Proceedings of the 42nd International Conference on Advanced Ceramics and Composites, Ceramic Engineering and Science Proceedings, Issue 3

Advancing Frontiers of Ceramics for Sustainable Societal Development: International Symposium in Honor of Dr. Mrityunjay Singh

PROGRESS IN POLYMER-DERIVED SiC-BASED FIBERS: IMPROVEMENT OF SURFACE ROUGHNESS

T. Ishikawa, K. Suwa, and R. Usukawa

Tokyo University of Science, Yamaguchi

1-1-1 Daigaku-Dori, Sanyo-Onoda, Yamaguchi 756-0884, Japan

ABSTRACT

Polymer-derived SiC-polycrystalline fiber (Tyranno SA) shows excellent heat-resistance up to 2000°C, and relatively high strength. However, to extend the application field, increase in the fiber’s strength is eagerly required. Up to now, through our research, the relationship between the strength and residual defects of the fiber, which were formed during the heat-treatment processes (degradation of raw fiber and sintering), has been clarified. In this paper, we addressed the relationship between the heat-treatment condition and the surface roughness of the obtained SiC-polycrystalline fiber, using three different raw fibers (Elementary ratio: Si1Al0.01C1.5O0.4˜0.5) and three types of carbon vessel (Open system, Partially closed system, and Closed system). With increase in the oxygen content in the raw fiber, the degradation during the heat-treatment process easily proceeded. This led to nearly stoichiometric composition of SiC crystal. And, higher oxygen content in the raw fiber and the closed system led to relatively high partial-pressure of SiO at the initial stage of the degradation process. In this case, the degradation reactions (SiO+2C=SiC+CO and SiO2+3C=SiC+2CO) in the inside of each filament became faster, and then the CO partial pressure at the surface region of each filament was found to be increased. In consequence, according to Le Chatelier’s principle, the surface degradation reaction and grain growth of formed SiC crystals would be considered to become slower. That is to say, using the raw fiber with higher oxygen content and closed system (highest CO content in the reactor), much smoother surface of the SiC-polycrystalline fiber could be achieved.

INTRODUCTION

Since the first precursor ceramics using polycarbosilane was developed, lots of polymer-derived SiC-base fibers have been developed. Through these developments, the heat-resistances of the SiC-based fibers were remarkably increased from 1200°C to 2000°C. Of these fibers, SiC-polycrystalline fibers (Tyranno SA, Hi-Nicalon Type S, and Sylramic) show the highest heat-resistance up to 2000°C, and then have been actively evaluated for aerospace applications as SiC/SiC composites [1-5]. However, to extend the application field, increase in the fiber’s strength is eagerly required. Up to now, through our research, the relationship between the strength and the residual defects contained in the fiber, which were formed during the production processes, has been clarified [6-9]. In these researches, we have proposed several new methods for reducing the residual defects, and demonstrated them using the conversion process from amorphous Si-Al-C-O fiber to SiC-polycrystalline fiber (Tyranno SA). Tyranno SA is produced by heat-treatment processes of amorphous Si-Al-C-O fiber which is synthesized from polyaluminocarbosilane. During the heat-treatment processes, a degradation of the Si-Al-C-O fiber and the subsequent sintering of the degraded fiber proceed as well, accompanied by a release of CO gas and compositional changes, to finally obtain the dense structure. Since these structural changes proceed in each filament, a strict control should be needed to minimize residual defects. As mentioned above, to reduce the residual defects, we proposed new conversion processes and demonstrated them. And then, using these new processes, the inside defects were remarkably reduced. In this case, by controlling the advantageous degradation-reaction and preventing the disappearance of gaseous SiO from each filament, residual carbon (one of residual defects) was remarkably reduced along with prevention of abnormal surface grain growth. In consequence, the surface roughness was relatively improved compared with the obtained surface using previous processes. However, the surface roughness was very sensitive depending on the atmospheric condition during the degradation process, and then obtaining much smoother surface was a relatively difficult problem. Smoother surface of the fiber is very important for obtaining good fibrous fracture behavior of ceramic matrix composites (CMCs) [10]. Accordingly, it should be important to clarify the relationship between the process condition and the surface roughness of the obtained fiber.

