Study on the Fracture Origin of SiC-polycrystalline Fiber

Transcription

1 Research Article Trends in Nanotechnology & Material Science Study on the Fracture Origin of SiC-polycrystalline Fiber Hiroshi Oda 1 and Toshihiro Ishikawa *2 1 UBE Industries Ltd, Kogushi, Ube, Yamaguchi, , Japan 2 Tokyo University of Science, Yamaguchi, Daigaku-Dori, Sanyo-Onoda, Yamaguchi , Japan * Corresponding author: Tokyo University of Science, Yamaguchi, Daigaku-Dori, Sanyo-Onoda, Yamaguchi , Japan, Tel: ; ishikawa@rs.tus. ac.jp Copyright: Hiroshi Oda and Toshihiro Ishikawa. This is an open-access article distributed under the terms of the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original author and source are credited. Citation: Hiroshi Oda, Toshihiro Ishikawa (2016) Study on the Fracture Origin of SiC-polycrystalline Fiber. Trends in Nanotechnology & Material Science 1: 1-5. Received Date: January 08, 2016 Accepted Date: January 20, 2016 Published Date: January 30, 2016 Abstract We established a new method to precisely capture the first fracture surface of the SiC-polycrystalline fiber with very high modulus (about 400 GPa). This is a very simple method using only liquid paraffin and powder papers during tensile testing. Using this method, every fracture origins, which dominate the strengths of the SiC-polycrystalline fibers, were correctly detected and analyzed by the use of FE-SEM, EDS, and TEM. The strength was found to be in inverse proportion to one-half power of the defect size. The defect was mainly residual carbon, different crystalline phase or micro-pore. And also, it was clarified that surface defects were much more sensitive than inside defects to the strengths. By controlling the heat-treatment processes (degradation and sintering processes), which cause the defects, surface roughness of the fiber was remarkably improved from 6.8 nm to 4.9 nm. Keywords: SiC-polycrystalline fiber, Strength, Defect,Fracture origin Introduction SiC-polycrystalline fibers are well-known as an excellent heat-resistant inorganic fiber which can withstand over 1500 o C even in air [1-5]. Accordingly, lots of research on the fine structure and characteristics of the fibers have been actively performed [6-8]. Furthermore, several types of composite materials using the SiC fibers have been evaluated aiming to various applications [9-11]. And also, representative aircraft engine manufacturers are expecting actual applications of the SiC-polycrystalline fibers for jet engines and land-based gas turbines in the near future. By the way, the mechanical properties of the composite materials are dominated by the fiber s strength. Hence, to extend the application field, an increase in the mechanical strengths of these fibers is eagerly required. Until several years ago, researches on some defects (residual carbon, and so on) contained or formed in the SiC-polycrystalline fibers have been carried out [12-15]. However, they could not directly make mention of the relationship between the initially contained defects and the fiber s strength. So, to achieve the further improvement of the fiber s strength, the formation mechanism of the defects has to be clarified, and the defects should be remarkably reduced. Of these SiC-polycrystalline fibers, Tyranno SA (UBE s SiC-polycrystalline fiber) is synthesized via further heat-treatment (~2000 o C) of an amorphous Si-Al-C-O fiber, which is synthesized from polyaluminocarbosilane [3]. During the aforementioned further heat-treatment, the degradation of the amorphous Si-Al-C-O fiber and the sintering of the degraded fiber proceed as well, accompanied by the release of CO gas and compositional changes, to finally obtain the dense SiC-polycrystalline fiber. Since these structural changes proceed in each filament, a strict control should be needed to minimize residual defects existing on the surface and in the inside of each filament. Actually, existence of some defects has been confirmed as mentioned above. Accordingly, by the decrease in the residual defects, much higher strengths will be expected. To achieve this objective, first, the correct detection of the residual defects and the realization of its influence on the fiber s strength are very important. In this paper, a new method to capture the first fracture surface, detection of the fracture origin (defect) of each filament, and the relationship between the defect and the fiber s strength will be discussed. And also, an experimental result on the remarkable improvement of the surface roughness will appear. 2. Experimental Procedure 2.1 Materials In this research, Tyranno SA (UBE s commercial SiC-polycrystalline fiber) and an amorphous Si-Al-C-O fiber (UBE s intermediate fiber) were used. The physical properties of Tyranno SA are shown in Table 1. Diameter μm 10 Number of filaments fil./yarn 800 Tex g/1000m 170 Tensile strength GPa 2.4 Tensile modulus GPa 380 Elongation % 0.7 Density g/cm Thermal conductivity W/mK 65 Coefficient of thermal expansion 10-6 /K 4.5 (RT-1000C) Chemical composition Table 1. Physical properties of Tyranno SA Si wt% 67 C 31 O <1 Ti - Zr - Al <2

