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Nanocrystalline eutectic Al–Si alloy produced by cryomilling
Materials Science and Engineering A 508 (2009) 43–49
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Nanocrystalline eutectic Al–Si alloy produced by cryomilling J. Milligan, R. Vintila, M. Brochu ∗ Aluminium Research Centre - REGAL, Department of Mining and Materials Engineering, McGill University, 3610 University St., Montreal, Quebec, Canada H3A 2B2
a r t i c l e
i n f o
Article history: Received 8 August 2008 Received in revised form 8 December 2008 Accepted 10 December 2008 Keywords: Aluminum Silicon Eutectic Cryomilling Nanocrystalline
a b s t r a c t A nanocrystalline Al–Si eutectic alloy was produced by mechanical attrition under cryogenic conditions. The resulting alloy was characterized using X-ray diffraction, scanning and transmission electron microscopy, and thermal heat treatments. The powder was milled for 4 h when the grain size reached 46 nm according to X-ray diffraction and 38 nm according to transmission electron microscopy examinations. An increase in solubility limit was achieved and was determined by measurements of the lattice parameter after milling, and quantified using Vegard’s law. The solubility reached after 4 h of milling time was 1.72 at.% of silicon in solid solution at room temperature. © 2008 Elsevier B.V. All rights reserved.
1. Introduction The mechanical properties of nanocrystalline materials have become of increasing interest because they are significantly different from those of the same material with coarse grains. Changes are seen in the mechanical properties such as hardness, strength and ductility, when the grain size is reduced to the nanometer scale [1]. The strength and plastic behaviour of metals, which strongly depends on dislocation activity, starts to differ substantially entering the nanometer regime because of the reduced availability of mobile dislocations in the smaller nanoscaled grains [2]. Nanocrystalline materials provide the opportunity to obtain properties that are unachievable with equilibrium materials. One proven synthesis technique available to produce nanocrystalline materials is mechanical milling [3]. Mechanical milling is a solidstate processing technique involving repeated welding, fracturing, and rewelding of powder particles in a low-energy ball mill to produce refined grain structures and particle sizes [4]. The milling action mixes the powders on an atomic scale leading to alloying of the powder with very fine grain sizes [5]. The mechanism of nanostructure formation during milling has been attributed to the formation of an array with a high density of dislocations, which annihilate and recombine to form high angle grain boundaries; the orientation of the grains then becomes random [4,6–8]. Materials that have a grain structure in the nanometric regime (<100 nm) have exhibited very high hardness and strength. Hard-
∗ Corresponding author. Tel.: +1 514 398 2354; fax: +1 514 398 4492. E-mail address: [email protected] (M. Brochu). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.12.017
ness values for nanocrystalline metals have been recorded up to 2–7 times higher than those of larger grained metals [2,3]. Lee et al. showed an increase of hardness in Al 5083 from 0.45 GPa up to 2 GPa when the grain size was decreased from 5 m to 200 nm [9]. Siegel and Fougere [10] summarized hardness values of 0.45 GPa for coarse grained copper which then increased to 1.2 GPa after a reduction in grain size to 25 nm. Coarse grained iron showed an increase of hardness from 1.3 GPa to 6.5 GPa once reaching a grain size of 17 nm. Nanocrystalline nickel showed an increase of hardness from 1.5 GPa to 6.5 GPa when the grain size was reduced to 20 nm from its coarse grained counterpart. Cryomilling, the mechanical attrition of powders within a cryogenic medium, has been proven to strengthen materials through refinement of the grain size and the dispersion of nanoscale particles [3]. A complete description of the process and mechanisms involved during milling was reported by Witkin and Lavernia [3]. The liquid nitrogen medium helps prevent oxidation and recrystallization during milling, which is advantageous given the large amount of free surfaces and dislocations being created during milling. Among the light elements, aluminum is becoming one of the most used elements because of its low density, high strength and numerous available alloys [5]. Previous work has been completed on the cryomilling of pure Al and aluminum alloys that are ductile [5,11,12]. Aluminum alloys that are used for wear resistance applications are based on the aluminum–silicon alloy system [13]. The silicon content in the aluminum alloy imparts excellent wear resistance, and the wear resistance can be increased with increasing refinement of the silicon particles [13]. There exist several techniques available to produce fine silicon particles and super
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J. Milligan et al. / Materials Science and Engineering A 508 (2009) 43–49
saturated solid solutions of Al–Si alloys. These methods depend on very rapid solidification and include methods such as chill casting [14], melt spinning [15], water quenching [16] and atomization [17]. In addition, super saturated solid solutions can be achieved through mechanical milling, but the property enhancements that are achievable are limited by the amount of alloying element that can be added in solution without the formation of a second phase [5]. By increasing the amount of alloying element in solid solution in the matrix, we can, for example, further increase the effect of solid solution strengthening of the alloy. Solid solubility extension also presents the potential that upon annealing of the metastable powder, the release of the atoms in solution may nucleate nanoscale precipitates. Cryomilling works were mainly targeted to structural alloys and no systematic work was devoted to the aluminum–silicon system prior to completion of this work. Aluminum–silicon powders have potential application as feedstock for coatings and claddings [18–20]. The present study involves the cryomilling of an aluminum silicon eutectic alloy for wear applications. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were used to characterize the morphology, and grain structure of the powder. X-ray diffraction (XRD) and XLAT software were used to measure the lattice parameter and calculate the variation of solubility of the silicon solute in the aluminum matrix. Thermal heat treatments were also preformed to observe the thermal stability, i.e. evolution of grain size distribution and solid solubility limit upon exposure of the powder at higher temperature.
