SiC nanocomposites by carbothermal reaction between Si3N4 and carbon

Acta Materialia 50 (2002) 4831–4840 www.actamat-journals.com

Synthesis and properties of porous Si3N4/SiC nanocomposites by carbothermal reaction between Si3N4 and carbon Jian-Feng Yang a,∗, Guo-Jun Zhang b, Naoki Kondo a, Tatsuki Ohji a a

Synergy Materials Research Center, AIST, Nagoya 463-8687, Japan b Synergy Ceramics Laboratory, FCRA, Nagoya 463-8687, Japan

Received 9 November 2001; received in revised form 7 June 2002; accepted 31 July 2002

Abstract Porous Si3N4/SiC nanocomposites with 20 and 50 vol% SiC have been fabricated by two-step sintering of powder mixture of α-Si3N4, carbon powder with a mean size of 13 nm and 5 wt% Y2O3. Nano-sized SiC particles were formed through reactions: (1) carbon and surface SiO2 on the Si3N4 particles, and (2) carbon and Si3N4 particles. A bonding between Si3N4 grains was formed by the reactions so the original dimensions of the compact remained unchanged after sintering, and high porosity of 50–70% was obtained. This method of fabricating porous ceramics can be referred to as restrained sintering by reaction bonding (RSRB). It was shown that, with the variation of porosity, the samples has flexural strength of 20–100 MPa, and compressive strength of twice as much as the flexural strength. The composites with high porosity exhibited good permeability, a large strain to failure and ‘pseudo-plastic’ deformation behavior during the compressive test.  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Sintering; Composites; Mechanical properties; Stress–strain relationship measurements

1. Introduction Recently, there has been great interest in the creation of porous materials for a range of applications at room or high temperatures as molten metal filters, lightweight cores for sandwich panels, radiant burners, catalyst supports, sensors, bone grafts, dust or soot collectors, heat

Corresponding author. Tel.: +81-52-739-0156; fax: +8152-739-0136. E-mail address: [email protected] (J.-F. Yang). ∗

exchangers, and so on [1–3]. The pores play the important roles in reducing the weight of structures and Young’s modulus, increasing the surface area, or filtering fluids. Control of the porosity and pore geometry and enhancement of the mechanical properties of the porous structures are critically important to improve the performance and the reliability of these structures. Silicon nitride (Si3N4) ceramics and its composites, no matter they are dense or porous, are potential materials used at high temperatures. Various processing techniques have been developed to fabricate porous Si3N4 ceramics and its composite

1359-6454/02/$22.00.  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 2 ) 0 0 3 5 0 - 6

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materials for structural or functional applications [4–11]. As examples, porous Si3N4 ceramics can be fabricated by either using the fugitive substance [4,5], or by controlled sintering of Si3N4 powder with low content of sintering aid [6,7], or partial hot-pressing [11]. Porous Si3N4 ceramics with a microstructure of rod-like β-Si3N4 grains shows superior mechanical properties, such as high strength, good thermal shock resistance, and high strain and damage tolerance [4–9]. Thus, an α→β phase transformation with little or no densification is preferred when fabricating porous Si3N4 ceramics by sintering. The densification by particle rearrangement begins at the liquid-phase formation temperature, followed by the α→β transformation through a solution–reprecipitation process [6]. Because densification by particle rearrangement usually occurs before the phase transformation, and the reconstructive α→β phase transition also provides a driving force for densification [7], ceramics with porosity of higher than 50% is difficult to obtain by sintering at high temperature with full phase transformation [6,7,11]. By using the fugitive substance, such as starch, it is possible to adjust the porosity by changing the starch content. However, the large sized fugitive substance degrades the shaping property of the powder mixture. On the other hand, it is difficult to produce porous Si3N4 ceramics with fine pores [4], due to the restriction of the particle size of fugitive substance. That how to increase the porosity while keep the fine pore size, good shaping behavior and sinterability is still a unresolved problem in the fabrication of porous ceramics. In a previous study, the present authors have demonstrated a successful means to fabricate porous Si3N4 ceramics with high porosity through sintering powder mixture of Si3N4, 1–5 vol% (0.7– 3.4 wt%) carbon and some sintering aid [12]. Reactions between SiO2 or Si3N4 and the added carbon powder occur in the ceramic green body [13]: 2C(s) ⫹ SiO2(s)→SiC(s) ⫹ 2CO(g)

