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Mechanical alloying of amorphous Al–SiO2 powders
Journal of Alloys and Compounds 456 (2008) 139–142
Mechanical alloying of amorphous Al–SiO2 powders Jinhui Wang ∗ Department of Physics, Shanghai Jiao Tong University, Shanghai 200240, PR China Received 13 January 2007; received in revised form 5 February 2007; accepted 5 February 2007 Available online 12 February 2007
Abstract Mechanical alloying (MA) of elemental Al and SiO2 mixtures were performed under argon and air atmospheres. Amorphization has been observed in the MA process under argon atmosphere, while MA under air atmosphere the starting materials are still present. The transformation from crystalline to amorphous state was examined using powder X-ray diffraction, transmission electron microscopy and differential thermal analysis. It is believed that the amorphization is due to the destablization of crystalline phase, induced by refinement of grain size, vitrification of pressure, and loss of long range of order in supersaturated solid solution due to interdiffusion. © 2007 Elsevier B.V. All rights reserved. PACS: 81.20.Ev; 61.43.Er; 61.46.Df; 81.05.Zx Keywords: Mechanical alloying; Amorphization; Aluminum; Silica
1. Introduction
2. Experimental procedure
Mechanical alloying (MA) has been one of the novel nonequilibrium methods to synthesize various interesting solid-state materials, such as amorphous, quasicrystalline and nanocrystalline alloys or composites [1]. Aluminum based composite materials reinforced by various high-temperature resistant ceramic fine particulates, such as SiO2 , SiC, Si3 N4 , TiC and other powders [1–2], have been synthesized by MA procedure, displaying a desirable enhance of mechanical properties at room and elevated temperatures. Meanwhile, MA method is used to fabricate amorphous Al alloys and composites. For instance, the amorphous phases of the Al–22.8 at.% Fe system were obtained by MA with an energy dose of 130 kJ/g [3], and ball milling Al–C mixtures can also lead to an amorphous state in the composition range of 28–50 at.% C [4]. In this work, it is reported that ball milling the mixtures of SiO2 and Al powders results in an amorphous phase after milling for 150 h (h), instead of a mixture of Si and ␣-Al2 O3 as mentioned in the literature [5]. It is found that the atmospheres during milling have a greatly impact on the final product.
Elemental Al (∼75 m) and SiO2 (∼45 m) mixtures with respect to molar ratio 1:1 were milled in a planetary ball mill at 200 rotations per minute under 0.15 MPa purified argon or air atmosphere. Vials and balls are made of harden stainless steel, and the volume of each vial is 100 ml, which contains 44 balls (10 mm in diameter). The ball-to-weight ratio is 10:1. The milling process was interrupted at different times, i.e., 0.3, 5, 10, 50, 100 and 150 h, and a small amount of powders was removed from the vial in a glove box or in air for characterization. Parts of as-prepared samples were annealed in a tube furnace under a flow of high-purified argon atmosphere. The temperatures are 1108 and 1273 K, respectively. The duration of annealing is half an hour, and then furnace cooling. We used X-ray diffraction (XRD, D/Dmax-RA, Rigaku) with Cu K␣ radiation (λ = 0.154 nm) and graphite filter to analyze the microstructure of as-milled samples. Scanning electron microscopy (SEM, S-2150, Hitachi) and transmission electron microscopy (TEM, JEM-2010, JEOL) were employed to observe particulate morphology and electron diffraction pattern. The thermal properties were examined using differential thermal analyzer (DTA, DTA 1700, PerkinElmer). The procedure was performed under a flow of nitrogen atmosphere and heating rate is 20 ◦ C/min.
∗
Tel.: +86 21 54742907. E-mail address: [email protected].
