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Microstructure and mechanical properties of Al–Zr nanocomposite materials
Materials Science and Engineering A 518 (2009) 100–107
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical properties of Al–Zr nanocomposite materials B. Srinivasarao a,b , C. Suryanarayana c , K. Oh-ishi b , K. Hono b,a,∗ a
Graduate School of Pure and Applied Sciences, University of Tsukuba, 1-2-1 Sengen, Tsukuba 305-0047, Japan National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Japan c Department of Mechanical, Materials and Aerospace Engineering, University of Central Florida, Orlando, FL 32816-2450, USA b
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Article history: Received 7 February 2009 Received in revised form 11 April 2009 Accepted 17 April 2009 Keywords: Aluminum alloy Nanocrystal Nanocomposites Mechanical alloying Spark plasma sintering
a b s t r a c t Three different processing routes were explored to develop Al–Zr nanocomposite alloys using mechanical alloying and spark plasma sintering methods. Depending on the route of milling adopted, the powder in the as-milled condition consisted of either a solid solution of Zr in Al or a mixture of Al-solid solution and Al3 Zr (L12 ) phases. The alloys after sintering consisted of Al and Al3 Zr (L12 ) with grain sizes of less than 100 nm. These nanocomposite alloys exhibited a high compressive strength of 1 GPa with 10% plasticity. The high strength observed in these alloys was explained on the basis of the retention of nanometer sized grains and also the fine dispersion of the L12 phase. On the other hand, the good amount of plasticity was explained to be due to excellent bonding between the powder particles and the presence of coarse Al grains in the matrix. © 2009 Elsevier B.V. All rights reserved.
1. Introduction The development of high strength heat resistant Al-based alloys (both monolithic and composite) has gained considerable interest recently [1–4]. For this purpose, transition metals such as Fe, Cr, Mo, Sc and Zr were added since these elements have a low solid solubility in Al, low diffusivity and additionally they form high melting temperature intermetallic phases with Al. Low solid solubility and diffusivity makes the growth of these intermetallic phases very difficult, making these alloys suitable for high temperature applications. Although the solid solubility of these transition elements in Al is limited, it can be significantly increased by rapid solidification or mechanical alloying methods and this supersaturated solid solution is stable without decomposition up to 320 ◦ C. Further heat treatment can lead to precipitation hardening with a large volume fraction of the precipitate, which enhances the strength of the alloy. Among these transition elements, Zr has the lowest diffusivity and also forms a metastable Al3 Zr (L12 ) phase which displays a very small lattice mismatch with the Al-matrix and therefore has high resistance for coarsening [5–9]. Rittner et al. [6] have synthesized nanocrystalline Al–Zr alloys by the inert gas condensation technique and observed a maximum hardness of 3 GPa. Subsequently Sasaki et al. [7] reported a high tensile strength of 800 MPa for ultrafine grained (260 nm) Al95.3 Zr4 Fe0.7 alloy prepared by physi-
∗ Corresponding author. E-mail address: [email protected] (K. Hono). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.04.032
cal vapor deposition technique. This alloy had also exhibited good high temperature mechanical properties. However, the samples processed by this method were only plates of less than 0.6 mm in thickness. Mechanical alloying (MA) is a process that can produce supersaturated solid solutions even in immiscible systems. The process involves repeated cold welding, fracturing, and rewelding of powder particles in a high-energy ball mill. This technique is capable of synthesizing supersaturated solid solutions, nanocrystalline alloys, amorphous alloys, and their composites with a variety of microstructural modifications [10–14]. Since this technique can be used to synthesize a large amount of powder at a time, it is an industrially viable process. The powders of supersaturated solid solutions or amorphous phase can be consolidated to bulk materials by subsequent consolidation processes like hot pressing [15–17], extrusion [18–22], and spark plasma sintering (SPS) [23–25]. Recently, Sasaki et al. [25] reported an ultrahigh strength bulk nanocrystalline Al–Fe alloy processed by MA and subsequent SPS process, which exhibited yield strength of higher than 1 GPa with excellent heat resistance. In line with the work, we have explored further possibility to process high strength heat resistant aluminum alloys using the Al–Zr system. 2. Experimental procedure Pure Al (99.99%), Pure Zr (99.9%) and home made (arc melted and crushed) Al3 Zr (D023 ) powders were used in this study. MA was conducted in a Fritsch P-4 planetary ball mill using 10 mm diameter
