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Size-dependent structure and properties of rapidly solidified aluminum alloy powders
Scripta Materialia 45 (2001) 1341±1347
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Size-dependent structure and properties of rapidly solidi®ed aluminum alloy powders S.J. Honga, C. Suryanarayanab*, and B.S. Chuna a
Rapidly Solidi®ed Materials Research Center, Chungnam National University, Taedok Science Town, Taejon 305-764, South Korea b Department of Mechanical, Materials and Aerospace Engineering, University of Central Florida, Orlando, FL 32816-2450, USA
Received 14 March 2001; accepted 3 September 2001
Abstract With decreasing powder particle size, formation of intermetallic compounds was suppressed in gasatomized Al±14wt.%Ni±14wt.%Mm alloy powder; powders with <26 lm particle size consisted of a-Al particles embedded in an amorphous matrix. The hardness decreased with increasing particle size and also on heat treatment due to particle coarsening and/or crystallization of the remnant amorphous phase. Ó 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Rapid solidi®cation; Nanocrystalline Al composites; X-ray diraction; Electron microscopy
Introduction Amorphous Al±Ln±TM (Ln lanthanide element and TM transition metal) alloys show very high tensile strengths of over 1000 MPa with adequate ductility [1±5]. The strength of the alloy can be further increased by about 50% by heat treatment to produce a nanocomposite material due to the formation of about 30 vol.% of ®ne 10±30 nm size fcc a-Al particles in the amorphous matrix. Several strengthening mechanisms have been suggested for the increased strength of the nanocomposite material [3,6±8]. Majority of the recent work on Al-base amorphous alloys has been on glass-forming ability, crystallization behavior, optimization of heat treatment to synthesize the desired microstructure, and evaluation of the structure and mechanical properties of amorphous and/or heat-treated alloys [9±13]. Since the cooling rate during rapid solidi®cation determines the constitution of the as-solidi®ed powder and consequently the
*
Corresponding author. Tel.: +1-407-823-6662; fax: +1-407-823-0208. E-mail address: [email protected] (C. Suryanarayana).
1359-6462/01/$ - see front matter Ó 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 6 2 ( 0 1 ) 0 1 1 6 6 - 6
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properties, one would assume that lot of information is available on this aspect. But, only a limited knowledge exists. For example, it has been reported that the cooling rate increases with a decrease in particle size [14]. It has also been reported that phase selection in Ti±Al alloys is dependent on particle size [15,16]. Further, it has been known for a long time that ZrO2 particles smaller than 30 nm in size stabilize the high-temperature tetragonal phase at room temperature [17]. Some other examples have also been reported in mechanically alloyed ultra®ne powders [18,19]. In this paper, we report on our results of the microstructure, hardness, and phase stability of gas-atomized Al±14wt.%Ni±14wt.%Mm (Mm misch metal) powder particles as a function of the powder particle size in the rapidly solidi®ed condition and also after heat treatment. Experimental procedure Rapidly solidi®ed Al±14wt.%Ni±14wt.%Mm (referred to as Al±14Ni±14Mm from now on) powders were prepared by gas atomization; the experimental details of which have been reported earlier [20]. The as-solidi®ed powder was classi®ed into dierent size fractions to examine the constitution and mechanical properties of the powders as a function of the powder size. The as-solidi®ed powder was normally divided into three groups with the particle size 75±90 lm (group A), 45±53 lm (group B), and <26 lm (group C). In some selected experiments, powders in the range of 63±75 lm were also examined. Microstructural analysis of the as-solidi®ed powder was conducted using scanning (SEM) and transmission electron microscopes (TEM). The crystallization behavior of the rapidly solidi®ed alloy was investigated using a dierential scanning calorimeter (DSC) at a heating rate of 10 K/min. The crystal structures of the dierent phases were characterized by X-ray diraction (XRD) using CuKa radiation. The microhardness was determined using a Shimadzu Vickers hardness tester with a 10 g load. Results and discussion Fig. 1a±c show microstructures of the as-solidi®ed powders as a function of the particle size. The larger particles (Fig. 1a) contained a needle-shaped phase with a size of about 20 lm, while the intermediate size particles (Fig. 1b) contained this needleshaped phase with a size of only about 5 lm. No such precipitates were seen in the ®nest size powders (Fig. 1c). Since the cooling rate during rapid solidi®cation increases with decreasing particle size, one can infer that the crystallization process in the smallest powder particle size has been suppressed due to the high cooling rate and large undercooling. Since the cooling rates of hypoeutectic Al±Si alloys under similar experimental conditions were measured to vary from about 102 to 104 K/s, depending on the powder size [13], we expect that the cooling rate experienced by the Al±14Ni±14Mm alloy
S.J. Hong et al./Scripta Materialia 45 (2001) 1341±1347
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Fig. 1. SEM micrographs of gas atomized Al±14Ni±14Mm alloy powders: (a) 75±90 lm, (b) 45±53 lm, and (c) <26 lm in diameter.
