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Chemical leaching of non-equilibrium Al(Ni-Fe) alloy powders produced by rod milling
Journal of Alloys and Compounds 456 (2008) 72–78
Chemical leaching of non-equilibrium Al(Ni-Fe) alloy powders produced by rod milling Hyun-Goo Kim a,∗ , Wha-Nam Myung b , K. Sumiyama c , K. Suzuki d a
Department of Physics Education, Chosun University, Gwangju 501-759, South Korea Department of Physics, Chonnam National University, Gwangju 500-757, South Korea c Department of Materials Science and Engineering, Nagoya Institute of Technology, Nagoya, Japan d Institute for Advanced Materials Research in Sendai, Tokai University in Tokyo, Japan b
Received 14 February 2006; received in revised form 1 February 2007; accepted 1 February 2007 Available online 6 February 2007
Abstract In this paper, chemical leaching of non-equilibrium Al0.6 (Ni0.5 Fe0.5 )0.4 powder by rod milling, is reported. X-ray diffractometry (XRD), transmission electron microscopy (TEM), differential scanning calorimetry (DSC), scanning electron microscopy (SEM), M¨ossbauer spectroscopy, and vibrating sample magnetometry (VSM) were used to characterize the rod-milled and leached specimens. After 200 h of milling, the AlNi bcc phase was detected in the XRD patterns. The crystallite size, calculated using Williamson–Hall and Scherrer methods, decreased rapidly for tm , less than 200 h, while the behavior of internal strain increases is the opposite. The leached specimens transformed to a ferromagnetic fcc Ni phase at above 400 ◦ C of annealing for 1 h. The peak temperature and the width of DSC curves decreased with increasing leaching temperature. The relative intensity of the magnetic lines of sextet in M¨ossbauer spectra is reduced and the spectrum collapses into a distribution of doublet as the milling time increased. When the specimen cooled from 750 ◦ C, the curve sharply increased to approximately 595, 580, and 542 ◦ C for leached specimens in a 30 wt.% KOH solution at room temperature, at 55 and at 85 ◦ C, respectively, indicating that the bcc phase has been transformed to a fcc phase. © 2007 Elsevier B.V. All rights reserved. Keywords: Rod milling; Non-equilibrium Al(Ni-Fe); Chemical leaching
1. Introduction In anticipation of unique or improved properties, very different from those of conventional coarse crystalline materials, considerable interest in producing and studying nanocrystalline materials has been created [1–4]. Recently, a technique of mechanically alloying pure elements has been proven to be one of the most effective methods for preparing nanocrystalline materials [5,6]. Nanometer order powders have found applications in catalysis, electromagnetic shielding, magnetic recording, refrigeration and in the processing of advanced engineering materials. In particular, nanoscale 3d-transition metal (Fe, Co and Ni) particles are expected to display unique magnetic and catalytic properties [7]. A combination of rod milling (RM) and chemical leaching has been used to prepare nanoscale metastable materials,
which represent many interesting properties [8–16]. The leaching treatment of Al–Ni alloys by an alkaline solution leads to removal of the major portion of the Al and retention of a “skeleton Ni” catalyst. Similarly, a nanocrystalline bcc Co phase is obtained by leaching out Al atoms from mechanically alloyed bcc Co40 Al60 [12]. After the non-equilibrium Al(Ni-Fe) powder had been leached and annealed, the bcc phase is transformed into the fcc phase [8]. Yamauchi et al. reported that it formed the metastable phase by alkaline leaching of Al of rapidly solidified Al3 (Nix Fe1−x ) alloys [17]. In this paper, the idea of leaching Al from a rod-milled Al0.6 (Ni0.5 Fe0.5 )0.4 alloy is used in a basic solution to obtain a nanocrystalline bcc Ni-Fe phase and to study its properties using XRD, DSC, SEM, TEM, M¨ossbauer spectroscope and vibrating sample magnetometer (VSM). 2. Experimental
∗
Corresponding author. Tel.: +82 62 230 7376. E-mail address: [email protected] (H.-G. Kim).
0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.02.002
Pure elemental powders of Al, Ni, and Fe were mixed in an argon glove box to provide a composition of Al0.6 (Ni0.5 Fe0.5 )0.4 , the nominal composition being
H.-G. Kim et al. / Journal of Alloys and Compounds 456 (2008) 72–78 expressed in atomic percent. The rod-to-powder weight ratio was 28:1. The stainless steel (SUS 304) vial together with SUS 304 rods (10 mm in diameter) was rotated with an angular velocity of 100 rev/min under purified argon gas. The vial was opened after 100, 200, 300 and 400 h (h, milling times (tm )) of processing, in an argon atmosphere. A small amount of sample powder was removed from the vial in a glove box for analysis and observation. The rod-milled powders (400 h of milling) were leached in a 30 wt.% KOH solution at approximately 20 ◦ C (room temperature, RT, L1), at 55 ◦ C (L2) and at 85 ◦ C (L3) to remove the Al atoms. After the leaching treatment, the powders were washed with distilled water until the pH value became the same as the distilled water. Then the powders were immersed in alcohol and dried in a glove box under argon atmosphere. This process is the conventional method for obtaining Raney Ni catalysts [18]. The leached L2 alloy powders were annealed in evacuated quartz tubes at 400 and 600 ◦ C for 1 h. The XRD was performed at room temperature by using Cu K␣ radiation. The scan speed was 1.5◦ min−1 and the sampling width was 0.05◦ . The heat treatment up to crystallization was monitored by DSC at a heating rate of 10 K min−1 in a flow of purified argon gas. The microstructure of the rod-milled powder and the leached specimens before and after annealing were observed by TEM using a 200 kV. SEM observations allowing for characterization of powder particles during RM have been performed as well. The M¨ossbauer spectroscopy was performed on all samples, at room temperature and at 14.4 keV with a Rh/Co57 source. A conventional VSM was employed to measure the magnetization from 20 up to 750 ◦ C in magnetic fields of 5 kOe. The composition of the alloys was determined using inductively coupled plasma-emission spectrometry (ICP).
