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Chemical leaching of rapidly solidified Al–Si binary alloys
Journal of Alloys and Compounds 396 (2005) 302–308
Chemical leaching of rapidly solidified Al–Si binary alloys I. Yamauchia,∗ , K. Takaharaa , T. Tanakab , K. Matsubarac a b
Department of Material Science and Engineering, School of Engineering of Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan Department of Mechanical Engineering, Faculty of Engineering, Osaka Sangyo University, 3-1-1 Nakagaito, Daito, Osaka 574-8530, Japan c Samsung Yokohama Research Institute, 2-7 Sugasawa-cho, Tsurumi-ku, Yokohama 230-0027, Japan Received 28 November 2004; received in revised form 27 December 2004; accepted 4 January 2005 Available online 1 February 2005
Abstract Various particulate precursors of Al100−x Six (x = 5–12) alloys were prepared by a rapid solidification process. The rapidly solidified structures of the precursors were examined by XRD, DSC and SEM. Most of Si atoms were dissolved into the ␣-Al(fcc) phase by rapid solidification though the solubility of Si in the ␣-Al phase is negligibly small in conventional solidification. In the case of 5 at.% Si alloy, a single ␣-Al phase was only formed. The amount of the primary Si phase increased with increase of Si content for the alloys beyond 8 at.% Si. Rapid solidification was effective to form super-saturated ␣-Al precursors. These precursors were chemically leached by using a basic solution (NaOH) or a hydrochloric acid (HCl) solution. All Al atoms were removed by a HCl solution as well as a NaOH solution. Granules of the Si phase were newly formed during leaching. The specific surface area was about 50–70 m2 /g independent of Si content. The leaching behavior in both solutions was slightly different. In the case of a NaOH solution, the shape of the precursor often degenerated after leaching. On the other hand, it was retained after leaching by a HCl solution. Fine Si particles precipitated in the ␣-Al phase by annealing of as-rapidly solidified precursors at 773 K for 7.2 × 103 s. In this case, it was difficult to obtain any products by NaOH leaching, but a few of Si particles were obtained by HCl leaching. Precipitated Si particles were dissolved by the NaOH solution. The X-ray diffraction patterns of leached specimens showed broad lines of the Si phase and its lattice constant was slightly larger than that of the pure Si phase. The microstructures of the leached specimens were examined by transmission electron microscopy. It showed that the leached specimens had a skeletal structure composed of slightly elongated particles of the Si phase and quite fine pores. The particle size was about 30–50 nm. It was of comparable order with that evaluated by Scherer’s equation. Leaching of rapidly solidified precursors was effective to form skeletal structures composed of fine Si particles with pores. © 2005 Elsevier B.V. All rights reserved. Keywords: Rapid solidification; Al–Ni–Ag; Skeletal material; Metastable phase
1. Introduction Skeletal materials composed of fine particles and pores can be obtained by applying a chemical leaching process to an aluminum base alloy [1–3]. In this process, aluminum atoms are removed from lattices sites of the alloy by chemical reaction and many vacancy sites are formed. The residual atoms on the lattice sites except Al atoms will rearrange and form a new lattice according to their chemical stability. When the mobility of these atoms is enhanced due to ∗
Corresponding author. Tel.: +81 6 6879 7524; fax: +81 6 6879 7536. E-mail address: [email protected] (I. Yamauchi).
0925-8388/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2005.01.006
the induced vacancies, these atoms can easily move together and may form some nucleus for a new phase. The mobile atoms around the nucleus will arrive on the new phase and make it grow. Thus the new phase becomes larger in size. However, the long distance movement of the atoms will be limited so that some of the atoms will form a new nucleus at another place. Finally, numerous fine particles of the new phase will be formed. The space between each granule will appear as fine pores. This is a principle of skeletal structure formation by chemical leaching. Several authors have been reported on the formation of non-equilibrium skeletal materials by applying chemical leaching to various rapidly solidified precursors [4–9].