(1)

3C(s) ⫹ Si3N4(s)→3SiC(s) ⫹ 2N2(g)

(2)

The boundary temperature for the conversion from Si3N4 to SiC has been estimated as ~1435°C [14], and the formation of SiC starts between

1400° and 1450°C [15]. The reactions were also used to fabricate SiC based composites [16,17]. Reaction (1) occurs more easily than does reaction (2) in the Si3N4–SiO2–C system [18]. After SiO2 on the Si3N4 surface is consumed, the remained carbon reacts with Si3N4, according to reaction (2). The Si3N4 grains would be bonded together by the reaction-formed SiC particles located at the grain boundaries, and the shrinkage of Si3N4 ceramics was decreased. On other hand, the reactions result in a weight loss of about 32% theoretically, thus it would be expected to increase the porosity simultaneously through the reactions during sintering. In another word, the porosity can partly be adjusted by the amount of carbon. In this investigation, high porous Si3N4/SiC nanocomposites with 20 and 50 vol% SiC were fabricated by the carbothermal reaction between Si3N4 and carbon. Effects of carbon addition on the sinterability, porosity, microstructure development, and mechanical properties were investigated.

2. Experimental procedures 2.1. Fabrication of porous Si3N4/SiC nanocomposites Porous Si3N4/SiC nanocomposites were fabricated using the reaction process between Si3N4 and carbon powders. The starting powders were highpurity Si3N4 powder (SN-E10, UBE Industries Ltd, Tokyo, Japan; α ratio: ⬎95%, mean particle size: 0.5 µm, main impurities by weight: O=1.6%; C⬍0.2%; Cl, Fe, Ca, and Al2O3⬍50 ppm), carbon powder (No. 2600, Mitsubishi Chemical Corp., Tokyo, Japan; mean particle size 13 nm), and sintering aid Y2O3 powder (99.9% purity, Shin-etsu Chemical Co, Ltd, Tokyo, Japan). For the fabrication of the Si3N4/SiC nanocomposites with 20 and 50 vol% SiC, which are volume fractions of solid parts of Si3N4 and SiC in the composites, the following compositions of powder mixture were used: 89.63 and 83.46 wt% Si3N4, 5.37 and 11.54wt% carbon, respectively, calculated according to the reactions (1) and (2). 5 wt% Y2O3 as the sintering additive was added in. The resultant SiC content after sintering maybe somewhat

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deviate from the calculation due to the difference in weight loss of individual composition. The above powder mixtures were firstly wet-milled in methanol for 24 h, using high-purity Si3N4 media. Then, the slurry was dried and sieved through a 250 µm screen. The powder mixtures were then uniaxially pressed under 10 MPa, to form rectangular bars measuring 55 × 35 × 10 mm. Some of the samples were then cold isostatic pressed (CIPed) under the pressure of 200 MPa. With regard to sample designation, for example, SC20C means that SiC content in the composite is 20 vol% and formed by CIPing. The green bodies were placed in a BN-coated graphite crucible. A pre-sintering of 1600°C/4 h in argon atmosphere of gas pressure 0.6 MPa was used to obtain SiC particles through the reactions. Then the samples were sintered in a graphite resistance furnace at 1750°, 1850° and 1900°C for 2 h in a nitrogen atmosphere of gas pressure 0.6 MPa with the temperature rising rate of 15°C/min. This two step method enhances the SiC formation and Si3N4 matrix sintering. 2.2. Characterization 2.2.1. Density The bulk density and open porosity of the sintered products was measured by the Archimedes displacement method using distilled water. The porosity was calculated from the relative density and theoretical density, which was calculated by the rule of mixtures. 2.2.2. Reaction process, phase composition, microstructure and pore size For investigating the reaction process, thermogravimetry-differential thermal analysis (TG-DTA, Model TG-DTA2020, MAC Science Co. Ltd, Yokohama, Japan) was performed at a heating rate of 10 K/min using the powder mixture. Crystalline phases in sintered samples were identified by Xray diffraction (XRD, Model RINT2000, Rigaku Co. Ltd, Tokyo, Japan) analysis at 40 kV and 100 mA using CuKα radiation, a step width of 0.02°, and a counting time of 0.3 s. The microstructures were characterized by scanning electron microscopy (SEM, Model S-5000, Hitachi Ltd,