powders is also in the range of 102 to 104 K/s, realizing that the smaller particles would have solidi®ed faster than the larger particles. The XRD pattern from the coarsest powder (Fig. 2a) consisted of many diraction peaks, which could be indexed to be a mixture of fcc a-Al, orthorhombic Al11 Ce3 with a 0:4389 nm, b 1:3025 nm, and c 1:0092 nm, orthorhombic Al3 Ni with a 0:6598 nm, b 0:7332 nm, and c 0:4862 nm, and orthorhombic Al11 La3 with a 0:4431 nm, b 1:3142 nm, and c 1:0132 nm. With decreasing powder particle size (Fig. 2b), the XRD pattern consisted of a-Al and orthorhombic Al11 Ce3 . However, the ®nest powder (Fig. 2c) is characterized by the presence of a-Al peaks superimposed on an amorphous phase (inferred from the weak diuse halo at 2h of 33±42°).
Fig. 2. XRD patterns of gas atomized Al±14Ni±14Mm alloy powders: (a) 75±90 lm, (b) 45±53 lm, and (c) <26 lm in diameter. Fig. 3. DSC curves of gas atomized Al±14Ni±14Mm alloy powders: (a) 75±90 lm, (b) 45±53 lm, and (c) <26 lm in diameter.
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From the above diraction patterns one can also notice that the a-Al peak positions have shifted to lower angle regions compared with that of the equilibrium phase. This shift in the peak positions is attributed to the eects of rapid solidi®cation in producing supersaturated solid solutions. Thus, the shift was maximum in the smallest powder particles (having solidi®ed most rapidly), and the least in the coarsest powder. Accordingly, the lattice parameter of the a-Al phase in these powders has varied from 0.4087 nm in the ®nest powder to 0.4073 nm in the coarsest powder. All these values are larger than that of the as-cast material, viz., 0.4055 nm [21], con®rming that the lattice parameter has increased with increasing solidi®cation rate due to supersaturation of Al with the solute elements. The Goldschmidt atomic radii of Ce and La are 0.182 and 0.187 nm, respectively, which are much larger than the value of 0.143 nm for Al [22]. Realizing that the atomic radius of Ni is smaller than that of Al, the increase in lattice parameter of a-Al seems to be related to the presence of larger amounts of Ce and La in the a-Al phase of the rapidly-solidi®ed ®ne powder than under equilibrium conditions. Fig. 3 shows the DSC traces of the three groups of rapidly solidi®ed alloy powders. Here we ®nd that the number of exothermic peaks is dierent in the dierent powder sizes. Two clear exothermic peaks±±one at 297 °C and the other at 355 °C±±can be seen from the powder of Group C. The second exothermic peak appears to be made up of two overlapping peaks; deconvolution of these two components was dicult. In the Group B powder, one notices only a small and broad peak at about 355 °C, again possibly made up of two overlapping peaks. The exothermic peak in the Group A powder was of very low intensity and barely visible, but at about the same temperature as the second peak in Group C powder. Thus, we can conclude that the powders that had solidi®ed most rapidly showed two exothermic peaks while those that solidi®ed relatively slowly showed only one exothermic peak. XRD patterns were recorded after the specimens were heated beyond the peak temperatures in the DSC traces. From a detailed analysis of these diraction patterns, it could be concluded that the ®rst exothermic peak (at 297 °C) was due to the precipitation of the a-Al phase from the amorphous matrix while the second peak (at 355 °C) was due to the precipitation of dierent intermetallics such as Al11 Ce3 , Al3 Ni, and Al11 La3 . Since the precipitating intermetallics have very similar characteristics (both structure and lattice parameters), it is only natural to expect that many of them will precipitate out at about the same temperature, thus