I. Yamauchi et al. / Journal of Alloys and Compounds 396 (2005) 302–308
The equilibrium phase diagram [10] of the binary Al–Si system shows that the mutual solubility of these elements is negligibly small at room temperature. In a hypo-eutectic Al–Si alloy, most of the Si atoms are usually solidified as a eutectic Si phase. Silicon is well known as an amphoteric element. Therefore, it was thought as an improper element for chemical leaching. There are only few reports on leaching for this system. A highly supersaturated ␣-Al(fcc) solid solution of Si can be obtained by rapid solidification [11]. The supersaturated ␣-Al solid solution means that the Si atoms substitutionally occupy the lattice sites of the ␣-Al phase. In this case, a new Si phase can be formed by leaching. In this study, we applied rapid solidification to obtain supersaturated Al–Si precursors and examined the structure and the morphology transition by leaching in detail.
2. Experimental procedures 2.1. Rapid solidification for precursors A series of Al100−x Six (x = 5–12) alloys was prepared from Al (99.99 mass%) and high purity Si by melting in Ar atmosphere. The melt was cast into a cylindrical metal mold of 10 mm in diameter. About 20 g of the alloy was remelted in a graphite nozzle with a small bore of 0.5 mm in diameter at its top end by high frequency induction and then the molten jet was ejected into a rotating water layer. Thus, rapidly solidified powders were obtained by using the rotating-wateratomization process [12]. The rotating water velocity was 42 m/s and the atomization conditions were the same for all experiments. The atomized powders were immediately collected with the water and dried in an evacuated chamber.
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Rapidly solidified thin ribbons were prepared by a conventional single roller process for some alloys. The annealed precursors were obtained by annealing the rapidly solidified powder at 773 K for 7.2 ks in vacuum. The structures of these specimens were examined by differential scanning calorimetry (DSC) at a heating rate of 0.25 K s−1 , SEM and XRD using a graphite monochromatic Cu K␣. For SEM observation, polished specimens were slightly etched by a dix Keller solution for 30 s. 2.2. Chemical leaching About 5 g of each rapidly solidified or annealed specimens was leached by a 20% NaOH basic solution or a 1:1 dilute hydrochloric acid (HCl) solution for 1.2 ks at 333 K. The leached specimen was collected by a Buchner’s funnel and washed several times with distilled water. The structure of the leached specimens was examined by XRD. In the case of the leached powder, the samples for TEM observation were prepared by molding them into plastic resin which then was cut to a quite thin plate by using a microtome with a diamond knife. The thickness of the plate was about 60 nm. On the other hand, the leached ribbon samples were directly examined by TEM without thinning. The specific surface was measured by the BET method. 3. Results and discussion 3.1. Effect of rapid solidification on as-solidified structure Fig. 1(a) and (b) shows as-rapidly solidified structures obtained by SEM of Al–5% Si and Al–12% Si alloys,
Fig. 1. Microstructures of as-solidified and annealed specimens of 5 at.% Si and 12 at.% Si alloys: (a) Al–5% Si RWA; (b) Al–12% Si RWA; (c) Al–5% Si heat treated; and (d) Al–12% Si heat treated.
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Fig. 2. X-ray diffraction patterns of rapidly solidified Al–Si alloys.
respectively. They show quite fine microstructures having white and dark contrasts. Their size increased with increase of the Si content. The X-ray diffraction patterns (Fig. 2) of these specimens show Si phase diffraction lines except for the Al–5% Si alloy where they were negligibly weak. The equilibrium diagram of the Al–Si binary system shows that the solubility of Si into the ␣-Al phase is negligibly small [10]. In conventional solidification processes, most of Si atoms in these hypo-eutectic alloys solidify as the eutectic Si phase. However, most of the Si atoms might be dissolved into the ␣-Al phase in a rapidly solidified 5 at.% Si alloy. In other alloys, a small part of Si atoms seems to solidify as Si phase. An exothermic peak which started at around 470 K was observed in each DSC curves of as-rapidly solidified specimens as shown in Fig. 3. The exothermic reaction was due to the precipitation of the Si phase from the supersaturated ␣-Al phase as described in a previous paper [11]. The total heat generated by the exothermic reaction was evaluated from the hatched area in Fig. 3. The area was maximum for Al–12% alloy and minimum for Al–5% Si alloy. Since the area is proportional to the amount of Si precipitates, the degree of the super-saturation increased with increase of the Si content. The microstructure change by annealing at 773 K for 7.2 ks is shown in Fig. 1(c) and (d). Fine and white granules were formed after annealing. Their number and size increased with increase of the Si content. The X-ray diffraction patterns of these specimens are shown in Fig. 4. Sharp and strong diffraction lines of the Si phase were observed for all specimens. It was confirmed that the fine and white granules in Fig. 1(c) and (d) are Si precipitates formed by annealing. The white contrast in Fig. 1(b) was formed by either microsegregation of Si in the primary ␣-Al phase during rapid solidification or the existence of a little Si phase. These results agreed with our previous results [11].
