deposition reaction than the inner (Fig. 5(b)). There is also residual pore intra-agglomeration as shown in Fig. 5(c). In some intra-agglomeration pores, the CVI Si3N4 whiskers also grow at the low super-saturation (Fig. 5(d)), which is around 5 m. The flexural strengths and dielectric constants of several Si3N4 materials are shown in Table 3. Compared with Fig. 5. Typical microstructure of Si3N4(P)/Si3N4 composites (a) cross-section; (b) composites edge coating; (c) intra-agglomerations; (d) CVI Si3N4 whisker.
Y Liu et al. Materials Science and Engineering A 475(2008)217-223 Table 3 Flexural strength and dielectric constant of CVI Si3 Nap/Si3 N4 and other Si Na materials lexural strength(MPa) Density (gcm-3) Dielectric constant 113.4 4.32(30MHz) CVI 104.5 4.27(30MHz) 104.3 4.13(30MHz) 4.34(30MHz) Reaction sintered Si3 N4 [31] 5.6(810GHz) SioN nano-composite [331 747(8.5GHz),7.14(35GHz) Electromagnetic window that of SiC(p/Sic composites, the flexural strength of CVI Acknowledgments NA(P/Si3N4 composite is low, only 100 MPa around. The low strength of Si3 N4(p/Si3 N4 composite has resulted from This work was supported by the Key Foundation of National the low density, 2.03-2.10g/cm-3, and high open pore ratio, Science in China(90405015)and National Elitist Youth Founda- about 12% of the composite. The microstructure of the tion in China(50425208) This work was also supported by the two types of composites also verified that the SiC(py/sic Doctorate Foundation of Northwestern Polytechnical University has more high strength due to the denser cross-section ( CX200505) Compared with other reports, CVI Si3N4(P/Si3 N4 compos- References Although reaction sintered Si3N4 composite has a high strength, [1) Mitomo Mamoru, Petzow Gunter, MRS Bull. 20(1995)I 500 MPa 31], the composites also has a high dielectric con- [23J. Barta, M Manela, R. Fischer, Mater. Sci. Eng. 71(1984)265. tant, which is disadvantageous for radome applications On the [3] J.H. She, D L Jiang, S.H. Tan, J.K.Guo, Key Eng Mater. 108-110(1995) other hand, CVI Si3 N4(p)/Si3 N4 composite has a higher strength and almost similar dielectric constant, compared with sintered 4 Y Hua, L. Zhang, L. Cheng, J. Wang, Mater. Sci. Eng. A 428(2006) nano Si3N4 [33]. Therefore, the CVI Si3N4(p)/Si3N4 compos- (5)R.Naslain, Phys IV France 123(2005)3-17 ite is a good candidate materials for high temperature radome (6)X D Wang, G.J. Qiao, Z.H. Jin, J.Am. Ceram Soc. 87(2004)565 [8] G.C. Dodds, R.A. Tanzilli. United States Patent, Appl. 5891815, 1999 4. Conclusions 9] I.G. Talmy, C.A. Martin, H.A. Deborah, et al. United States Patent, Appl 5573986,1996 [10] K. Strecker, S Ribeiro, R. Oberacker, M.J. Hoffmann, Int J Refract Met. (1) The designed particle preform exist with two kinds of pores H Mater.22(2004)169 with size of 500-800 um inter-agglomerations and 5-10 um [11] U Paik, H C. Park, S.C. Choi, C.G. Ha, J.W. Kim, Y.G. Jung, Mater. Sci. intra-agglomerations, respectively. The agglomerate parti EngA334(2002)267 cle preform is uniform and porous, with density 1.29 of [12] H.J. Choi, Yw.Kim, M.Mitomo,TNishimura,JHLee,DY.Kim,Scripta g/cm and pore ratio 59.6%. The gas diffusion mechanism [13] M.J. Hoffmann, A. Geyer, R. Oberacker, J. Eur. Ceram Soc. 19(13-14) through inter-agglomerations pores belongs to