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Tokyo, Japan). The pore size distribution was measured by mercury porosimetry. 2.2.3. Liquid permeability A liquid suction test was used to measure the liquid permeability of the obtained porous Si3N4/SiC nanocomposites using water [19]. The square samples for the permeability test were machined to a size of 10 × 10 × 15 mm (height) by diamond wheel of 600 mesh, and then the samples were carefully cleaned and dried. 2.2.4. Flexural strength The specimens were machined into test bars of 36 × 4 × 3 mm, for bending strength measurement according to JIS R1601 (Japanese Industrial Standard). All surfaces of the test bars were finally ground on an 800-grit diamond wheel, and the edges were beveled. The directions of both the diamond grinding and the beveling were parallel to that of the bar length. The three-point bending strength was measured on specimen bars with a span of 30 mm at a cross-head speed of 0.5 mm/min, using universal testing machine (Autograph AG-10TC, Shimadzu Corp., Kyoto, Japan). Each final value was averaged over six measurements. 2.2.5. Compressive test Square Si3N4/SiC composite samples of 12.5 mm length and 5 × 5 mm section area having a porous structure were subjected to uniaxial compressive loading. The tests were conducted at a cross-head speed of 0.5 mm/min on the same universal testing machine as that used for the flexural strength test. The edges of each traction surface were chamfered using 600-grit SiC paper before testing. Compressive strength was obtained by dividing the peak load by the cross-sectional area of the sample. Each final value was averaged over 3–6 measurements. 3. Results and discussion 3.1. Densification According to reaction (1) and (2), considering 1.6% oxygen on the surface of raw Si3N4 particles,

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weight loss after sintering are 9.28 and 18.82% theoretically for the composites with 20 and 50 vol% SiC, respectively. The measured weight loss, relative density and porosity of sintered bodies were shown in Table 1. As can be seen in Table 1, weight loss of sample sintered at 1650° and 1750°C was lower than the theoretical value, while that of the samples sintered at 1850° and 1900°C was almost the same. The low weight loss at low temperature of 1650° and 1750°C was come from a fact that a mass of the carbon remained after the sintering and the reaction was not completed, indicating that it need a high temperature to transfer carbon to SiC for the bulk samples. The shrinkage of all samples after sintering was within 1%, indicated that the original dimensions of the powder compact remain virtually unchanged during sintering. The low shrinkage indicated a significant effect of carbon addition on the densification behavior, and the shrinkage was apparently restrained by the reaction between carbon and Si3N4. The densification of Si3N4 ceramics using Y2O3–Al2O3 oxide as sintering additives begin at temperature above 1400°C [20], at which the glass phase formed and particle rearrangement was the main densification mechanism. The reaction between carbon and Si3N4 (reaction (2)) also

begin at this temperature, and prior to it, reaction between carbon and surface SiO2 (reaction (1)) begins. The occurrence of these reactions were at surface of Si3N4 particles, and the reactant located either at their surface or between the Si3N4 particle, which result in a reaction bonding between Si3N4 grains. It was well known that glass phase play an important role in the Si3N4 particle rearrangement, and the densification of Si3N4 ceramics by liquid phase sintering dependent on the characteristics of glass phase, such as amount and viscosity [20]. At relatively low temperature of lower than 1600°C, the viscosity of glass phase is high, so the densification is limited [20]. As the reactions begin almost simultaneously with the glass formation, the bonding of the Si3N4 particles by the reaction formed SiC particles is very likely to obstacle the movement of Si3N4 particles and restrain their rearrangement. With increase in sintering temperature, the increased viscosity of glass phase indicated the densification tendency, however, as the bonding among Si3N4 grains was strong, hardly did the grain rearrangement take place. Finally low sintering shrinkage (within 1%) and high porosity after sintering were resulted in. This method of producing porous ceramics can be defined as the restrained sintering by reaction bonding (RSRB).