explaining the broad and overlapping peaks. It has been known that amorphous alloys can crystallize in one of the three possible ways±±(a) primary, (b) polymorphic, or (c) eutectic [23]. Primary crystallization involves precipitation of a solid solution phase followed by crystallization of the remnant amorphous phase either in a polymorphic (i.e. without change in composition) or a eutectic (i.e. simultaneous formation of two crystalline phases) mode. Thus the nature of the crystalline phase formed depends on the composition of the amorphous phase and its volume fraction on the amount of the amorphous phase available for transformation. In the present investigation primary crystallization, followed by polymorphic crystallization, seems to have occurred in the Al±14Ni±14Mm alloy. The DSC curves presented in Fig. 3 can be interpreted in the following way. The observation of two exothermic peaks in the ®nest powder size suggests that the as-
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solidi®ed material contains mostly the amorphous phase (perhaps with some amount of the a-Al phase as indicated by the XRD patterns). The low intensity of the ®rst exothermic peak suggests that the a-Al phase has formed from the amorphous phase. If the solidi®cation rate were to be higher, then the as-solidi®ed powder would have been fully amorphous and consequently, the peak at 297 °C would have a larger area. But, if the rapidly solidi®ed powder were to contain a signi®cant amount of the a-Al phase, then the concentration of the remnant amorphous phase will be such as to directly precipitate out the intermetallics. Such a behavior is observed in powders of groups A and B (Fig. 3a and b), where only the second peak is observed at about 355 °C. Since the slowest solidi®ed powder (group A) contains very little of the amorphous phase, the peak is hardly observable and its area is very small. From the combined studies of microstructure, DSC and XRD on the crystallization behavior of rapidly solidi®ed Al±14Ni±14Mm powders, it can be concluded that the solidi®cation rate of the powder determines the crystallization behavior and hence the microstructure. From these studies, the crystallization sequence in this alloy system can be summarized as: Amorphous ! Amorphous' a-Al ! a-Al Al11 Ce3 Al11 La3 Al3 Ni where Amorphous' represents the amorphous phase with a dierent composition after precipitation of the a-Al phase from the original amorphous phase. Fig. 4(a) shows the bright-®eld TEM micrograph and the selected area diraction (SAD) pattern from Group B powders. A similar set for Group C powders is shown in Fig. 4b. The SAD pattern from the Group C powder (<26 lm) consists of a diuse halo on which some sharp diraction spots, corresponding to the a-Al phase, are superimposed; con®rming the result obtained from the XRD patterns. The bright-®eld electron micrograph from this powder shows a-Al particles of 10±30 nm size embedded in the amorphous matrix, whereas much coarser particles of 40±80 nm with a complex microstructure are observed in the powders of Group B. The EDS spectrum from some of the particles in Fig. 4a shows the presence of 76.6at.%Al and 23.4at.%Ni, and combined with the XRD patterns, these particles were identi®ed as Al3 Ni. Fig. 5 shows the room temperature Vickers microhardness of the Al±14Ni±14Mm powders in the as-solidi®ed condition and also of the powder heat treated at 400 °C for
Fig. 4. Bright-®eld images and SAD patterns of gas atomized Al±14Ni±14Mm alloy powders. The powder particle sizes are (a) 45±53 lm, and (b) <26 lm.
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Fig. 5. Room temperature Vickers microhardness of the powders in the as-solidi®ed condition and after heat treatment at 400 °C for 1 h.