3. Result and discussion Fig. 1 presents the change in the X-ray diffraction patterns of the powders with tm : 0, 100, 200, 300 and 400 h. Directly after milling, a significant decrease in the relative intensities of the Ni, the Fe and the Al peaks were detected, and after 200 h, only the bcc-type AlNi phase was detected in the XRD patterns. However, the XRD peak for 2θ = 44.65◦ at 0 h of milling is asymmetric after 100 h of RM and slightly shifted toward
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Fig. 1. X-ray diffraction patterns of rod-milled Al0.6 (Ni0.5 Fe0.5 ) 0.4 for different milling times.
the lower 2θ side gradually with increasing tm , indicating that Fe atoms are arranged into the bcc-type AlNi phase. The lattice constants of peak for 2θ = 44.65◦ are 2.862, 2.862, 2.87999, 2.87999, and 2.88925 nm for 0, 100, 200, 300, and 400 h of milling, respectively, gradually increasing with increasing tm (the inset of Fig. 1). The value of the lattice constants for 400 h was similar with that of 2.888 nm for AlNi. The evolutions of crystallite size and internal strain with the tm were evaluated using Williamson–Hall [19–21] and Scherrer methods [22]. Fig. 2(a) represents a βcos θ/λ versus sin θ/λ plot for powder with 100 h of milling. The slope of the straight line gives the amount of internal strain, which comes out to be 4.28 × 10−3 . The intercept on the βcos θ/λ axis gives the crystallite size as 11.15 nm. Fig. 2(b) shows the evolution of crystallite
Fig. 2. (a) Williamson–Hall plot showing βcos θ/λ vs. sin θ/λ for the powder of 100 h for milling. (b) Evolution of crystallite size and internal strain with the milling time obtained from Williamson–Hall method for the Al0.6 (Ni0.5 Fe0.5 )0.4 . (c) Crystallite size obtained applying the Scherrer formula to different peaks of XRD patterns for the powder with 100 h of milling time. (d) Evolution of crystallite size with the milling time obtained from Scherrer method.
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size with tm . For tm , less than 200 h crystallite size decreases rapidly, while the behavior of internal strain increases is the opposite. This behavior of powders is explained as formation of defects such as dislocation and/or stacking faults during the milling. The internal strain increased in 100 h then decreased until reaching 400 h. Also, we have determined the crystallite size using the Scherrer formula given by Dh k l = kλ/βcos θ, where Dh k l is the crystallite size estimated from the (h k l) line, k the Scherrer constant, β the half width, λ the X-ray wavelength, and θ is the diffraction angle. Fig. 2(c) shows the average crystallite size for the milled powder of 100 h considering all diffraction peaks, calculating their width and finally, applying the Scherrer formula at all. Fig. 2(d) shows the decreased of average crystallite size with increasing tm and the crystallite size decreases rapidly for tm , less than 200 h. The crystallite sizes calculated by Williamson–Hall method are higher than those obtained by Scherrer formula. This difference could be due to the fact of
peak broadening associated to internal strain is not considered in the Scherrer model. A SEM was employed to follow the changes in the size and shape of the Al0.6 (Ni0.5 Fe0.5 )0.4 powders at different tm ’s and of the leached specimen produced at 400 h (Fig. 3). The powders before milling were irregular, with blocky, bar, wild berries shapes (Fig. 3(a)). The initial powders produced at 100 h had agglomerated to form composite units averaging approximately 85 m in diameter (Fig. 3(b)); they were blocky or plate-like. Upon increasing tm , the agglomerated particles broke down, and after 200 h of milling (Fig. 3(c)) the diameter of the agglomerated particles demonstrated an averaged decrease of approximately 5–10 m, the shape of the powders are irregular with globe and/or flake-like morphology. After 300 h of milling, the particles become smaller and more uniform in shape (Fig. 3(d)), with an average particle size of approximately 1.5–4.0 m. After 400 h of milling, the particles become more
Fig. 3. SEM images of Al0.6 (Ni0.5 Fe0.5 )0.4 after the following milling times: (a) 0 h, (b) 100 h, (c) 200 h, (d) 300 h, (e) 400 h, and (f) for leached L2 specimen of 400 h of milling.
