OARS - Open Access Rewards System
DOI : 10.2240/azojomo0286

Sintering of Silicon Nitride Nano-Powders Prepared by High-Energy Mechanical Milling

Toshiyuki Nishimura, Xin Xu, Naoto Hirosaki, Yoshinobu Yamamoto and Hidehiko Tanaka

Copyright AD-TECH; licensee AZoM.com Pty Ltd.
This is an AZo Open Access Rewards System (AZo-OARS) article distributed under the terms of the AZo-OARS https://www.azom.com/oars.asp which permits unrestricted use provided the original work is properly cited but is limited to non-commercial distribution and reproduction.
AZojomo (ISSN 1833-122X) Volume 6 November 2010

Topics Covered

Abstract
Keywords
Introduction
Experimental Procedure
Results and Discussion
     Effect of Metallic Aluminum as Grinding Additive
     Effect of Sintering Additives
Conclusions
Acknowledgements
References
Contact Details

Abstract

Silicon nitride powder with sintering additives was ground by high-energy milling to produce nanometer sized powder. This powder was densified by pulsed electric current sintering (SPS) for short time to prevent grain growth. Addition of metallic aluminum as grinding additive improved effciency of high-energy milling, but densification of powder with higher amount of Al started at higher temperatures and did not completed at the same sintering temperature that a powder with low amount of Al completely densified. Densification of powder with Y2O3-MgO started at lower temperatures than that of powder with Y2O3-Al2O3 and with Y2O3.

Keywords

Silicon Nitride, High Energy Milling, SPS, Grinding Additives, Sintering Additives

Introduction

Improvement of mechanical properties, e.g., strength, superplasticity, is expected by reducing grain size of ceramics, from micrometer scale to nanometer scale. A silicon nitride is a candidate for structural use, because it has high strength even at high temperature and high toughness. Fine-grained ceramics show superplasticity [1] and superplasticity of silicon nitride ceramics has been investigated for both scientific interest and technological application [2-7]. The superplasticity can be applied to plastic forming in complex shape. Raw silicon nitride powders were prepared by vapor phase reaction [2], attrition milling of commercial sub-micrometer powder [3-5] and removing large particles from sub-micrometer powder [6,7]. Grain size of ceramics from these powders was smaller than 500 nm. Grain size of hot-pressed silicon nitride ceramic from fine-β-powder was about 200 nm and the lowest superplastic temperature of the ceramic was 1450°C [6,7]. Superplastic deformation of a sialon ceramic from attrition milled powder was also reported at 1450°C [5]. Further lowering the superplastic temperature is needed, in particularfor, for industrial applications. Nano-sized raw powder is necessary for fabrication of nano-ceramics. We fabricated silicon nitride nano-powder from sub-micrometer powder by high-energy milling, and nano-ceramics were fabricated by rapid sintering of the nano-powder [8]. Effect of grinding additive in high-energy milling was reported [9]. In this paper, effect of metallic aluminum as a grinding additive on high-energy milling and sintering is investigated. Recently, we measured superplasticity of ceramics from high-energy milled powder with variety of sintering additives [10]. Effect of sintering additives on sintering is also investigated in conjunction with the superplasticity.

Experimental Procedure

Commercial sub-micrometer silicon nitride powder (SN-E10, Ube Industries Ltd., Tokyo, Japan) was used to investigate the effect of metallic aluminum addition. The powder was mixed in ethanol with 5 mol% of Y2O3 (purity: 99.9%, Shin-Etsu-Chemical Co. Ltd., Tokyo, Japan) and 2 mol% Al2O3 (AKP-20, Sumitomo Chemical Co. Ltd., Tokyo, Japan) by conventional planetary ball mill. The mixed powder was milled in nitrogen atmosphere with metallic aluminum particle (purity: 99.9%, average grain size: about 20µm, Kojundo Chemical Lab. Co., Ltd., Saitama, Japan) by high-energy mill (PM1200, Seishin Enterprise Co. Ltd., Tokyo, Japan). The milling time was 4 h. The milled powders were then sintered by spark plasma sintering machine (SPS-1030, Sumitomo Coal Mining Co. Ltd., Tokyo, Japan) for 5 min in nitrogen atmosphere. Heating rate was about 300 K/min and holding time was 5 min. for preventing grain growth.