1 h under argon atmosphere. The Group B (45±53 lm) powder showed higher hardness than Group A powder (75±90 lm), mostly due to the ®ne size of the precipitates. The hardness decreased also with heat treatment. Kim et al. [7] have reported that the mechanical properties of amorphous Al±TM± RE (RE rare-earth metal) alloys are determined by solute concentration in the amorphous matrix and the precipitated particle size according to the relation: rpa rsc rppt where rpa is the strength of the partially amorphous material, rsc is the strength of the matrix due to solute concentration, and rppt is the additional strength due to the formation of the precipitates. If the precipitates are extremely ®ne, the strength of the alloy is essentially determined by the strength of the precipitate phase, because rppt is close to the theoretical strength value. From this it can be easily understood that the strength (or hardness) will decrease when the particle size increases either due to particle coarsening or crystallization of the remaining amorphous phase or both. Further, since nanocrystalline materials cannot contain many dislocations, they behave like an almost perfect (defect-free) material with a high strength [19]. Thus, the Hall±Petch equation, which explains the grain size dependence of strength on the grain size may not be fully applicable to nanocrystalline materials since it is very dicult to generate and propagate dislocations in ®ne particles. The critical grain size below which the strength of aluminum does not change signi®cantly occurs at about 60 nm, while a rapid reduction in strength is expected to occur above this size. Conclusions Gas-atomized Al±14Ni±14Mm alloy powders were divided into three groups, viz., 75±90 lm, 45±53 lm and <26 lm, in order to identify the eect of powder particle size on the microstructure and transformation behavior. With decreasing particle size,
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precipitation of intermetallic compounds such as Al3 Ni, Al11 Ce3 and Al11 La3 was suppressed, and only a-Al precipitated from the amorphous matrix in the ®nest powders. With increasing powder size from <26 lm to 45±53 lm, the a-Al precipitate size increased from about 10±30 nm to 40±80 nm; consequently, the hardness decreased with increasing powder size. Also, heat treating of the powders caused a reduction in the hardness due to particle coarsening and/or crystallization of the remaining amorphous phase.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
Inoue, A., Ohtera, K., Tsai, A. P., & Masumoto, T. (1988). Jpn J Appl Phys 27, L280. Inoue, A., Ohtera, K., Tsai, A. P., & Masumoto, T. (1988). Jpn J Appl Phys 27, L479. Kim, Y. H., Inoue, A., & Masumoto, T. (1991). Mater Trans JIM 32, 331. Higashi, K., Mukai, T., Uoya, A., Inoue, A., & Masumoto, T. (1995). Mater Trans JIM 36, 1467. Davis, L. A., Das, S. K., Li, J. C. M., & Zedalis, M. S. (1994). Int J Rapid Solidif 8, 73. Zhong, Z. C., Jiang, X. Y., & Greer, A. L. (1997). Mater Sci Eng A 226±228, 531. Kim, H. S., Warren, P. J., Cantor, B., & Lee, H. R. (1999). Nanostructured Mater 11, 241. Kim, H. S., & Hong, S. I. (1999). Acta Mater 47, 2059. Inoue, A. (1998). Prog Mater Sci 43, 365. Takeda, Y., Hayashi, T., Odani, Y., Amano, N., & Kuroishi, N. (1988). Modern Dev Powder Met 18±21, 553. Groza, J. R. (1999). In C. Suryanarayana (Ed.). Non-Equilibrium Processing of Materials (p. 347). Oxford: Pergamon. Ohtera, K., Inoue, A., Terabayashi, T., Nagahama, H., & Masumoto, T. (1992). Mater Trans JIM 33, 775. Yoon, G. M., Kim, Y. H., Jung, B. H., & Kim, I. B. (1996). J Korean Inst Met & Mater 34, 696. Suryanarayana, C. (1991). Processing of Metals. In R.W. Cahn (Ed.). Materials Science and TechnologyÐA Comprehensive Treatment (Vol. 15, p. 57). Weinheim, Germany: VCH. Graves, J. A., Perepezko, J. H., Ward, C. H., & Froes, F. H. (1987). Scripta Met 21, 567. McCollough, C., Valencia, J. J., Levi, C. G., & Mehrabian, R. (1987). Scripta Met 21, 1341. Garvie, R. C. (1978). J Phys Chem 82, 218. Suryanarayana, C. (2001). Prog Mater Sci 46, 1. Suryanarayana, C. (1995). Int Mater Rev 40, 41. Hong, S. J., Kim, T. S., Suryanarayana, C., Cho, D. H., & Chun, B. S. (2001). J Mater Synth Proc 9, 39. Mondolfo, L. F. (1976). Structure and Properties of Aluminum Alloys. London: Butterworths. Goldschmidt, H. J. (1967). Interstitial Alloys. London: Butterworths. K oster, U., & Herold, U. (1981). In H. J. Guntherodt & H. Beck (Eds.). Glassy Alloys I (p. 225). Berlin: Springer.