H.-G. Kim et al. / Journal of Alloys and Compounds 456 (2008) 72–78
Fig. 4. X-ray diffraction patterns of (a) rod-milled powder, and of specimens leached under (b) L1, (c) L2, and (d) L3 conditions.
regular with an average size of approximately 2.0 m (Fig. 3(e)). The expanded SEM micrograph taken at 400 h is presented in the inset of Fig. 3(e). Fig. 3(f) presents an SEM micrograph of a leached Al0.6 (Ni0.5 Fe0.5 )0.4 specimen originally produced at 400 h. The size and the shape of the particles before (Fig. 3(e)) and after (Fig. 3(f)) leaching are almost identical. Nevertheless, the shapes of the leached powders were shown to be somewhat pressed because of the removal of Al by leaching. Fig. 4 presents the X-ray diffraction patterns of the rod-milled powder (curve a) and leached specimens under L1 (curve b), L2 (curve c), and L3 (curve d) conditions. The leaching of Al atoms under L1 and L2 conditions did not induce significant changes in diffraction patterns, even though the Al atoms had been removed. The peak position, 2θ (max) of the nanoscale crystalline phase with broad bcc peak linearly shifts to the low angle side with increasing leaching temperatures. The peak of a Fe3 O4 compound was observed at L3 condition. Therefore, we used L2 condition for leaching in this experiment. The Al content of the leached L1, L2, and L3 specimens determined by ICP is 12.3, 10.7, and 6.54 wt.%, respectively. The leaching temperature is known to have very significant effects on the structural and magnetic properties. Fig. 5 presents the X-ray diffraction patterns of the rod-milled powder (Fig. 5(a)) and leached L2 specimen before (Fig. 5(b)) and after annealing at 400 ◦ C (Fig. 5(c)) and 600 ◦ C (Fig. 5(d)) for 1 h. Leaching of the Al atoms from the bcc rod-milled powder does not induce significant change in the diffraction pattern, even though the Al atoms are removed. For the specimen heated to 400 ◦ C for 1 h, after passing the landing point between first and second DSC Peak, the fcc Bragg peak is observed as shown in Fig. 5(c). After annealed at 600 ◦ C for 1 h, the bcc phase was transformed to the fcc Ni phase, accompanied by a change in the magnetic properties [13]. The DSC traces of the leached specimens obtained with a heating rate 10 ◦ C/min are presented in Fig. 6. Two broad but distinct exothermic peaks except peak of 148.5 ◦ C of L1 are observed. The first peak temperatures, Tp1 , are 264.1, 260.4, and 238.3 ◦ C for L1, L2, and L3, respectively. The second peak temperatures, Tp2 , are 518.8, 489.0, and 457.3 ◦ C for L1, L2, and
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Fig. 5. X-ray diffraction patterns of (a) rod-milled powder, (b) leached L2 specimen, (c) leached L2 specimen after annealing at 400 ◦ C for 1 h, and (d) leached L2 specimen after annealing at 600 ◦ C for 1 h.
L3, respectively. The second peak corresponds to the structural transformation from the bcc to fcc structure accompanied by the change in the magnetic properties [13]. The peak temperatures and the width of DSC curves decreased with increasing leaching temperature. The weak low-temperature peak of L1 does not correlate with any structural changes. The room temperature M¨ossbauer spectra of rod-milled powder and leached L2 specimen are presented in Fig. 7. For the unmilled powder of the Al0.6 (Ni0.5 Fe0.5 )0.4 , the spectrum shows the presence of a typical sextet corresponding to the ferromagnetism of iron existing in the powder (Fig. 7(a)). As the tm increased, the relative intensity of the magnetic lines is reduced and the spectrum collapses into a distribution of doublet more than 200 h of milling (Fig. 7(c)), meaning that the alloying at the atomic level was completed at RM (Fig. 7(e)). The width and the intensity of the doublet for leached-specimen becomes a little wider and changes from that of the rod-milled powder, but total intensity is the same (Fig. 7(f)).
Fig. 6. DSC traces of leached specimens under different leaching conditions: (a) L1, (b) L2, and (c) L3 conditions.
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Fig. 8 presents the temperature dependence of magnetization, M for the leached L1, L2, and L3 specimens between RT and 750 ◦ C in an applied magnetic field, H, of approximately 5 kOe. The curve shapes for the L1 and L2 conditions are similar. On cooling of the specimen from 750 ◦ C, the curve sharply increased to approximately 595, 580 ◦ C for L1 specimen (Fig. 8(a)) and L2 specimen (Fig. 8(b)), respectively, indicating that the bcc phase had been transformed to the ferromagnetic fcc phase. However, the curve for L3 specimen sharply increased at approximately 468 and 542 ◦ C (Fig. 8(c)). Therefore, the point transformed
Fig. 7. (a–f) Room-temperature M¨ossbauer spectra of Al0.6 (Ni0.5 Fe0.5 )0.4 powder milled for various times. Fig. 8. Temperature dependence of the magnetization of leached specimen for temperature cycle RT → 750 ◦ C → RT: (a) L1, (b) L2, and (c) L3 conditions.
H.-G. Kim et al. / Journal of Alloys and Compounds 456 (2008) 72–78
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Fig. 9. Bright-field, dark-field images, and corresponding selected-area diffraction patterns of (a) rod-milled powder, (b) leached L2 specimen, (c) leached L2 specimen after annealing at 400 ◦ C for 1 h, and (d) leached L2 specimen after annealing at 600 ◦ C for 1 h.