a Cauchy–Gaussian profile of the reflection [21,22]: ı2 (2)
K = L tan2 o
ı(2) tan o sin o
+ 16e2
(1)
in which K is a constant taken as 1, L is the grain size, e is the lattice strain, and ı(2) = B(1 − b2 /B2 ), b and B are the breadths in radians of the same peak in the reference and experimental patterns, respectively. Lattice parameters were determined using XLAT software which uses a least squares fit for the precise refinement of cell constants. Tungsten was used as the diffraction standard for the XRD testing. Bright field and dark field transmission electron micrographs were acquired using a Philips CM200 microscope with a beam energy of 200 keV. Grain sizes were measured from dark field micrographs taking the average of six measurements of diameter for each grain. Grain size distributions were also measured from the varying width of the diffraction rings of the selected area diffraction patterns according to Eq. (2) [23]: D=
rd r
(2)
where D is the grain size, r is the radius of the ring, d is the interplanar spacing of the respective ring, and r is the width of the ring. To verify the thermal stability of the nanostructured powders, samples were heat treated in a tube furnace under a high-purity argon atmosphere at 100 ◦ C, 200 ◦ C, and 300 ◦ C for 1 h and were characterized for solid solubility and grain size. 3. Results and discussion
2. Experimental procedures The starting powder was a −325 mesh, atomized eutectic aluminum silicon alloy from Valimet Inc. The composition was measured by ICP and is presented in Table 1. The starting powder has an average particle size of 32 m and a starting grain size of 2 m. The cryogenic milling experiments were completed using a Union Process 1-S attrition mill in a stainless steel vial under a rotational speed of 180 rpm. Stainless steel balls, with a diameter of 4.85 mm, were used as grinding media and a powder to ball weight ratio of 1:32 was used. To prevent adhesion of the powder to the balls, tank and attritor, and to control the fracturing events, 0.25 wt% of stearic acid (CH3 (CH2 )16 CO2 H) was added as a process control agent. The powder was milled for 4 h. Liquid nitrogen was continuously added to the vial to maintain constant slurry during milling, and to maintain a constant temperature of −196 ◦ C. Particle size distributions of the powder after 0, 2, and 4 h of milling were measured using a HORIBA LA-920 particle size analyzer. The evolution of the morphology of the powder after various stage of milling was characterized using secondary electron micrographs acquired using a Hitachi VP-SEM 3000SN microscope. Grain size and lattice parameter measurements were made with X-ray diffraction using Phillips PW1070 diffractometer (Cu K␣ = 1.54056 Å). A pattern was obtained from 35◦ to 110◦ with a step size of 0.005◦ and a dwell time of 0.2 s/step. The grain size and lattice strain were determined from the XRD peak broadening using Eq. (1), assuming Table 1 The chemical composition of the starting powder measured by ICP. Element
Amount (wt%)
Al Si Fe Cu Mg Mn Cr Zn
Bal. 12.07 0.1166 0.00613 0.0097 0.0053 0.00004 0.0021
Fig. 1 shows the powder morphology evolution throughout the milling process, which changes from spherical to flakes. The initial morphology of the master alloy was spherical from the atomization process as depicted in Fig. 1a. During ball milling the powder particles are continuously being flattened, cold welded and fractured. Due to the initial powder matrix being soft, their tendency to flatten and form larger particles is high during the initial stages of milling (shown in Fig. 1b); as the process continues the particles become work hardened and eventually fracture [4] (shown in Fig. 1c). The average particle size distribution of the powder throughout the milling process was measured. Fig. 2 shows the variation of the particle size with milling time. The average starting particle size was 32 m. After 2 h of milling, the average particle size increased to 61 m, and then decreased to 18 m after 4 h of milling time. After 2 h of milling time, the average particle size increases as the particles are flattened, which can be seen in Fig. 1b. When fracture occurs the particle size of the powder then decreases [24]. The results obtained are showing a faster reduction in particle size than what is generally observed for fully metallic powders. In particular, a reported minimum milling time of 10 h is required to yield similar particle size in an Al–Mg system [11] as seen in Fig. 2. The formation of AlN and SiN are thermodynamically favorable during the milling process, however high resolution examination was not undertaken to determine the existence and characteristics of these phases after milling. It is believed that the presence of the Si phase accelerates the deformation and fracture of the powders either through the formation of nitrides during milling or through the presence of a second hard phase in the mill. Similar results have been observed by Chung et al. [6], where they added a small volume fraction of AlN into pure Ni powders. A particle size reduction of 63% was observed after 8 h of milling time with the addition of the hard phase to the Ni powders. In the Al–Si system compared to the Al–Mg system a 52% reduction in particle size was observed after 4 h of milling time. The powder was scanned using X-ray diffraction to determine the grain size, lattice strain and lattice parameter evolution during cryomilling. Fig. 3 shows the starting powder used as a reference, where both the aluminum and silicon phases are observable. The
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Fig. 1. Morphology of the powder (a) prior to milling (b) after 2 h of milling and (c) after 4 h of milling time.
aluminum and silicon phases correspond to JCPDS patterns 040787 and 05-0565, respectively. Broadening of the fcc-Al peaks is observed in the XRD spectrum after 2 h and 4 h of milling time. With increasing milling time, the peaks broaden and loose intensity as the grain size and lattice strain evolve during milling. The broadening of the Al peaks allows the grain size and lattice strain to be determined based on the FWHM of the peaks. Table 2 lists the
Fig. 2. Change in particle size versus milling time.