Fig. 3. DSC curves of rapidly solidified Al–Si alloys.
3.2. Crystal structure transition obtained by chemical leaching Examples of the external appearance of the specimens obtained by a NaOH or HCl solution leaching of as-rapidly solidified precursors is shown in Fig. 5(a)–(c). The leached specimens were often too fragile to retain their particulate shape after NaOH solution leaching. On the other hand, their external shape was usually retained after HCl solution leaching. Fig. 6 shows the X-ray diffraction patterns of these leached specimens obtained by using either a NaOH or a HCl solution. There were no lines of the Al(fcc) phase and only broad Si lines were observed. Similar results were also obtained for 12% Si and 5% Si specimens. Thus, it was concluded that the
Fig. 4. X-ray diffraction patterns of annealed Al–Si alloys at 773 K for 7.2 ks.
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Fig. 6. Effect of leaching solution on as-leached structure.
Fig. 5. Change of external appearance by leaching: (a) before leaching; (b) after leaching by NaOH solution; and (c) after leaching by HCl solution.
particles shown in Fig. 5(b) and (c) were composed of the Si phase. The diffraction pattern of pure Si powder was included into this figure for comparison with the leached ones. The diffraction patterns of the specimens leached by a HCl solution were slightly broader than those by NaOH leaching. The peak positions of both specimens were shifted toward lower diffraction angle. The expansion of the lattice constant was 0.3–0.4%. It is an interesting results but the reason for this is not clear.
As silicon is an amphoteric element, the Si phase formed by leaching will also dissolve into the NaOH solution. It is reasonable to expect that no products are formed after leaching. However, the specimen leached by a NaOH solution still retained its shape after leaching as shown in Fig. 5(b). In the case of an annealed precursor, it was impossible to obtain any materials after NaOH solution leaching. The leaching solution was clear after leaching by a NaOH solution. The reason why the Si phase obtained from as-rapidly solidified precursor was stable in a NaOH solution is left as a future work. One possibility is as follows. A very thin silicon dioxide film may have covered the surface and protected its internal part from further attack by the NaOH solution. On the other hand, the solution after leaching was muddy for the HCl solution. In this case, a few residues could be collected from the solution. The X-ray diffraction patterns of the residues showed only sharp Si lines as shown in Fig. 7. In this case, Si free ␣-Al phase was only leached. The Si lines were those of the fine Si precipitates in the annealed precursor. The above-mentioned Si phase was not newly formed but it was obtained by extracting the ␣-Al phase from a Al–Si duplex material. The broad peaks of the Si phase in Figs. 6 and 7 suggest that the grain size of the as-leached specimens was quite fine. From Scherer’s equation, the grain size was evaluated to be about 10–40 nm as shown in Table 1. The evaluated grain size was rather finer for HCl leaching. By assuming each grain as an independent fine particle, a skeletal structure with high specific surface area can be described. Table 2 shows the summary of the specific surface area obtained by the BET method. The specific surface area was quite high and it suggested that a skeletal structure was formed. If the
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Fig. 8(a) and (b) shows microstructures of as-leached 5% Si alloys by TEM. The specimens for TEM observation were prepared by microtome thinning technique. Slightly elongated granules were observed and their size was 30–50 nm. The selected area diffraction patterns of these areas showed only Si diffraction patterns. It was concluded that these granules consisted the Si phase. Thus, fine Si granules were newly formed by leaching of the supersaturated Al–Si precursor. The size was slightly smaller for HCl solution leached specimens. It qualitatively agreed with the result shown in Table 1. Similar results were obtained for other alloys. Fig. 8 shows that the granules are distributed inhomogeneously and formed a network structure like a eutectic cell. These morphologies are different from a
Fig. 7. Effect of heat treatment on X-ray diffraction pattern of 5% Si alloys leached by HCl solution. Table 1 Evaluated particle size from XRD patterns evaluated by Scherer’s equation Alloys
NaOH
HCl
Al–5% Si Al–8% Si Al–10% Si Al–12% Si
42 31 36 38
12 20 25 26
shape of the Si particle is spherical, the specific surface area σ (m2 /g) can be expressed as σ=
6 dρ
(1)
where d and ρ are the diameter (m) and the density (g/m3 ) of the Si sphere, respectively. By taking 65 as a typical value of σ and 2.3 × 10−6 as ρ in Eq. (1), the value of d was evaluated to be about 40 × 10−9 m (=40 nm). This is the same order of the grain size evaluated by X-ray diffraction as shown in Table 1. Although the effect of the shape factor of the particles and the surface smoothness on the specific surface area should be accounted for, it suggest that Si particles of a few tens nanometers were formed by leaching. Table 2 Specific surface area (m2 /g) Alloys
NaOH
HCl
Al–5% Si Al–8% Si Al–10% Si Al–12% Si
64 75 66 47
67 68 64 73
Fig. 8. TEM observation of leached 5% Si specimens by NaOH and HCl solution: (a) leached by NaOH solution; and (b) leached by HCl solution.