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from the bcc phase to ferromagnetic fcc phase decreased with increasing leaching temperature. In order to understand the microstructure, TEM observations were carried out. Fig. 9(a) presents a bright-field image (BFI), dark-field image (DFI), and the corresponding specific-area diffraction (SAD) pattern of the rod-milled powder. The electron micrographs reveal nanometer-sized structures, and the corresponding SAD pattern demonstrate the typical bcc structure. The leached L2 specimen topotactically kept its bcc structure even though most of the Al atoms were removed, as shown in Fig. 9(b). However, the crystallite size of the leached specimen was smaller and more uniform than that of the rod-milled powder. The electron micrograph and the SAD pattern of the L2 specimen annealed at 400 ◦ C for 1 h, corresponding to Fig. 5(c), is presented in Fig. 9(c). The nanoscale crystallites and SAD pattern were almost identical to that of the leached L2 specimen. The electron micrograph and the SAD pattern of the L2 specimen annealed at 600 ◦ C for 1 h is presented in Fig. 9(d). The initial nanoscale crystallites are now relatively larger crystals. The corresponding SAD pattern present several diffraction rings and spots, which are attributed to the fcc structure, in accordance with the results of XRD. 4. Conclusions The nanocrystalline Al0.6 (Ni0.5 Fe0.5 )0.4 powder and the leached L1, L2, and L3 specimens have been prepared using a combination of rod milling and leaching of Al in a basic solution. After 200 h of milling, the bcc-type AlNi phase was detected in the XRD patterns. The strongest line of the XRD peak slightly shifted toward the lower 2θ side gradually, with increasing tm . The crystallite size, calculated by two different methods, decreased rapidly for tm , less than 200 h, while the behavior of internal strain increases is the opposite. The lattice constant of strongest peak for 2θ = 44.65◦ at 0 h of milling gradually increased with increasing tm . The strongest line of the XRD peaks linearly shifts to the low angle side with increasing leaching temperatures. The leached specimens transform to a ferromagnetic fcc Ni phase at above 400 ◦ C of annealing for 1 h. The peak temperature and the width of DSC curves decreased with increasing leaching temperature. The relative intensity of the magnetic lines of sextet in M¨ossbauer spectra is reduced and the spectrum collapses into a distribution of doublet
as the milling time increased. On cooling of the specimen from 750 ◦ C, the curve sharply increased at approximately 595, 580, and 542 ◦ C for L1, L2 and L3 specimens, respectively, and the point transformed from bcc phase to ferromagnetic fcc phase decreased with increasing leaching temperature. Acknowledgement This study was supported by research funds from Chosun University, 2005. References [1] R.P. Andres, R.S. Averback, W.L. Brus, W.A. Goddard III, A. Kaldor, S.G. Louie, M. Moskovits, P.S. Oeercy, R.W. Siegel, F. Spaepen, Y. Wang, J. Mater. Res. 4 (1989) 704. [2] R.W. Siegel, Processing of Metals and Alloys, VCH, Weinheim, 1991, p. 583. [3] D. Oleszak, H. Matyja, Nanostructured Mater. 6 (1995) 425. [4] H.G. Kim, K. Sumiyama, K. Suzuki, J. Alloys Compd. 239 (1996) 88. [5] P.H. Shingu, B. Hung, S.R. Nishitani, S. Nasu, Trans. Jpn. Inst. Mat. Suppl. 29 (1988) 3. [6] H.G. Kim, J.Y. Park, S. Yamamuro, K. Sumiyama, K. Suzuki, Mater. Sci. Eng. A217/218 (1996) 269. [7] R.L. Whetten, D.M. Cox, D.J. Trevor, A. Kaldor, Surf. Sci. 156 (1985) 8. [8] H.G. Kim, W.N. Myung, K. Sumiyama, K. Suzuki, J. Korean Phys. Soc. 31-1 (1997) 189. [9] B.H. Zeifert, J. Salmines, J.A. Hern´andez, R. Reynoso, N. Nava, J.G. Caba˜nas-Moreno, G. Aguilar R´ıos, Mater. Lett. 43 (2000) 244. [10] E. Ivanov, G. Golubkova, T. Grigorieva, React. Solids 8 (1990) 73. [11] E. Ivanov, S.A. Makhlouf, K. Sumiyama, H. Yamauchi, K. Suzuki, G. Golubkova, J. Alloys Compd. 185 (1992) 25. [12] S.A. Makhlouf, E. Ivanov, K. Sumiyama, K. Suzuki, J. Alloys Compd. 189 (1992) 117. [13] S.A. Makhlouf, K. Sumiyama, E. Ivanov, H. Yamauchi, T. Hihara, K. Suzuki, Mater. Sci. Eng. A181/A182 (1994) 1184. [14] H.G. Kim, W.N. Myung, Inter. J. Non-Equilib. Process 11 (2000) 271. [15] I. Yamauchi, K. Takahara, T. Tanaka, K. Matsubara, J. Alloys Compd. 396 (2005) 302. [16] H.G. Kim, W.N. Myung, K. Sumiyama, K. Suzuki, J. Alloys Compd. 398 (2005) 74. [17] I. Yamauchi, I. Ohnaka, M. Murata, J. Jpn. Inst. Metals 56 (1992) 1385. [18] P. Fouilloux, Appl. Catal. 8 (1983) 1. [19] G.K. Williamson, W.H. Hall, Acta Met. 1 (1953) 22. [20] M.E. Rabanal, A. V´arez, B. Levenfeld, J.M. Torralba, J. Mater. Process. Technol. 143 (2003) 470. [21] W.H. Hall, Proc. Phys. A62 (1949) 741. [22] B.D. Cullity, Elements of X-ray diffraction, Addison-Wesley, 1978, p. 99.