Table 3 The change in lattice parameter and solid solubility during milling. Milling time (h)
Lattice parameter (Å)
Si in S.S. (at.%)
0 2 4
4.0494 ± 0.0005 4.0479 ± 0.0019 4.0439 ± 0.0016
0 0.47 1.72
values for the grain size and lattice strain for the different milling times. The grain size of the Al–Si alloy was reduced to 46 nm after 4 h of milling time with a lattice strain of 0.09% which is comparable to the grain size and strain reported for cryomilling of pure aluminum, where a 47 nm grain size and a lattice strain of 0.10% strain were obtained after 4 h of milling [12]. The lattice parameter of the aluminum matrix was measured using the shifts in the XRD spectrum. According to Vegard’s law the lattice parameter of an alloy varies linearly with the solid solubility. The solid solubility increases with decreasing lattice parameter of aluminum (Al = 0.4094 nm) to the lattice parameter of an imaginary cubic closed packed silicon (Si = 0.3731 nm) [15,25]. Fig. 4 depicts the Vegard’s law (bold line) relationship for the aluminum–silicon system. The results on the extension of the solubility limit of silicon in the aluminum matrix are presented in Table 3 and a maximum solubility of 1.72 at.% was obtained during cryomilling. It is worth mentioning that the equilibrium solid solution limit for this system
Table 2 The average grain size and lattice strain in nc Al and nc Al–12Si. Material
Milling time (h)
Grain size (nm)
Lattice strain (%)
T (K)
Reference
Al–12Si Al–12Si Al
2 4 4
105 46 47
0.22 0.09 0.10
80 80 90
This paper This paper [12]
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J. Milligan et al. / Materials Science and Engineering A 508 (2009) 43–49
Fig. 3. XRD spectrums after various milling times for the Al–12Si powder.
is 1.5 at.% at 577 ◦ C and 0% at room temperature [26]. Fig. 4 also contains comparable results on the extension of the solubility limit obtained using a SPEX 8000 shaker mill after 17 h of milling with a ball to powder ratio of 10:1. Their results are showing an increase in solubility limit up to 1.3 at.% of silicon when a starting powder containing 10 at.% of silicon was used [5]. It was also reported that the maximum solubility limit achieved with a starting solute content of 30 at.% silicon was approximately 4.5 at.% [5]. Based on the results of solid solubility extension (SSE) in the Al–Si obtained in this work and from literature data, the SSE extension seems to be linearly proportional to the solute concentration in the Al–Si system. To further study the thermal stability of the SSE, heat treatments at 50 ◦ C, 75 ◦ C, and 100 ◦ C for 1 h were performed. Fig. 5 presents the modified aluminum silicon phase diagram containing the data of the decomposition of the metastable solid solubility as a function of temperature for the alloy Al–12Si. After a heat treatment of 1 h at room temperature, approximately 1 at.% of silicon remains in solid solution in the aluminum. The solid solubility decreased down to 0.72 at.% after 1 h at 50 ◦ C, down to 0.16 at.% after 1 h at 75 ◦ C, and finally to complete dissociation after 1 h of heat treatment at 100 ◦ C.
TEM analysis of the cryomilled powder was completed to observe the grain structure after milling. The milling time used for this powder was sufficient to produce a random dispersion of small grains through the sample which can be observed in the bright and dark field micrographs shown in Fig. 6a and b. Fig. 6c shows the selected area diffraction pattern of the cryomilled powder showing a ring pattern indicating that the powder has reached a nanoscaled grain size. The rings corresponding to aluminum and silicon are completely distinguishable. The grain size distribution measurements from the TEM dark field micrographs are presented in Fig. 6d and show an average grain size of 38.8 nm. The microstructure consists mainly of small nanometer grains with some well-dispersed ultra fine grains. In order to characterize the respective grain size distribution of the Al and Si grains, as this is a limitation in the differentiation arising from the dark field images, their respective grain size distribution was measured using the width of the diffraction rings. Fig. 7 illustrates the grain size distributions measured
Fig. 4. Solid solubility versus lattice parameter for the Al–Si system according to Vegard’s law.
Fig. 5. The aluminum–silicon phase diagram with the solid solubility of nano-Al–Si (starting powder Al–12Si).
J. Milligan et al. / Materials Science and Engineering A 508 (2009) 43–49
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Fig. 6. TEM bright field (a), dark field (b), and diffraction pattern (c) after 4 h of milling. (d) The distribution of grains measured from the dark field images.
from the rings in the diffraction pattern shown in Fig. 6c. The grain sizes were calculated according to Eq. (2), to distinguish the size of the aluminum grains versus the silicon grains in the cryomilled powder. Fig. 7a shows the distribution of the aluminum grains determined from the (1 1 1) and (2 0 0) rings. An average aluminum grain size of 26 nm was measured. The distribution of the silicon grains were determined from the (1 1 1), (2 2 0) and (3 1 1) rings and have an average of 19 nm (seen in Fig. 7b). The grain size measured is dependant on the instantaneous ring width, larger widths indicating a smaller grain. As the number of grains
increases the few large grains that remain will produce a small ring width that may be convoluted beneath the section of the ring produced by a smaller grain at the same rotation in the selected area diffraction pattern. As such a uniform grain size distribution is required to apply this measurement technique. Combining the two distributions together (shown in Fig. 7c), the weighted average of the aluminum and silicon grains between the grain sizes of 0.1–70 nm is 24 nm, which agrees with the average grain size measured in the dark field images in the same range. This distribution is comparable to that found for pure aluminum that
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J. Milligan et al. / Materials Science and Engineering A 508 (2009) 43–49
Fig. 7. Distribution of the aluminum grains (a), the silicon grains (b) and aluminum + silicon grains (c) measured from the ring pattern after 4 h of milling time.
has been cryomilled [12,27]. The distributions measured show that during milling there is a simultaneous refinement of both the aluminum and silicon phases present; this effect was also observed by Chung et al. [6] during milling of a composite system. The thermal stability of the powder was investigated through heat treatments of the powder up to 300 ◦ C to determine the effect of temperature on the grain size and lattice parameter. Table 4 presents the results of the heat treatments on the powder that was milled for 4 h. The lattice parameter of the powder after 1 h of annealing at 100 ◦ C has returned to that of pure aluminum which indicates that the SSE achieved during milling has been annihilated. This is observed in the samples annealed at 200 ◦ C and 300 ◦ C as well. Significant grain growth was observed from the XRD calculations. The grain size increased from 47 nm to 80 nm after 1 h of annealing at 100 ◦ C with a reduction of lattice strain to 0.04%. The
Fig. 8. Distribution of the grains measured from the dark field images after (a) annealing at 100 ◦ C, (b) 200 ◦ C, and (c) 300 ◦ C for 1 h.