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was similarly leached and then directly observed by TEM without any thinning. Fig. 9(a) and (b) shows the effect of microtome thinning on the observed structure morphology. In Fig. 9(a) that was obtained by microtome thinning, a clear network was observed and there were only few granules in the interior between the network. On the other hand, a similar network was observed with a rather dark contrast as shown in Fig. 9(b). The size and shape of the network were similar to those of Fig. 9(a). In this case, fine granules were observed in the interior part of the network. The size and shape of the internal granules were almost the same as those of the network. Similar results were also obtained for other alloys. As the rapid solidification effect on the structure was almost the same in both particle and ribbon specimens, the above morphological difference may be caused during the thinning process. The mounted particles within the plastic resin were cut into slices of about 50 nm in thickness by microtome. Each of the interior Si granules seemed to be kept in place by weak bonding. The granules were too hard and brittle to be cut by a diamond knife. So most of the granules may have dropped from the mounted resin during the cutting process. The slightly bound granules on the network remained after cutting. Although the observed TEM morphology of leached powder precursor was slightly different with the true morphology, the size and shape of the granules on the network seemed to be the same as those of the interior. 4. Conclusions
Fig. 9. Removal of some Si granules in interiors of the network during thin foil preparation for TEM observation by a microtome: (a) foil prepared by microtome; and (b) direct observation without thinning (both samples are prepared from rapidly solidified 5 at.% Si thin ribbon): (a) sample was prepared by microtome thinning and (b) sample was directly prepared from leached ribbon.
previous result [9] where homogeneously distributed granules were observed. To examine the reason why there were no Si granules in the interior of the network, some rapidly solidified ribbon precursors were prepared by a conventional single roller melt spinning method. In the case of a particles specimens of 200 m in diameter, this was too thick for TEM observation. Therefore, a thinning technique by a microtome was used in this experiment. On the other hand, some regions were thin enough for TEM observation in a ribbon specimen due to an inhomogeneous thickness distribution that was generally present in ribbon specimens. The ribbon precursor
(1) Highly saturated ␣-Al precursors were obtained by rapid solidification for Al100−x Six (x = 5–12) alloys. (2) Almost all Al atoms were removed by a HCl solution as well as by a NaOH solution. (3) A granular Si phase was newly formed during leaching. The specific surface area was about 50–70 m2 /g independent of Si content. (4) The X-ray diffraction patterns of leached specimens showed broad lines of the Si phase and its lattice constant was slightly larger than that of the pure Si phase. (5) TEM observation showed that the leached specimens had a skeletal structure composed of slightly elongated particles of the Si phase and quite fine pores. The granule size was about 30–50 nm.
Acknowledgement The authors would like to thank Mr. K. Izumi for the BET measurement.
References [1] F. Matsuda, Chemtech 7 (1977) 306. [2] D.J. Young, M.S. Wainwright, R.B. Anderson, J. Catal. 64 (1980) 116.
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[3] M.L. Bakker, D.J. Young, M.S. Wainwright, J. Mater. Sci. 23 (1988) 3921. [4] I. Ohnaka, I. Yamauchi, M. Itaya, J. Jpn. Inst. Met. 56 (1992) 973. [5] I. Yamauchi, I. Ohnaka, Y. Ohashi, J. Jpn. Inst. Met. 57 (1993) 1385. [6] I. Ohnaka, I. Yamauchi, Mater. Sci. Eng. A 181/182 (1994) 1190. [7] I. Yamauchi, M. Ohmori, I. Ohnaka, Powder Powder Mater. 43 (1996) 573.