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Size-dependent structure and properties of rapidly solidi®ed aluminum alloy powders S.J. Honga, C. Suryanarayanab*, and B.S. Chuna a
Rapidly Solidi®ed Materials Research Center, Chungnam National University, Taedok Science Town, Taejon 305-764, South Korea b Department of Mechanical, Materials and Aerospace Engineering, University of Central Florida, Orlando, FL 32816-2450, USA
Received 14 March 2001; accepted 3 September 2001
Abstract With decreasing powder particle size, formation of intermetallic compounds was suppressed in gasatomized Al±14wt.%Ni±14wt.%Mm alloy powder; powders with <26 lm particle size consisted of a-Al particles embedded in an amorphous matrix. The hardness decreased with increasing particle size and also on heat treatment due to particle coarsening and/or crystallization of the remnant amorphous phase. Ó 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Rapid solidi®cation; Nanocrystalline Al composites; X-ray diraction; Electron microscopy
Introduction Amorphous Al±Ln±TM (Ln lanthanide element and TM transition metal) alloys show very high tensile strengths of over 1000 MPa with adequate ductility [1±5]. The strength of the alloy can be further increased by about 50% by heat treatment to produce a nanocomposite material due to the formation of about 30 vol.% of ®ne 10±30 nm size fcc a-Al particles in the amorphous matrix. Several strengthening mechanisms have been suggested for the increased strength of the nanocomposite material [3,6±8]. Majority of the recent work on Al-base amorphous alloys has been on glass-forming ability, crystallization behavior, optimization of heat treatment to synthesize the desired microstructure, and evaluation of the structure and mechanical properties of amorphous and/or heat-treated alloys [9±13]. Since the cooling rate during rapid solidi®cation determines the constitution of the as-solidi®ed powder and consequently the
*
Corresponding author. Tel.: +1-407-823-6662; fax: +1-407-823-0208. E-mail address: [email protected] (C. Suryanarayana).
1359-6462/01/$ - see front matter Ó 2001 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 6 2 ( 0 1 ) 0 1 1 6 6 - 6
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S.J. Hong et al./Scripta Materialia 45 (2001) 1341±1347
properties, one would assume that lot of information is available on this aspect. But, only a limited knowledge exists. For example, it has been reported that the cooling rate increases with a decrease in particle size [14]. It has also been reported that phase selection in Ti±Al alloys is dependent on particle size [15,16]. Further, it has been known for a long time that ZrO2 particles smaller than 30 nm in size stabilize the high-temperature tetragonal phase at room temperature [17]. Some other examples have also been reported in mechanically alloyed ultra®ne powders [18,19]. In this paper, we report on our results of the microstructure, hardness, and phase stability of gas-atomized Al±14wt.%Ni±14wt.%Mm (Mm misch metal) powder particles as a function of the powder particle size in the rapidly solidi®ed condition and also after heat treatment. Experimental procedure Rapidly solidi®ed Al±14wt.%Ni±14wt.%Mm (referred to as Al±14Ni±14Mm from now on) powders were prepared by gas atomization; the experimental details of which have been reported earlier [20]. The as-solidi®ed powder was classi®ed into dierent size fractions to examine the constitution and mechanical properties of the powders as a function of the powder size. The as-solidi®ed powder was normally divided into three groups with the particle size 75±90 lm (group A), 45±53 lm (group B), and <26 lm (group C). In some selected experiments, powders in the range of 63±75 lm were also examined. Microstructural analysis of the as-solidi®ed powder was conducted using scanning (SEM) and transmission electron microscopes (TEM). The crystallization behavior of the rapidly solidi®ed alloy was investigated using a dierential scanning calorimeter (DSC) at a heating rate of 10 K/min. The crystal structures of the dierent phases were characterized by X-ray diraction (XRD) using CuKa radiation. The microhardness was determined using a Shimadzu Vickers hardness tester with a 10 g load. Results and discussion Fig. 1a±c show microstructures of the as-solidi®ed powders as a function of the particle size. The larger particles (Fig. 1a) contained a needle-shaped phase with a size of about 20 lm, while the intermediate size particles (Fig. 1b) contained this needleshaped phase with a size of only about 5 lm. No such precipitates were seen in the ®nest size powders (Fig. 1c). Since the cooling rate during rapid solidi®cation increases with decreasing particle size, one can infer that the crystallization process in the smallest powder particle size has been suppressed due to the high cooling rate and large undercooling. Since the cooling rates of hypoeutectic Al±Si alloys under similar experimental conditions were measured to vary from about 102 to 104 K/s, depending on the powder size [13], we expect that the cooling rate experienced by the Al±14Ni±14Mm alloy
S.J. Hong et al./Scripta Materialia 45 (2001) 1341±1347
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Fig. 1. SEM micrographs of gas atomized Al±14Ni±14Mm alloy powders: (a) 75±90 lm, (b) 45±53 lm, and (c) <26 lm in diameter.