Here, we describe the formation mechanism of the surface structure and the relationship between the process condition and the surface roughness of the obtained SiC-polycrystalline fiber.

EXPERIMENTAL

The SiC-polycrystalline fiber (Tyranno SA) was synthesized by heat-treatment up to 1900°C of an amorphous Si-Al-C-O fiber, which is synthesized from polyaluminocarbosilane. The polyaluminocarbosilane was synthesized by a reaction of polycarbosilane with tetra-butoxyaluminum at 300°C in nitrogen atmosphere. A spun fiber was obtained by melt-spinning of the polyaluminocarbosilane, and then the spun fiber was cured at around 200°C in air. The cured fiber was fired at around 1300°C in nitrogen atmosphere to obtain the amorphous Si-Al-C-O fiber. The Si-Al-C-O fiber was composed of SiC fine crystals, oxide phases (estimated forms: SiO2, AlOx), and excess carbons. By the way, as mentioned above, since in this synthesis we used polyaluminocarbosilane which was synthesized by the reaction of polycarbosilane and tetra-butoxyaluminum, we presumed that the aluminum existed as some oxide forms in the Si-Al-C-O fiber. In the next step, the amorphous Si-Al-C-O fiber was heat-treated up to around 1500°C in argon gas atmosphere. During the heat-treatment, by the existence of the oxide phase and excess carbon in the fiber, the amorphous Si-Al-C-O fiber was degraded accompanied by a release of CO gas to obtain a porous degraded fiber. The porous degraded fiber was composed of a nearly stoichiometric SiC composition containing small amount of aluminum. In this case, since a part of the aluminum contained in the amorphous Si-Al-C-O fiber was found to vaporize as some oxide materials during the heat-treatment process, consequently a very small amount of aluminum (less than 1wt%) was contained in the degraded fiber. By the existence of the small amount of aluminum, at the next step, an effective sintering proceeded in each degraded filament composed of the nearly stoichiometric SiC crystals during further heat-treatment up to 2000°C in argon atmosphere. The production scheme of the polymer-derived SiC polycrystalline fiber using the Si-Al-C-O fiber as the raw fiber is shown in Fig.1. As mentioned in our previous papers, degradation reactions of the amorphous Si-Al-C-O fiber enclosed in red frame in this figure (Fig.1) strongly affects the final fine-structure (Cross-section and Surface) of the SiC-polycrystalline fiber. Especially, reaction condition concerning CO gas content during the degradation process is most important [6, 9]. Accordingly, in this research we adopted three types of reaction vessel (Open system, Partially closed system, and Closed system) made of carbon shown in Fig.2. For the heat-treatment (degradation reaction and sintering) of the Si-Al-C-O fiber, we used Super High Temperature Inert Gas Furnace (NEWTONIAN Pascal-40, Produced by NAGANO) under argon gas flow (1 L/min). The size of the heating zone (made of graphite and C/C composites) is 35 mm in diameter and 40 mm in height. Several types of raw fibers (about 10 mg) (Elementary ratio: Si1Al0.01C1.5O0.4˜0.5) were used and located in each vessel. The programing rate and the maximum temperature were 400°C/min and 2000°C, respectively.

Figure 1 The production scheme of the SiC-polycrystalline fiber using a raw Si-Al-C-O fiber

Figure 2 Experimental condition for research on fiber’s surface roughness

The surfaces and cross sections of the obtained fibers were observed using a field emission scanning electron microscope (FE-SEM), model JSM-700F (JEOL, Ltd.). Parts of surface region and inside of the several samples were sharpened by an etching machine using focused ion beam (FIB), and then the fine structures were observed by the transmission electron microscope (TEM), model JEM-2100F (JEOL, Ltd.). Surface roughness was observed using Atomic Force Microscope (AFM), model AFM 5000II (Hitachi, Ltd.).

RESULTS AND DISCUSSION

Morphological changes during the degradation process

As mentioned before, for obtaining the SiC-polycrystalline fiber, at the first step, the amorphous Si-Al-C-O fiber was heat-treated up to 1500°C in Ar gas atmosphere. During the heat-treatment process, by the existence of the oxide phase and excess carbon in the fiber, the amorphous Si-Al-C-O fiber was degraded accompanied by a release of CO gas to obtain a porous degraded fiber. This degradation of the Si-Al-C-O fiber proceeds mainly by the following two types of reactions.