grain size continued to increase to 149 nm (0.03% lattice strain) and 217 nm (0.02% lattice strain) at annealing temperatures of 200 ◦ C and 300 ◦ C, respectively. Fig. 8 contains the grain size distributions of the grains after annealing for 1 h at 100 ◦ C, 200 ◦ C, and 300 ◦ C measured using dark field micrographs. The average grain size according to the TEM micrographs increased to 64 nm after annealing at 100 ◦ C, 71 nm after annealing at 200 ◦ C and 91 nm after annealing at 300 ◦ C, respectively. Zhou et al. measured the thermal stability of pure nanocrystalline Al at various temperatures [27]. It is hypothesized that the lower grain size stability is due to the
Table 4 The grain size, lattice parameter and solid solution for 1 h heat treatments at various temperatures. Condition
Grain size (nm)
Lattice strain (%)
After milling Annealed 100 ◦ C Annealed 200 ◦ C Annealed 300 ◦ C
47 80 149 217
0.09 0.04 0.03 0.02
± ± ± ±
0.009 0.004 0.003 0.002
Lattice parameter (Å) 0.40439 0.40498 0.40515 0.40502
± ± ± ±
0.0016 0.0009 0.0005 0.0010
S.S. (at.%) 1.72 0.00 0.00 0.00
J. Milligan et al. / Materials Science and Engineering A 508 (2009) 43–49
shorter milling time. With the shorter time there is less reaction with the liquid nitrogen to form the AlN particles which pin grain boundaries and promote good thermal stability of the grains [3].
[2] [3] [4] [5]
4. Conclusions
[6]
Nanostructured Al–Si alloy was produced by mechanical attrition in a liquid nitrogen slurry. After 4 h of milling time the grain size reached 46 nm with a lattice strain of 0.09% according to the XRD spectrum and this was verified by TEM. The grain size determined from TEM dark field images was 38.8 nm and the diffraction pattern was found to have formed rings. The width of the rings was used to calculate a distribution of the aluminum and silicon grains showing an average grain size of 26 nm and 19 nm, respectively. Extension of the solubility limit of silicon in aluminum was achieved during cryomilling up to 1.72 at.% in solid solution after 4 h of milling. This extension of the solubility limit can be attributed to the excellent mixing on the atomic level that occurs during the welding, fracturing and rewelding of the particles throughout milling. The extended solubility was removed after 1 h of annealing at 100 ◦ C and grain growth was observed when the temperature is further increased.
[7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]
Acknowledgments [23]
The author would like to thank McGill University, REGAL and NSERC for their financial support of this project. References [1] M.A. Meyers, A. Mishra, D.J. Benson, Prog. Mater. Sci. 51 (2006) 427–556.
[24] [25] [26]
[27]
49
E. Bonetti, L. Pasquini, E. Sampaolesi, Nanostruct. Mater. 9 (1997) 611. D.B. Witkin, E.J. Lavernia, Prog. Mater. Sci. 51 (2006) 1–60. C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1–184. C.R. Clark, C. Suryanarayana, F.H. Froes, P/M Particul. Mater. 14–17 (1995) 135–143. K.H. Chung, J. He, D.H. Shin, J.M. Schoenung, Mater. Sci. Eng. A 356 (2003) 23–31. H.J. Fecht, Nanostruct. Mater. 6 (1995) 33–42. C.C. Koch, Rev. Adv. Mater. Sci. 5 (2003) 91–99. Z. Lee, F. Zhou, R.Z. Valiev, E.J. Lavernia, S.R. Nutt, Scr. Mater. 51 (2004) 209–214. R.W. Siegel, G.E. Fougere, Nanostruct. Mater. 6 (1995) 205–216. F. Zhou, R. Rodriguez, E.J. Lavernia, Mater. Sci. Forum 386–388 (2002) 409–414. F. Zhou, D. Witkin, S.R. Nutt, E.J. Lavernia, Mater. Sci. Eng. A 375–377 (2004) 917–921. J.R. Davis, Davis & Associates, Aluminum and Aluminum Alloys, ASM International, 1993. O. Uzun, T. Karaaslan, M. Keskin, Turk. J. Phys. 25 (2001) 455–466. A. Bendijk, R. Delhez, L. Katgerman, Th.H. De keijser, E.J. Mittemeijer, N.M. Van Der Pers, J. Mater. Sci. 15 (1980) 2803–2810. M.C. Gui, J. Jia, G.S. Song, Q.C. Li, J. Mater. Sci. Technol. 15 (1999) 225–228. M. Rajabi, A. Simchi, M. Vahidi, P. Davami, J. Alloys Compd. 466 (2008) 111–118. P. Volovitch, J.E. Masse, A. Fabre, L. Barrallier, W. Saikaly, Surf. Coat. Technol. 202 (2008) 4901–4914. T.J. Konno, R. Sinclair, Mater. Chem. Phys. 35 (1993) 99–113. Y. Gao, C. Wang, Q. Lin, H. Liu, M. Yao, Surf. Coat. Technol. 201 (2006) 2701–2706. R.A. Varin, J. Bystrzycki, A. Calka, Intermetallics 7 (1999) 785–796. H.P. Klug, L. Alexander, X-ray Diffraction Procedures for Polycrystalline and Amorphous Materials, 2nd ed., Wiley, New York, 1974. D.B. Williams, C.B. Carter, Transmission Electron Microscopy: A Textbook for Materials Science, Springer, 1996. ˜ F. Sun, P. Rojas, A. Zúniga, E.J. Lavernia, Mater. Sci. Eng. A 430 (2006) 90–97. W.B. Pearson, The Crystal Chemistry and Physics of Metals and Alloys, Wiley Interscience, New York, 1972, p. 135. B. Predel, O. Madelung, Landolt-Bornstein, Group IV Physical Chemistry – Phase Equilibria, Crystallographic and Thermodynamic Data of Binary Alloys, vol. 5 – Electronic Materials and Semiconductors, Springer, Verlag, 1998. F. Zhou, J. Lee, S. Dallek, E.J. Lavernia, J. Mater. Res. 16 (2001) 3451–3458.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Nanocrystalline eutectic Al–Si alloy produced by cryomilling J. Milligan, R. Vintila, M. Brochu ∗ Aluminium Research Centre - REGAL, Department of Mining and Materials Engineering, McGill University, 3610 University St., Montreal, Quebec, Canada H3A 2B2
a r t i c l e
i n f o
Article history: Received 8 August 2008 Received in revised form 8 December 2008 Accepted 10 December 2008 Keywords: Aluminum Silicon Eutectic Cryomilling Nanocrystalline
a b s t r a c t A nanocrystalline Al–Si eutectic alloy was produced by mechanical attrition under cryogenic conditions. The resulting alloy was characterized using X-ray diffraction, scanning and transmission electron microscopy, and thermal heat treatments. The powder was milled for 4 h when the grain size reached 46 nm according to X-ray diffraction and 38 nm according to transmission electron microscopy examinations. An increase in solubility limit was achieved and was determined by measurements of the lattice parameter after milling, and quantified using Vegard’s law. The solubility reached after 4 h of milling time was 1.72 at.% of silicon in solid solution at room temperature. © 2008 Elsevier B.V. All rights reserved.