0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.02.027
3. Results and discussion Fig. 1 shows room-temperature XRD patterns at different milling times. The starting materials were milled under argon atmosphere. Obviously, with increasing milling time from 0.3 to 100 h, the peaks related to SiO2 (quartz phase) and Al still exist. However, after milling for 150 h, the intensity of reflections
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J. Wang / Journal of Alloys and Compounds 456 (2008) 139–142
Fig. 1. XRD patterns for Al–SiO2 after different milling times under argon atmosphere.
greatly decreases. The pattern displays a broad diffraction maximum without detectable sharp diffraction lines, indicating the formation of an amorphous phase. In comparison, when the milling process is carried out under air environment, as depicted in Fig. 2, even after milling for 150 h, the corresponding XRD patterns still reveal distinguishable peaks, which can be indexed with SiO2 and Al phases, respectively. Moreover, the trace of Al oxides or elemental Si cannot be detected in both milling atmospheres, probably implying that the redox reaction 4Al + 3SiO2 → 3Si + 2Al2 O3 [5] does not take place. Furthermore, it can be seen that the initial sharp diffraction lines are considerably broadened with increasing milling time. The result is more distinct in the inset of Fig. 3. It is mainly due to the refinement of the crystallite size. Meanwhile, an increase
Fig. 2. XRD patterns for Al–SiO2 after different milling times under air atmosphere.
Fig. 3. The average grain size D of Al and SiO2 crystallites as a function of milling time. Inset shows the SiO2 (1 0 1) reflection profiles at different milling times.
of atomic-level strain caused by repeated collisions and coldwelds during milling plays a role on line broadening to certain extent. Roughly, the average crystallite size D can be estimated by Scherrer equation [6]: D=
0.91λ B(2θ) cos θ
where B(2θ) is the full width at half maximum (in radians) of the diffraction peak, corrected for Cu K␣ and instrumental broadening. λ is the wave length of incident X-rays and θ is the diffraction angle. Fig. 3 displays that D of SiO2 and Al crystallites rapidly decreases in the primarily milling stage, and then almost remains unchanged as the milling time is larger than 50 h. Moreover, the D of SiO2 and Al crystallites in the samples milled under argon atmosphere are smaller than that milled under air environment, respectively. For instance, after milling for 50 h, the D of SiO2 is about 28 and 50 nm under argon and air atmosphere, respectively. In the same time, the D of Al is about 16 and 33 nm. It is also found that in the inset of Fig. 3, the SiO2 (1 0 1) reflection peaks slightly shift to lower diffraction angle side (about 0.04◦ ) as the milling time is greater than 10 h. In contrast, the (1 1 1) reflections of Al shift to larger diffraction angle side (about 0.02◦ ). The shifts of diffraction peaks may be attributed to the interdiffusion of heterogeneous atoms. The particulate morphology of the Al–SiO2 mixtures milled for 150 h under argon atmosphere is shown in Fig. 4a. The ball milling for 150 h forms irregular particles with grain size up to 5 m. We also observed micrograph of a fine particle by means of TEM, as indicated in Fig. 4b. It is obvious that the particle (about 400 nm in size) is composed of several smaller grains. Fig. 4c shows the corresponding select area diffraction pattern (SADP) of the particle. A broad diffraction halo without continuous rings of crystallites, displays a typical feature of an amorphous phase, and thus confirming the XRD results mentioned above.
J. Wang / Journal of Alloys and Compounds 456 (2008) 139–142
141
Fig. 4. SEM image of particles after 150 h of milling under argon atmosphere (a), TEM micrograph of a fine particle (b), and selected area diffraction pattern (c).
Fig. 5 reveals DTA traces of the sample after 150 h of milling under argon atmosphere. The sample has a broad exothermic peak at temperatures ranging from 693 to 1093 K, and a narrow exothermic peak from 1123 to 1273 K. The temperature of the maximum value of peak and enthalpy (the area under the exothermic peak) are 875 K, 89 Cal/g and 1205 K, 35 Cal/g, respectively. In order to understand the origin of these exothermic peaks, a small amount of as-milled amorphous sample was annealed at 1108 and 1273 K for half an hour, respectively. The powders were fired in a flow of purified argon gas to avoid oxidization. The XRD patterns of the annealed powders, as shown in Fig. 6, indicates that after annealing at the temperature of 1108 K, a small amount of Al2 SiO5 and SiO2 phases emerges, besides amorphous component. With increasing annealing temperature up to 1273 K, the volume percent of crystalline phases
Fig. 5. DTA trace of powders after 150 h of milling under argon atmosphere.
become greater. Owing to repeated welding and fracturing during milling, numerous faults and dislocations are produced in the particles, the broad exothermic peak at low temperature stage may mainly originate from strain relaxation, defect recovery and grain growth. The broad exothermic peak also exists in the DSC scan for partially amorphous silicon and other nanocrystalline powders [7,8]. The contribution arising from the crystallization of small amount of amorphous powders should be less if duration of annealing reduces. The exothermic peak at the elevated temperature may be almost attributed to the crystallization of Al–SiO2 amorphous phases. The mechanism of crystal-to-amorphous transition has been discussed in the literature [1], several models, such as local melting of materials followed by rapid solidification [9], solid-state reaction similar to the thin film diffusion couples [10], pressure-
Fig. 6. XRD patterns of annealed samples.
