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Shape memory characteristics of TiNi casting alloys made by using self-propagating high-temperature synthesis
Materials Science and Engineering A 438–440 (2006) 675–678
Shape memory characteristics of TiNi casting alloys made by using self-propagating high-temperature synthesis K. Kitamura a,∗ , T. Kuchida b , T. Inaba b , M. Tokuda b , Y. Yoshimi c a
Nagano National College of Technology, Department of Mechanical Engineering, 716 Tokuma, Nagano, Nganao, Japan b Mie University, Tsu, Japan c Yoshimi Inc., Obu, Japan Received 8 May 2005; received in revised form 25 January 2006; accepted 17 February 2006
Abstract Self-propagating high-temperature synthesis (SHS) is a new method of making an ingot of TiNi. This method shows little gravity segregation. The purpose of this study is to investigate the difference of the shape memory characteristics of the TiNi casting alloys made from a SHS ingot and from a conventional melt cast ingot. The samples used in this study were rods made by centrifugal casting. Differential scanning calorimetry (DSC), X-ray diffraction, and tensile test were used to examine the shape memory characteristics on the samples. The heat treatment conditions were 773 K–1.8 ks and 1073 K–3.6 ks, respectively. The DSC samples were both ends (the top and bottom area) of the rod samples. The results of the XRD measurements showed that TiNi phase was obtained in all the samples. In contrast, the result of the DSC test showed that more gravity segregation effect happened in the melt cast sample than in the SHS sample. As the conclusion of this study, gravity segregation had little effect in the SHS ingot sample. © 2006 Elsevier B.V. All rights reserved. Keywords: TiNi; Self-propagating high-temperature synthesis; Shape memory alloy; Casting alloy
1. Introduction Near-equiatomic TiNi alloys are used in various fields because of their excellent shape memory characteristics, superelasticity, good corrosion resistance and high biocompatibility. These alloys are applied to sensors, actuators and medical devices because of these features. Generally, this alloy is made through a melting method such as vacuum arc melting, which requires multiple re-melts to achieve sufficient homogeneity and it has gravity segregation regionally. On the other hand, selfpropagating high-temperature synthesis (SHS) is a synthesis approach without gravity segregation. Ti and Ni powders were used in the SHS method. Powder metallurgy (PM) creates the chemical composition of the alloy as designed because of the dependence of the temperature range of martensite ↔ austenite transformation on nickel content [1–4]. In a recent study, TiNi alloy was synthesized using many PM techniques such as SHS [5–16], plasma sintering [17], element powder metallurgy
∗
Corresponding author. E-mail address: [email protected] (K. Kitamura).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.02.172
(EPM) [18], hot isostatic pressing (HIP) [19] and pre-alloy powder metallurgy [20]. Currently, porous structure TiNi alloy using SHS [10–16] is making its way in bone implantation because it can combine with the new bone tissue in the body. Researchers have been doing studies about the shape memory characteristics of SHS ingot. There is also a thesis about the cast of the alloy in a dental research [21–23]. The purpose of this paper is to prepare cast TiNi alloy from melt method and SHS ingots and to study the effect of gravity segregation on the samples. 2. Experimental Ti–50.8 at.% Ni ingots were prepared by melting and SHS methods. SHS is the synthesis of TiNi alloys in a wave of the combustion that propagates over starting reactive mixture of Ti and Ni powder owing to layer-by-layer heat transfer [5,6]. An SHS ingot made by KCM Corporation was used in this research. All the samples were rod shape alloy made by centrifugal casting method. The size of all the samples was 2 mm in diameter and 45 mm in length. The samples were annealed at 773 K for 1.8 ks and 1073 K for 3.6 ks followed by water quenching. The surfaces of the sample were finally polished by emery paper. The
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K. Kitamura et al. / Materials Science and Engineering A 438–440 (2006) 675–678