Chemical leaching of non-equilibrium Al(Ni-Fe) alloy powders produced by rod milling Hyun-Goo Kim a,∗ , Wha-Nam Myung b , K. Sumiyama c , K. Suzuki d a
Department of Physics Education, Chosun University, Gwangju 501-759, South Korea Department of Physics, Chonnam National University, Gwangju 500-757, South Korea c Department of Materials Science and Engineering, Nagoya Institute of Technology, Nagoya, Japan d Institute for Advanced Materials Research in Sendai, Tokai University in Tokyo, Japan b
Received 14 February 2006; received in revised form 1 February 2007; accepted 1 February 2007 Available online 6 February 2007
Abstract In this paper, chemical leaching of non-equilibrium Al0.6 (Ni0.5 Fe0.5 )0.4 powder by rod milling, is reported. X-ray diffractometry (XRD), transmission electron microscopy (TEM), differential scanning calorimetry (DSC), scanning electron microscopy (SEM), M¨ossbauer spectroscopy, and vibrating sample magnetometry (VSM) were used to characterize the rod-milled and leached specimens. After 200 h of milling, the AlNi bcc phase was detected in the XRD patterns. The crystallite size, calculated using Williamson–Hall and Scherrer methods, decreased rapidly for tm , less than 200 h, while the behavior of internal strain increases is the opposite. The leached specimens transformed to a ferromagnetic fcc Ni phase at above 400 ◦ C of annealing for 1 h. The peak temperature and the width of DSC curves decreased with increasing leaching temperature. The relative intensity of the magnetic lines of sextet in M¨ossbauer spectra is reduced and the spectrum collapses into a distribution of doublet as the milling time increased. When the specimen cooled from 750 ◦ C, the curve sharply increased to approximately 595, 580, and 542 ◦ C for leached specimens in a 30 wt.% KOH solution at room temperature, at 55 and at 85 ◦ C, respectively, indicating that the bcc phase has been transformed to a fcc phase. © 2007 Elsevier B.V. All rights reserved. Keywords: Rod milling; Non-equilibrium Al(Ni-Fe); Chemical leaching
1. Introduction In anticipation of unique or improved properties, very different from those of conventional coarse crystalline materials, considerable interest in producing and studying nanocrystalline materials has been created [1–4]. Recently, a technique of mechanically alloying pure elements has been proven to be one of the most effective methods for preparing nanocrystalline materials [5,6]. Nanometer order powders have found applications in catalysis, electromagnetic shielding, magnetic recording, refrigeration and in the processing of advanced engineering materials. In particular, nanoscale 3d-transition metal (Fe, Co and Ni) particles are expected to display unique magnetic and catalytic properties [7]. A combination of rod milling (RM) and chemical leaching has been used to prepare nanoscale metastable materials,
which represent many interesting properties [8–16]. The leaching treatment of Al–Ni alloys by an alkaline solution leads to removal of the major portion of the Al and retention of a “skeleton Ni” catalyst. Similarly, a nanocrystalline bcc Co phase is obtained by leaching out Al atoms from mechanically alloyed bcc Co40 Al60 [12]. After the non-equilibrium Al(Ni-Fe) powder had been leached and annealed, the bcc phase is transformed into the fcc phase [8]. Yamauchi et al. reported that it formed the metastable phase by alkaline leaching of Al of rapidly solidified Al3 (Nix Fe1−x ) alloys [17]. In this paper, the idea of leaching Al from a rod-milled Al0.6 (Ni0.5 Fe0.5 )0.4 alloy is used in a basic solution to obtain a nanocrystalline bcc Ni-Fe phase and to study its properties using XRD, DSC, SEM, TEM, M¨ossbauer spectroscope and vibrating sample magnetometer (VSM). 2. Experimental
∗
Corresponding author. Tel.: +82 62 230 7376. E-mail address: [email protected] (H.-G. Kim).
0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.02.002
Pure elemental powders of Al, Ni, and Fe were mixed in an argon glove box to provide a composition of Al0.6 (Ni0.5 Fe0.5 )0.4 , the nominal composition being
H.-G. Kim et al. / Journal of Alloys and Compounds 456 (2008) 72–78 expressed in atomic percent. The rod-to-powder weight ratio was 28:1. The stainless steel (SUS 304) vial together with SUS 304 rods (10 mm in diameter) was rotated with an angular velocity of 100 rev/min under purified argon gas. The vial was opened after 100, 200, 300 and 400 h (h, milling times (tm )) of processing, in an argon atmosphere. A small amount of sample powder was removed from the vial in a glove box for analysis and observation. The rod-milled powders (400 h of milling) were leached in a 30 wt.% KOH solution at approximately 20 ◦ C (room temperature, RT, L1), at 55 ◦ C (L2) and at 85 ◦ C (L3) to remove the Al atoms. After the leaching treatment, the powders were washed with distilled water until the pH value became the same as the distilled water. Then the powders were immersed in alcohol and dried in a glove box under argon atmosphere. This process is the conventional method for obtaining Raney Ni catalysts [18]. The leached L2 alloy powders were annealed in evacuated quartz tubes at 400 and 600 ◦ C for 1 h. The XRD was performed at room temperature by using Cu K␣ radiation. The scan speed was 1.5◦ min−1 and the sampling width was 0.05◦ . The heat treatment up to crystallization was monitored by DSC at a heating rate of 10 K min−1 in a flow of purified argon gas. The microstructure of the rod-milled powder and the leached specimens before and after annealing were observed by TEM using a 200 kV. SEM observations allowing for characterization of powder particles during RM have been performed as well. The M¨ossbauer spectroscopy was performed on all samples, at room temperature and at 14.4 keV with a Rh/Co57 source. A conventional VSM was employed to measure the magnetization from 20 up to 750 ◦ C in magnetic fields of 5 kOe. The composition of the alloys was determined using inductively coupled plasma-emission spectrometry (ICP).