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440C stainless steel balls in a 440C stainless steel vial under argon atmosphere at a speed of 300 rpm. All of the powder handling was carried out in a glove box having Ar atmosphere. Three different methods were employed to make the nanocrystalline composite powders. Route I: Pure elemental Al and Zr powders (resulting total alloy compositions of 0.5, 1.5 and 3 at.% Zr) were milled with a ball to powder weight ratio of 10:1 for a milling time of 300 h to obtain a solid solution of Zr in Al. Route II: Pure elemental Al powder and D023 -Al3 Zr were used (total alloy composition were 2.1, 3, 4 and 4.5 at.% Zr) with a ball to powder weight ratio of 10:1 for a milling time of 120 h to synthesize a solid solution of Zr in Al. Route III: Prealloyed L12 -Al3 Zr was first prepared by mixing stoichiometric amounts of pure Al and Zr powders and milling them for 10 h. Subsequently, different amounts of this resultant powder (10, 25 and 40 wt.%) were mixed with pure Al powder and milled with a ball to powder weight ratio of 20:1 for 5 h to make a nanocomposite. (The total alloy compositions after converting into atomic percentages are 1.63, 4.33 and 7.37 at.% Zr, respectively.) Ethanol (1.5–4 ml or 6–15 wt.%) was used as a process control agent (PCA) in all the experiments. Following milling, the powder was packed in a tungsten carbide die (coated with liquid BN spray) with an inner diameter of 10 mm and compacted by an SPS machine (Sumitomo Coal Mining Company, Model 1050) under a vacuum of <10−3 Pa. Sintering was carried out as follows: at first a stress of 445 MPa was applied, and then the material was heated to desired temperature (i.e., 400–490 ◦ C) at a heating rate of 80 ◦ C/min and held for 10 min. After finishing sintering the load was removed and the sample was water cooled. The final sintered compact had dimensions of 5 mm thickness and 10 mm diameter. The surface of the compact was cleaned by grinding on emery paper to mirror finish and then the density was measured using the Archimedes method with water as a liquid medium. The density of all the compacts was found to be >97% of the theoretical value. Compression tests were performed at a quasi-static strain rate of 1 × 10−4 s−1 in an Instron machine using rectangular samples of 2 mm width, 2 mm breadth and 4 mm height. X-ray diffraction (XRD) analysis was carried out with a Rigaku RINT-2500 X-ray diffractometer using Cu K␣ radiation. Both scanning electron microscopy (1540EsB CrossBeam) and transmission electron microscopy (TEM) (Philips CM200 and TECNAI G2 F30) techniques were used to characterize the microstructural features of the alloys. 3. Results and discussion Figs. 1(a) and 2(a) show X-ray diffraction profile of ball milled powders in Route I and Route II, respectively. In both routes, either the elemental zirconium (in Route I) or intermetallic D023 -Al3 Zr (in Route II) is gradually dissolved into the Al-matrix with increase in milling time to form a solid solution. Note that the amount of the D023 -Al3 Zr added in Route II was such that the total Zr content in the alloy was 3 at.%. The solid solution was formed after 300 h of milling in Route I, while it took only 120 h in Route II. This suggests that the time required for the elemental Zr to completely dissolve was much longer than that for the intermetallic phase. This is because interdiffusion between different phases predominantly occurs only when the grain size of the two phases reaches a critical value during the MA process [11]. Since Al3 Zr is a brittle material compared to elemental Zr, it is easy for the intermetallic phase to reach the minimum critical grain size in a very short time, and hence diffusion occurs quickly. It is also possible that the initial formation of ZrH2 (formed due to reaction between PCA and Zr) phase in the
Fig. 1. Typical XRD profiles of powders made by Route I. (a) Evolution of Al-solid solution during milling of Al–0.5 at.% Zr and (b) summary of the various alloys after 300 h of milling.