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J. Wang / Journal of Alloys and Compounds 456 (2008) 139–142
induced and crystallite-refinement-induced amorphization [7], have been proposed. However, Schwarz et al. reported that the peak temperature increase of powders during milling is only about 38 K (above the average processing temperature), well below the melting temperature of the starting materials [11]. Thus, the local melting model seems not true. As for Al–SiO2 system, the amorphization after milling for 150 h under argon atmosphere is due to the destabilization of crystalline phase. In one approach, during MA, the accumulation of structure defects, such as vacancies, dislocations, grain boundaries and distortion of lattice, considerably enhances the Gibbs free energy with decreasing grain size. If the free energy of crystal increases beyond the amorphous state, a crystal to amorphous transition can occur [12]. The refinement of grain will elevates the free energy due to the increase of grain boundary density. For example, Fechet et al. reported that the excess enthalpy of about 30% of the heat of fusion is stored in the nanocrystalline metals after high-energy ball milling [13]. The critical grain sizes for solid-state amorphization of various pure metalloids and metals, have been successfully predicted by Zhang et al. [14]. In addition, the pressure of about 109 Pa exists between the contacting surfaces during collisions for ball to ball and ball to vial. The powders are subjected to rapidly loading and unloading rates, may have not enough time turning to equilibrium state, thus the transformation to metastably amorphous state can be expected [1,7]. In the other approach, the shifts of diffraction peak in the inset of Fig. 3 may indicate that Al atoms diffuse into silica matrix. The elastic mismatch energy stored in a supersaturate solid solution will increase when the long-range chemical order is destroyed. Once the mismatch energy is higher than a critical value, the transformation for crystal to amorphous phase also possibly takes place [15]. When ball milling was performed under air atmosphere, the case seems different. Although the trace of Al oxide has not been evidenced in the XRD patterns, very fine Aluminum oxide particles will emerge due to the high oxygen affinity to aluminum. Generally, aluminum oxide layer is always present on the surface of Al particles, and incessantly produces during milling in air atmosphere. The introduction of Al2 O3 particles into Al matrix may form so-called Al-based metal matrix composites. This experimental process is similar to the low cost bonding of technique of hot-cold rolling, though aluminum oxide is produced by anodizing of pure aluminum using sulfuric acid [16]. The Al-based particulate reinforced metal matrix composites are an important class of engineered materials and show substantial advantages such as increased stiffness, strength, creep and wear resistance, and have superior performance at elevated temperatures. For example, the hardness of aluminum matrix composites increases from 634.1 to 1401.4 MPa when the volume fraction of Al2 O3 particles changes from 0 to 30% [17]. With the addition of small amount (∼1 vol.%) of nanometric alumina particulate in the Al, the yield strength rises approximately
from 70 to 120 GPa [18]. The reinforcement can be explained by several models [19], such as the shear lag theory assuming load transfer between the particle and the matrix at the particle–matrix interface, dislocation-particles interactions and grain boundaries strengthening. Therefore, it appears that the obtained Al2 O3 particles during milling in air atmosphere will disperse into the milling mixtures, and then reinforce its strength. Under the applied stress, Al2 O3 particulates and increasing amount of grain boundaries (due to significant grain refinement) acts as obstacles to the dislocation movement. The grain sizes of SiO2 and Al particles are difficult to decrease during MA in air than that in argon atmosphere. The free energy arising from the accumulation of grain boundaries and other structure defects may be lower than that of amorphous phase because the grain sizes of particles are larger than the critical size required for solid-state amorphization, thus, the destabilization of crystalline phase will not occur. 