3. Result and discussion Fig. 1 presents the change in the X-ray diffraction patterns of the powders with tm : 0, 100, 200, 300 and 400 h. Directly after milling, a significant decrease in the relative intensities of the Ni, the Fe and the Al peaks were detected, and after 200 h, only the bcc-type AlNi phase was detected in the XRD patterns. However, the XRD peak for 2θ = 44.65◦ at 0 h of milling is asymmetric after 100 h of RM and slightly shifted toward
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Fig. 1. X-ray diffraction patterns of rod-milled Al0.6 (Ni0.5 Fe0.5 ) 0.4 for different milling times.
the lower 2θ side gradually with increasing tm , indicating that Fe atoms are arranged into the bcc-type AlNi phase. The lattice constants of peak for 2θ = 44.65◦ are 2.862, 2.862, 2.87999, 2.87999, and 2.88925 nm for 0, 100, 200, 300, and 400 h of milling, respectively, gradually increasing with increasing tm (the inset of Fig. 1). The value of the lattice constants for 400 h was similar with that of 2.888 nm for AlNi. The evolutions of crystallite size and internal strain with the tm were evaluated using Williamson–Hall [19–21] and Scherrer methods [22]. Fig. 2(a) represents a βcos θ/λ versus sin θ/λ plot for powder with 100 h of milling. The slope of the straight line gives the amount of internal strain, which comes out to be 4.28 × 10−3 . The intercept on the βcos θ/λ axis gives the crystallite size as 11.15 nm. Fig. 2(b) shows the evolution of crystallite
Fig. 2. (a) Williamson–Hall plot showing βcos θ/λ vs. sin θ/λ for the powder of 100 h for milling. (b) Evolution of crystallite size and internal strain with the milling time obtained from Williamson–Hall method for the Al0.6 (Ni0.5 Fe0.5 )0.4 . (c) Crystallite size obtained applying the Scherrer formula to different peaks of XRD patterns for the powder with 100 h of milling time. (d) Evolution of crystallite size with the milling time obtained from Scherrer method.
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size with tm . For tm , less than 200 h crystallite size decreases rapidly, while the behavior of internal strain increases is the opposite. This behavior of powders is explained as formation of defects such as dislocation and/or stacking faults during the milling. The internal strain increased in 100 h then decreased until reaching 400 h. Also, we have determined the crystallite size using the Scherrer formula given by Dh k l = kλ/βcos θ, where Dh k l is the crystallite size estimated from the (h k l) line, k the Scherrer constant, β the half width, λ the X-ray wavelength, and θ is the diffraction angle. Fig. 2(c) shows the average crystallite size for the milled powder of 100 h considering all diffraction peaks, calculating their width and finally, applying the Scherrer formula at all. Fig. 2(d) shows the decreased of average crystallite size with increasing tm and the crystallite size decreases rapidly for tm , less than 200 h. The crystallite sizes calculated by Williamson–Hall method are higher than those obtained by Scherrer formula. This difference could be due to the fact of
peak broadening associated to internal strain is not considered in the Scherrer model. A SEM was employed to follow the changes in the size and shape of the Al0.6 (Ni0.5 Fe0.5 )0.4 powders at different tm ’s and of the leached specimen produced at 400 h (Fig. 3). The powders before milling were irregular, with blocky, bar, wild berries shapes (Fig. 3(a)). The initial powders produced at 100 h had agglomerated to form composite units averaging approximately 85 m in diameter (Fig. 3(b)); they were blocky or plate-like. Upon increasing tm , the agglomerated particles broke down, and after 200 h of milling (Fig. 3(c)) the diameter of the agglomerated particles demonstrated an averaged decrease of approximately 5–10 m, the shape of the powders are irregular with globe and/or flake-like morphology. After 300 h of milling, the particles become smaller and more uniform in shape (Fig. 3(d)), with an average particle size of approximately 1.5–4.0 m. After 400 h of milling, the particles become more
Fig. 3. SEM images of Al0.6 (Ni0.5 Fe0.5 )0.4 after the following milling times: (a) 0 h, (b) 100 h, (c) 200 h, (d) 300 h, (e) 400 h, and (f) for leached L2 specimen of 400 h of milling.