powders is also in the range of 102 to 104 K/s, realizing that the smaller particles would have solidi®ed faster than the larger particles. The XRD pattern from the coarsest powder (Fig. 2a) consisted of many diraction peaks, which could be indexed to be a mixture of fcc a-Al, orthorhombic Al11 Ce3 with a 0:4389 nm, b 1:3025 nm, and c 1:0092 nm, orthorhombic Al3 Ni with a 0:6598 nm, b 0:7332 nm, and c 0:4862 nm, and orthorhombic Al11 La3 with a 0:4431 nm, b 1:3142 nm, and c 1:0132 nm. With decreasing powder particle size (Fig. 2b), the XRD pattern consisted of a-Al and orthorhombic Al11 Ce3 . However, the ®nest powder (Fig. 2c) is characterized by the presence of a-Al peaks superimposed on an amorphous phase (inferred from the weak diuse halo at 2h of 33±42°).
Fig. 2. XRD patterns of gas atomized Al±14Ni±14Mm alloy powders: (a) 75±90 lm, (b) 45±53 lm, and (c) <26 lm in diameter. Fig. 3. DSC curves of gas atomized Al±14Ni±14Mm alloy powders: (a) 75±90 lm, (b) 45±53 lm, and (c) <26 lm in diameter.
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From the above diraction patterns one can also notice that the a-Al peak positions have shifted to lower angle regions compared with that of the equilibrium phase. This shift in the peak positions is attributed to the eects of rapid solidi®cation in producing supersaturated solid solutions. Thus, the shift was maximum in the smallest powder particles (having solidi®ed most rapidly), and the least in the coarsest powder. Accordingly, the lattice parameter of the a-Al phase in these powders has varied from 0.4087 nm in the ®nest powder to 0.4073 nm in the coarsest powder. All these values are larger than that of the as-cast material, viz., 0.4055 nm [21], con®rming that the lattice parameter has increased with increasing solidi®cation rate due to supersaturation of Al with the solute elements. The Goldschmidt atomic radii of Ce and La are 0.182 and 0.187 nm, respectively, which are much larger than the value of 0.143 nm for Al [22]. Realizing that the atomic radius of Ni is smaller than that of Al, the increase in lattice parameter of a-Al seems to be related to the presence of larger amounts of Ce and La in the a-Al phase of the rapidly-solidi®ed ®ne powder than under equilibrium conditions. Fig. 3 shows the DSC traces of the three groups of rapidly solidi®ed alloy powders. Here we ®nd that the number of exothermic peaks is dierent in the dierent powder sizes. Two clear exothermic peaks±±one at 297 °C and the other at 355 °C±±can be seen from the powder of Group C. The second exothermic peak appears to be made up of two overlapping peaks; deconvolution of these two components was dicult. In the Group B powder, one notices only a small and broad peak at about 355 °C, again possibly made up of two overlapping peaks. The exothermic peak in the Group A powder was of very low intensity and barely visible, but at about the same temperature as the second peak in Group C powder. Thus, we can conclude that the powders that had solidi®ed most rapidly showed two exothermic peaks while those that solidi®ed relatively slowly showed only one exothermic peak. XRD patterns were recorded after the specimens were heated beyond the peak temperatures in the DSC traces. From a detailed analysis of these diraction patterns, it could be concluded that the ®rst exothermic peak (at 297 °C) was due to the precipitation of the a-Al phase from the amorphous matrix while the second peak (at 355 °C) was due to the precipitation of dierent intermetallics such as Al11 Ce3 , Al3 Ni, and Al11 La3 . Since the precipitating intermetallics have very similar characteristics (both structure and lattice parameters), it is only natural to expect that many of them will precipitate out at about the same temperature, thus explaining the broad