case of Route I, has slowed down the diffusion of Zr atoms into the Al lattice to form the solid solution [26]. In case of Route I, the elemental Zr was reacting with the PCA and forming as ZrH2 and also because of the long milling times the amount of PCA needed is very high (i.e., 4 ml). Whereas in the case of Route II, the formation of ZrH2 is not observed and also the milling time was less (i.e., 120 instead of 300 h) due to which the amount of PCA needed to avoid the agglomeration of powders is less (i.e., minimum of 1.5 ml). Fig. 1(b) shows the XRD patterns of the blended elemental powders with different Zr contents processed for 300 h in Route I. It is to be noted that up to 3 at.% Zr could be dissolved in Al to form the solid solution after 300 h of milling. Note that the alloys with higher Zr contents exhibit broader diffraction peaks, suggesting finer grain size with Zr content. The average grain size of the solid solution phase obtained after 300 h of milling in Route I was calculated from the XRD profiles by Voigt function [27] as 67, 62 and 42 nm for 0.5, 1.5 and 3 at.% Zr, respectively. The structural evolution in the powders with different amounts of Al3 Zr (leading to different total Zr contents in the alloy) milled for 120 h following the Route II process is presented in Fig. 2(b) for the overall Zr content of 2.1, 3, 4.0, and 4.7 at.%. Note that the solid solution formation after 120 h is observed up to 4 at.% Zr. Beyond this composition, the metastable L12 -Al3 Zr phase appeared after 120 h of milling. Because of the presence of metastable phase and overlap of the peak positions it is difficult to calculate the grain size from the XRD plot, instead it was calculated from the dark field electron microscopy images of the powder particles. Typical TEM
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Fig. 2. Typical XRD profiles of powders made by Route II. (a) Evolution of Al-solid solution during milling of Al–3 at.% Zr and (b) summary of the various alloys after 120 h of milling.
micrographs of BF and DF was shown in Fig. 3 for 2.1 at.% Zr alloy powder after 120 h of milling and the figure shows a nanocrystalline solid solution which is consistent with Fig. 2(b). The average grain sizes were 48, 32, 35 and 30 nm respectively for 2.1, 3, 4 and 4.7 at.% Zr alloys. This result is in consistent with the previous reports that the grain size of milled powders is smaller for the higher the solute content in the alloys [10,11]. This is because the alloy with the higher solute content becomes stronger and harder due to solid solution strengthening. But the smaller grain size for the 3 at.% Zr was due to the use of a higher amount of PCA, viz., 3 ml ethanol whereas only 1.5 ml PCA was used for all other powders in this route. In Route III, at first L12 -Al3 Zr alloy powder was prepared by mixing pure elemental Zr and Al with 1.5 ml PCA and Fig. 4(a) shows the XRD pattern of the powder after 10 h of milling. From the figure, it can be seen that Al, L12 -Al3 Zr and ZrH2 phases were present and no Zr peaks were observed. The average grain size of the L12 Al3 Zr phase was estimated to be about 9 nm from the dark field TEM images. Secondly, this powder was mixed with Al in different amounts (10, 25 and 40 wt.%) and milled together for extra 5 h with an extra 0.5 ml of PCA and the XRD profiles were shown in Fig. 4(b). It may be noted from the figure that the powder contains Al, L12 -Al3 Zr and ZrH2 phases. The absence of elemental Zr in these powders suggests that all the Zr was completely used up to form the L12 -Al3 Zr and ZrH2 phases. Alternately, it is possible that its volume fraction is so low that it can be hardly detected by the XRD method. Fig. 5 shows the TEM bright field and corresponding dark field image of Al–25 wt.% Al3 Zr nanocomposite powder. The dark
Fig. 3. Typical TEM micrograph of 2.1 at.% Zr alloy powder after 120 h of milling (Route II). (a) Bright field micrograph and (b) dark field micrograph.
field image clearly shows two regions where clusters of nano-Al3 Zr clustered grains (also shown by broken lines in Fig. 5(a)) in one region and the second region contains Al-matrix grains. The matrix grain sizes were found to be 90, 62 and 48 nm for 10, 25 and 40 wt.% Al3 Zr additions. Fig. 6 represents the XRD patterns of the bulk samples consolidated by SPS at different temperatures. To achieve good high temperature mechanical properties it is important that the alloy contains the metastable L12 -Al3 Zr phase than the equilibrium D023 -Al3 Zr from the solid solution in the sintered compact because high creep resistance can be obtained if the matrix/precipitate interface has smaller lattice mismatch [28]. Fig. 6(a) and (b) shows XRD profiles of the sintered samples containing 3 at.% Zr prepared by Route I (300 h) and 2.1 at.% Zr prepared by Route II (120 h), respectively. From these two figures it is clear that L12 -Al3 Zr phase precipitated from the supersaturated solid solution during sintering up to a critical temperature, above which the equilibrium D023 -Al3 Zr phase had formed. This critical temperature was 490 ◦ C for the Al–3 at.% Zr alloy processed by Route I and 450 ◦ C for the
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Fig. 4. Typical XRD profiles of powders made by Route III. (a) L12 -Al3 Zr phase after 10 h of milling and (b) summary of the various alloys after 15 h of milling.