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Fig. 4. X-ray diffraction patterns of (a) rod-milled powder, and of specimens leached under (b) L1, (c) L2, and (d) L3 conditions.
regular with an average size of approximately 2.0 m (Fig. 3(e)). The expanded SEM micrograph taken at 400 h is presented in the inset of Fig. 3(e). Fig. 3(f) presents an SEM micrograph of a leached Al0.6 (Ni0.5 Fe0.5 )0.4 specimen originally produced at 400 h. The size and the shape of the particles before (Fig. 3(e)) and after (Fig. 3(f)) leaching are almost identical. Nevertheless, the shapes of the leached powders were shown to be somewhat pressed because of the removal of Al by leaching. Fig. 4 presents the X-ray diffraction patterns of the rod-milled powder (curve a) and leached specimens under L1 (curve b), L2 (curve c), and L3 (curve d) conditions. The leaching of Al atoms under L1 and L2 conditions did not induce significant changes in diffraction patterns, even though the Al atoms had been removed. The peak position, 2θ (max) of the nanoscale crystalline phase with broad bcc peak linearly shifts to the low angle side with increasing leaching temperatures. The peak of a Fe3 O4 compound was observed at L3 condition. Therefore, we used L2 condition for leaching in this experiment. The Al content of the leached L1, L2, and L3 specimens determined by ICP is 12.3, 10.7, and 6.54 wt.%, respectively. The leaching temperature is known to have very significant effects on the structural and magnetic properties. Fig. 5 presents the X-ray diffraction patterns of the rod-milled powder (Fig. 5(a)) and leached L2 specimen before (Fig. 5(b)) and after annealing at 400 ◦ C (Fig. 5(c)) and 600 ◦ C (Fig. 5(d)) for 1 h. Leaching of the Al atoms from the bcc rod-milled powder does not induce significant change in the diffraction pattern, even though the Al atoms are removed. For the specimen heated to 400 ◦ C for 1 h, after passing the landing point between first and second DSC Peak, the fcc Bragg peak is observed as shown in Fig. 5(c). After annealed at 600 ◦ C for 1 h, the bcc phase was transformed to the fcc Ni phase, accompanied by a change in the magnetic properties [13]. The DSC traces of the leached specimens obtained with a heating rate 10 ◦ C/min are presented in Fig. 6. Two broad but distinct exothermic peaks except peak of 148.5 ◦ C of L1 are observed. The first peak temperatures, Tp1 , are 264.1, 260.4, and 238.3 ◦ C for L1, L2, and L3, respectively. The second peak temperatures, Tp2 , are 518.8, 489.0, and 457.3 ◦ C for L1, L2, and
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Fig. 5. X-ray diffraction patterns of (a) rod-milled powder, (b) leached L2 specimen, (c) leached L2 specimen after annealing at 400 ◦ C for 1 h, and (d) leached L2 specimen after annealing at 600 ◦ C for 1 h.
L3, respectively. The second peak corresponds to the structural transformation from the bcc to fcc structure accompanied by the change in the magnetic properties [13]. The peak temperatures and the width of DSC curves decreased with increasing leaching temperature. The weak low-temperature peak of L1 does not correlate with any structural changes. The room temperature M¨ossbauer spectra of rod-milled powder and leached L2 specimen are presented in Fig. 7. For the unmilled powder of the Al0.6 (Ni0.5 Fe0.5 )0.4 , the spectrum shows the presence of a typical sextet corresponding to the ferromagnetism of iron existing in the powder (Fig. 7(a)). As the tm increased, the relative intensity of the magnetic lines is reduced and the spectrum collapses into a distribution of doublet more than 200 h of milling (Fig. 7(c)), meaning that the alloying at the atomic level was completed at RM (Fig. 7(e)). The width and the intensity of the doublet for leached-specimen becomes a little wider and changes from that of the rod-milled powder, but total intensity is the same (Fig. 7(f)).
Fig. 6. DSC traces of leached specimens under different leaching conditions: (a) L1, (b) L2, and (c) L3 conditions.
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Fig. 8 presents the temperature dependence of magnetization, M for the leached L1, L2, and L3 specimens between RT and 750 ◦ C in an applied magnetic field, H, of approximately 5 kOe. The curve shapes for the L1 and L2 conditions are similar. On cooling of the specimen from 750 ◦ C, the curve sharply increased to approximately 595, 580 ◦ C for L1 specimen (Fig. 8(a)) and L2 specimen (Fig. 8(b)), respectively, indicating that the bcc phase had been transformed to the ferromagnetic fcc phase. However, the curve for L3 specimen sharply increased at approximately 468 and 542 ◦ C (Fig. 8(c)). Therefore, the point transformed
Fig. 7. (a–f) Room-temperature M¨ossbauer spectra of Al0.6 (Ni0.5 Fe0.5 )0.4 powder milled for various times. Fig. 8. Temperature dependence of the magnetization of leached specimen for temperature cycle RT → 750 ◦ C → RT: (a) L1, (b) L2, and (c) L3 conditions.
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Fig. 9. Bright-field, dark-field images, and corresponding selected-area diffraction patterns of (a) rod-milled powder, (b) leached L2 specimen, (c) leached L2 specimen after annealing at 400 ◦ C for 1 h, and (d) leached L2 specimen after annealing at 600 ◦ C for 1 h.