The porous degraded fiber was composed of a nearly stoichiometric SiC composition containing small amount of aluminum (less than 1 wt%). By the existence of the small amount of aluminum, at the next step, an effective sintering proceeded in each degraded filament composed of the nearly stoichiometric SiC crystals during further heat-treatment up to 2000°C in Ar gas atmosphere. And then, the dense SiC-polycrystalline fiber was obtained. The morphological changes of each filament during the further heat-treatment are shown in Fig.3.

Figure 3 Morphological changes during the degradation and sintering processes

As can be seen from this figure (Fig.3), the degradation proceeded from outside to inside of the amorphous Si-Al-C-O fiber. And, regarding the SiC-crystalline size of the obtained sintered fiber, the surface SiC-crystals were relatively small compared with the inside crystals. This phenomenon was considered to be caused by the atmospheric condition (Especially: CO gas content) during the degradation process. That is to say, the degradation reactions (SiO2 + 3C = SiC + 2CO and SiO + 2C = SiC + CO) are strongly dominated by the CO gas content in the reactor. According to Le Chatelier’s principle, the higher the CO content becomes, the slower the reaction becomes. Anyway, the abovementioned degradation proceeds in the inside of each filament accompanied by a release of CO gas. So, the inside of each filament is saturated by the formed CO gas, and the surplus CO gas is ejected from the surface region to the outside. Furthermore, on the surface region of each filament, some boundary layer composed of CO gas must be formed. By these changes, in consequence, some CO gas distribution would be formed from the inside to the surface region of each filament. Accordingly, the degradation in the inside of each Si-Al-C-O filament was considered to proceed as shown in Fig.4.

Figure 4 General degradation in the inside of each Si-Al-C-O filament

As can be seen from this figure (Fig.4), the first degradation reaction smoothly proceeds near the surface by relatively low CO content compared with that in the inside. So, the degradation proceeds from outside to inside. By the way, it is considered that the CO content at the surface region is relatively high compared with that in the inside of each filament. Accordingly, it is estimated that at the surface region, the degradation slowly proceeds to consequently prevent the SiC-crystalline grain growth. In the next section, we will address the relationship between the surface degradation reaction and the surface roughness of the obtained SiC-polycrystalline fiber.

Change in the surface roughness of the SiC-polycrystalline fiber

As mentioned before, both the degradation reaction and grain growth of the formed SiC crystals are strongly affected by the CO content in the reactor during the degradation reaction. As shown in Fig.2, we used three different types of vessel (Open system, Partially-open system, and Closed system) and several types of raw fiber (Elementary ratio: Si1Al0.01C1.5O0.4˜0.5) for changing actual CO content in the reactor. Under our reaction condition, we calculated that the maximum CO content in the reactor was 73vol% when we used both the closed system and the raw fiber with highest oxygen content (Si1Al0.01C1.5O0.5), whereas the minimum CO content was 0.2vol% when we used both the open system and the other raw fiber with lowest oxygen content (Si1Al0.01C1.5O0.4). The surface structures (SEM images) of the obtained SiC-polycrystalline fibers are shown in Fig.5. These fibers were obtained by heat-treatment of the different types of Si-Al-C-O fiber at 1900°C in Ar using the three different types of vessels. As can be seen from this figure (Fig.5), higher oxygen content of the raw fiber and closed system led to the much smoother surface.

Figure 5 Changes in the surface structures (SEM images) using different types of raw fiber and vessels

These results are closely related to the CO content in the reactor during the degradation reaction. The much smoother surface was obtained using the raw fiber with highest oxygen content (Si1Al0.01C1.5O0.5) and the closed system. As can be seen from these results, the surface roughness is effectively controllable by changing the degradation conditions (Especially; CO content in the reactor). The most important factors for change in the CO content in the reactor are (1) Oxygen content of the raw Si-Al-C-O fiber, and (2) Reactor system (Open system, Partially-open system, and Closed system). By change in the combination of these factors, different degrees of surface roughness could be obtained as can be seen from Fig.5. Some phenomena caused by increase in the oxygen content of the raw Si-Al-C-O fiber are shown as follows along with some differences caused by change in the system (Open system, Partially-open system, and Closed system). As the oxygen content in the law fiber increases, the degradation reaction easily occurs. It leads to nearly stoichiometric composition of the degraded fiber. In this case, the partial pressure of CO gas at the surface region of the fiber becomes higher compared with that of the inside. This leads to lower reaction rate at the surface region of the fiber. Consequently, increase in the oxygen content of the raw Si-Al-C-O fiber led to decrease in the SiC crystalline size at the surface region. This means getting smooth surface. That is to say, higher oxygen content of the raw Si-Al-C-O fiber and closed system cause relatively higher CO content in the reactor during the degradation reaction, and then the consequent higher CO partial pressure at the surface region reduces the degradation reaction, that results in getting smooth surface.