2 2.2 Tensile test The tensile test was performed by a monofilament method using a universal testing machine (Orientec Corporation) Tensilon UTM-II-20 (25mm gauge length and 2mm/min cross-head speed at room temperature). Before the tensile testing, the diameter of each filament was measured with a digital microscope, model VHX-5000 (KEYENCE Corporation). Each filament was glued on the paper holder with elastic adhesive as shown in Figure 1. Before starting this research, there were some difficulties to catch the first fracture surface of the SiC-polycrystalline fiber, because of very fine diameter and high modulus of Tyranno SA. To remove these difficulties, we established a new simple method to prevent the scatter of the destroyed fragments. The tensile specimen was held between two thin papers containing liquid paraffin (Figure 1). These thin papers didn t touch both chucks of the tensile testing machine, so that the tensile testing could be performed without their undesirable effects. Using this simple method, we could catch all broken fragments, and then we could effectively obtain the first fracture surfaces. Paper holder Elastic adhesive Gauge length 25mm Tensile direction Tensile direction Tyranno SA fiber Thin film with liquid paraffin Figure 2: Weibull plot for the strengths of Tyranno SA. 3.2 Characterization of the fracture surface As mentioned before, we developed a new method for capturing the all fragments during the monofilament tensile testing of Tyranno SA, and then effectively observed their first fracture surface. Using FE-SEM, we observed the first fracture surface and detected the fracture origin. Figure 3 shows the typical fracture surface of Tyranno SA. Some river-like cracks propagating from the fracture origin can be seen as shown in Figure 3. This type of crack-propagation is named River Pattern. As mentioned above, the gathering point of the river-like cracks is considered as the fracture origin, which is the weakest point in the tensile specimen. Because, the specimen must be destroyed at the weakest point during the tensile testing. So, we studied on the fracture origin in detail using FE-SEM, and then the defect s size was measured. Furthermore, the fine structures of the fracture origins were investigated in detail using FE-SEM and electron distribution spectroscopy (EDS). Figure 1: The specimen s geometry during the tensile testing 2.3 Observation of the fracture origin The fracture surfaces of captured fragments were effectively observed using a field emission scanning electron microscope (FE-SEM), model JSM-700F (JEOL, Ltd.). After that, we sharpened the part of fracture origin by an etching machine using focused ion beam (FIB), and picked it up as a sample for the transmission electron microscope (TEM), model JEM-2100F (JEOL, Ltd.). TEM was used for observing the crystalline phase abnormally existing at the fracture origin. 3. Results and Discussion 3.1 Tensile test and Weibull analysis The results of the tensile test were arranged by a Weibull plot in Figure 2. These plots include the results of 200 specimens. In the region of ln (σ) >0.6, there is linear relationship. Otherwise, in the region lower than 0.6 of ln (σ), it shows some inclination. This indicates that multiple fracture mode had occurred. That is to say, in the lower strength region, there might be several types of fracture origins. So, the characterization of these fracture origins is very important for the improvement of the fiber s strength. In the next section, the results of characterization of the fracture surface and the fracture origin will be discussed. Figure 3: Typical fracture surface of Tyranno SA, showing the river pattern and the fracture origin. Next, we would like to show two types of fracture patterns. First one is the fracture surface of the relatively weaker filament. And another one is that of the relatively stronger filament. Figure 4 shows both fracture surfaces, from which different types of River Patterns can be observed. In case of former pattern (relatively lower strength), the river-like cracks propagated from the surface region. That is to say, in this case, the fracture origin existed in the surface region. And also, this type of fracture origin was relatively large. On the other hand, in the case of later pattern (relatively higher strength), the river-like crack propagated from the internal small point. In this case, it was considered that the fracture origin existed in the inside of the filament. Both fracture origins were considered as a remarkable Defect. From these observations, in the case of this type of SiC-polycrystalline fiber, it was decided that the larger defect existing in