and overlapping peaks. It has been known that amorphous alloys can crystallize in one of the three possible ways±±(a) primary, (b) polymorphic, or (c) eutectic [23]. Primary crystallization involves precipitation of a solid solution phase followed by crystallization of the remnant amorphous phase either in a polymorphic (i.e. without change in composition) or a eutectic (i.e. simultaneous formation of two crystalline phases) mode. Thus the nature of the crystalline phase formed depends on the composition of the amorphous phase and its volume fraction on the amount of the amorphous phase available for transformation. In the present investigation primary crystallization, followed by polymorphic crystallization, seems to have occurred in the Al±14Ni±14Mm alloy. The DSC curves presented in Fig. 3 can be interpreted in the following way. The observation of two exothermic peaks in the ®nest powder size suggests that the as-
S.J. Hong et al./Scripta Materialia 45 (2001) 1341±1347
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solidi®ed material contains mostly the amorphous phase (perhaps with some amount of the a-Al phase as indicated by the XRD patterns). The low intensity of the ®rst exothermic peak suggests that the a-Al phase has formed from the amorphous phase. If the solidi®cation rate were to be higher, then the as-solidi®ed powder would have been fully amorphous and consequently, the peak at 297 °C would have a larger area. But, if the rapidly solidi®ed powder were to contain a signi®cant amount of the a-Al phase, then the concentration of the remnant amorphous phase will be such as to directly precipitate out the intermetallics. Such a behavior is observed in powders of groups A and B (Fig. 3a and b), where only the second peak is observed at about 355 °C. Since the slowest solidi®ed powder (group A) contains very little of the amorphous phase, the peak is hardly observable and its area is very small. From the combined studies of microstructure, DSC and XRD on the crystallization behavior of rapidly solidi®ed Al±14Ni±14Mm powders, it can be concluded that the solidi®cation rate of the powder determines the crystallization behavior and hence the microstructure. From these studies, the crystallization sequence in this alloy system can be summarized as: Amorphous ! Amorphous' a-Al ! a-Al Al11 Ce3 Al11 La3 Al3 Ni where Amorphous' represents the amorphous phase with a dierent composition after precipitation of the a-Al phase from the original amorphous phase. Fig. 4(a) shows the bright-®eld TEM micrograph and the selected area diraction (SAD) pattern from Group B powders. A similar set for Group C powders is shown in Fig. 4b. The SAD pattern from the Group C powder (<26 lm) consists of a diuse halo on which some sharp diraction spots, corresponding to the a-Al phase, are superimposed; con®rming the result obtained from the XRD patterns. The bright-®eld electron micrograph from this powder shows a-Al particles of 10±30 nm size embedded in the amorphous matrix, whereas much coarser particles of 40±80 nm with a complex microstructure are observed in the powders of Group B. The EDS spectrum from some of the particles in Fig. 4a shows the presence of 76.6at.%Al and 23.4at.%Ni, and combined with the XRD patterns, these particles were identi®ed as Al3 Ni. Fig. 5 shows the room temperature Vickers microhardness of the Al±14Ni±14Mm powders in the as-solidi®ed condition and also of the powder heat treated at 400 °C for
Fig. 4. Bright-®eld images and SAD patterns of gas atomized Al±14Ni±14Mm alloy powders. The powder particle sizes are (a) 45±53 lm, and (b) <26 lm.
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Fig. 5. Room temperature Vickers microhardness of the powders in the as-solidi®ed condition and after heat treatment at 400 °C for 1 h.