[8] I. Yamauchi, T. Mase, T. Kajiwara, M. Saraoka, J. Alloys Compd. 348 (1/2) (2002) 270. [9] I. Yamauchi, H. Kawamura, K. Nakano, T. Tanaka, J. Alloys Compd. 387 (2005) 187. [10] T.B. Massalski (Ed.), Binary Alloy Phase Diagram, vol. 1, ASM, 1986, p. 760. [11] I. Yamauchi, I. Ohnaka, S. Kawamoto, T. Fukusako, Trans. JIM 27 (1986) 187–194. [12] I. Ohnaka, T. Fukusako, H. Tsutsumi, J. Jpn. Inst. Met. 46 (1982) 1095.
Chemical leaching of rapidly solidified Al–Si binary alloys I. Yamauchia,∗ , K. Takaharaa , T. Tanakab , K. Matsubarac a b
Department of Material Science and Engineering, School of Engineering of Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan Department of Mechanical Engineering, Faculty of Engineering, Osaka Sangyo University, 3-1-1 Nakagaito, Daito, Osaka 574-8530, Japan c Samsung Yokohama Research Institute, 2-7 Sugasawa-cho, Tsurumi-ku, Yokohama 230-0027, Japan Received 28 November 2004; received in revised form 27 December 2004; accepted 4 January 2005 Available online 1 February 2005
Abstract Various particulate precursors of Al100−x Six (x = 5–12) alloys were prepared by a rapid solidification process. The rapidly solidified structures of the precursors were examined by XRD, DSC and SEM. Most of Si atoms were dissolved into the ␣-Al(fcc) phase by rapid solidification though the solubility of Si in the ␣-Al phase is negligibly small in conventional solidification. In the case of 5 at.% Si alloy, a single ␣-Al phase was only formed. The amount of the primary Si phase increased with increase of Si content for the alloys beyond 8 at.% Si. Rapid solidification was effective to form super-saturated ␣-Al precursors. These precursors were chemically leached by using a basic solution (NaOH) or a hydrochloric acid (HCl) solution. All Al atoms were removed by a HCl solution as well as a NaOH solution. Granules of the Si phase were newly formed during leaching. The specific surface area was about 50–70 m2 /g independent of Si content. The leaching behavior in both solutions was slightly different. In the case of a NaOH solution, the shape of the precursor often degenerated after leaching. On the other hand, it was retained after leaching by a HCl solution. Fine Si particles precipitated in the ␣-Al phase by annealing of as-rapidly solidified precursors at 773 K for 7.2 × 103 s. In this case, it was difficult to obtain any products by NaOH leaching, but a few of Si particles were obtained by HCl leaching. Precipitated Si particles were dissolved by the NaOH solution. The X-ray diffraction patterns of leached specimens showed broad lines of the Si phase and its lattice constant was slightly larger than that of the pure Si phase. The microstructures of the leached specimens were examined by transmission electron microscopy. It showed that the leached specimens had a skeletal structure composed of slightly elongated particles of the Si phase and quite fine pores. The particle size was about 30–50 nm. It was of comparable order with that evaluated by Scherer’s equation. Leaching of rapidly solidified precursors was effective to form skeletal structures composed of fine Si particles with pores. © 2005 Elsevier B.V. All rights reserved. Keywords: Rapid solidification; Al–Ni–Ag; Skeletal material; Metastable phase
1. Introduction Skeletal materials composed of fine particles and pores can be obtained by applying a chemical leaching process to an aluminum base alloy [1–3]. In this process, aluminum atoms are removed from lattices sites of the alloy by chemical reaction and many vacancy sites are formed. The residual atoms on the lattice sites except Al atoms will rearrange and form a new lattice according to their chemical stability. When the mobility of these atoms is enhanced due to ∗
Corresponding author. Tel.: +81 6 6879 7524; fax: +81 6 6879 7536. E-mail address: [email protected] (I. Yamauchi).
0925-8388/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2005.01.006
the induced vacancies, these atoms can easily move together and may form some nucleus for a new phase. The mobile atoms around the nucleus will arrive on the new phase and make it grow. Thus the new phase becomes larger in size. However, the long distance movement of the atoms will be limited so that some of the atoms will form a new nucleus at another place. Finally, numerous fine particles of the new phase will be formed. The space between each granule will appear as fine pores. This is a principle of skeletal structure formation by chemical leaching. Several authors have been reported on the formation of non-equilibrium skeletal materials by applying chemical leaching to various rapidly solidified precursors [4–9].