1. Introduction The mechanical properties of nanocrystalline materials have become of increasing interest because they are significantly different from those of the same material with coarse grains. Changes are seen in the mechanical properties such as hardness, strength and ductility, when the grain size is reduced to the nanometer scale [1]. The strength and plastic behaviour of metals, which strongly depends on dislocation activity, starts to differ substantially entering the nanometer regime because of the reduced availability of mobile dislocations in the smaller nanoscaled grains [2]. Nanocrystalline materials provide the opportunity to obtain properties that are unachievable with equilibrium materials. One proven synthesis technique available to produce nanocrystalline materials is mechanical milling [3]. Mechanical milling is a solidstate processing technique involving repeated welding, fracturing, and rewelding of powder particles in a low-energy ball mill to produce refined grain structures and particle sizes [4]. The milling action mixes the powders on an atomic scale leading to alloying of the powder with very fine grain sizes [5]. The mechanism of nanostructure formation during milling has been attributed to the formation of an array with a high density of dislocations, which annihilate and recombine to form high angle grain boundaries; the orientation of the grains then becomes random [4,6–8]. Materials that have a grain structure in the nanometric regime (<100 nm) have exhibited very high hardness and strength. Hard-
∗ Corresponding author. Tel.: +1 514 398 2354; fax: +1 514 398 4492. E-mail address: [email protected] (M. Brochu). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.12.017
ness values for nanocrystalline metals have been recorded up to 2–7 times higher than those of larger grained metals [2,3]. Lee et al. showed an increase of hardness in Al 5083 from 0.45 GPa up to 2 GPa when the grain size was decreased from 5 m to 200 nm [9]. Siegel and Fougere [10] summarized hardness values of 0.45 GPa for coarse grained copper which then increased to 1.2 GPa after a reduction in grain size to 25 nm. Coarse grained iron showed an increase of hardness from 1.3 GPa to 6.5 GPa once reaching a grain size of 17 nm. Nanocrystalline nickel showed an increase of hardness from 1.5 GPa to 6.5 GPa when the grain size was reduced to 20 nm from its coarse grained counterpart. Cryomilling, the mechanical attrition of powders within a cryogenic medium, has been proven to strengthen materials through refinement of the grain size and the dispersion of nanoscale particles [3]. A complete description of the process and mechanisms involved during milling was reported by Witkin and Lavernia [3]. The liquid nitrogen medium helps prevent oxidation and recrystallization during milling, which is advantageous given the large amount of free surfaces and dislocations being created during milling. Among the light elements, aluminum is becoming one of the most used elements because of its low density, high strength and numerous available alloys [5]. Previous work has been completed on the cryomilling of pure Al and aluminum alloys that are ductile [5,11,12]. Aluminum alloys that are used for wear resistance applications are based on the aluminum–silicon alloy system [13]. The silicon content in the aluminum alloy imparts excellent wear resistance, and the wear resistance can be increased with increasing refinement of the silicon particles [13]. There exist several techniques available to produce fine silicon particles and super
44
J. Milligan et al. / Materials Science and Engineering A 508 (2009) 43–49
saturated solid solutions of Al–Si alloys. These methods depend on very rapid solidification and include methods such as chill casting [14], melt spinning [15], water quenching [16] and atomization [17]. In addition, super saturated solid solutions can be achieved through mechanical milling, but the property enhancements that are achievable are limited by the amount of alloying element that can be added in solution without the formation of a second phase [5]. By increasing the amount of alloying element in solid solution in the matrix, we can, for example, further increase the effect of solid solution strengthening of the alloy. Solid solubility extension also presents the potential that upon annealing of the metastable powder, the release of the atoms in solution may nucleate nanoscale precipitates. Cryomilling works were mainly targeted to structural alloys and no systematic work was devoted to the aluminum–silicon system prior to completion of this work. Aluminum–silicon powders have potential application as feedstock for coatings and claddings [18–20]. The present study involves the cryomilling of an aluminum silicon eutectic alloy for wear applications. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were used to characterize the morphology, and grain structure of the powder. X-ray diffraction (XRD) and XLAT software were used to measure the lattice parameter and calculate the variation of solubility of the silicon solute in the aluminum matrix. Thermal heat treatments were also preformed to observe the thermal stability, i.e. evolution of grain size distribution and solid solubility limit upon exposure of the powder at higher temperature.