transi- (1999)2359 tion diffusion and Knudsen diffusion intra-agglomerations [14] J.S. Park, Y. Katoh, A Kohyama, J K. Lee, J.J. Sha, H.K. Yoon, J Nucl pores Mater.329-333(2004)558 (2)SiC()/SiC and Sis N4(p)/Si3 N4 composites were fabricated [15K. Biswas, G Rixecker, EAldinger,J.Eur.CeramSoc.23(2003)1099 using the designed particle preform and chemical vapor 17]K. Strecker, M.J. Hoffmann, J. Eur. Ceram Soc. 25 (2005)80 infiltration technique. There existed large amounts of [18]QWHuang, L.H. Zhu, Mater. Lett. 59(2005)1732. silicon carbide and silicon nitride in the inter-and intra- [19] T. Tani, Compos. Part A-Appl S 30(1999)419. agglomerations pores, although there are some residual [20] P. Bhandhubanyong, T. Akhadejdamrong, JMater. Process. Technol. 63 pores ()The flexural strength of SiC(p /SiC composite changes with [21 R. Naslain, R Paller, X. Bourrat, G. Vignoles, Key Eng Mater. 159-160 molding-agglomerations pressure and molding-preforms [22]R. Naslain, Compos. Sci. Technol. 64(2004) pressure. The maximum of the strength was 284 MPa, and [23]N. Igawaa, T. Taguchi, T Nozawa, L L. Snead, T Hinoki, J C. McLaughlin, the ratio of the retained strength was 95.4% at 1600C, and Y. Katoh, S Jitsukawa, A. Kohyama, J. Phys. Chem. Solids 66(2005) fracture roughness was 7. 11 MPam. The Si3N4(p/Si3 N4 composite had an acceptable strength, 113. 4 MPa and low 4] Y Liu, L Cheng, L Zhang, et al., J Inorg. Mater. 20(5)(2005)979(in dielectric constant. about 4.2-4.3 [25] Y. Liu, L Cheng, L. Zhang, et al., J. Univ Sci. Technol. B, 2007, in press
222 Y. Liu et al. / Materials Science and Engineering A 475 (2008) 217–223 Table 3 Flexural strength and dielectric constant of CVI Si3N4p/Si3N4 and other Si3N4 materials Materials Flexural strength (MPa) Density (g cm−3) Dielectric constant CVI Si3N4p/Si3N4 113.4 2.08 4.32 (30 MHz) 104.5 2.03 4.27 (30 MHz) 104.3 2.09 4.13 (30 MHz) 104.3 2.15 4.34 (30 MHz) Reaction sintered Si3N4 [31] 500 – 5.6 (8–10 GHz) Sintered nano-Si3N4 [32] 89 2.29 4.8–5.7 SiON nano-composite [33] 190 – 4.78–5.00 SiBAlON composite [8] – – 7.47 (8.5 GHz), 7.14 (35 GHz) Electromagnetic window [9] 85 – 4.03 that of SiC(p)/SiC composites, the flexural strength of CVI Si3N4(p)/Si3N4 composite is low, only 100 MPa around. The low strength of Si3N4(p)/Si3N4 composite has resulted from the low density, 2.03–2.10 g/cm−3, and high open pore ratio, about 12% of the composite. The microstructure of the two types of composites also verified that the SiC(p)/SiC has more high strength due to the denser cross-section morphology. Compared with other reports, CVI Si3N4(p)/Si3N4 composite has an acceptable strength and low dielectric constant. Although reaction sintered Si3N4 composite has a high strength, 500 MPa [31], the composites also has a high dielectric constant, which is disadvantageous for radome applications. On the other hand, CVI Si3N4(p)/Si3N4 composite has a higher strength and almost similar dielectric constant, compared with sintered nano Si3N4 [33]. Therefore, the CVI Si3N4(p)/Si3N4 composite is a good candidate materials for high temperature radome applications. 