I. Yamauchi et al. / Journal of Alloys and Compounds 396 (2005) 302–308
The equilibrium phase diagram [10] of the binary Al–Si system shows that the mutual solubility of these elements is negligibly small at room temperature. In a hypo-eutectic Al–Si alloy, most of the Si atoms are usually solidified as a eutectic Si phase. Silicon is well known as an amphoteric element. Therefore, it was thought as an improper element for chemical leaching. There are only few reports on leaching for this system. A highly supersaturated ␣-Al(fcc) solid solution of Si can be obtained by rapid solidification [11]. The supersaturated ␣-Al solid solution means that the Si atoms substitutionally occupy the lattice sites of the ␣-Al phase. In this case, a new Si phase can be formed by leaching. In this study, we applied rapid solidification to obtain supersaturated Al–Si precursors and examined the structure and the morphology transition by leaching in detail.
2. Experimental procedures 2.1. Rapid solidification for precursors A series of Al100−x Six (x = 5–12) alloys was prepared from Al (99.99 mass%) and high purity Si by melting in Ar atmosphere. The melt was cast into a cylindrical metal mold of 10 mm in diameter. About 20 g of the alloy was remelted in a graphite nozzle with a small bore of 0.5 mm in diameter at its top end by high frequency induction and then the molten jet was ejected into a rotating water layer. Thus, rapidly solidified powders were obtained by using the rotating-wateratomization process [12]. The rotating water velocity was 42 m/s and the atomization conditions were the same for all experiments. The atomized powders were immediately collected with the water and dried in an evacuated chamber.
303
Rapidly solidified thin ribbons were prepared by a conventional single roller process for some alloys. The annealed precursors were obtained by annealing the rapidly solidified powder at 773 K for 7.2 ks in vacuum. The structures of these specimens were examined by differential scanning calorimetry (DSC) at a heating rate of 0.25 K s−1 , SEM and XRD using a graphite monochromatic Cu K␣. For SEM observation, polished specimens were slightly etched by a dix Keller solution for 30 s. 2.2. Chemical leaching About 5 g of each rapidly solidified or annealed specimens was leached by a 20% NaOH basic solution or a 1:1 dilute hydrochloric acid (HCl) solution for 1.2 ks at 333 K. The leached specimen was collected by a Buchner’s funnel and washed several times with distilled water. The structure of the leached specimens was examined by XRD. In the case of the leached powder, the samples for TEM observation were prepared by molding them into plastic resin which then was cut to a quite thin plate by using a microtome with a diamond knife. The thickness of the plate was about 60 nm. On the other hand, the leached ribbon samples were directly examined by TEM without thinning. The specific surface was measured by the BET method. 3. Results and discussion 3.1. Effect of rapid solidification on as-solidified structure Fig. 1(a) and (b) shows as-rapidly solidified structures obtained by SEM of Al–5% Si and Al–12% Si alloys,
Fig. 1. Microstructures of as-solidified and annealed specimens of 5 at.% Si and 12 at.% Si alloys: (a) Al–5% Si RWA; (b) Al–12% Si RWA; (c) Al–5% Si heat treated; and (d) Al–12% Si heat treated.
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Fig. 2. X-ray diffraction patterns of rapidly solidified Al–Si alloys.