size of the samples used for XRD and DSC measurements was 3.0 mm × 1 mm. The tensile samples were used without cutting off cast samples. The phase constituent was determined by XRD analysis of Philips, X’Pert-MPD PW3050 at room temperature. The transformation behavior during cooling and heating was investigated using a differential scanning calorimeter, Shimazu DSC-50. The cooling and heating rates were 10 K/min. Mechanical properties of the shape memory alloys were measured using a tensile machine, Shimazu AG-IS. The distortion gauge was used for the tensile test of the sample. 3. Results and discussions Fig. 1 shows the XRD patterns of the melt cast samples heat-treated at 773 K–1.8 ks and 873 K–3.6 ks. The lower curve shows the sample heat-treated at 773 K–1.8 ks, and the upper curve shows the sample heat-treated at 873 K–3.6 ks. In the TiNi alloy both parent phase and martensite phase appeared. It indicated that two precipitates, Ti2 Ni and TiNi3 , were formed in the sample. It also indicated that the heat treatment temperature dependency did not appear in these samples. Fig. 2 shows the XRD patterns of the SHS method samples heat-treated at 773 K–1.8 ks and 873 K–3.6 ks. In the TiNi alloy both parent phase and martensite phase appeared. It indicated that two precipitates, Ti2 Ni and TiNi3 , were formed in the sample. The martensite peak was weaker than that of the melt sample. This result shows that the volume fraction of the martensite phase is low. The heat-treatment temperature dependency did not appear in these samples. The XRD peaks from the precipitates were weaker than the melt method sample. Fig. 3 shows the DSC cooling curves of melt method sample heat-treated at 773 K–1.8 ks. The solid and dashed curve of upper sample and lower sample show a three-stage transformation from the parent (B2) to M (martensite)-phase upon cooling. In comparison with lower sample, upper sample is low on the second martensite transformation peak. The difference of these DSC curves is primarily due to Ni content that has a stronger effect on gravity segregation.
Fig. 1. X-ray diffraction patterns of melting method specimen at 773 K–1.8 ks and 873 K–3.6 ks heat treatment.
Fig. 2. X-ray diffraction patterns of SHS specimen after 773 K–1.8 ks and 873 K–3.6 ks heat treatment.
Fig. 4 shows the DSC heating curves for melt method sample heat-treated at 773 K–1.8 ks. The solid and dashed curve of upper sample and lower sample shows a three-stage reversetransformation from the M-phase to B2 upon heating. In comparison with lower sample, upper sample is low on the third reverse-martensite transformation peak. Fig. 5 shows the DSC cooling curves for SHS sample heattreated at 773 K–1.8 ks. The solid curves of the upper sample show a three-stage transformation from the B2 to M-phase upon cooling. The second martensite transformation peak is very small. The dashed curves of the lower sample show only twostage transformation from the B2 to M-phase. The influence of gravity segregation is little or none. Fig. 6 shows the DSC heating curves for SHS sample heattreated at 773 K–1.8 ks. The solid and dashed curve shows a two-stage reverse-transformation from the M-phase to B2 upon heating. In comparison with lower sample, upper sample is low on the first reverse-martensite transformation peak. Fig. 7 shows the correspondence of the transformation and the reverse-transformation peak that appears to DSC curve. Line (a) in Fig. 7 shows transformation peak, and Line (a ) in Fig. 7
Fig. 3. DSC curves of melt method specimens upon cooling.
K. Kitamura et al. / Materials Science and Engineering A 438–440 (2006) 675–678
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Fig. 4. DSC curves of melt method specimens upon heating. Fig. 7. Correspondence of the transformation peaks apeear to DSC.
shows reverse-transformation peaks of DSC curves. The transformation peak and the reverse-transformation peak correspond. Lines (b) and (b ) in Fig. 7 show the result of reversely transformed peak after the second martensite transformation ends. The second martensite transformation peak was clarified to the second reverse-martensite transformation peak from this result. Lines (c) and (c ) in Fig. 7 show the result of the cooling and heating between 200 and 360 K. As a result, the third martensite transformation peak was clarified to the first reverse-martensite transformation peak. The correspondence of each transformation peak and the reverse-transformation peak was clarified from these results. Solid lines in Fig. 8 show the DSC cooling and heating curves for melt method sample heat-treated at 873 K–3.6 ks, and dashed Fig. 5. DSC curves of SHS specimens upon cooling.
Fig. 6. DSC curves of SHS specimens upon heating.
Fig. 8. DSC curves of melting method specimen and SHS specimen upon cooling and heating.
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K. Kitamura et al. / Materials Science and Engineering A 438–440 (2006) 675–678
E = 34 GPa and that the induced martensitic phase transformation stress at σ = 225 MPa. An excellent shape memory characteristic appeared at the SHS sample as well as melt method sample. 4. Conclusions
Fig. 9. Stress–strain curves of melt method specimen at room temperature.
Using both the SHS TiNi ingot and melting TiNi ingot the rod samples were successfully made by centrifugal casting method, and all the rod samples had shape memory characteristics. However, the result of the DSC measurement showed that the sample made by melting method had more gravity segregation effect than the sample made by the SHS method. If the SHS sample is used for a centrifugal casting, an excellent shape memory characteristic with low gravity segregation of a casting material can be produced. References [1] [2] [3] [4] [5] [6] [7]
Fig. 10. Stress–strain curves of SHS specimen at room temperature.
lines in Fig. 8 show SHS sample heat-treated at the same condition. In the melt method sample multi stage transformation and reverse-transformation peaks appeared. In the SHS sample only one large transformation and one large reverse-transformation peak show on cooling and heating. This result shows that the internal structure is not fully homogenized through the solution treatment process. Fig. 9 shows the stress–strain relation for tension test of melt method sample at room temperature. It is observed that in the first stage of the loading process this material shows the apparent elastic deformation martensitic phase with Young’s modulus E = 37 GPa and the induce martensitic phase transformation stress at σ = 215 MPa. Fig. 10 shows the stress–strain relation for tension test of melt method sample at room temperature. It is observed that in the first stage of the loading process this material shows the apparent elastic deformation martensitic phase with Young’s modulus
[8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
K. Otsuka, X. Ren, Intermetallics 7 (1999) 511–528. H. Funakubo, Shape Memory Alloys, Gorden & Breach, New York, 1987. D. Starosvetsky, I. Gotman, Biomaterials 22 (2001) 1853–1859. A. Kapanen, J. Ryh¨anen, A. Danilov, J. Tuukkanen, Biomaterials 22 (2001) 2475–2480. Y. Kaieda, M. Otaguchi, N. Oguro, T. Oie, S. Shite, M. Hatakeyama, MRS Int. Mtg. Adv. Mats. 9 (1989) 623–628. Y. Kaieda, Adv. Synth. Process. Composites Adv. Ceram. 56 (1995) 28–37. C.L. Chu, B. Li, S.D. Wang, S.G. Zhang, X.X. Yang, Z.D. Yin, Trans. Nonferr. Met. Soc. China 7 (1997) 84–87. N. Bertolino, M. Monagheddu, A. Tacca, P. Giuliani, C. Zanotti, U. Anselmi Tamburini, Intermetallics 11 (2003) 41–49. C.L. Yeh, W.Y. Sung, J. Alloys Compd. 376 (2004) 79–88. B.Y. Li, L.J. Rong, Y.Y. Li, V.E. Gjunter, Acta Mater. 48 (2000) 3895–3904. B.-Y. Li, L.-J. Rong, Y.Y. Li, V.E. Gjunter, Intermetallics 8 (2000) 881–884. C.L. Chu, C.Y. Chung, P.H. Lin, Mater. Sci. Eng. A 392 (2005) 106–111. S.L. Zhu, X.J. Yang, F. Hu, S.H. Deng, Z.D. Cui, Mater. Lett. 58 (2004) 2369–2373. B.-Y. Li, L.-J. Rong, Y.-Y. Li, Scripta Mater. 44 (2001) 823–827. C.Y. Chung, C.L. Chu, S.D. Wang, Mater. Lett. 58 (2004) 1683–1686. A. Biswas, Acta Mater. 53 (2005) 1415–1425. M.D. Meneese, D.C. Lagoudas, T.C. Pollock, Mater. Sci. Eng. A 280 (2000) 334–348. S.M. Green, D.M. Grant, N.R. Kelly, Powder Metall. 40 (1) (1997) 43–47. M.D. McNeese, D.C. Material, T.C. Pollock, Mater. Sci. Eng. A 280 (2000) 334–348. M. Igharo, J.V. Wood, Powder Metall. 29 (1986) 37. Y. Soga, H. Doi, E. Kobayashi, T. Yoneyama, H. Hamanaka, J. Jpn. Soc. Dent. Mater. Devices 16 (1997) 150. H. Doi, E. Kobayashi, T. Yoneyama, H. Hamanaka, J. Jpn. Soc. Dent. Mater. Devices 19 (2000) 62. J. Kasuga, T. Yoneyama, H. Doi, E. Kobayashi, O. Fukushima, H. Hamanaka, J. Jpn. Soc. Dent. Mater. Devices 22 (2003) 139.
Shape memory characteristics of TiNi casting alloys made by using self-propagating high-temperature synthesis K. Kitamura a,∗ , T. Kuchida b , T. Inaba b , M. Tokuda b , Y. Yoshimi c a
Nagano National College of Technology, Department of Mechanical Engineering, 716 Tokuma, Nagano, Nganao, Japan b Mie University, Tsu, Japan c Yoshimi Inc., Obu, Japan Received 8 May 2005; received in revised form 25 January 2006; accepted 17 February 2006
Abstract Self-propagating high-temperature synthesis (SHS) is a new method of making an ingot of TiNi. This method shows little gravity segregation. The purpose of this study is to investigate the difference of the shape memory characteristics of the TiNi casting alloys made from a SHS ingot and from a conventional melt cast ingot. The samples used in this study were rods made by centrifugal casting. Differential scanning calorimetry (DSC), X-ray diffraction, and tensile test were used to examine the shape memory characteristics on the samples. The heat treatment conditions were 773 K–1.8 ks and 1073 K–3.6 ks, respectively. The DSC samples were both ends (the top and bottom area) of the rod samples. The results of the XRD measurements showed that TiNi phase was obtained in all the samples. In contrast, the result of the DSC test showed that more gravity segregation effect happened in the melt cast sample than in the SHS sample. As the conclusion of this study, gravity segregation had little effect in the SHS ingot sample. © 2006 Elsevier B.V. All rights reserved. Keywords: TiNi; Self-propagating high-temperature synthesis; Shape memory alloy; Casting alloy
1. Introduction Near-equiatomic TiNi alloys are used in various fields because of their excellent shape memory characteristics, superelasticity, good corrosion resistance and high biocompatibility. These alloys are applied to sensors, actuators and medical devices because of these features. Generally, this alloy is made through a melting method such as vacuum arc melting, which requires multiple re-melts to achieve sufficient homogeneity and it has gravity segregation regionally. On the other hand, selfpropagating high-temperature synthesis (SHS) is a synthesis approach without gravity segregation. Ti and Ni powders were used in the SHS method. Powder metallurgy (PM) creates the chemical composition of the alloy as designed because of the dependence of the temperature range of martensite ↔ austenite transformation on nickel content [1–4]. In a recent study, TiNi alloy was synthesized using many PM techniques such as SHS [5–16], plasma sintering [17], element powder metallurgy
∗
Corresponding author. E-mail address: [email protected] (K. Kitamura).
0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.02.172
(EPM) [18], hot isostatic pressing (HIP) [19] and pre-alloy powder metallurgy [20]. Currently, porous structure TiNi alloy using SHS [10–16] is making its way in bone implantation because it can combine with the new bone tissue in the body. Researchers have been doing studies about the shape memory characteristics of SHS ingot. There is also a thesis about the cast of the alloy in a dental research [21–23]. The purpose of this paper is to prepare cast TiNi alloy from melt method and SHS ingots and to study the effect of gravity segregation on the samples. 2. Experimental Ti–50.8 at.% Ni ingots were prepared by melting and SHS methods. SHS is the synthesis of TiNi alloys in a wave of the combustion that propagates over starting reactive mixture of Ti and Ni powder owing to layer-by-layer heat transfer [5,6]. An SHS ingot made by KCM Corporation was used in this research. All the samples were rod shape alloy made by centrifugal casting method. The size of all the samples was 2 mm in diameter and 45 mm in length. The samples were annealed at 773 K for 1.8 ks and 1073 K for 3.6 ks followed by water quenching. The surfaces of the sample were finally polished by emery paper. The
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K. Kitamura et al. / Materials Science and Engineering A 438–440 (2006) 675–678
size of the samples used for XRD and DSC measurements was 3.0 mm × 1 mm. The tensile samples were used without cutting off cast samples. The phase constituent was determined by XRD analysis of Philips, X’Pert-MPD PW3050 at room temperature. The transformation behavior during cooling and heating was investigated using a differential scanning calorimeter, Shimazu DSC-50. The cooling and heating rates were 10 K/min. Mechanical properties of the shape memory alloys were measured using a tensile machine, Shimazu AG-IS. The distortion gauge was used for the tensile test of the sample. 3. Results and discussions Fig. 1 shows the XRD patterns of the melt cast samples heat-treated at 773 K–1.8 ks and 873 K–3.6 ks. The lower curve shows the sample heat-treated at 773 K–1.8 ks, and the upper curve shows the sample heat-treated at 873 K–3.6 ks. In the TiNi alloy both parent phase and martensite phase appeared. It indicated that two precipitates, Ti2 Ni and TiNi3 , were formed in the sample. It also indicated that the heat treatment temperature dependency did not appear in these samples. Fig. 2 shows the XRD patterns of the SHS method samples heat-treated at 773 K–1.8 ks and 873 K–3.6 ks. In the TiNi alloy both parent phase and martensite phase appeared. It indicated that two precipitates, Ti2 Ni and TiNi3 , were formed in the sample. The martensite peak was weaker than that of the melt sample. This result shows that the volume fraction of the martensite phase is low. The heat-treatment temperature dependency did not appear in these samples. The XRD peaks from the precipitates were weaker than the melt method sample. Fig. 3 shows the DSC cooling curves of melt method sample heat-treated at 773 K–1.8 ks. The solid and dashed curve of upper sample and lower sample show a three-stage transformation from the parent (B2) to M (martensite)-phase upon cooling. In comparison with lower sample, upper sample is low on the second martensite transformation peak. The difference of these DSC curves is primarily due to Ni content that has a stronger effect on gravity segregation.
Fig. 1. X-ray diffraction patterns of melting method specimen at 773 K–1.8 ks and 873 K–3.6 ks heat treatment.
Fig. 2. X-ray diffraction patterns of SHS specimen after 773 K–1.8 ks and 873 K–3.6 ks heat treatment.
Fig. 4 shows the DSC heating curves for melt method sample heat-treated at 773 K–1.8 ks. The solid and dashed curve of upper sample and lower sample shows a three-stage reversetransformation from the M-phase to B2 upon heating. In comparison with lower sample, upper sample is low on the third reverse-martensite transformation peak. Fig. 5 shows the DSC cooling curves for SHS sample heattreated at 773 K–1.8 ks. The solid curves of the upper sample show a three-stage transformation from the B2 to M-phase upon cooling. The second martensite transformation peak is very small. The dashed curves of the lower sample show only twostage transformation from the B2 to M-phase. The influence of gravity segregation is little or none. Fig. 6 shows the DSC heating curves for SHS sample heattreated at 773 K–1.8 ks. The solid and dashed curve shows a two-stage reverse-transformation from the M-phase to B2 upon heating. In comparison with lower sample, upper sample is low on the first reverse-martensite transformation peak. Fig. 7 shows the correspondence of the transformation and the reverse-transformation peak that appears to DSC curve. Line (a) in Fig. 7 shows transformation peak, and Line (a ) in Fig. 7
Fig. 3. DSC curves of melt method specimens upon cooling.
K. Kitamura et al. / Materials Science and Engineering A 438–440 (2006) 675–678
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Fig. 4. DSC curves of melt method specimens upon heating. Fig. 7. Correspondence of the transformation peaks apeear to DSC.
shows reverse-transformation peaks of DSC curves. The transformation peak and the reverse-transformation peak correspond. Lines (b) and (b ) in Fig. 7 show the result of reversely transformed peak after the second martensite transformation ends. The second martensite transformation peak was clarified to the second reverse-martensite transformation peak from this result. Lines (c) and (c ) in Fig. 7 show the result of the cooling and heating between 200 and 360 K. As a result, the third martensite transformation peak was clarified to the first reverse-martensite transformation peak. The correspondence of each transformation peak and the reverse-transformation peak was clarified from these results. Solid lines in Fig. 8 show the DSC cooling and heating curves for melt method sample heat-treated at 873 K–3.6 ks, and dashed Fig. 5. DSC curves of SHS specimens upon cooling.
Fig. 6. DSC curves of SHS specimens upon heating.
Fig. 8. DSC curves of melting method specimen and SHS specimen upon cooling and heating.
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E = 34 GPa and that the induced martensitic phase transformation stress at σ = 225 MPa. An excellent shape memory characteristic appeared at the SHS sample as well as melt method sample. 4. Conclusions
Fig. 9. Stress–strain curves of melt method specimen at room temperature.
Using both the SHS TiNi ingot and melting TiNi ingot the rod samples were successfully made by centrifugal casting method, and all the rod samples had shape memory characteristics. However, the result of the DSC measurement showed that the sample made by melting method had more gravity segregation effect than the sample made by the SHS method. If the SHS sample is used for a centrifugal casting, an excellent shape memory characteristic with low gravity segregation of a casting material can be produced. References [1] [2] [3] [4] [5] [6] [7]
Fig. 10. Stress–strain curves of SHS specimen at room temperature.
lines in Fig. 8 show SHS sample heat-treated at the same condition. In the melt method sample multi stage transformation and reverse-transformation peaks appeared. In the SHS sample only one large transformation and one large reverse-transformation peak show on cooling and heating. This result shows that the internal structure is not fully homogenized through the solution treatment process. Fig. 9 shows the stress–strain relation for tension test of melt method sample at room temperature. It is observed that in the first stage of the loading process this material shows the apparent elastic deformation martensitic phase with Young’s modulus E = 37 GPa and the induce martensitic phase transformation stress at σ = 215 MPa. Fig. 10 shows the stress–strain relation for tension test of melt method sample at room temperature. It is observed that in the first stage of the loading process this material shows the apparent elastic deformation martensitic phase with Young’s modulus
[8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
K. Otsuka, X. Ren, Intermetallics 7 (1999) 511–528. H. Funakubo, Shape Memory Alloys, Gorden & Breach, New York, 1987. D. Starosvetsky, I. Gotman, Biomaterials 22 (2001) 1853–1859. A. Kapanen, J. Ryh¨anen, A. Danilov, J. Tuukkanen, Biomaterials 22 (2001) 2475–2480. Y. Kaieda, M. Otaguchi, N. Oguro, T. Oie, S. Shite, M. Hatakeyama, MRS Int. Mtg. Adv. Mats. 9 (1989) 623–628. Y. Kaieda, Adv. Synth. Process. Composites Adv. Ceram. 56 (1995) 28–37. C.L. Chu, B. Li, S.D. Wang, S.G. Zhang, X.X. Yang, Z.D. Yin, Trans. Nonferr. Met. Soc. China 7 (1997) 84–87. N. Bertolino, M. Monagheddu, A. Tacca, P. Giuliani, C. Zanotti, U. Anselmi Tamburini, Intermetallics 11 (2003) 41–49. C.L. Yeh, W.Y. Sung, J. Alloys Compd. 376 (2004) 79–88. B.Y. Li, L.J. Rong, Y.Y. Li, V.E. Gjunter, Acta Mater. 48 (2000) 3895–3904. B.-Y. Li, L.-J. Rong, Y.Y. Li, V.E. Gjunter, Intermetallics 8 (2000) 881–884. C.L. Chu, C.Y. Chung, P.H. Lin, Mater. Sci. Eng. A 392 (2005) 106–111. S.L. Zhu, X.J. Yang, F. Hu, S.H. Deng, Z.D. Cui, Mater. Lett. 58 (2004) 2369–2373. B.-Y. Li, L.-J. Rong, Y.-Y. Li, Scripta Mater. 44 (2001) 823–827. C.Y. Chung, C.L. Chu, S.D. Wang, Mater. Lett. 58 (2004) 1683–1686. A. Biswas, Acta Mater. 53 (2005) 1415–1425. M.D. Meneese, D.C. Lagoudas, T.C. Pollock, Mater. Sci. Eng. A 280 (2000) 334–348. S.M. Green, D.M. Grant, N.R. Kelly, Powder Metall. 40 (1) (1997) 43–47. M.D. McNeese, D.C. Material, T.C. Pollock, Mater. Sci. Eng. A 280 (2000) 334–348. M. Igharo, J.V. Wood, Powder Metall. 29 (1986) 37. Y. Soga, H. Doi, E. Kobayashi, T. Yoneyama, H. Hamanaka, J. Jpn. Soc. Dent. Mater. Devices 16 (1997) 150. H. Doi, E. Kobayashi, T. Yoneyama, H. Hamanaka, J. Jpn. Soc. Dent. Mater. Devices 19 (2000) 62. J. Kasuga, T. Yoneyama, H. Doi, E. Kobayashi, O. Fukushima, H. Hamanaka, J. Jpn. Soc. Dent. Mater. Devices 22 (2003) 139.