Table 1 Summary of weight loss, relative density, and porosity of the sintered porous Si3N4/SiC nanocomposite samples Sample

Relative density of green body

SC20

0.397

SC20C

0.551

SC50

0.407

SC50C

0.573

Sintering Weight loss (%) temperature (°C) 1650 1750 1850 1900 1650 1750 1850 1900 1650 1750 1850 1900 1650 1750 1850 1900

7.2 7.9 9.0 9.4 7.3 7.6 9.0 9.4 12.7 13.4 18.0 18.5 14.0 13.9 17.8 18.3

Relative density after sintering 0.387 0.384 0.360 0.359 0.489 0.488 0.498 0.494 0.363 0.359 0.311 0.314 0.455 0.458 0.447 0.442

Total porosity

0.613 0.616 0.640 0.641 0.511 0.512 0.502 0.506 0.637 0.641 0.689 0.686 0.545 0.542 0.553 0.558

Open porosity

0.611 0.615 0.630 0.640 0.493 0.511 0.475 0.429 0.625 0.619 0.686 0.679 0.524 0.478 0.488 0.535

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Due to the low shrinkage and high weight loss during the reactions mentioned above, the relative densities after sintering were even lower than the relative densities of the green bodies. The higher content the SiC reaction-formed, the lower the relative density and the higher the porosity were. The porosities were also shown in Table 1. With the combination of the SiC content changing and degree of the CIPing, the final relative densities were adjustable from 30 to 50%. The pores were almost in open type. As pointed out in JIS R1634, the open porosity was limited to be lower than 10%, however, our careful surface water removal ensured the validity of the Archimedes method in measurement of high porous ceramics. Comparing with the density obtained by weight:volume ratio, the difference between the resultant density was within 1%, indicated that the results obtained from the Archimedes method were believable. 3.2. Differential thermal analysis TG-DTA was performed on the powder mixture of specimen with 20 vol% SiC, and the curve is shown in Fig. 1. Because of the temperature limitation of the TG-DTA apparatus, the analysis was stopped at 1700°C. It can be concluded from this curve that the reaction begins to take place at about 1400°C, and with further rising in temperature a sharp decrease in the TG curve and a sharp

Fig. 1. Thermogravimetry-differential thermal analysis of the powder mixture of sample 20% SiC at a heating rate of 10°C/min in Ar.

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increase in the DTA curve occurs. On the other hand, it can be seen that the reaction took place gradually in view of the broad exothermic peak on the curve. 3.3. X-ray diffraction The XRD patterns for samples sintered at 1750°, 1850°, and 1900°C are shown in Fig. 2. It can be seen that the reaction between α-Si3N4 and carbon could not take place completely when sintered at 1750°C, so that a mass of α-Si3N4 remained in the sample. In addition, the existence of carbon would also prohibit the α–β Si3N4 phase transformation. This result is not coincident with the TG-DTA analysis discussed in the above section, probably due to that the sample used for the TG-DTA analysis was powder and that for the XRD was the sintered compact. At high temperatures of 1850° and 1900°C, the reaction took place completely, so that β-Si3N4, 3C-SiC and 2H-SiC could be detected. Beside the Si3N4 and SiC phases, some other phases such as Y2Si3O3N4 and YNSiO2 formed from the sintering additive, surface SiO2 on the Si3N4 particles and Si3N4, were also detected. Due to the amorphous type of the starting reactant C there is no peak in the patterns. 3.4. Microstructure Fig. 3 shows the microstructure of green body of SC20 and samples SC20 and SC50 sintered at various temperatures. The micrograph of green body exhibited very fine particles composed of fine α-Si3N4 particles and ultra-fine carbon particles (Fig. 3a). At high temperature, transformation from α-Si3N4 to β-Si3N4 occurs, and the β-Si3N4 can be easily identified in SEM microstructure due to its elongated grain morphology. The SiC particle should be very fine due to the use of the carbon particle with the size of 13 nm. Heated at 1700°C, fine β-Si3N4 elongated grains exhibited (Fig. 3b), however, equiaxial α-Si3N4 and carbon particles still remained, due to the incomplete reaction, which also was indicated by the XRD result. With the increasing the sintering temperature to 1850° or 1900°C, grain growth of β-Si3N4 was shown,