Changes in actual surface roughness of the obtained SiC-polycrystalline fibers, which were synthesized by heat-treatment at 1900°C in argon gas atmosphere using different raw fibers and different vessels, are shown in Fig.6. As can be seen from this figure (Fig.6), the surface roughness could be controlled by change in both oxygen content of the raw Si-Al-C-O fiber and the reaction vessel. In this case, for obtaining A1, B1, C1 in Fig.6, the raw fiber composed of Si1Al0.01C1.5O0.4 was used for the synthesis. And, for obtaining A2, B2, C2 and for obtaining A3, B3, C3, the raw fibers composed of Si1Al0.01C1.5O0.45 and Si1Al0.01C1.5O0.5 were used, respectively. As can be seen from Fig.6, we could control the surface roughness from 67.99nm (maximum value) to 9.05nm (minimum value).

Figure 6 Change in the surface roughness of the SiC-polycrystalline fiber synthesized from different raw Si-Al-C-O fiber with different oxygen content at 1900°C in argon atmosphere. For obtaining A1, B1, C1, the raw fiber composed of Si1Al0.01C1.5O0.4 was used. For obtaining A2, B2, C2 and for obtaining A3, B3, C3, the raw fibers composed of Si1Al0.01C1.5O0.45 and Si1Al0.01C1.5O0.5 were used, respectively.

In this research, we used degradation process of the amorphous raw fiber (Si-Al-C-O fiber) accompanied by a release of CO gas and the subsequent sintering process, and showed the controllable SiC crystalline size constructing the obtained SiC-polycrystalline fiber by changing the CO gas partial pressure in the reactor. In consequence, we could control the surface roughness of the SiC-polycrystalline fiber using CO gas released from the raw fiber. However, this means that an intentional change in CO gas partial pressure in the reaction vessel can lead to preferable crystalline structure.

Fig.7 shows an improvement result regarding the surface roughness by achieving the adjustment of the CO gas partial pressure in the reactor during the degradation reaction. In this case, using Le Chatelier’s principle, we accelerated the following reaction (SiO2+3C=SiC+2CO) to consequently increase the CO gas partial pressure at the surface region of the fiber at initial degradation process. Regarding this degradation condition, we reported the detailed content in the previous paper [6].

Figure 7 Improvement of the surface roughness by adjustment of CO content in the reactor

CONCLUSIONS

We clarified the relationship between the heat-treatment condition and the surface roughness of the obtained SiC-polycrystalline fiber, using three different raw fibers (Elementary ratio: Si1Al0.01C1.5O0.4˜0.5) and three types of carbon vessel (Open system, Partially closed system, and Closed system). With increase in the oxygen content in the raw fiber, the degradation during the heat-treatment process easily proceeded accompanied by a release of relatively high concentration of CO gas. When we used the raw fiber composed of Si1Al0.01C1.5O0.5 and closed system, much smoother surface of the obtained SiC-polycrystalline fiber could be achieved. In this case, the degradation reactions (SiO+2C=SiC+CO and SiO2+3C=SiC+2CO) at the first stage in the inside of each filament became faster, and then the CO partial pressure at the surface region of each filament was found to be increased. In consequence, according to Le Chatelier’s principle, the surface degradation reaction and grain growth of formed SiC crystals would be considered to become slower.

ACKNOWLEDGMENT

This study was funded by a Grant from NEDO (New Energy and Industrial Technology Development Organization) via Ube Industries, Ltd. We gratefully acknowledge this financial support.