Al–2.1 at.% Zr alloy processed by Route II. A similar observation was reported in the case of rapidly solidified Al–Zr alloys [5,29] and vapor deposited Al–Ti–Cr alloys [30], which can be understood in terms of the lower activation energy required to form the metastable L12 phase due to coherent precipitation with a small interfacial energy than the incoherent D023 phase. Moon and Lee [14] reported that the thermal stability of the L12 -Al3 Zr phase was
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increased when a ternary element such as Cu, Ni and Mn is added. In case of Route I process, because of large milling time and because the milling media is 440C stainless steel, there is a possibility that Ni and Mn from the milling media could be present in the milled powder. These additions might have made the L12 -Al3 Zr phase stable up to higher temperatures (viz., 490 ◦ C). Fig. 6(c) and (d) shows the phases present in the sintered alloys with different Zr concentrations made by Route I and Route II, respectively. The XRD patterns show the presence of metastable L12 phase after sintering at 470 and 400 ◦ C, respectively. These results should be compared with that of Fig. 6(e), which shows the typical XRD pattern of the Al–25 wt.% Al3 Zr alloy sintered at 450 ◦ C; this figure reveals a number of phases, viz., Al, Al3 Zr (L12 ) ZrC and iron aluminides. ZrC was formed by the reaction between the carbon present in the PCA and Zr, whereas the iron aluminides are due to contamination of iron from the milling media. Nanocrystalline L12 -Al3 Zr has very high Vickers hardness value approaching 10 GPa [31] whereas Al–Zr solid solution was less than 400 Hv (Route I and Route II). Because of the high hardness of L12 phase in Route III the iron contamination will be higher even though the milling time is smaller (15 h) than Route I and Route II processes (120 and 300 h). Fig. 7 shows the SEM backscattered micrographs of sintered compacts prepared by Route I and Route II. The black regions in Fig. 7(a) and (b) correspond to pure Al formed during SPS and the grey region corresponds to the milled powder particles bonded together. All the compacts made by Route I and subsequent sintering (up to 490 ◦ C) show poor bonding between the particles and lot of processing defects such as porosity, cracks and voids indicated by white arrows in Fig. 7(a) whereas for Route II processed and sintered samples the bonding between the particles has greatly improved and very few isolated porosity and voids were present (Fig. 7(b)). The major cause for the poor particle bonding in Route I could be due to the high powder particle strength; a result of contamination from the milling media because of long milling times (300 h) and the use of large amount of PCA (4 ml ethanol). Both these result in incorporation of oxide, carbide and other ceramic particles into the metallic matrix and consequently increase the strength of the powder. And because of the high particle strength, the particles are not sufficiently compressed during the sintering process to establish full contact between them. The presence of pure Al at the particle boundaries is considered to be due to the partial melting because of the local joule heating at the particle interfaces and is commonly observed in spark plasma sintered aluminum alloys [23–25]. During the solidification, Zr is rejected from Al due to the limited solubility, and hence coarse grains of pure Al are formed at the particle inter-
Fig. 5. TEM micrograph of Al–25 wt.% (L12 )Al3 Zr powder particle after milling. (a) Bright field image and (b) dark field image.
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Fig. 6. Typical XRD patterns of sintered samples (a) and (c) for Route I, and, (b) and (d) for Route II and (e) for Route III process.
faces. The partial melting of Al results in the density values of >97% of the theoretical value, because the melt fills the spacing between the particles. Fig. 7(c) and (d) shows the SEM backscattered high magnification images taken from Fig. 7(a) and (b), respectively. The white contrast in Fig. 7(c) and (d) indicates the presence of Al3 Zr (L12 ) phase and Al13 Fe4 phase whereas the black contrast is due to oxide particles i.e., Al2 O3 mostly came from polishing. These Al3 Zr particles are in the size range of 10–50 nm. Fig. 8 shows the SEM micrographs of the samples containing different amounts of Al3 Zr (L12 ) made by Route III and after sintering at 450 ◦ C. The figure shows a uniform microstructure and no pro-
cessing defects such as cracks, debonding and porosity are seen. The white contrast in these figures corresponds to Al3 Zr (L12 ) and Al13 Fe4 which are uniformly dispersed throughout the matrix and most of them are well below 2 m in size. Fig. 9(a) shows the bright field electron micrograph of Al–3 at.% Zr alloy made by Route II with 3 ml of ethanol as PCA. The selected area diffraction pattern (SAED) (shown as an inset) indicates the presence of Al3 Zr (L12 ) and Al-phases, confirming the results obtained by XRD, Fig. 6(d). Mostly the second phase particles were of equiaxed grain size forming at the grain boundaries of the matrix Al. Fig. 9(b) gives an idea of the matrix grain size
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Fig. 7. Typical SEM backscattered micrographs of sintered samples. (a) 3 at.% Zr (Route I) sintered at 470 ◦ C, (b) 4 at.% Zr (Route II) sintered at 400 ◦ C (c and d) High magnification back scatter images taken from (a) and (b), respectively.