4. Conclusion In short, it is found that ball milling Al–SiO2 blends under argon atmosphere results in an amorphous phase, which was confirmed by XRD, SADP and DTA measurements. The amorphization is possibly due to the destabilization of crystalline phase, induced by a combination of factors including refinement of grain size, high pressure exerted to powders during repeated collision, and elastic mismatched energy stored in the supersaturated solid solution. However, which mechanism dominated the transformation of crystal to amorphous phase is still left to investigate. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1, and therein references. H.L. Rizkalla, A. Abdulwahed, J. Mater. Proc. Tech. 56 (1996) 398. V.I. Fadeeva, A.V. Leonov, Mater. Sci. Eng. A 206 (1996) 90. N.Q. Wu, J.M. Wu, G.X. Wang, Z.Z. Li, J. Alloys Compd. 260 (1997) 121. P. Matteazzi, G. Le Ca¨er, J. Am. Ceram. Soc. 75 (1992) 2749. P. Scherrer, Gott, Nachr. 2 (1918) 98. T.D. Shen, C.C. Koch, T.L. McCormick, R.J. Nemanich, J.Y. Huang, J.G. Huang, J. Mater. Res. 10 (1995) 139. J. Eckert, J.C. Holzer, C.E. Krill III, W.L. Johnson, J. Mater. Res. 7 (1992) 1751. A.E. Yermakov, E.E. Yurchikov, V.A. Barinov, Fiz. Met. Metalloved 52 (1981) 1184. R.B. Schwarz, R.R. Petrich, J. Non-cryst. Solids 76 (1985) 281. R.B. Schwarz, C.C. Koch, Appl. Phys. Lett. 49 (1986) 146. J.S.C. Jang, C.C. Koch, J. Mater. Res. 5 (1990) 498. H.L. Fecht, E. Hellstern, Z. Fu, W.L. Johnson, Metall. Trans. 21A (1990) 2333. Y.H. Zhao, J. Non-Cryst. Solids 352 (2006) 5578. A.W. Weeber, H. Bakker, Physica B 153 (1988) 93. X.X. Yu, W.B. Lee, Compos.: Part A 31 (2000) 245. H. Huang, M.B. Bush, Mater. Sci. Eng. A 232 (1997) 63. Y. Kang, S.L. Chan, Mater. Chem. Phys. 85 (2004) 438. G.M. Owolabi, A.G. Odeshi, M.N. K. Singh, M.N. Bassim, Mater. Sci. Eng. A, in press.
Mechanical alloying of amorphous Al–SiO2 powders Jinhui Wang ∗ Department of Physics, Shanghai Jiao Tong University, Shanghai 200240, PR China Received 13 January 2007; received in revised form 5 February 2007; accepted 5 February 2007 Available online 12 February 2007
Abstract Mechanical alloying (MA) of elemental Al and SiO2 mixtures were performed under argon and air atmospheres. Amorphization has been observed in the MA process under argon atmosphere, while MA under air atmosphere the starting materials are still present. The transformation from crystalline to amorphous state was examined using powder X-ray diffraction, transmission electron microscopy and differential thermal analysis. It is believed that the amorphization is due to the destablization of crystalline phase, induced by refinement of grain size, vitrification of pressure, and loss of long range of order in supersaturated solid solution due to interdiffusion. © 2007 Elsevier B.V. All rights reserved. PACS: 81.20.Ev; 61.43.Er; 61.46.Df; 81.05.Zx Keywords: Mechanical alloying; Amorphization; Aluminum; Silica
1. Introduction
2. Experimental procedure
Mechanical alloying (MA) has been one of the novel nonequilibrium methods to synthesize various interesting solid-state materials, such as amorphous, quasicrystalline and nanocrystalline alloys or composites [1]. Aluminum based composite materials reinforced by various high-temperature resistant ceramic fine particulates, such as SiO2 , SiC, Si3 N4 , TiC and other powders [1–2], have been synthesized by MA procedure, displaying a desirable enhance of mechanical properties at room and elevated temperatures. Meanwhile, MA method is used to fabricate amorphous Al alloys and composites. For instance, the amorphous phases of the Al–22.8 at.% Fe system were obtained by MA with an energy dose of 130 kJ/g [3], and ball milling Al–C mixtures can also lead to an amorphous state in the composition range of 28–50 at.% C [4]. In this work, it is reported that ball milling the mixtures of SiO2 and Al powders results in an amorphous phase after milling for 150 h (h), instead of a mixture of Si and ␣-Al2 O3 as mentioned in the literature [5]. It is found that the atmospheres during milling have a greatly impact on the final product.