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Fig. 2. XRD patterns of the samples obtained at various temperatures. (a) 20% SiC. (b) 50% SiC.

and no carbon particles can be seen (Fig. 3c and d). Similar to the dense Si3N4/SiC nanocomposites, the high SiC content resulted in the restrained βSi3N4 grain growth [21], as can be seen in Fig. 3(e) and (f) for SC50 samples. It can be seen that the SiC particles obtained from reaction between carbon and Si3N4 has a very fine size of about 100– 200 nm, which was almost independent on sintering temperature. The SiC particles locate around the β-Si3N4 grains and are suggested to prohibit the densification of the Si3N4 matrix, as analyzed in previous section. The microstructure contains a large fraction of pores as can be seen in Fig. 3. The pore size distribution in the sample with 50% SiC is shown in Fig. 4 indicating the very fine pore size. The samples sintered at different temperatures have almost the same pore size, but samples with low porosity (CIPed before sintering) show a fine pore size distribution. 3.5. Liquid permeability

Fig. 3. SEM micrographs of Si3N4/SiC nanocomposites of SC20 sample: (a) green body, (b) heated at 1750°C, (c) heated at 1850°C, and (d) heated at 1900°C, and SC50 sample (e) heated at 1850°C, and (f) 1900°C.

The liquid permeability of a porous material is closely related to the open porosity and pore size. Fig. 5 shows the liquid permeability of the two samples of SC50 sintered at 1850° and 1900°C, and SC50C sintered at 1850°C. It can be seen that

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gas permeability, the trend of decreased permeability with low open porosity and fine pore size should be the same. 3.6. Flexural strength

Fig. 4. Pore size distributions of porous Si3N4/SiC composites with 50 vol% SiC. The samples are sintered at 1850° and 1900°C respectively, and the CIPed sample was sintered at 1900°C.

The flexural strength is almost independent on the sintering temperature, while it is closely related with the material density (porosity). The flexural strength as a function of the volume fraction of porosity is shown in Fig. 6. The flexural strength is affected by the carbon contents (or the SiC contents), low addition of carbon black results in a considerable high strength. This is probably attributed to the formation of a high fraction of fibrous β-Si3N4 grains microstructure in sample with low carbon content (also low SiC content). The failure trends followed the exponential equation proposed first by Ryshekewitch and Duckworth [22,23]: sf ⫽ s0e-bP

(3)

where σf is the strength of the porous structure in compression, σ0 is the strength of the nonporous structure, P stands for the volume fraction of porosity, and b is an empirical constant. Ryshekewitch determined that the value of b varies between 6 and 9 at different loading conditions for zirconia and alumina ceramics [22]. Flexural strengths of samples with 20 and 50

Fig. 5. Liquid (water) permeability of the porous Si3N4/SiC nanocomposites with 50 vol% SiC.

the SC50 samples sintered both at 1850° and 1900°C have a high liquid permeability of 1.03 and 1.14 kg/m2 s1/2. This result is coincident with the data shown in Table 1, that is, these two materials have almost the same open porosity, and the sample SC50 sintered at 1900°C has larger pore size. Relative low liquid permeability for the SC50C sample is considered to be from the low porosity comparing with the SC50 samples. Although the liquid permeability is something different from the

Fig. 6. Flexural strength against volume fraction of porosity for the porous Si3N4/SiC nanocomposites.

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vol% are fitted in Eq. (3). The equations can be given as: 20%SiC: sf ⫽ 992e⫺4.5548P

(4)

50%SiC: sf ⫽ 1755e⫺6.4791P

(5)

The two equations are also plotted in the Fig. 6, and it can be seen that there is a rapid drop of the strength for the sample with 50 vol% SiC, comparing with the strength drop of samples with 20 vol% SiC. The σ0 values for each case are unrealistic, and is not surely the strength of the dense sample, as the σ0 values were obtained from the exponential fitting, only using strength date of high porous ceramics. Together with the low strength of porous ceramics with 50% SiC, the high σ0 value for the 50% SiC samples also indicated the rapid drop of the strength.