a Cauchy–Gaussian profile of the reflection [21,22]: ı2 (2)
K = L tan2 o
ı(2) tan o sin o
+ 16e2
(1)
in which K is a constant taken as 1, L is the grain size, e is the lattice strain, and ı(2) = B(1 − b2 /B2 ), b and B are the breadths in radians of the same peak in the reference and experimental patterns, respectively. Lattice parameters were determined using XLAT software which uses a least squares fit for the precise refinement of cell constants. Tungsten was used as the diffraction standard for the XRD testing. Bright field and dark field transmission electron micrographs were acquired using a Philips CM200 microscope with a beam energy of 200 keV. Grain sizes were measured from dark field micrographs taking the average of six measurements of diameter for each grain. Grain size distributions were also measured from the varying width of the diffraction rings of the selected area diffraction patterns according to Eq. (2) [23]: D=
rd r
(2)
where D is the grain size, r is the radius of the ring, d is the interplanar spacing of the respective ring, and r is the width of the ring. To verify the thermal stability of the nanostructured powders, samples were heat treated in a tube furnace under a high-purity argon atmosphere at 100 ◦ C, 200 ◦ C, and 300 ◦ C for 1 h and were characterized for solid solubility and grain size. 3. Results and discussion
2. Experimental procedures The starting powder was a −325 mesh, atomized eutectic aluminum silicon alloy from Valimet Inc. The composition was measured by ICP and is presented in Table 1. The starting powder has an average particle size of 32 m and a starting grain size of 2 m. The cryogenic milling experiments were completed using a Union Process 1-S attrition mill in a stainless steel vial under a rotational speed of 180 rpm. Stainless steel balls, with a diameter of 4.85 mm, were used as grinding media and a powder to ball weight ratio of 1:32 was used. To prevent adhesion of the powder to the balls, tank and attritor, and to control the fracturing events, 0.25 wt% of stearic acid (CH3 (CH2 )16 CO2 H) was added as a process control agent. The powder was milled for 4 h. Liquid nitrogen was continuously added to the vial to maintain constant slurry during milling, and to maintain a constant temperature of −196 ◦ C. Particle size distributions of the powder after 0, 2, and 4 h of milling were measured using a HORIBA LA-920 particle size analyzer. The evolution of the morphology of the powder after various stage of milling was characterized using secondary electron micrographs acquired using a Hitachi VP-SEM 3000SN microscope. Grain size and lattice parameter measurements were made with X-ray diffraction using Phillips PW1070 diffractometer (Cu K␣ = 1.54056 Å). A pattern was obtained from 35◦ to 110◦ with a step size of 0.005◦ and a dwell time of 0.2 s/step. The grain size and lattice strain were determined from the XRD peak broadening using Eq. (1), assuming Table 1 The chemical composition of the starting powder measured by ICP. Element
Amount (wt%)
Al Si Fe Cu Mg Mn Cr Zn
Bal. 12.07 0.1166 0.00613 0.0097 0.0053 0.00004 0.0021
Fig. 1 shows the powder morphology evolution throughout the milling process, which changes from spherical to flakes. The initial morphology of the master alloy was spherical from the atomization process as depicted in Fig. 1a. During ball milling the powder particles are continuously being flattened, cold welded and fractured. Due to the initial powder matrix being soft, their tendency to flatten and form larger particles is high during the initial stages of milling (shown in Fig. 1b); as the process continues the particles become work hardened and eventually fracture [4] (shown in Fig. 1c). The average particle size distribution of the powder throughout the milling process was measured. Fig. 2 shows the variation of the particle size with milling time. The average starting particle size was 32 m. After 2 h of milling, the average particle size increased to 61 m, and then decreased to 18 m after 4 h of milling time. After 2 h of milling time, the average particle size increases as the particles are flattened, which can be seen in Fig. 1b. When fracture occurs the particle size of the powder then decreases [24]. The results obtained are showing a faster reduction in particle size than what is generally observed for fully metallic powders. In particular, a reported minimum milling time of 10 h is required to yield similar particle size in an Al–Mg system [11] as seen in Fig. 2. The formation of AlN and SiN are thermodynamically favorable during the milling process, however high resolution examination was not undertaken to determine the existence and characteristics of these phases after milling. It is believed that the presence of the Si phase accelerates the deformation and fracture of the powders either through the formation of nitrides during milling or through the presence of a second hard phase in the mill. Similar results have been observed by Chung et al. [6], where they added a small volume fraction of AlN into pure Ni powders. A particle size reduction of 63% was observed after 8 h of milling time with the addition of the hard phase to the Ni powders. In the Al–Si system compared to the Al–Mg system a 52% reduction in particle size was observed after 4 h of milling time. The powder was scanned using X-ray diffraction to determine the grain size, lattice strain and lattice parameter evolution during cryomilling. Fig. 3 shows the starting powder used as a reference, where both the aluminum and silicon phases are observable. The
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Fig. 1. Morphology of the powder (a) prior to milling (b) after 2 h of milling and (c) after 4 h of milling time.
aluminum and silicon phases correspond to JCPDS patterns 040787 and 05-0565, respectively. Broadening of the fcc-Al peaks is observed in the XRD spectrum after 2 h and 4 h of milling time. With increasing milling time, the peaks broaden and loose intensity as the grain size and lattice strain evolve during milling. The broadening of the Al peaks allows the grain size and lattice strain to be determined based on the FWHM of the peaks. Table 2 lists the
Fig. 2. Change in particle size versus milling time.