respectively. They show quite fine microstructures having white and dark contrasts. Their size increased with increase of the Si content. The X-ray diffraction patterns (Fig. 2) of these specimens show Si phase diffraction lines except for the Al–5% Si alloy where they were negligibly weak. The equilibrium diagram of the Al–Si binary system shows that the solubility of Si into the ␣-Al phase is negligibly small [10]. In conventional solidification processes, most of Si atoms in these hypo-eutectic alloys solidify as the eutectic Si phase. However, most of the Si atoms might be dissolved into the ␣-Al phase in a rapidly solidified 5 at.% Si alloy. In other alloys, a small part of Si atoms seems to solidify as Si phase. An exothermic peak which started at around 470 K was observed in each DSC curves of as-rapidly solidified specimens as shown in Fig. 3. The exothermic reaction was due to the precipitation of the Si phase from the supersaturated ␣-Al phase as described in a previous paper [11]. The total heat generated by the exothermic reaction was evaluated from the hatched area in Fig. 3. The area was maximum for Al–12% alloy and minimum for Al–5% Si alloy. Since the area is proportional to the amount of Si precipitates, the degree of the super-saturation increased with increase of the Si content. The microstructure change by annealing at 773 K for 7.2 ks is shown in Fig. 1(c) and (d). Fine and white granules were formed after annealing. Their number and size increased with increase of the Si content. The X-ray diffraction patterns of these specimens are shown in Fig. 4. Sharp and strong diffraction lines of the Si phase were observed for all specimens. It was confirmed that the fine and white granules in Fig. 1(c) and (d) are Si precipitates formed by annealing. The white contrast in Fig. 1(b) was formed by either microsegregation of Si in the primary ␣-Al phase during rapid solidification or the existence of a little Si phase. These results agreed with our previous results [11].
Fig. 3. DSC curves of rapidly solidified Al–Si alloys.
3.2. Crystal structure transition obtained by chemical leaching Examples of the external appearance of the specimens obtained by a NaOH or HCl solution leaching of as-rapidly solidified precursors is shown in Fig. 5(a)–(c). The leached specimens were often too fragile to retain their particulate shape after NaOH solution leaching. On the other hand, their external shape was usually retained after HCl solution leaching. Fig. 6 shows the X-ray diffraction patterns of these leached specimens obtained by using either a NaOH or a HCl solution. There were no lines of the Al(fcc) phase and only broad Si lines were observed. Similar results were also obtained for 12% Si and 5% Si specimens. Thus, it was concluded that the
Fig. 4. X-ray diffraction patterns of annealed Al–Si alloys at 773 K for 7.2 ks.
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Fig. 6. Effect of leaching solution on as-leached structure.
Fig. 5. Change of external appearance by leaching: (a) before leaching; (b) after leaching by NaOH solution; and (c) after leaching by HCl solution.
particles shown in Fig. 5(b) and (c) were composed of the Si phase. The diffraction pattern of pure Si powder was included into this figure for comparison with the leached ones. The diffraction patterns of the specimens leached by a HCl solution were slightly broader than those by NaOH leaching. The peak positions of both specimens were shifted toward lower diffraction angle. The expansion of the lattice constant was 0.3–0.4%. It is an interesting results but the reason for this is not clear.
As silicon is an amphoteric element, the Si phase formed by leaching will also dissolve into the NaOH solution. It is reasonable to expect that no products are formed after leaching. However, the specimen leached by a NaOH solution still retained its shape after leaching as shown in Fig. 5(b). In the case of an annealed precursor, it was impossible to obtain any materials after NaOH solution leaching. The leaching solution was clear after leaching by a NaOH solution. The reason why the Si phase obtained from as-rapidly solidified precursor was stable in a NaOH solution is left as a future work. One possibility is as follows. A very thin silicon dioxide film may have covered the surface and protected its internal part from further attack by the NaOH solution. On the other hand, the solution after leaching was muddy for the HCl solution. In this case, a few residues could be collected from the solution. The X-ray diffraction patterns of the residues showed only sharp Si lines as shown in Fig. 7. In this case, Si free ␣-Al phase was only leached. The Si lines were those of the fine Si precipitates in the annealed precursor. The above-mentioned Si phase was not newly formed but it was obtained by extracting the ␣-Al phase from a Al–Si duplex material. The broad peaks of the Si phase in Figs. 6 and 7 suggest that the grain size of the as-leached specimens was quite fine. From Scherer’s equation, the grain size was evaluated to be about 10–40 nm as shown in Table 1. The evaluated grain size was rather finer for HCl leaching. By assuming each grain as an independent fine particle, a skeletal structure with high specific surface area can be described. Table 2 shows the summary of the specific surface area obtained by the BET method. The specific surface area was quite high and it suggested that a skeletal structure was formed. If the