Elemental Al (∼75 m) and SiO2 (∼45 m) mixtures with respect to molar ratio 1:1 were milled in a planetary ball mill at 200 rotations per minute under 0.15 MPa purified argon or air atmosphere. Vials and balls are made of harden stainless steel, and the volume of each vial is 100 ml, which contains 44 balls (10 mm in diameter). The ball-to-weight ratio is 10:1. The milling process was interrupted at different times, i.e., 0.3, 5, 10, 50, 100 and 150 h, and a small amount of powders was removed from the vial in a glove box or in air for characterization. Parts of as-prepared samples were annealed in a tube furnace under a flow of high-purified argon atmosphere. The temperatures are 1108 and 1273 K, respectively. The duration of annealing is half an hour, and then furnace cooling. We used X-ray diffraction (XRD, D/Dmax-RA, Rigaku) with Cu K␣ radiation (λ = 0.154 nm) and graphite filter to analyze the microstructure of as-milled samples. Scanning electron microscopy (SEM, S-2150, Hitachi) and transmission electron microscopy (TEM, JEM-2010, JEOL) were employed to observe particulate morphology and electron diffraction pattern. The thermal properties were examined using differential thermal analyzer (DTA, DTA 1700, PerkinElmer). The procedure was performed under a flow of nitrogen atmosphere and heating rate is 20 ◦ C/min.
∗
Tel.: +86 21 54742907. E-mail address: [email protected].
0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.02.027
3. Results and discussion Fig. 1 shows room-temperature XRD patterns at different milling times. The starting materials were milled under argon atmosphere. Obviously, with increasing milling time from 0.3 to 100 h, the peaks related to SiO2 (quartz phase) and Al still exist. However, after milling for 150 h, the intensity of reflections
140
J. Wang / Journal of Alloys and Compounds 456 (2008) 139–142
Fig. 1. XRD patterns for Al–SiO2 after different milling times under argon atmosphere.
greatly decreases. The pattern displays a broad diffraction maximum without detectable sharp diffraction lines, indicating the formation of an amorphous phase. In comparison, when the milling process is carried out under air environment, as depicted in Fig. 2, even after milling for 150 h, the corresponding XRD patterns still reveal distinguishable peaks, which can be indexed with SiO2 and Al phases, respectively. Moreover, the trace of Al oxides or elemental Si cannot be detected in both milling atmospheres, probably implying that the redox reaction 4Al + 3SiO2 → 3Si + 2Al2 O3 [5] does not take place. Furthermore, it can be seen that the initial sharp diffraction lines are considerably broadened with increasing milling time. The result is more distinct in the inset of Fig. 3. It is mainly due to the refinement of the crystallite size. Meanwhile, an increase
Fig. 2. XRD patterns for Al–SiO2 after different milling times under air atmosphere.
Fig. 3. The average grain size D of Al and SiO2 crystallites as a function of milling time. Inset shows the SiO2 (1 0 1) reflection profiles at different milling times.
of atomic-level strain caused by repeated collisions and coldwelds during milling plays a role on line broadening to certain extent. Roughly, the average crystallite size D can be estimated by Scherrer equation [6]: D=
0.91λ B(2θ) cos θ
where B(2θ) is the full width at half maximum (in radians) of the diffraction peak, corrected for Cu K␣ and instrumental broadening. λ is the wave length of incident X-rays and θ is the diffraction angle. Fig. 3 displays that D of SiO2 and Al crystallites rapidly decreases in the primarily milling stage, and then almost remains unchanged as the milling time is larger than 50 h. Moreover, the D of SiO2 and Al crystallites in the samples milled under argon atmosphere are smaller than that milled under air environment, respectively. For instance, after milling for 50 h, the D of SiO2 is about 28 and 50 nm under argon and air atmosphere, respectively. In the same time, the D of Al is about 16 and 33 nm. It is also found that in the inset of Fig. 3, the SiO2 (1 0 1) reflection peaks slightly shift to lower diffraction angle side (about 0.04◦ ) as the milling time is greater than 10 h. In contrast, the (1 1 1) reflections of Al shift to larger diffraction angle side (about 0.02◦ ). The shifts of diffraction peaks may be attributed to the interdiffusion of heterogeneous atoms. The particulate morphology of the Al–SiO2 mixtures milled for 150 h under argon atmosphere is shown in Fig. 4a. The ball milling for 150 h forms irregular particles with grain size up to 5 m. We also observed micrograph of a fine particle by means of TEM, as indicated in Fig. 4b. It is obvious that the particle (about 400 nm in size) is composed of several smaller grains. Fig. 4c shows the corresponding select area diffraction pattern (SADP) of the particle. A broad diffraction halo without continuous rings of crystallites, displays a typical feature of an amorphous phase, and thus confirming the XRD results mentioned above.