Fig. 8. Strain to failure of samples with different SiC content and porosity under compression.

3.7. Compressive test The results of the failure stress and strain to failure under compression test are shown in Fig. 7 and Fig. 8. The compressive strength is almost twice as much as the flexural strength. The higher density, such as for CIPed sample, resulted in the higher compressive strength. Typical stress–strain plot of some selected samples is shown in Fig. 9. It can be seen from this figure that the relatively high dense material (SC50C) exhibits a strong elastic response, while the low dense material (SC50) Fig. 9. Stress and strain plot under uniaxial loading for SC50 and SC50C sintered at 1850°C.

Fig. 7. Failure stress of samples with different SiC content and porosity under compression.

exhibits a pseudo-plastic response [24]. This pseudo-plastic deformation behavior is resulted from an increased strain to failure as shown in Fig. 9. The high strain to failure of near 6% was obtained for the SC50 sample. The high strain to failure is a result of high porosity of about 70% for the sample SC50. When the load reached the peak value, the samples begun cracking, but the further propagation of the crack was inhibited by the porous structure. Thus, for the samples SC20, SC20C and SC50C, the failure of samples in the compressive test was sudden breaking, however, for the sample SC50, the failure

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originated from some pores and broken gradually and bulk breaking could not be found as can be seen in Fig. 10. The remained debris resisted the load until full failure occurred, which was similar to the plastic deformation in metallic material. It is apparent that the ‘pseudo-plastic’ deformation behavior in compressive test is closely related with the porosity. For porous ceramics with high porosity, pore–crack interaction should be taken into account when explain the ‘pseudo-plastic’ deformation behavior. According to Rice’s analysis [25], under the compressive loading crack generation and propagation occur primarily due to local tensile stress from pores or other heterogeneity. Compressive failure of brittle materials is typically the result of cumulative damage from a number of crack generation–linkage–propagation events rather than a single weak link source, and stress concentrations from pores are seen as sources of crack generation, or propagation in compressive loading. Thus for the porous Si3N4/SiC nanocomposites with relatively low porosity of 50%, the pore distance is large and a linkage of cracks results in a large propagation of the cracks. As a result, catastrophic failure occurs after reaching the material strength. On the other hand, for the porous Si3N4/SiC nanocomposites with relatively high porosity of 70%, the pore distance is small and a linkage of pores results in a small propagation of the cracks. As a result, the crack propagation is short and does not result in the catastrophic failure suddenly. After reaching

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the material strength, the specimen is destroyed to several pieces along the direction parallel with compressive loading. With the load increase, the remained debris of the sample still resists the load, but a decreased load value to the deformation exhibits, which is not similar to the plastic deformation behavior in the metallic material. 4. Conclusions With restrained sintering by reaction bonding (RSRB) method, Si3N4/SiC nanocomposites with high porosity were fabricated using carbothermal reaction between Si3N4 and carbon. The investigation of this kind of materials leads to the following results: 1. Controlled porosity is obtainable by varying the carbon content and green density. Due to the near-to-zero shrinkage after sintering, the densities of the samples after sintering were lower than the green density, and samples with porosity of 苲50–70% were obtained. 2. The samples exhibit the microstructure that is composed of fibrous Si3N4 grains and nanosized SiC particles, fine pore size, and a good permeability due to the high porosity. 3. The flexural strength increases with the decrease in porosity, and is between 20 and 100 MPa. The samples with small fraction of SiC particles have relatively high flexural strength due to the high fraction of fibrous Si3N4 grains in microstructure. 4. The compressive strength is about twice as much as the flexural strength and the samples with high porosity exhibit obvious ‘pseudo-plasti’ deformation behavior under uniaxial compression. This ‘pseudo-plastic’ deformation behavior can be explained by the pore–crack interaction, as the high porous ceramics exhibited the short crack propagation during the fracture. Acknowledgements

Fig. 10. Photograph of fracture progress in the ‘pseudo-plastic’ behavior region.

This work was supported by AIST, METI, Japan, as part of the Synergy Ceramics Project.

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The authors are members of the Joint Research Consortium of Synergy Ceramics.

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