3 the surface region is much more sensitive to the tensile strength, compared with the smaller internal defect. Figure 4: Two types of fracture surfaces of Tyranno SA fibers with relatively (a) lower strength and (b) higher strength. By the use of our new method for capturing all broken fragments of the tensile specimen, we could also obtain the perfect pairs of the broken surfaces. The SEM images of these pairs are shown in Figure 5, where the fracture origins existing in the both upper- and lower- broken surfaces are shown in the dotted circles. As can be seen from this figure, both upperand lower- surfaces fitted nicely. Of these, the right filament was broken at the grain boundary in a relatively larger defect existing in the surface region. In this fracture origin, it was considered that some interfacial secondary phase probably existed at the grain boundary. So, this type of filament showed a relatively low strength. Regarding the existence of that type of different phases in the fracture origin, we will discuss in detail in the later section. Figure 5: A typical pairs of fracture surfaces of Tyranno SA containing fracture origin. Figure 6: Relationship between the defect size and the tensile strength of Tyranno SA. 3.3 Characterization of the fracture origin Figure 7 shows EDS mapping of fracture surface of average Tyranno SA with middle strength. Some agglomeration was partly observed by detailed elementary analysis. In the case of this filament, it was estimated that the fracture was initiated from this anisotropic region, where the carbon content was slightly large. Basically, Tyranno SA is a nearly stoichiometric SiC-polycrystalline fiber. That is to say, this fiber ought to be composed of the stoichiometric composition of SiC crystal. By the way, the precursor polymer for preparing Tyranno SA was a polyaluminocarbosilane synthesized by the reaction of polycarbosilane with tetra-butoxyaluminum at 300oC in N2 atmosphere. The aforementioned anisotropy of the composition (larger carbon content) was caused from the carbon-rich precursor. In the production process of Tyranno SA, to achieve the stoichiometric composition of SiC crystal, the following degradation reactions of an intermediate fiber (Si-Al-C-O fiber) and the sintering of the degraded fiber effectively occur. (1) SiO2 + 3C = SiC + 2CO ( G < 0 over 1522oC) Main-reaction (2) SiO + 2C = SiC + CO ( G < 0 at all temperatures range) Sub-reaction Though these degradation reactions have to be absolutely performed for achieving the stoichiometric composition, partly anisotropic reaction in some filaments must result in the aforementioned carbon-rich regions. Next, we would like to show the relationship between the defect size and the tensile strength of Tyranno SA along with the SEM images of the fracture surfaces (Figure 6). As can be seen from this figure, the strength of the fiber increases with a decrease in the defect size. Besides, the fracture origin of the relatively weaker filament existed in the surface region. And also, mischief by the internal fine defects of each filament was found to be relatively small. Regarding these results, the strengths are found to follow the well-known law which shows that the strength is decreased in inverse proportion to one-half power of the defect size as follows. (σ: tensile strength, K1C: fracture toughness, a: defect size) excelyticspublishers.com Figure 7: The results of EDS mapping of Tyranno SA at the fracture surfaces. Hiroshi Oda

4 Next, we investigated the fine structure of both a peculiar fracture origin and a normal fracture origin using TEM micrograph. The obtained results were shown in Figure 8-a and Figure 8-b. In these figures, the fracture origins were shown in the dotted circles. We sharpened these parts of fracture origins by an etching machine using focused ion beam (FIB), and picked them up as samples for the transmission electron microscope (TEM). And also, we observed the selected-area diffraction (SAD) patterns of these specimens. From these diffraction patterns, in this peculiar fracture origin (Figure 8-a) an alpha-sic crystal (hexagonal crystalline structure) existed, and the crystal existing in the normal fracture origin (Figure 7-b) was a beta-sic crystal (cubic crystalline structure). By the way, concerning two types of SiC crystals, an alphasic crystal has a cleavage plane, so that the region containing relatively larger alpha-sic crystals must show a weaker property, compared with normal regions composed of beta-sic crystals. Actually, the filament of Figure 8-a showed the lower strength (3.1 GPa) compared with the other filament of Figure 8-b (4.1 GPa). From these results, for obtaining the uniform, higher strength, the aforementioned anisotropic regions should be reduced under an improved production conditions. treatment reactions. Figure 9: Schematic interior structure of Tyranno SA containing several types of structural imperfections. Figure 8-a: Transmission electron micrograph and SAD patterns of Tyranno SA with a peculiar surface region. By strictly controlling the heat-treatment conditions, we could experimentally improve the surface roughness from 6.8 nm to 4.9 nm. About 