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from the bcc phase to ferromagnetic fcc phase decreased with increasing leaching temperature. In order to understand the microstructure, TEM observations were carried out. Fig. 9(a) presents a bright-field image (BFI), dark-field image (DFI), and the corresponding specific-area diffraction (SAD) pattern of the rod-milled powder. The electron micrographs reveal nanometer-sized structures, and the corresponding SAD pattern demonstrate the typical bcc structure. The leached L2 specimen topotactically kept its bcc structure even though most of the Al atoms were removed, as shown in Fig. 9(b). However, the crystallite size of the leached specimen was smaller and more uniform than that of the rod-milled powder. The electron micrograph and the SAD pattern of the L2 specimen annealed at 400 ◦ C for 1 h, corresponding to Fig. 5(c), is presented in Fig. 9(c). The nanoscale crystallites and SAD pattern were almost identical to that of the leached L2 specimen. The electron micrograph and the SAD pattern of the L2 specimen annealed at 600 ◦ C for 1 h is presented in Fig. 9(d). The initial nanoscale crystallites are now relatively larger crystals. The corresponding SAD pattern present several diffraction rings and spots, which are attributed to the fcc structure, in accordance with the results of XRD. 4. Conclusions The nanocrystalline Al0.6 (Ni0.5 Fe0.5 )0.4 powder and the leached L1, L2, and L3 specimens have been prepared using a combination of rod milling and leaching of Al in a basic solution. After 200 h of milling, the bcc-type AlNi phase was detected in the XRD patterns. The strongest line of the XRD peak slightly shifted toward the lower 2θ side gradually, with increasing tm . The crystallite size, calculated by two different methods, decreased rapidly for tm , less than 200 h, while the behavior of internal strain increases is the opposite. The lattice constant of strongest peak for 2θ = 44.65◦ at 0 h of milling gradually increased with increasing tm . The strongest line of the XRD peaks linearly shifts to the low angle side with increasing leaching temperatures. The leached specimens transform to a ferromagnetic fcc Ni phase at above 400 ◦ C of annealing for 1 h. The peak temperature and the width of DSC curves decreased with increasing leaching temperature. The relative intensity of the magnetic lines of sextet in M¨ossbauer spectra is reduced and the spectrum collapses into a distribution of doublet
as the milling time increased. On cooling of the specimen from 750 ◦ C, the curve sharply increased at approximately 595, 580, and 542 ◦ C for L1, L2 and L3 specimens, respectively, and the point transformed from bcc phase to ferromagnetic fcc phase decreased with increasing leaching temperature. Acknowledgement This study was supported by research funds from Chosun University, 2005. References [1] R.P. Andres, R.S. Averback, W.L. Brus, W.A. Goddard III, A. Kaldor, S.G. Louie, M. Moskovits, P.S. Oeercy, R.W. Siegel, F. Spaepen, Y. Wang, J. Mater. Res. 4 (1989) 704. [2] R.W. Siegel, Processing of Metals and Alloys, VCH, Weinheim, 1991, p. 583. [3] D. Oleszak, H. Matyja, Nanostructured Mater. 6 (1995) 425. [4] H.G. Kim, K. Sumiyama, K. Suzuki, J. Alloys Compd. 239 (1996) 88. [5] P.H. Shingu, B. Hung, S.R. Nishitani, S. Nasu, Trans. Jpn. Inst. Mat. Suppl. 29 (1988) 3. [6] H.G. Kim, J.Y. Park, S. Yamamuro, K. Sumiyama, K. Suzuki, Mater. Sci. Eng. A217/218 (1996) 269. [7] R.L. Whetten, D.M. Cox, D.J. Trevor, A. Kaldor, Surf. Sci. 156 (1985) 8. [8] H.G. Kim, W.N. Myung, K. Sumiyama, K. Suzuki, J. Korean Phys. Soc. 31-1 (1997) 189. [9] B.H. Zeifert, J. Salmines, J.A. Hern´andez, R. Reynoso, N. Nava, J.G. Caba˜nas-Moreno, G. Aguilar R´ıos, Mater. Lett. 43 (2000) 244. [10] E. Ivanov, G. Golubkova, T. Grigorieva, React. Solids 8 (1990) 73. [11] E. Ivanov, S.A. Makhlouf, K. Sumiyama, H. Yamauchi, K. Suzuki, G. Golubkova, J. Alloys Compd. 185 (1992) 25. [12] S.A. Makhlouf, E. Ivanov, K. Sumiyama, K. Suzuki, J. Alloys Compd. 189 (1992) 117. [13] S.A. Makhlouf, K. Sumiyama, E. Ivanov, H. Yamauchi, T. Hihara, K. Suzuki, Mater. Sci. Eng. A181/A182 (1994) 1184. [14] H.G. Kim, W.N. Myung, Inter. J. Non-Equilib. Process 11 (2000) 271. [15] I. Yamauchi, K. Takahara, T. Tanaka, K. Matsubara, J. Alloys Compd. 396 (2005) 302. [16] H.G. Kim, W.N. Myung, K. Sumiyama, K. Suzuki, J. Alloys Compd. 398 (2005) 74. [17] I. Yamauchi, I. Ohnaka, M. Murata, J. Jpn. Inst. Metals 56 (1992) 1385. [18] P. Fouilloux, Appl. Catal. 8 (1983) 1. [19] G.K. Williamson, W.H. Hall, Acta Met. 1 (1953) 22. [20] M.E. Rabanal, A. V´arez, B. Levenfeld, J.M. Torralba, J. Mater. Process. Technol. 143 (2003) 470. [21] W.H. Hall, Proc. Phys. A62 (1949) 741. [22] B.D. Cullity, Elements of X-ray diffraction, Addison-Wesley, 1978, p. 99.