J. Wang / Journal of Alloys and Compounds 456 (2008) 139–142
141
Fig. 4. SEM image of particles after 150 h of milling under argon atmosphere (a), TEM micrograph of a fine particle (b), and selected area diffraction pattern (c).
Fig. 5 reveals DTA traces of the sample after 150 h of milling under argon atmosphere. The sample has a broad exothermic peak at temperatures ranging from 693 to 1093 K, and a narrow exothermic peak from 1123 to 1273 K. The temperature of the maximum value of peak and enthalpy (the area under the exothermic peak) are 875 K, 89 Cal/g and 1205 K, 35 Cal/g, respectively. In order to understand the origin of these exothermic peaks, a small amount of as-milled amorphous sample was annealed at 1108 and 1273 K for half an hour, respectively. The powders were fired in a flow of purified argon gas to avoid oxidization. The XRD patterns of the annealed powders, as shown in Fig. 6, indicates that after annealing at the temperature of 1108 K, a small amount of Al2 SiO5 and SiO2 phases emerges, besides amorphous component. With increasing annealing temperature up to 1273 K, the volume percent of crystalline phases
Fig. 5. DTA trace of powders after 150 h of milling under argon atmosphere.
become greater. Owing to repeated welding and fracturing during milling, numerous faults and dislocations are produced in the particles, the broad exothermic peak at low temperature stage may mainly originate from strain relaxation, defect recovery and grain growth. The broad exothermic peak also exists in the DSC scan for partially amorphous silicon and other nanocrystalline powders [7,8]. The contribution arising from the crystallization of small amount of amorphous powders should be less if duration of annealing reduces. The exothermic peak at the elevated temperature may be almost attributed to the crystallization of Al–SiO2 amorphous phases. The mechanism of crystal-to-amorphous transition has been discussed in the literature [1], several models, such as local melting of materials followed by rapid solidification [9], solid-state reaction similar to the thin film diffusion couples [10], pressure-
Fig. 6. XRD patterns of annealed samples.
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J. Wang / Journal of Alloys and Compounds 456 (2008) 139–142
induced and crystallite-refinement-induced amorphization [7], have been proposed. However, Schwarz et al. reported that the peak temperature increase of powders during milling is only about 38 K (above the average processing temperature), well below the melting temperature of the starting materials [11]. Thus, the local melting model seems not true. As for Al–SiO2 system, the amorphization after milling for 150 h under argon atmosphere is due to the destabilization of crystalline phase. In one approach, during MA, the accumulation of structure defects, such as vacancies, dislocations, grain boundaries and distortion of lattice, considerably enhances the Gibbs free energy with decreasing grain size. If the free energy of crystal increases beyond the amorphous state, a crystal to amorphous transition can occur [12]. The refinement of grain will elevates the free energy due to the increase of grain boundary density. For example, Fechet et al. reported that the excess enthalpy of about 30% of the heat of fusion is stored in the nanocrystalline metals after high-energy ball milling [13]. The critical grain sizes for solid-state amorphization of various pure metalloids and metals, have been successfully predicted by Zhang et al. [14]. In addition, the pressure of about 109 Pa exists between the contacting surfaces during collisions for ball to ball and ball to vial. The powders are subjected to rapidly loading and unloading rates, may have not enough time turning to equilibrium state, thus the transformation to metastably amorphous state can be expected [1,7]. In the other approach, the shifts of diffraction peak in the inset of Fig. 3 may indicate that Al atoms diffuse into silica matrix. The elastic mismatch energy stored in a supersaturate solid solution will increase when the long-range chemical order is destroyed. Once the mismatch energy is higher than a critical