Table 3 The change in lattice parameter and solid solubility during milling. Milling time (h)
Lattice parameter (Å)
Si in S.S. (at.%)
0 2 4
4.0494 ± 0.0005 4.0479 ± 0.0019 4.0439 ± 0.0016
0 0.47 1.72
values for the grain size and lattice strain for the different milling times. The grain size of the Al–Si alloy was reduced to 46 nm after 4 h of milling time with a lattice strain of 0.09% which is comparable to the grain size and strain reported for cryomilling of pure aluminum, where a 47 nm grain size and a lattice strain of 0.10% strain were obtained after 4 h of milling [12]. The lattice parameter of the aluminum matrix was measured using the shifts in the XRD spectrum. According to Vegard’s law the lattice parameter of an alloy varies linearly with the solid solubility. The solid solubility increases with decreasing lattice parameter of aluminum (Al = 0.4094 nm) to the lattice parameter of an imaginary cubic closed packed silicon (Si = 0.3731 nm) [15,25]. Fig. 4 depicts the Vegard’s law (bold line) relationship for the aluminum–silicon system. The results on the extension of the solubility limit of silicon in the aluminum matrix are presented in Table 3 and a maximum solubility of 1.72 at.% was obtained during cryomilling. It is worth mentioning that the equilibrium solid solution limit for this system
Table 2 The average grain size and lattice strain in nc Al and nc Al–12Si. Material
Milling time (h)
Grain size (nm)
Lattice strain (%)
T (K)
Reference
Al–12Si Al–12Si Al
2 4 4
105 46 47
0.22 0.09 0.10
80 80 90
This paper This paper [12]
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Fig. 3. XRD spectrums after various milling times for the Al–12Si powder.
is 1.5 at.% at 577 ◦ C and 0% at room temperature [26]. Fig. 4 also contains comparable results on the extension of the solubility limit obtained using a SPEX 8000 shaker mill after 17 h of milling with a ball to powder ratio of 10:1. Their results are showing an increase in solubility limit up to 1.3 at.% of silicon when a starting powder containing 10 at.% of silicon was used [5]. It was also reported that the maximum solubility limit achieved with a starting solute content of 30 at.% silicon was approximately 4.5 at.% [5]. Based on the results of solid solubility extension (SSE) in the Al–Si obtained in this work and from literature data, the SSE extension seems to be linearly proportional to the solute concentration in the Al–Si system. To further study the thermal stability of the SSE, heat treatments at 50 ◦ C, 75 ◦ C, and 100 ◦ C for 1 h were performed. Fig. 5 presents the modified aluminum silicon phase diagram containing the data of the decomposition of the metastable solid solubility as a function of temperature for the alloy Al–12Si. After a heat treatment of 1 h at room temperature, approximately 1 at.% of silicon remains in solid solution in the aluminum. The solid solubility decreased down to 0.72 at.% after 1 h at 50 ◦ C, down to 0.16 at.% after 1 h at 75 ◦ C, and finally to complete dissociation after 1 h of heat treatment at 100 ◦ C.
TEM analysis of the cryomilled powder was completed to observe the grain structure after milling. The milling time used for this powder was sufficient to produce a random dispersion of small grains through the sample which can be observed in the bright and dark field micrographs shown in Fig. 6a and b. Fig. 6c shows the selected area diffraction pattern of the cryomilled powder showing a ring pattern indicating that the powder has reached a nanoscaled grain size. The rings corresponding to aluminum and silicon are completely distinguishable. The grain size distribution measurements from the TEM dark field micrographs are presented in Fig. 6d and show an average grain size of 38.8 nm. The microstructure consists mainly of small nanometer grains with some well-dispersed ultra fine grains. In order to characterize the respective grain size distribution of the Al and Si grains, as this is a limitation in the differentiation arising from the dark field images, their respective grain size distribution was measured using the width of the diffraction rings. Fig. 7 illustrates the grain size distributions measured
Fig. 4. Solid solubility versus lattice parameter for the Al–Si system according to Vegard’s law.
Fig. 5. The aluminum–silicon phase diagram with the solid solubility of nano-Al–Si (starting powder Al–12Si).
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Fig. 6. TEM bright field (a), dark field (b), and diffraction pattern (c) after 4 h of milling. (d) The distribution of grains measured from the dark field images.
from the rings in the diffraction pattern shown in Fig. 6c. The grain sizes were calculated according to Eq. (2), to distinguish the size of the aluminum grains versus the silicon grains in the cryomilled powder. Fig. 7a shows the distribution of the aluminum grains determined from the (1 1 1) and (2 0 0) rings. An average aluminum grain size of 26 nm was measured. The distribution of the silicon grains were determined from the (1 1 1), (2 2 0) and (3 1 1) rings and have an average of 19 nm (seen in Fig. 7b). The grain size measured is dependant on the instantaneous ring width, larger widths indicating a smaller grain. As the number of grains
increases the few large grains that remain will produce a small ring width that may be convoluted beneath the section of the ring produced by a smaller grain at the same rotation in the selected area diffraction pattern. As such a uniform grain size distribution is required to apply this measurement technique. Combining the two distributions together (shown in Fig. 7c), the weighted average of the aluminum and silicon grains between the grain sizes of 0.1–70 nm is 24 nm, which agrees with the average grain size measured in the dark field images in the same range. This distribution is comparable to that found for pure aluminum that
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Fig. 7. Distribution of the aluminum grains (a), the silicon grains (b) and aluminum + silicon grains (c) measured from the ring pattern after 4 h of milling time.
has been cryomilled [12,27]. The distributions measured show that during milling there is a simultaneous refinement of both the aluminum and silicon phases present; this effect was also observed by Chung et al. [6] during milling of a composite system. The thermal stability of the powder was investigated through heat treatments of the powder up to 300 ◦ C to determine the effect of temperature on the grain size and lattice parameter. Table 4 presents the results of the heat treatments on the powder that was milled for 4 h. The lattice parameter of the powder after 1 h of annealing at 100 ◦ C has returned to that of pure aluminum which indicates that the SSE achieved during milling has been annihilated. This is observed in the samples annealed at 200 ◦ C and 300 ◦ C as well. Significant grain growth was observed from the XRD calculations. The grain size increased from 47 nm to 80 nm after 1 h of annealing at 100 ◦ C with a reduction of lattice strain to 0.04%. The
Fig. 8. Distribution of the grains measured from the dark field images after (a) annealing at 100 ◦ C, (b) 200 ◦ C, and (c) 300 ◦ C for 1 h.