β-silicon nitride powder (NP500 grade, Denki Kagaku Kogyo Co., Tokyo, Japan) was used to investigate the effect of sintering additives. The sintering additives selected are 7.85 mass% Y2O3 and 1.42 mass% Al2O3 (as denoted as Y-A), 5 mass% Y2O3 and 2 mass% MgO (high-purity grade, Wako Pure Chemical Industries, Osaka, Japan) (as denoted as Y-M), and 10 mass% Y2O3 (as denoted as Y). All starting powder mixtures were high-energy milled for 4 h with a rotation speed of 475 rpm. Y-M, Y-A and Y were sintered at 1600°C for 2 min, 1600°C for 5 min and 1620°C for 5 min, respectively.

Results and Discussion

Effect of Metallic Aluminum as Grinding Additive

High-energy milled powders composed of a-silicon nitride, yttrium oxide and aluminum oxide with metallic aluminum were analyzed with powder XRD. Peak height of silicon nitride decreased and peaks of yttrium oxide and metallic aluminum diminished. The peak height of β-silicon nitride decreased with increasing of ball/powder ratio in high energy milling [9], which suggested that milling effect can be represented by peak height reduction. Relative intensity of milled silicon nitride powder to intensity of powder before milling was calculated and relationship between the relative intensity and aluminum content are shown in Fig. 1. A relative intensity in Fig. 1 is an average value of the relative intensity of (201), (102), (210), (321) and (322) peaks of a-silicon nitride. The relative intensity decreases with increasing of aluminum content, which means that higher aluminum content is effective for grinding of silicon nitride.

Differential thermal analysis was conducted for a silicon nitride mixed powder with 5 mol% Y2O3 and 2 mol% Al2O3 (powder A), powder A with metallic aluminum (powder B) and milled powder B (powder C) (Fig. 2). An endotherm is observed at about 680°C in DTA curve of powder B, which may correspond to melt of metallic aluminum. The endotherm is not observed in DTA curve of powder C and a slight exotherm is observed at about 650°C. AlN was formed during sintering of powder C at 1400°C [9]. The metallic aluminum may react with nitrogen without melting in powder C.

Figure 1. Effect of aluminum additive content on high-energy milling of silicon nitride.

Figure 2. DTA result of high-energy milled silicon nitride powder

Figure 3 shows displacement and load change during sintering. Applied load is not constant in standard SPS system, even if load control is set at a constant value. At the beginning of heating, displacement decreases and load increases by thermal expansion of punches and then load decreases suddenly. After that, displacement increases quickly. The load decreases or keeps a constant value during the increase of displacement and when the increase of displacement almost stops, the load begins to increase. These are typical change of displacement and load in SPS of silicon nitride. The load starts decreasing at about 1100 °C for a milled powder with 5.3 mass% Al. The displacement starts increasing at about 1200 °C and the increase almost completes at 1600 °C. The load decreasing temperature is about 1230°C for milled powder with 10.9 mass% Al. The displacement starts increasing at about 1350°C and still increases during holding step at 1600°C. In the case of milled powder with 22.0 mass% Al, the load starts decreasing at about 1330°C, the displacement starts increasing at about 1380°C and the final displacement is much lower than that of other powders. This result suggests that lower aluminum content is favorable for low temperature sintering, when fine-grained silicon nitride is fabricated.

Figure 3. Temperature, displacement and load in SPS of silicon nitride.

Figure 4 shows effect of sintering temperature on density of high-energy milled silicon nitride powder. Higher density is obtained from a powder with lower aluminum content at 1500°C. Density from a powder with 5.3 mass % Al is almost same as that from a powder with 10.9 mass% Al at 1550 and 1600°C. Nano-ceramic was obtained just from the powder with 5.3 mass% Al by sintering at 1550°C [9]. Considering the Figs. 1, 3 and 4, 5.3 mass% Al is most effective in both milling and sintering, and 1550°C is enough for densification.

Figure 4. Effect of aluminum addition on density of silicon nitride ceramics from high-energy milled powder.

Effect of Sintering Additives

Silicon nitride nano-ceramics were fabricated from powders Y-A, Y-M and Y and superplastic deformation was measured. The silicon nitride from powder Y-M deformed at the highest strain rate at a same temperature and under a same stress condition [10]. In this section, densification of the three silicon nitride ceramics is compared.

The powder Y-A can be densified at 1600°C and density of ceramic from powder Y-M is slightly lower than that from Y-A, probably because sintering time of Y-M is 2 min. Density of Y is lower than that of Y-A and Y-M (Fig. 5). Y2O3 added silicon nitride has been investigated as heat resistant material and dense material was obtained by sintering at 1750°C or higher for a several hours [11-13]. Dense nano-ceramic was obtained from powder Y by sintering at 1620°C for 5 min in this work, which is a typical example of effectiveness of the high-energy milling for improving sinterability of silicon nitride.