REFERENCES

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M.Takeda, A.Urano, J.Sakamoto, and Y.Imai, Microstructure and oxidative degradation behavior of silicon carbide fiber Hi-Nicalon type S, Journal of Nuclear Materials, 258-263 (1998) 1594-1599.

T.Ishikawa, Advances in Inorganic Fibers, Advanced Polymer Science (Springer-Vrlag Berlin Heidelberg) 178 (2005) 109-144.

J.J.Sha, T.Nozawa, J.S.Park, Y.Katoh, and A.Kohyama, Effect of heat treatment on the tensile strength and creep resistance of advanced SiC fibers, Journal of Nuclear Materials, 329-333 (2004) 592-596.

K.Itatani, K.Hattori, D.Harima, M.Aizawa, and I.Okada, Mechanical and thermal properties of silicon-carbide composites fabricated with short Tyranno Si-Zr-C-O fiber, Journal of Materials Science, 36 (2001) 3679-3686.

H.Oda and T.Ishikawa, Microstructure and mechanical properties of SiC-polycrystalline fiber and new defect-controlling process, International Journal of Applied Ceramic Technology, 14 (2017) 1031-1040.

T.Ishikawa and H.Oda, Defect control of SiC polycrystalline fiber synthesized from polyaluminocarbosilane, Journal of European Ceramic Society, 35 (2016) 3657-3662.

T.Ishikawa and H.Oda, Structural control aiming for high-performance SiC polycrystalline fiber, Journal of the Korean Ceramic Society, 53(6) (2016) 615-621.

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C.Sauder, A.Brusson, and J.Lamon, Influence of interface characteristics on the mechanical properties of Hi-Nicalon type-S or Tyranno-SA3 fiber-reinforced SiC/SiC minicomposites, International Journal of Applied Ceramic Technology, 7(3) (2010) 291-303.

THE INFLUENCE OF CASTING-CALENDERING PROCESS ON THE MICROSTRUCTURE OF PURE Al2O3 CERAMIC SUBSTRATE

S. X. Wang, H. F. Lan, W. J. Wang, Y. J. Huang and S. J. Li

College of Engineering, Shantou University

Shantou, Guangdong, 515063, China

ABSTRACT

The forming process of ceramic green tape has heavy influence on the performance of ceramic substrate. Aimed on the application for electric package, a new technology of casting-calendering process for preparing pure Al2O3 substrate was proposed in this paper. Micro and nano multi-scale mixed ceramic powders were used as raw material. The ceramic green tape with micro-nano hierarchical structure was prepared by casting-calendering process. Finally, the pure Al2O3 ceramic substrates were sintered under 1300 to 1600°C, and the microstructure of the ceramic substrates prepared by casting-calendering process and casting process was compared. Experimental results showed that the pure Al2O3 ceramic substrate sintered at 1500°C with no additives exhibits good densification degree of 92.5%. The thermal conductivity of the pure Al2O3 ceramic substrate prepared by casting-calendering process and sintered at 1500°C was 17.5W/(m·K), which was much higher than that by tape casting process under the same sintering condition.

INTRODUCTION

Due to the advantages of high thermal conductivity, good insulation, compatible thermal expansion coefficient with the chip, ceramic substrate has been widely used in electronic components, integrated circuits and electronic packaging fields. In recent years, with the rapid development of electronic packaging technology, packaging devices is becoming smaller and smaller and the packing density is becoming higher and higher. Thus, the heat density in the integrated circuit device is increasing. If the heat in the electronic device couldn't be dissipated quickly, it will not only reduce the reliability and performance of electronic devices, but also cause operation failure. According to the statistics1, more than half of the failure of electronic device was caused by heat. Overheating has become the bottleneck that restricts the development of electronic device technology. Improving the thermal performance of the substrate has become the focus of the electronic packaging field².

Traditionally, the electronic packaging substrate mainly includes plastic substrate, metal substrate and ceramic substrate. Plastic substrate is the most commonly used substrate at present, but the thermal conductivity and reliability of which is too low. Metal substrate has high thermal conductivity, but its CET is several times higher than the chip, which will produce high thermal stress between the metal substrate and the chip during operation. As a result, it is easy to cause the chip peeled off or fracture failure³. Ceramic substrate has high thermal conductivity, compatible CET with chips and good insulation, which is a preeminent packaging substrate material with excellent comprehensive performance. However, ceramic materials are usually difficult to melt because they are compounds that have a combination of ionic and covalent bonds, thus the sintering temperature is usually very high. Besides the production process is complex, and the production equipment is expensive⁴. At present, the cost of traditional high performance ceramic substrate by hot press process is very high. The thermal conductivity of LTCC (Low temperature co-fired ceramics) substrate decreased sharply due to the great amount of glass additive added⁵. For electronic packaging substrates, it is always expected to have high thermal conductivity, low dielectric constant and low dielectric loss.