4. Conclusions (1) The designed particle preform exist with two kinds of pores with size of 500–800m inter-agglomerations and 5–10 m intra-agglomerations, respectively. The agglomerate particle preform is uniform and porous, with density 1.29 of g/cm3 and pore ratio 59.6%. The gas diffusion mechanism through inter-agglomerations pores belongs to transition diffusion and Knudsen diffusion intra-agglomerations pores. (2) SiC(p)/SiC and Si3N4(p)/Si3N4 composites were fabricated using the designed particle preform and chemical vapor infiltration technique. There existed large amounts of silicon carbide and silicon nitride in the inter- and intraagglomerations pores, although there are some residual pores. (3) The flexural strength of SiC(p)/SiC composite changes with molding-agglomerations pressure and molding-preforms pressure. The maximum of the strength was 284 MPa, and the ratio of the retained strength was 95.4% at 1600 ◦C, and fracture roughness was 7.11 MPa m1/2. The Si3N4(p)/Si3N4 composite had an acceptable strength, 113.4 MPa and low dielectric constant, about 4.2–4.3. Acknowledgments This work was supported by the Key Foundation of National Science in China (90405015) and National Elitist Youth Foundation in China (50425208). This work was also supported by the Doctorate Foundation of Northwestern Polytechnical University (CX200505). References [1] Mitomo Mamoru, Petzow Gunter, MRS Bull. 20 (1995) 19–20. [2] J. Barta, M. Manela, R. Fischer, Mater. Sci. Eng. 71 (1984) 265. [3] J.H. She, D.L. Jiang, S.H. Tan, J.K. Guo, Key Eng. Mater. 108–110 (1995) 45–52. [4] Y. Hua, L. Zhang, L. Cheng, J. Wang, Mater. Sci. Eng. A 428 (2006) 346–350. [5] R. Naslain, J. Phys. IV France 123 (2005) 3–17. [6] X.D. Wang, G.J. Qiao, Z.H. Jin, J. Am. Ceram. Soc. 87 (2004) 565. [7] G. Magnani, BeltramiG, G.L. Minoccari, L. Pilotti, J. Eur. Ceram. Soc. 21 (2001) 633. [8] G.C. Dodds, R.A. Tanzilli. United States Patent, Appl. 5891815, 1999. [9] I.G. Talmy, C.A. Martin, H.A. Deborah, et al. United States Patent, Appl. 5573986, 1996. [10] K. Strecker, S. Ribeiro, R. Oberacker, M.J. Hoffmann, Int. J. Refract. Met. H Mater. 22 (2004) 169. [11] U. Paik, H.C. Park, S.C. Choi, C.G. Ha, J.W. Kim, Y.G. Jung, Mater. Sci. Eng. A 334 (2002) 267. [12] H.J. Choi, Y.W. Kim, M. Mitomo, T. Nishimura, J.H. Lee, D.Y. Kim, Scripta Mater. 50 (2004) 1203. [13] M.J. Hoffmann, A. Geyer, R. Oberacker, J. Eur. Ceram. Soc. 19 (13–14) (1999) 2359. [14] J.S. Park, Y. Katoh, A. Kohyama, J.K. Lee, J.J. Sha, H.K. Yoon, J. Nucl. Mater. 329–333 (2004) 558. [15] K. Biswas, G. Rixecker, F. Aldinger, J. Eur. Ceram. Soc. 23 (2003) 1099. [16] K.S. Cho, H.J. Choi, J.G. Lee, Y.W. Kim, Ceram. Int. 24 (1998) 299. [17] K. Strecker, M.J. Hoffmann, J. Eur. Ceram. Soc. 25 (2005) 801. [18] Q.W. Huang, L.H. Zhu, Mater. Lett. 59 (2005) 1732. [19] T. Tani, Compos. Part A-Appl. S 30 (1999) 419. [20] P. Bhandhubanyong, T. Akhadejdamrong, J. Mater. Process. Technol. 63 (1997) 277–280. [21] R. Naslain, R. Pailler, X. Bourrat, G. Vignoles, Key Eng. Mater. 159–160 (1999) 359. [22] R. Naslain, Compos. Sci. Technol. 64 (2004) 155. [23] N. Igawaa, T. Taguchi, T. Nozawa, L.L. Snead, T. Hinoki, J.C. McLaughlin, Y. Katoh, S. Jitsukawa, A. Kohyama, J. Phys. Chem. Solids 66 (2005) 551. [24] Y. Liu, L. Cheng, L. Zhang, et al., J. Inorg. Mater. 20 (5) (2005) 979 (in Chinese). [25] Y. Liu, L. Cheng, L. Zhang, et al., J. Univ. Sci. Technol. B, 2007, in press.
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