1 h under argon atmosphere. The Group B (45±53 lm) powder showed higher hardness than Group A powder (75±90 lm), mostly due to the ®ne size of the precipitates. The hardness decreased also with heat treatment. Kim et al. [7] have reported that the mechanical properties of amorphous Al±TM± RE (RE rare-earth metal) alloys are determined by solute concentration in the amorphous matrix and the precipitated particle size according to the relation: rpa rsc rppt where rpa is the strength of the partially amorphous material, rsc is the strength of the matrix due to solute concentration, and rppt is the additional strength due to the formation of the precipitates. If the precipitates are extremely ®ne, the strength of the alloy is essentially determined by the strength of the precipitate phase, because rppt is close to the theoretical strength value. From this it can be easily understood that the strength (or hardness) will decrease when the particle size increases either due to particle coarsening or crystallization of the remaining amorphous phase or both. Further, since nanocrystalline materials cannot contain many dislocations, they behave like an almost perfect (defect-free) material with a high strength [19]. Thus, the Hall±Petch equation, which explains the grain size dependence of strength on the grain size may not be fully applicable to nanocrystalline materials since it is very dicult to generate and propagate dislocations in ®ne particles. The critical grain size below which the strength of aluminum does not change signi®cantly occurs at about 60 nm, while a rapid reduction in strength is expected to occur above this size. Conclusions Gas-atomized Al±14Ni±14Mm alloy powders were divided into three groups, viz., 75±90 lm, 45±53 lm and <26 lm, in order to identify the eect of powder particle size on the microstructure and transformation behavior. With decreasing particle size,
S.J. Hong et al./Scripta Materialia 45 (2001) 1341±1347
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precipitation of intermetallic compounds such as Al3 Ni, Al11 Ce3 and Al11 La3 was suppressed, and only a-Al precipitated from the amorphous matrix in the ®nest powders. With increasing powder size from <26 lm to 45±53 lm, the a-Al precipitate size increased from about 10±30 nm to 40±80 nm; consequently, the hardness decreased with increasing powder size. Also, heat treating of the powders caused a reduction in the hardness due to particle coarsening and/or crystallization of the remaining amorphous phase.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
Inoue, A., Ohtera, K., Tsai, A. P., & Masumoto, T. (1988). Jpn J Appl Phys 27, L280. Inoue, A., Ohtera, K., Tsai, A. P., & Masumoto, T. (1988). Jpn J Appl Phys 27, L479. Kim, Y. H., Inoue, A., & Masumoto, T. (1991). Mater Trans JIM 32, 331. Higashi, K., Mukai, T., Uoya, A., Inoue, A., & Masumoto, T. (1995). Mater Trans JIM 36, 1467. Davis, L. A., Das, S. K., Li, J. C. M., & Zedalis, M. S. (1994). Int J Rapid Solidif 8, 73. Zhong, Z. C., Jiang, X. Y., & Greer, A. L. (1997). Mater Sci Eng A 226±228, 531. Kim, H. S., Warren, P. J., Cantor, B., & Lee, H. R. (1999). Nanostructured Mater 11, 241. Kim, H. S., & Hong, S. I. (1999). Acta Mater 47, 2059. Inoue, A. (1998). Prog Mater Sci 43, 365. Takeda, Y., Hayashi, T., Odani, Y., Amano, N., & Kuroishi, N. (1988). Modern Dev Powder Met 18±21, 553. Groza, J. R. (1999). In C. Suryanarayana (Ed.). Non-Equilibrium Processing of Materials (p. 347). Oxford: Pergamon. Ohtera, K., Inoue, A., Terabayashi, T., Nagahama, H., & Masumoto, T. (1992). Mater Trans JIM 33, 775. Yoon, G. M., Kim, Y. H., Jung, B. H., & Kim, I. B. (1996). J Korean Inst Met & Mater 34, 696. Suryanarayana, C. (1991). Processing of Metals. In R.W. Cahn (Ed.). Materials Science and TechnologyÐA Comprehensive Treatment (Vol. 15, p. 57). Weinheim, Germany: VCH. Graves, J. A., Perepezko, J. H., Ward, C. H., & Froes, F. H. (1987). Scripta Met 21, 567. McCollough, C., Valencia, J. J., Levi, C. G., & Mehrabian, R. (1987). Scripta Met 21, 1341. Garvie, R. C. (1978). J Phys Chem 82, 218. Suryanarayana, C. (2001). Prog Mater Sci 46, 1. Suryanarayana, C. (1995). Int Mater Rev 40, 41. Hong, S. J., Kim, T. S., Suryanarayana, C., Cho, D. H., & Chun, B. S. (2001). J Mater Synth Proc 9, 39. Mondolfo, L. F. (1976). Structure and Properties of Aluminum Alloys. London: Butterworths. Goldschmidt, H. J. (1967). Interstitial Alloys. London: Butterworths. K oster, U., & Herold, U. (1981). In H. J. Guntherodt & H. Beck (Eds.). Glassy Alloys I (p. 225). Berlin: Springer.