grain size continued to increase to 149 nm (0.03% lattice strain) and 217 nm (0.02% lattice strain) at annealing temperatures of 200 ◦ C and 300 ◦ C, respectively. Fig. 8 contains the grain size distributions of the grains after annealing for 1 h at 100 ◦ C, 200 ◦ C, and 300 ◦ C measured using dark field micrographs. The average grain size according to the TEM micrographs increased to 64 nm after annealing at 100 ◦ C, 71 nm after annealing at 200 ◦ C and 91 nm after annealing at 300 ◦ C, respectively. Zhou et al. measured the thermal stability of pure nanocrystalline Al at various temperatures [27]. It is hypothesized that the lower grain size stability is due to the
Table 4 The grain size, lattice parameter and solid solution for 1 h heat treatments at various temperatures. Condition
Grain size (nm)
Lattice strain (%)
After milling Annealed 100 ◦ C Annealed 200 ◦ C Annealed 300 ◦ C
47 80 149 217
0.09 0.04 0.03 0.02
± ± ± ±
0.009 0.004 0.003 0.002
Lattice parameter (Å) 0.40439 0.40498 0.40515 0.40502
± ± ± ±
0.0016 0.0009 0.0005 0.0010
S.S. (at.%) 1.72 0.00 0.00 0.00
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shorter milling time. With the shorter time there is less reaction with the liquid nitrogen to form the AlN particles which pin grain boundaries and promote good thermal stability of the grains [3].
[2] [3] [4] [5]
4. Conclusions
[6]
Nanostructured Al–Si alloy was produced by mechanical attrition in a liquid nitrogen slurry. After 4 h of milling time the grain size reached 46 nm with a lattice strain of 0.09% according to the XRD spectrum and this was verified by TEM. The grain size determined from TEM dark field images was 38.8 nm and the diffraction pattern was found to have formed rings. The width of the rings was used to calculate a distribution of the aluminum and silicon grains showing an average grain size of 26 nm and 19 nm, respectively. Extension of the solubility limit of silicon in aluminum was achieved during cryomilling up to 1.72 at.% in solid solution after 4 h of milling. This extension of the solubility limit can be attributed to the excellent mixing on the atomic level that occurs during the welding, fracturing and rewelding of the particles throughout milling. The extended solubility was removed after 1 h of annealing at 100 ◦ C and grain growth was observed when the temperature is further increased.
[7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]
Acknowledgments [23]
The author would like to thank McGill University, REGAL and NSERC for their financial support of this project. References [1] M.A. Meyers, A. Mishra, D.J. Benson, Prog. Mater. Sci. 51 (2006) 427–556.
[24] [25] [26]
[27]
49
E. Bonetti, L. Pasquini, E. Sampaolesi, Nanostruct. Mater. 9 (1997) 611. D.B. Witkin, E.J. Lavernia, Prog. Mater. Sci. 51 (2006) 1–60. C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1–184. C.R. Clark, C. Suryanarayana, F.H. Froes, P/M Particul. Mater. 14–17 (1995) 135–143. K.H. Chung, J. He, D.H. Shin, J.M. Schoenung, Mater. Sci. Eng. A 356 (2003) 23–31. H.J. Fecht, Nanostruct. Mater. 6 (1995) 33–42. C.C. Koch, Rev. Adv. Mater. Sci. 5 (2003) 91–99. Z. Lee, F. Zhou, R.Z. Valiev, E.J. Lavernia, S.R. Nutt, Scr. Mater. 51 (2004) 209–214. R.W. Siegel, G.E. Fougere, Nanostruct. Mater. 6 (1995) 205–216. F. Zhou, R. Rodriguez, E.J. Lavernia, Mater. Sci. Forum 386–388 (2002) 409–414. F. Zhou, D. Witkin, S.R. Nutt, E.J. Lavernia, Mater. Sci. Eng. A 375–377 (2004) 917–921. J.R. Davis, Davis & Associates, Aluminum and Aluminum Alloys, ASM International, 1993. O. Uzun, T. Karaaslan, M. Keskin, Turk. J. Phys. 25 (2001) 455–466. A. Bendijk, R. Delhez, L. Katgerman, Th.H. De keijser, E.J. Mittemeijer, N.M. Van Der Pers, J. Mater. Sci. 15 (1980) 2803–2810. M.C. Gui, J. Jia, G.S. Song, Q.C. Li, J. Mater. Sci. Technol. 15 (1999) 225–228. M. Rajabi, A. Simchi, M. Vahidi, P. Davami, J. Alloys Compd. 466 (2008) 111–118. P. Volovitch, J.E. Masse, A. Fabre, L. Barrallier, W. Saikaly, Surf. Coat. Technol. 202 (2008) 4901–4914. T.J. Konno, R. Sinclair, Mater. Chem. Phys. 35 (1993) 99–113. Y. Gao, C. Wang, Q. Lin, H. Liu, M. Yao, Surf. Coat. Technol. 201 (2006) 2701–2706. R.A. Varin, J. Bystrzycki, A. Calka, Intermetallics 7 (1999) 785–796. H.P. Klug, L. Alexander, X-ray Diffraction Procedures for Polycrystalline and Amorphous Materials, 2nd ed., Wiley, New York, 1974. D.B. Williams, C.B. Carter, Transmission Electron Microscopy: A Textbook for Materials Science, Springer, 1996. ˜ F. Sun, P. Rojas, A. Zúniga, E.J. Lavernia, Mater. Sci. Eng. A 430 (2006) 90–97. W.B. Pearson, The Crystal Chemistry and Physics of Metals and Alloys, Wiley Interscience, New York, 1972, p. 135. B. Predel, O. Madelung, Landolt-Bornstein, Group IV Physical Chemistry – Phase Equilibria, Crystallographic and Thermodynamic Data of Binary Alloys, vol. 5 – Electronic Materials and Semiconductors, Springer, Verlag, 1998. F. Zhou, J. Lee, S. Dallek, E.J. Lavernia, J. Mater. Res. 16 (2001) 3451–3458.