Al2O3 ceramic is one of the most commonly used substrate materials in the integrated circuits packaging field. High purity Al2O3 ceramic substrate has good insulative and dielectric properties. The Al2O3 ceramic is regarded as the most competitive materials in the low-cost ceramic substrate field⁶. Hot-pressing is the main process for preparing high density Al2O3 substrate. Though much lower than that of the pure AlN substrate, the thermal conductivity of pure Al2O3 substrate is still up to 25W/(m·K), which could meet the requirements of electronic packaging substrate⁷. However, the high pure Al2O3 substrate requires high sintering temperature and high pressure equipment, which is difficult to produce in large scale by the tunnel kiln and roller kiln. Though lower than that of AlN substrate, the production cost of Al2O3 substrate is still very high. It is difficult to widely apply in high power integrated circuit and LED lighting field.

In order to reduce the manufacturing cost of Al2O3 substrate, many scholars have done lots of research on low temperature sintered substrate field in recent years. Liu Ming8 prepared the borosilicate glass/Al2O3 composites with the thermal conductivity of 2.89W/(m·K), the dielectric constant of 7.82, the dielectric loss of 0.53 × 10−3 (10 MHz) by low-temperature sintering process at 875°C. Chen et al9 prepared alumina ceramic composites using SiO2−B2O3−CaO−MgO glass additives with the dielectric constant of 7.3, the dielectric loss of 1.15×10−3 at 875°C, its thermal conductivity is still only 3.56 W/(m·K).

The Al2O3 ceramics sintered in low temperature usually added a large amount of glass phases in order to improve the relative density. The commercialized alumina ceramic substrates by low-temperature co-fired usually contained above 50vol% glass¹⁰. Because of the barrier effect of glass phase, the thermal conductivity of alumina ceramic substrate by low-temperature co-fired is generally only 2-4W/(m·K) even at high density¹¹. Induja et al12 developed Al2O3/BBSZ composites by adding glass of 35Bi2O3: 32ZnO: 27B2O3: 6SiO2, its dielectric constant is 11.3, its dielectric loss is 0.001 (1 MHz), and its thermal conductivity is 7.2W/(m·K). But it is still far lower than that of pure Al2O3 ceramic sintered at high temperature (about 25W/(m·K). Therefore, it still needs to make great efforts for preparing alumina ceramic substrate with high thermal conductivity under low sintering temperature.

Aimed on application for high power electric package, a new technology of casting-calendering process for preparing pure Al2O3 substrate was proposed in this research. Micro and nano multi-scale mixed ceramic powder was used as raw material. The ceramic green tape with micro-nano hierarchical structure was prepared by casting-calendering process. Finally, the pure Al2O3 substrates were sintered at 1300 to 1600°C.

EXPERIMENTAL

Experimental Materials

The main raw materials of the experiments are listed as following: micro scale α-Al2O3 powder (D50=5.68 μm, 99.9%, Jiankun chemical Co. Ltd., Hefei, China), nano scale α-Al2O3 powder (D50=50nm, 99.9%, Deke Daojin Co. Ltd., Beijing, China), anhydrous ethanol (AR, Tianjin Damao Chemical Co. Ltd., Tianjin, China), butanone (AR, Tianjin Damao Chemical Co. Ltd., Tianjin, China), polyvinyl butyral (PVB, AR, Shandong Haoyao New Material Co. Ltd., Shandong, China), three butyl phosphate (AR, Tianjin Damao Chemical Co. Ltd., Tianjin, China), dibutyl phthalate (DBP, AR, Guangzhou Chemical Co. Ltd., Guangzhou, China), polyethylene glycol (AR, Guangzhou Chemical Co. Ltd., Guangzhou, China).

Preparation of Samples

In this experiment, ceramic green tape was prepared