value, the transformation for crystal to amorphous phase also possibly takes place [15]. When ball milling was performed under air atmosphere, the case seems different. Although the trace of Al oxide has not been evidenced in the XRD patterns, very fine Aluminum oxide particles will emerge due to the high oxygen affinity to aluminum. Generally, aluminum oxide layer is always present on the surface of Al particles, and incessantly produces during milling in air atmosphere. The introduction of Al2 O3 particles into Al matrix may form so-called Al-based metal matrix composites. This experimental process is similar to the low cost bonding of technique of hot-cold rolling, though aluminum oxide is produced by anodizing of pure aluminum using sulfuric acid [16]. The Al-based particulate reinforced metal matrix composites are an important class of engineered materials and show substantial advantages such as increased stiffness, strength, creep and wear resistance, and have superior performance at elevated temperatures. For example, the hardness of aluminum matrix composites increases from 634.1 to 1401.4 MPa when the volume fraction of Al2 O3 particles changes from 0 to 30% [17]. With the addition of small amount (∼1 vol.%) of nanometric alumina particulate in the Al, the yield strength rises approximately
from 70 to 120 GPa [18]. The reinforcement can be explained by several models [19], such as the shear lag theory assuming load transfer between the particle and the matrix at the particle–matrix interface, dislocation-particles interactions and grain boundaries strengthening. Therefore, it appears that the obtained Al2 O3 particles during milling in air atmosphere will disperse into the milling mixtures, and then reinforce its strength. Under the applied stress, Al2 O3 particulates and increasing amount of grain boundaries (due to significant grain refinement) acts as obstacles to the dislocation movement. The grain sizes of SiO2 and Al particles are difficult to decrease during MA in air than that in argon atmosphere. The free energy arising from the accumulation of grain boundaries and other structure defects may be lower than that of amorphous phase because the grain sizes of particles are larger than the critical size required for solid-state amorphization, thus, the destabilization of crystalline phase will not occur. 4. Conclusion In short, it is found that ball milling Al–SiO2 blends under argon atmosphere results in an amorphous phase, which was confirmed by XRD, SADP and DTA measurements. The amorphization is possibly due to the destabilization of crystalline phase, induced by a combination of factors including refinement of grain size, high pressure exerted to powders during repeated collision, and elastic mismatched energy stored in the supersaturated solid solution. However, which mechanism dominated the transformation of crystal to amorphous phase is still left to investigate. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1, and therein references. H.L. Rizkalla, A. Abdulwahed, J. Mater. Proc. Tech. 56 (1996) 398. V.I. Fadeeva, A.V. Leonov, Mater. Sci. Eng. A 206 (1996) 90. N.Q. Wu, J.M. Wu, G.X. Wang, Z.Z. Li, J. Alloys Compd. 260 (1997) 121. P. Matteazzi, G. Le Ca¨er, J. Am. Ceram. Soc. 75 (1992) 2749. P. Scherrer, Gott, Nachr. 2 (1918) 98. T.D. Shen, C.C. Koch, T.L. McCormick, R.J. Nemanich, J.Y. Huang, J.G. Huang, J. Mater. Res. 10 (1995) 139. J. Eckert, J.C. Holzer, C.E. Krill III, W.L. Johnson, J. Mater. Res. 7 (1992) 1751. A.E. Yermakov, E.E. Yurchikov, V.A. Barinov, Fiz. Met. Metalloved 52 (1981) 1184. R.B. Schwarz, R.R. Petrich, J. Non-cryst. Solids 76 (1985) 281. R.B. Schwarz, C.C. Koch, Appl. Phys. Lett. 49 (1986) 146. J.S.C. Jang, C.C. Koch, J. Mater. Res. 5 (1990) 498. H.L. Fecht, E. Hellstern, Z. Fu, W.L. Johnson, Metall. Trans. 21A (1990) 2333. Y.H. Zhao, J. Non-Cryst. Solids 352 (2006) 5578. A.W. Weeber, H. Bakker, Physica B 153 (1988) 93. X.X. Yu, W.B. Lee, Compos.: Part A 31 (2000) 245. H. Huang, M.B. Bush, Mater. Sci. Eng. A 232 (1997) 63. Y. Kang, S.L. Chan, Mater. Chem. Phys. 85 (2004) 438. G.M. Owolabi, A.G. Odeshi, M.N. K. Singh, M.N. Bassim, Mater. Sci. Eng. A, in press.