Figure 5. Effect of sintering temperature on density of silicon nitride ceramics.

Densification behavior during sintering of the three materials from powder Y-A, Y-M and Y is precisely investigated to clarify the difference in compressive deformation. Typical changes in displacement and load were explained in Fig. 3. A densification starting temperature (DST) is defined as a temperature at which the load decreases or the displacement increases sharply. Figure 6 shows the DST obtained from load and displacement. The both DSTs from load and displacement of Y-M are lower than those of Y-A and Y. The DST from load of Y-A is same as that of Y and that from displacement of Y-A is rather higher than that of Y. This result suggests that powder Y-M forms liquid phase during sintering at lower temperature than Y-A and Y, and also melting temperature of grain boundary in Y-M is lower than that of Y-A and Y. Grain boundary sliding is dominant in superplastic deformation at high stress region [14], therefore, viscosity of grain boundary phase in silicon nitride Y-M might be the lowest within the three. This result supports that viscosity of liquid phase in the three ceramics is Y > Y-A > Y-M at 1500°C [10].

Figure 6. Effect of sintering additives on densification starting temperature of silicon nitride powder.

Conclusions

Effect of grinding additive and sintering additives on high-energy milling and sintering was investigated. More metallic aluminum was better for milling, but less aluminum is favorable for low temperature sintering. Precise control is necessary to fabricate dense nano-ceramics. Selection of sintering additives is important for lowering superplastic temperature of silicon nitride nano-ceramics. The sintering additives affect liquid phase formation in sintering and also viscosity of grain boundary phase in superplastic deformation.

Acknowledgements

This work was partly supported by Grant-in-Aids for Scientific Research No. 17360345 (Section (B)) from the Ministry of Education, Culture, Sports, Science and Technology of Japan.

References

1. J. G. Wang and R. Raj, J. Am. Ceram. Soc., 67 (1984) 399-409.
2. F. Wakai, Y. Kodama, S. Sakaguchi, N. Murayama, K. Izaki and K. Niihara, Nature, 344,.421-423 (1990)
3. I. Chen and S. Hwang, J. Am. Ceram. Soc., 75 (1992) 1073-1079.
4. P. Burger, R. Duclos and J. Crampon, J. Am. Ceram. Soc., 80 (1997) 879-885.
5. A. Rosenflanz and I. Chen, J. Am. Ceram. Soc., 80 (1997) 1341-52.
6. M. Mitomo, H. Hirotsuru, H. Suematsu and T. Nishimura, J. Am. Ceram. Soc., 78 (1995) 211-214.
7. G. Zhan, M. Mitomo, T. Nishimura, R. Xie, T. Sakuma and Y. Ikuhara, J. Am. Ceram. Soc., 83 (2000) 841-847.
8. X. Xu, T. Nishimura, N. Hirosaki, R. Xie, Y. Zhu, Y. Yamamoto, and H. Tanaka, J. Am. Ceram. Soc. 88 (2005) 934-937.
9. T. Nishimura, X. Xu, N. Hirosaki, K. Kimoto, Y. Yamamoto and H. Tanaka, Key Eng. Mater. 287 (2005) 156-159.
10. X. Xu, T. Nishimura, N. Hirosaki, R. Xie, Y. Yamamoto and H. Tanaka, J. Am. Ceram. Soc. (in press).
11. G. E. Gazza, J. Am. Ceram. Soc., 56, 662 (1973).
12. A. Tsuge, K. Nishida and M. Komatsu, J. Am. Ceram. Soc., 58 (1975) 323-326.
13. J. T. Smith and C. L. Quackenbush, Ceram. Bull., 59 (1980) 529-532, 537.
14. X. Xu, T. Nishimura, N. Hirosaki, R. Xie, Y. Yamamoto and H. Tanaka, Acta Mater. 54 (2006) 255-262.

Contact Details

Toshiyuki NISHIMURA, Xin XU, Naoto HIROSAKI, Yoshinobu YAMAMOTO and Hidehiko TANAKA
Advanced Materials Laboratory, National Institute for Materials Science
1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan

This paper was also published in print form in "Advances in Technology of Materials and Materials Processing", 10[2] (2008) 89-94.

Tell Us What You Think

Do you have a review, update or anything you would like to add to this article?

Leave your feedback
Submit