20% smoother surfaces could be obtained. The result of this improvement is shown in Figure 10. Figure 8-b: Transmission electron micrograph and SAD patterns of Tyranno SA with relatively higher strength. As mentioned above, present Tyranno SA has contained several types of structural imperfections. The schematic interior structure of the present SiC-polycrystalline fiber is shown in Figure 9. These structural imperfections must be remained or formed during the production process. As shown in section 3.3, Tyranno SA was produced by a further heat-treatment (degradation reactions and sintering) of an intermediate Si-Al-C-O fiber. In particular, the degradation reaction accompanied by the release of CO gas has to be strictly controlled for reducing the interior defects. And also, the aforementioned sub-reaction (SiO+2C=SiC+CO) contains a gaseous SiO reactant, whose partial pressure remarkably appears over 1150oC. Accordingly, we have to prevent the vaporization of the SiO gas from each filament before the aforementioned sub-reaction proceeds, as the disappearance of SiO gas leads to increase in the residual carbon in each filament. Under these considerations, to increase in the tensile strength by reducing the residual defects, we have to strictly control the Ar gas flow and atmospheric conditions in the reactor during the heatexcelyticspublishers.com Figure 10: The result of the improvement of the surface roughness. 4. Conclusions For effectively capturing all of the fractured fragments of the SiCpolycrystalline fiber with very high modulus, we established a new simple using only liquid paraffin and powder papers during tensile testing. Using this method, it became possible to capture every fracture origins, which dominate the strengths of the SiC-polycrystalline fibers. By a detailed analysis of the fracture origins using FE-SEM, EDS, and TEM, in the Hiroshi Oda

5 fracture origins, the existence of excess carbon, different crystalline phase, or micro-pore was observed. And also, the strength was found to be in inverse proportion to one-half power of the defect size. Furthermore, it was clarified that surface defects were much more sensitive than inside defects to the strengths. And finally, by controlling the heat-treatment processes (degradation and sintering processes), which cause the defects; surface roughness of the fiber was remarkably improved from 6.8 nm to 4.9 nm. References 1. Flores O, Bordia RK, Nestler D, Krenkel W, Motz G (2014) Ceramic Fibers Based on SiC and SiCN Systems: Current Research, Development, and Commercial Status, Advanced Engineering Materials 16: Colombo P, Mera G, Riedel R, Soraru GD (2013) Polymer-Derived Ceramics: 40 Years of Research and Innovation in Advanced Ceramics, Ceramic Science and Technology: Applications Edited by Ralf Riedel and I-Wei Chen 4: Ishikawa T, Kohtoku Y, Kumagawa K, Yamamura T, Nagasawa T (1998) Highstrength alkali-resistant sintered SiC fibre stable to 2200oC, Nature, 391: Takeda M, Urano A, Sakamoto J, Imai Y (1998) Microstructure and oxidative degradation behavior of silicon carbide fiber Hi-Nicalon type S, Journal of Nuclear Materials Ishikawa T (2005) Advances in Inorganic Fibers. Advanced Polymer Science (Springer-Vrlag Berlin Heidelberg) 178: Huguet-Garcia J, Jankowiak A, Miro S, Gosset D, Serruys Y, et al. (2015) Study of the Ion-Irradiation Behavior of Advanced SiC Fibers by Raman Spectroscopy and Transmission Electron Microscopy. Journal of American Ceramic Society, 98: Huguet-Garcia J, Jankowiak A, Miro S, Vandenberghe T, Grygiel C, (2015) Journal of Materials Research. 30: Kondo S, Hinoki T, Nonaka M, Ozawa K (2015) Irradiation-induced shrinkage of highly crystalline SiC fibers. Acta Materialia 83: Idris MI, Konishi H, Imai M, Yoshida K, Yano T (2015) Neutron Irradiation Swelling of SiC and SiCf/SiC for Advanced Nuclear Applications. Energy Procedia 71: Ortona A, Fend T, Yu HW, Raju K, Fitriani P, et al. (2015) Tubular Si-infiltrated SiCf/SiC composites for solar receiver application Part 1: Fabrication by replica and electrophoretic deposition. Solar Energy Materials and Solar Cells 132: Katoh Y, Nozawa T, Shih C, Ozawa K, Koyanagi T, et al. (2015) High-dose neutron irradiation of Hi-Nicalon Type S silicon carbide composites. Part 2: Mechanical and physical properties. Journal of Nuclear Materials 462: Ho CY, Tsai SC, Lin HT, Chen FR, Kai JJ (2013) Microstructural investigation of Si-ion-irradiated single crystal 3C-SiC and SA-Tyrannohex SiC fiber-bonded composite at high temperatures. Journal of Nuclear Materials. 443: Sauder C, Lamon J (2007) Tensile Creep Behavior of SiC-Based Fibers with a Low Oxygen Content. Journal of American Ceramic Society 90: Silvestroni L, Fabbriche DD, Sciti D (2015) Tyranno SA3 fiber-zrb2 composites. Part 1: Microstructure and densification. Materials and Design 65: Sha JJ, Park JS, Hinoki T, Kohyama A (2007) Tensile behavior and microstructural characterization of SiC fibers under loading. Materials Science and Engineering A 456: