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Ti50Cu25Ni20Sn5 bulk metallic glass fabricated by powder consolidation
Materials Letters 61 (2007) 4591 – 4594 www.elsevier.com/locate/matlet
Ti50Cu25Ni20Sn5 bulk metallic glass fabricated by powder consolidation P.P. Choi a,⁎, J.S. Kim b , O.T.H. Nguyen b , Y.S. Kwon b a
b
Nano-materials Research Center, Korea Institute of Science and Technology, P.O. Box 131, Cheongryang, Seoul 130–650, Republic of Korea Research Center for Machine Parts and Materials Processing, School of Materials Science and Engineering, University of Ulsan, Namgu Mugeo 2-Dong, San 29, Ulsan 680–749, Republic of Korea Received 16 October 2006; accepted 24 February 2007 Available online 3 March 2007
Abstract A bulk metallic glass of Ti50Cu25Ni20Sn5 nominal composition was produced via a powder metallurgical route, namely the preparation of glassy powders by mechanical alloying followed by their consolidation by spark-plasma sintering. Samples were characterized with respect to their structure and thermal properties before and after sintering. A bulk glassy sample, having nearly full density and containing only a small fraction of the intermetallic NiTi2 phase, could be obtained after careful selection of sintering parameters. © 2007 Elsevier B.V. All rights reserved. Keywords: Titanium alloys; Powder consolidation; Sintering; Bulk amorphous materials
1. Introduction For more than a decade numerous bulk metallic glasses (BMGs) have been developed, which possess properties superior to those of their crystalline counterparts such as high strength, a large elastic limit, corrosion and wear resistance, etc [1]. Owing to these properties, BMGs are considered to have great potential for the application as structural materials. A common way of BMG fabrication is rapidly solidifying an alloy in liquid state, for instance by mold-casting. For systems showing high glass-forming ability such as Zr-based BMGs, samples of more than 10 mm in thickness have been obtained in this way [2]. Nevertheless, the shapes of rapidly quenched glassy alloys are usually restricted to thin rods or even to ribbons or powders due to a minimum cooling-rate required for super-cooling without undergoing devitrification. To overcome such size and shape restrictions, powder metallurgical processing techniques have been explored as alternative fabrication methods of BMGs [3–7]. The existence of a super-cooled liquid region, where a metallic glass exhibits Newtonian viscous flow, enables glassy powder to be consolidated to bulk form. Ti-based BMGs have been reported to exhibit excellent mechanical properties [8–10]. Among them, Ti50Cu25Ni20Sn5 ⁎ Corresponding author. Tel.: +82 2 958 6673; fax: +82 2 958 5529. E-mail address: [email protected] (P.P. Choi). 0167-577X/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2007.02.066
has an unusually large super-cooled liquid region (ΔT = 60 K [9]), making this alloy a promising candidate for powder consolidation. Spark-plasma sintering (SPS), which combines conventional electric current sintering with hot-pressing, is a relatively novel but highly efficient method of rapidly sintering hard-to-sinter materials at low temperatures [11]. Therefore, it shows outstanding potential for consolidation of amorphous powder into bulk samples. In the present study, suitable experimental conditions for preparing a bulk amorphous Ti50Cu25Ni20Sn5 alloy by means of mechanical alloying (MA) and SPS have been systematically explored. 2. Experimental Elemental powders of Ti, Cu, Ni and Sn (purity ≥ 99.9%) were mixed to the composition of Ti50Cu25Ni20Sn5 and mechanically alloyed in a high-energy planetary ball mill (AGO-2). Ball milling was performed at a rotational velocity of 300 rpm under protective Ar atmosphere, using milling tools made of hardened steel. A ball to powder weight ratio of 200 g:10 g was chosen. Structural characterization of the samples was done by X-ray diffraction (XRD). Finally obtained amorphous powders were additionally analyzed with a 3D atom probe (3DAP), where 3DAP samples were prepared by focused ion beam milling. Details about 3DAP analysis and the
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Fig. 1. XRD patterns of samples after MA for (a) 1 h, (b) 2 h, (c) 5 h, (d) 20 h, and (e) after MA for 20 h followed by SPS at 723 K and 500 MPa for 3 min.
apparatus used in this study are described elsewhere [12]. The thermal properties of amorphous samples were investigated by isochronous and isothermal differential scanning calorimetry (DSC) analyses. Powder consolidation by SPS was carried out at 698 and 723 K and an applied pressure of 500 MPa, using a mold and punches made of tungsten carbide. 3. Results and discussion Fig. 1(a)–(e) show the structural evolution of powder samples with progressing milling time. While only elemental peaks can be detected after 1 h of MA, a halo peak appears at 2Θ = 36–46° after 2 h. A gradual decrease in the intensity of the elemental peaks with increasing
Fig. 3. DSC curves obtained at a heating rate of 20 K/min: (a) a sample after MA for 20 h, and (b) a sample after MA for 20 h followed by SPS at 723 K and 500 MPa for 3 min.
milling time can be observed. After milling for 5 h, the only crystalline phases that can be detected are elemental Ti and Cu and the intermetallic NiTi2 phase. The transformation into a fully amorphous structure is completed after 20 h of MA. The elemental distribution maps as detected by 3DAP (see Fig. 2(a)) indicate chemical homogeneity within the amorphous alloy. This observation can be quantitatively confirmed by statistical χ2 tests (compare Fig. 2 (b)). The measured concentration frequency distributions do not significantly differ from binomial distributions. χ2 values do not exceed critical χ2a values corresponding to a significance level of 95%. Besides the components of the alloy, small traces of Fe and oxygen impurities (not shown in Fig. 2) that were incorporated into the powder during ball milling, were detected with a content of 0.4 and 0.2 at.%, respectively. It should also be mentioned that the composition of the amorphous
Fig. 2. 3DAP data of a sample after MA for 20 h: (a) elemental distribution maps, (b) concentration frequency distributions.
P.P. Choi et al. / Materials Letters 61 (2007) 4591–4594
Fig. 4. Polished cross-sectional view of samples spark-plasma sintered at: (a) 698 K and (b) 723 K for 3 min under an applied pressure of 500 MPa.
alloy, as measured by 3DAP (Ti46.5Cu20.7Ni27.1Sn5.1Fe0.4 O0.2), significantly deviates from the nominal composition. In comparison, a value of Ti47.4Cu27.4Ni20.5Sn4.7 was measured by inductively coupled plasma (ICP) spectroscopy. The reduced Ti concentration, which is detected in both 3DAP and ICP analyses, can be attributed to a compositional shift caused by preferential sticking of Ti powder particles to the milling tools in the initial stages of ball milling. The increase in Ni and decrease in Cu concentration, as detected by 3DAP, are most likely artifacts resulting from the lower field-evaporation field strength of Cu compared to Ni [13]. The DSC curve of an amorphous powder sample, obtained for isochronous annealing at a heating rate of 20 K/min, is shown in Fig. 3(a). It is seen that glass transition and crystallization occur at Tg = 709 K and Tx = 749 K, respectively. The product of primary crystallization was identified as NiTi2 from XRD data, as was reported for an amorphous Ti50Cu23Ni20Sn7 alloy [14]. In comparison with rapidly quenched Ti50Cu25Ni20Sn5 samples, for which Tg = 710 K and Tx = 770 K were reported [9], the Tx and ΔT values of the mechanically powder are significantly lower. Such a difference in thermal stability is related to the compositional shift and incorporation of impurities occurring during MA. SPS at optimum sintering parameters (at 723 K, applying a pressure of 500 MPa for 3 min) yields a mostly amorphous sample containing only a small fraction of the NiTi2 phase, as revealed by XRD (see Fig. 1 (e)). The presence of NiTi2 crystallites gives a reasonable explanation for the change in the DSC traces before and after SPS (see Fig. 3 (a) and (b)). The sintered sample exhibits a lower crystallization temperature and enthalpy (Tx = 739 K, ΔH ≈ − 69.5 J/g) than the initial powder
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(Tx = 749 K, ΔH ≈ −75.5 J/g). NiTi2 grains in the sintered sample are expected to serve as preferential nucleation sites for crystallization and thus lower the activation energy and onset temperature of crystallization. Assuming that the released crystallization enthalpy is proportional to the volume fraction of the amorphous phase, the volume fraction of the NiTi2 phase is estimated to be about 10%. Optical micrographs of polished cross-sections of sintered samples are shown in Fig. 4. A careful selection of sintering parameters was crucial for obtaining cylindrical bulk amorphous samples (10 mm in diameter and 5 mm in length). First and foremost, the sintering temperature was chosen to be within the super-cooled liquid region (between Tg and Tx), where a metallic glass exhibits Newtonian viscous flow and enables powder consolidation while crystallization is mostly retarded. To determine a suitable value for the sintering time, isothermal DSC analyses at temperatures within the super-cooled liquid region were performed. At 723 K, the maximum sintering temperature at which the samples still remained mostly amorphous after SPS, yielded an incubation period of 3 min. Therefore, the chosen sintering time was 3 min in order to avoid crystallization during SPS. It was found that glassy samples could only be obtained for sintering at temperatures ≤ 723 K, more than 25 K below Tx. Choosing higher sintering temperatures, in the range from 733 K to 743 K, resulted in completely crystallized samples. Short incubation periods at these temperatures can be regarded as one of the reasons for the occurrence of crystallization. Furthermore, crystallization can also be ascribed to errors made when measuring the sample temperature. Since the thermocouple cannot be directly placed into the center of the sample but only into a hole drilled into the mold, a significant difference between the measured and actual sample temperature is expected. In an attempt to completely avoid crystallization, a sample was sintered at 698 K. However, this sample exhibited low density (≈92%), containing numerous pores of up to 70 μm in size (compare Fig. 4 (a)) due to an insufficient viscous flow during SPS. Compacts with the highest relative density (≈ 99%) could be obtained upon sintering at 723 K and 500 MPa (see Fig. 4 (b)) with a small fraction of NiTi2 (≈ 10 vol.%), as mentioned before.
4. Conclusions A Ti50Cu25Ni20Sn5 BMG, showing nearly full density, was fabricated by spark-plasma sintering of mechanically alloyed powders. Despite choosing a sintering temperature more than 25 K lower than the onset temperature of primary crystallization a small fraction of the intermetallic NiTi2 phase was formed. Such an observation may be ascribed to a deviation of the SPS temperature measured by the thermocouple from the actual sample temperature. Other possible causes of partial crystallization are the short incubation periods of crystallization in combination with delayed sample cooling after the sintering process. A comparative study of the mechanical properties of the sintered compacts with those of rapidly quenched samples is the subject of on-going work. Acknowledgements We would like to thank the 3D Atom Probe group at the University of Göttingen in Germany for granting access to their apparatus. Financial support by the Korean Ministry of Commerce, Industry and Energy through the ReMM of the University of Ulsan is greatly acknowledged.
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P.P. Choi et al. / Materials Letters 61 (2007) 4591–4594
References [1] A. Inoue, Acta Mater. 48 (2000) 279. [2] A. Peker, W.L. Johnson, Appl. Phys. Lett. 63 (1993) 2342. [3] G.M. Dougherty, G.J. Shiftlet, S.J. Poon, Acta Metall. Mater. 42 (1994) 2275. [4] Y. Kawamura, H. Kato, A. Inoue, T. Masumoto, Appl. Phys. Lett. 67 (1995) 2008. [5] M.H. Lee, D.H. Bae, W.T. Kim, D.H. Kim, E. Rozhkova, P.B. Wheelock, D.J. Sordelet, J. Non-Cryst. Solids 315 (2003) 89. [6] D.J. Sordelet, E. Rozhkova, M.F. Besser, M.J. Kramer, J. Non-Cryst. Solids 317 (2003) 137. [7] K. Ozaki, T. Nishio, A. Matsumoto, K. Kobayashi, Mater. Sci. Eng. A 375–377 (2004) 857.
[8] K. Amiya, N. Nishiyama, A. Inoue, T. Masumoto, Mater. Sci. Eng., A Struct. Mater.: Prop. Microstruct. Process. 179−180 (1994) 692. [9] T. Zhang, A. Inoue, Mater. Trans., JIM 39 (1998) 1001. [10] J.M. Park, H.J. Chang, K.H. Han, W.T. Kim, D.H. Kim, Scr. Mater. 53 (2005) 1. [11] M. Tokita, J. Soc. Powder Technol., Jpn. 30 (1995) 790. [12] D. Blavette, B. Deconihout, A. Bostel, J. Sarrau, M. Bouet, A. Menand, Rev. Sci. Instrum. 64 (1993) 2911. [13] M.K. Miller, A. Cerezo, M.G. Hetherington, G.D.W. Smith, Atom Probe Field Ion Microscopy, Clarendon Press, Oxford, 1996. [14] D.V. Louzguine, A. Inoue, Scr. Mater. 43 (2000) 371.
Ti50Cu25Ni20Sn5 bulk metallic glass fabricated by powder consolidation P.P. Choi a,⁎, J.S. Kim b , O.T.H. Nguyen b , Y.S. Kwon b a
b
Nano-materials Research Center, Korea Institute of Science and Technology, P.O. Box 131, Cheongryang, Seoul 130–650, Republic of Korea Research Center for Machine Parts and Materials Processing, School of Materials Science and Engineering, University of Ulsan, Namgu Mugeo 2-Dong, San 29, Ulsan 680–749, Republic of Korea Received 16 October 2006; accepted 24 February 2007 Available online 3 March 2007
Abstract A bulk metallic glass of Ti50Cu25Ni20Sn5 nominal composition was produced via a powder metallurgical route, namely the preparation of glassy powders by mechanical alloying followed by their consolidation by spark-plasma sintering. Samples were characterized with respect to their structure and thermal properties before and after sintering. A bulk glassy sample, having nearly full density and containing only a small fraction of the intermetallic NiTi2 phase, could be obtained after careful selection of sintering parameters. © 2007 Elsevier B.V. All rights reserved. Keywords: Titanium alloys; Powder consolidation; Sintering; Bulk amorphous materials
1. Introduction For more than a decade numerous bulk metallic glasses (BMGs) have been developed, which possess properties superior to those of their crystalline counterparts such as high strength, a large elastic limit, corrosion and wear resistance, etc [1]. Owing to these properties, BMGs are considered to have great potential for the application as structural materials. A common way of BMG fabrication is rapidly solidifying an alloy in liquid state, for instance by mold-casting. For systems showing high glass-forming ability such as Zr-based BMGs, samples of more than 10 mm in thickness have been obtained in this way [2]. Nevertheless, the shapes of rapidly quenched glassy alloys are usually restricted to thin rods or even to ribbons or powders due to a minimum cooling-rate required for super-cooling without undergoing devitrification. To overcome such size and shape restrictions, powder metallurgical processing techniques have been explored as alternative fabrication methods of BMGs [3–7]. The existence of a super-cooled liquid region, where a metallic glass exhibits Newtonian viscous flow, enables glassy powder to be consolidated to bulk form. Ti-based BMGs have been reported to exhibit excellent mechanical properties [8–10]. Among them, Ti50Cu25Ni20Sn5 ⁎ Corresponding author. Tel.: +82 2 958 6673; fax: +82 2 958 5529. E-mail address: [email protected] (P.P. Choi). 0167-577X/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2007.02.066
has an unusually large super-cooled liquid region (ΔT = 60 K [9]), making this alloy a promising candidate for powder consolidation. Spark-plasma sintering (SPS), which combines conventional electric current sintering with hot-pressing, is a relatively novel but highly efficient method of rapidly sintering hard-to-sinter materials at low temperatures [11]. Therefore, it shows outstanding potential for consolidation of amorphous powder into bulk samples. In the present study, suitable experimental conditions for preparing a bulk amorphous Ti50Cu25Ni20Sn5 alloy by means of mechanical alloying (MA) and SPS have been systematically explored. 2. Experimental Elemental powders of Ti, Cu, Ni and Sn (purity ≥ 99.9%) were mixed to the composition of Ti50Cu25Ni20Sn5 and mechanically alloyed in a high-energy planetary ball mill (AGO-2). Ball milling was performed at a rotational velocity of 300 rpm under protective Ar atmosphere, using milling tools made of hardened steel. A ball to powder weight ratio of 200 g:10 g was chosen. Structural characterization of the samples was done by X-ray diffraction (XRD). Finally obtained amorphous powders were additionally analyzed with a 3D atom probe (3DAP), where 3DAP samples were prepared by focused ion beam milling. Details about 3DAP analysis and the
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Fig. 1. XRD patterns of samples after MA for (a) 1 h, (b) 2 h, (c) 5 h, (d) 20 h, and (e) after MA for 20 h followed by SPS at 723 K and 500 MPa for 3 min.
apparatus used in this study are described elsewhere [12]. The thermal properties of amorphous samples were investigated by isochronous and isothermal differential scanning calorimetry (DSC) analyses. Powder consolidation by SPS was carried out at 698 and 723 K and an applied pressure of 500 MPa, using a mold and punches made of tungsten carbide. 3. Results and discussion Fig. 1(a)–(e) show the structural evolution of powder samples with progressing milling time. While only elemental peaks can be detected after 1 h of MA, a halo peak appears at 2Θ = 36–46° after 2 h. A gradual decrease in the intensity of the elemental peaks with increasing
Fig. 3. DSC curves obtained at a heating rate of 20 K/min: (a) a sample after MA for 20 h, and (b) a sample after MA for 20 h followed by SPS at 723 K and 500 MPa for 3 min.
milling time can be observed. After milling for 5 h, the only crystalline phases that can be detected are elemental Ti and Cu and the intermetallic NiTi2 phase. The transformation into a fully amorphous structure is completed after 20 h of MA. The elemental distribution maps as detected by 3DAP (see Fig. 2(a)) indicate chemical homogeneity within the amorphous alloy. This observation can be quantitatively confirmed by statistical χ2 tests (compare Fig. 2 (b)). The measured concentration frequency distributions do not significantly differ from binomial distributions. χ2 values do not exceed critical χ2a values corresponding to a significance level of 95%. Besides the components of the alloy, small traces of Fe and oxygen impurities (not shown in Fig. 2) that were incorporated into the powder during ball milling, were detected with a content of 0.4 and 0.2 at.%, respectively. It should also be mentioned that the composition of the amorphous
Fig. 2. 3DAP data of a sample after MA for 20 h: (a) elemental distribution maps, (b) concentration frequency distributions.
P.P. Choi et al. / Materials Letters 61 (2007) 4591–4594
Fig. 4. Polished cross-sectional view of samples spark-plasma sintered at: (a) 698 K and (b) 723 K for 3 min under an applied pressure of 500 MPa.
alloy, as measured by 3DAP (Ti46.5Cu20.7Ni27.1Sn5.1Fe0.4 O0.2), significantly deviates from the nominal composition. In comparison, a value of Ti47.4Cu27.4Ni20.5Sn4.7 was measured by inductively coupled plasma (ICP) spectroscopy. The reduced Ti concentration, which is detected in both 3DAP and ICP analyses, can be attributed to a compositional shift caused by preferential sticking of Ti powder particles to the milling tools in the initial stages of ball milling. The increase in Ni and decrease in Cu concentration, as detected by 3DAP, are most likely artifacts resulting from the lower field-evaporation field strength of Cu compared to Ni [13]. The DSC curve of an amorphous powder sample, obtained for isochronous annealing at a heating rate of 20 K/min, is shown in Fig. 3(a). It is seen that glass transition and crystallization occur at Tg = 709 K and Tx = 749 K, respectively. The product of primary crystallization was identified as NiTi2 from XRD data, as was reported for an amorphous Ti50Cu23Ni20Sn7 alloy [14]. In comparison with rapidly quenched Ti50Cu25Ni20Sn5 samples, for which Tg = 710 K and Tx = 770 K were reported [9], the Tx and ΔT values of the mechanically powder are significantly lower. Such a difference in thermal stability is related to the compositional shift and incorporation of impurities occurring during MA. SPS at optimum sintering parameters (at 723 K, applying a pressure of 500 MPa for 3 min) yields a mostly amorphous sample containing only a small fraction of the NiTi2 phase, as revealed by XRD (see Fig. 1 (e)). The presence of NiTi2 crystallites gives a reasonable explanation for the change in the DSC traces before and after SPS (see Fig. 3 (a) and (b)). The sintered sample exhibits a lower crystallization temperature and enthalpy (Tx = 739 K, ΔH ≈ − 69.5 J/g) than the initial powder
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(Tx = 749 K, ΔH ≈ −75.5 J/g). NiTi2 grains in the sintered sample are expected to serve as preferential nucleation sites for crystallization and thus lower the activation energy and onset temperature of crystallization. Assuming that the released crystallization enthalpy is proportional to the volume fraction of the amorphous phase, the volume fraction of the NiTi2 phase is estimated to be about 10%. Optical micrographs of polished cross-sections of sintered samples are shown in Fig. 4. A careful selection of sintering parameters was crucial for obtaining cylindrical bulk amorphous samples (10 mm in diameter and 5 mm in length). First and foremost, the sintering temperature was chosen to be within the super-cooled liquid region (between Tg and Tx), where a metallic glass exhibits Newtonian viscous flow and enables powder consolidation while crystallization is mostly retarded. To determine a suitable value for the sintering time, isothermal DSC analyses at temperatures within the super-cooled liquid region were performed. At 723 K, the maximum sintering temperature at which the samples still remained mostly amorphous after SPS, yielded an incubation period of 3 min. Therefore, the chosen sintering time was 3 min in order to avoid crystallization during SPS. It was found that glassy samples could only be obtained for sintering at temperatures ≤ 723 K, more than 25 K below Tx. Choosing higher sintering temperatures, in the range from 733 K to 743 K, resulted in completely crystallized samples. Short incubation periods at these temperatures can be regarded as one of the reasons for the occurrence of crystallization. Furthermore, crystallization can also be ascribed to errors made when measuring the sample temperature. Since the thermocouple cannot be directly placed into the center of the sample but only into a hole drilled into the mold, a significant difference between the measured and actual sample temperature is expected. In an attempt to completely avoid crystallization, a sample was sintered at 698 K. However, this sample exhibited low density (≈92%), containing numerous pores of up to 70 μm in size (compare Fig. 4 (a)) due to an insufficient viscous flow during SPS. Compacts with the highest relative density (≈ 99%) could be obtained upon sintering at 723 K and 500 MPa (see Fig. 4 (b)) with a small fraction of NiTi2 (≈ 10 vol.%), as mentioned before.
4. Conclusions A Ti50Cu25Ni20Sn5 BMG, showing nearly full density, was fabricated by spark-plasma sintering of mechanically alloyed powders. Despite choosing a sintering temperature more than 25 K lower than the onset temperature of primary crystallization a small fraction of the intermetallic NiTi2 phase was formed. Such an observation may be ascribed to a deviation of the SPS temperature measured by the thermocouple from the actual sample temperature. Other possible causes of partial crystallization are the short incubation periods of crystallization in combination with delayed sample cooling after the sintering process. A comparative study of the mechanical properties of the sintered compacts with those of rapidly quenched samples is the subject of on-going work. Acknowledgements We would like to thank the 3D Atom Probe group at the University of Göttingen in Germany for granting access to their apparatus. Financial support by the Korean Ministry of Commerce, Industry and Energy through the ReMM of the University of Ulsan is greatly acknowledged.
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P.P. Choi et al. / Materials Letters 61 (2007) 4591–4594
References [1] A. Inoue, Acta Mater. 48 (2000) 279. [2] A. Peker, W.L. Johnson, Appl. Phys. Lett. 63 (1993) 2342. [3] G.M. Dougherty, G.J. Shiftlet, S.J. Poon, Acta Metall. Mater. 42 (1994) 2275. [4] Y. Kawamura, H. Kato, A. Inoue, T. Masumoto, Appl. Phys. Lett. 67 (1995) 2008. [5] M.H. Lee, D.H. Bae, W.T. Kim, D.H. Kim, E. Rozhkova, P.B. Wheelock, D.J. Sordelet, J. Non-Cryst. Solids 315 (2003) 89. [6] D.J. Sordelet, E. Rozhkova, M.F. Besser, M.J. Kramer, J. Non-Cryst. Solids 317 (2003) 137. [7] K. Ozaki, T. Nishio, A. Matsumoto, K. Kobayashi, Mater. Sci. Eng. A 375–377 (2004) 857.
[8] K. Amiya, N. Nishiyama, A. Inoue, T. Masumoto, Mater. Sci. Eng., A Struct. Mater.: Prop. Microstruct. Process. 179−180 (1994) 692. [9] T. Zhang, A. Inoue, Mater. Trans., JIM 39 (1998) 1001. [10] J.M. Park, H.J. Chang, K.H. Han, W.T. Kim, D.H. Kim, Scr. Mater. 53 (2005) 1. [11] M. Tokita, J. Soc. Powder Technol., Jpn. 30 (1995) 790. [12] D. Blavette, B. Deconihout, A. Bostel, J. Sarrau, M. Bouet, A. Menand, Rev. Sci. Instrum. 64 (1993) 2911. [13] M.K. Miller, A. Cerezo, M.G. Hetherington, G.D.W. Smith, Atom Probe Field Ion Microscopy, Clarendon Press, Oxford, 1996. [14] D.V. Louzguine, A. Inoue, Scr. Mater. 43 (2000) 371.