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Fig. 8(a) and (b) shows microstructures of as-leached 5% Si alloys by TEM. The specimens for TEM observation were prepared by microtome thinning technique. Slightly elongated granules were observed and their size was 30–50 nm. The selected area diffraction patterns of these areas showed only Si diffraction patterns. It was concluded that these granules consisted the Si phase. Thus, fine Si granules were newly formed by leaching of the supersaturated Al–Si precursor. The size was slightly smaller for HCl solution leached specimens. It qualitatively agreed with the result shown in Table 1. Similar results were obtained for other alloys. Fig. 8 shows that the granules are distributed inhomogeneously and formed a network structure like a eutectic cell. These morphologies are different from a
Fig. 7. Effect of heat treatment on X-ray diffraction pattern of 5% Si alloys leached by HCl solution. Table 1 Evaluated particle size from XRD patterns evaluated by Scherer’s equation Alloys
NaOH
HCl
Al–5% Si Al–8% Si Al–10% Si Al–12% Si
42 31 36 38
12 20 25 26
shape of the Si particle is spherical, the specific surface area σ (m2 /g) can be expressed as σ=
6 dρ
(1)
where d and ρ are the diameter (m) and the density (g/m3 ) of the Si sphere, respectively. By taking 65 as a typical value of σ and 2.3 × 10−6 as ρ in Eq. (1), the value of d was evaluated to be about 40 × 10−9 m (=40 nm). This is the same order of the grain size evaluated by X-ray diffraction as shown in Table 1. Although the effect of the shape factor of the particles and the surface smoothness on the specific surface area should be accounted for, it suggest that Si particles of a few tens nanometers were formed by leaching. Table 2 Specific surface area (m2 /g) Alloys
NaOH
HCl
Al–5% Si Al–8% Si Al–10% Si Al–12% Si
64 75 66 47
67 68 64 73
Fig. 8. TEM observation of leached 5% Si specimens by NaOH and HCl solution: (a) leached by NaOH solution; and (b) leached by HCl solution.
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was similarly leached and then directly observed by TEM without any thinning. Fig. 9(a) and (b) shows the effect of microtome thinning on the observed structure morphology. In Fig. 9(a) that was obtained by microtome thinning, a clear network was observed and there were only few granules in the interior between the network. On the other hand, a similar network was observed with a rather dark contrast as shown in Fig. 9(b). The size and shape of the network were similar to those of Fig. 9(a). In this case, fine granules were observed in the interior part of the network. The size and shape of the internal granules were almost the same as those of the network. Similar results were also obtained for other alloys. As the rapid solidification effect on the structure was almost the same in both particle and ribbon specimens, the above morphological difference may be caused during the thinning process. The mounted particles within the plastic resin were cut into slices of about 50 nm in thickness by microtome. Each of the interior Si granules seemed to be kept in place by weak bonding. The granules were too hard and brittle to be cut by a diamond knife. So most of the granules may have dropped from the mounted resin during the cutting process. The slightly bound granules on the network remained after cutting. Although the observed TEM morphology of leached powder precursor was slightly different with the true morphology, the size and shape of the granules on the network seemed to be the same as those of the interior. 4. Conclusions
Fig. 9. Removal of some Si granules in interiors of the network during thin foil preparation for TEM observation by a microtome: (a) foil prepared by microtome; and (b) direct observation without thinning (both samples are prepared from rapidly solidified 5 at.% Si thin ribbon): (a) sample was prepared by microtome thinning and (b) sample was directly prepared from leached ribbon.
previous result [9] where homogeneously distributed granules were observed. To examine the reason why there were no Si granules in the interior of the network, some rapidly solidified ribbon precursors were prepared by a conventional single roller melt spinning method. In the case of a particles specimens of 200 m in diameter, this was too thick for TEM observation. Therefore, a thinning technique by a microtome was used in this experiment. On the other hand, some regions were thin enough for TEM observation in a ribbon specimen due to an inhomogeneous thickness distribution that was generally present in ribbon specimens. The ribbon precursor
(1) Highly saturated ␣-Al precursors were obtained by rapid solidification for Al100−x Six (x = 5–12) alloys. (2) Almost all Al atoms were removed by a HCl solution as well as by a NaOH solution. (3) A granular Si phase was newly formed during leaching. The specific surface area was about 50–70 m2 /g independent of Si content. (4) The X-ray diffraction patterns of leached specimens showed broad lines of the Si phase and its lattice constant was slightly larger than that of the pure Si phase. (5) TEM observation showed that the leached specimens had a skeletal structure composed of slightly elongated particles of the Si phase and quite fine pores. The granule size was about 30–50 nm.
Acknowledgement The authors would like to thank Mr. K. Izumi for the BET measurement.
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