which is well below 100 nm. The average matrix grain sizes calculated from the dark field images are 65, 58 and 45 nm for 0.5, 1.5 and 3 at.% Zr alloys, respectively for Route I and 85, 45, 48 and 35 nm for 2.1, 3, 4 and 4.5 at.% Zr alloys made by Route II, respectively. The average grain size of the composites made by Route I and Route II are well below 100 nm even after sintering and this small grain size is expected to confer high strength on these materials. Fig. 10(a) shows the microstructure of sintered compact of
Al–25 wt.% Al3 Zr alloy in which the dark regions correspond to cluster of Al3 Zr grains. While the clusters have sizes of typically ≤1 m, the average grain size in the Al3 Zr cluster was found to be about 10 nm, clearly seen from Fig. 10(b). The average matrix grain sizes were 200, 120 and 50 nm for 10, 25 and 40 wt.% Al3 Zr alloys, respectively. Fig. 11 shows the typical stress–strain curves obtained by compression test for the samples prepared by Routes I, II and III. For
Fig. 8. Typical SEM micrographs of sintered samples sintered at 450 ◦ C (Route III process). (a) 10 wt.%, (b) 25 wt.% and (c) 40 wt.% Al3 Zr.
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Fig. 9. Typical TEM micrograph of 3 at.% Zr sintered compact (Route II). (a) Bright field image, (b) dark field image showing nanocrystalline grain size of Al.
the case of Route I a high compressive strength of 1 GPa could be achieved in all of the compositions but they broke into pieces immediately after yielding and sometimes in the elastic regime itself (to remove the effect of brittle phase Al3 Zr-D023 on compression properties sintered samples of 480 ◦ C was used because D023 form starting from 490 ◦ C as observed from Fig. 6). This type of failure can be attributed to the improper bonding (many powder particles are barely touching each other) between the particles and also due to the presence of lot of defects (Fig. 7(a)). The defects oriented perpendicular to the compression axis will be closed during the test, but at other angles they will initiate cracks and these cracks will grow faster because of the high stress levels and therefore lead to fracture. For the Route II processed samples, the compression testing (Fig. 11(b)) results show some plasticity (∼10%). This improvement in plasticity is due to the improvement in bonding between the particles and a significant reduction of process defects (Fig. 7(b)). A high strength of more than 1 GPa was observed for 3 at.% Zr and this is because of the use of a large amount of PCA (i.e., 3 ml whereas only 1.5 ml was used in other cases) which will increase the amount of carbon, oxygen and also will affect the contamination level from the milling media. From the XRD patterns as
well as low magnification TEM studies, these carbides and oxides were not observed suggesting that these were present either in small sizes or still dissolved in the matrix. This could be due to rapid sintering by the SPS technique (the total sintering time was less than 20 min). Fig. 11(c) indicates that Route III processed nanocomposites exhibit good plasticity and this plasticity had decreased when the amount of Al3 Zr (L12 ) phase increased. All of the compression data in this study show strain softening behavior which could be due to either small size (i.e., <100 nm) or because of the use pure Al-matrix (Route III process). The high strength found in the Route III processed alloys is due to the high hardness of L12 -Al3 Zr clusters [31]. From the above results Route II and Route III processes are the most suitable methods for scaling up, because of the minimized processing time, reasonable compressive strength and plasticity. Currently, scaling up of these alloys by extrusion and their high temperature properties are under investigation. So far, we obtained the plastic strain of 10% for the strength level of 1 GPa only in the compression test. Further optimization of the alloy compositions, processing routes and microstructures is necessary to obtain tensile ductility.
Fig. 10. Typical TEM micrograph of Al–25 wt.% Al3 Zr sintered compact (Route III). (a) Bright field micrograph showing L12 clusters and matrix and (b) dark field image showing L12 grains.
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Acknowledgments One of the authors (BSR) acknowledges the National Institute for Materials Science for the provision of a NIMS Graduate Research Assistantship. C. S. is grateful to the Japan Society for Promotion of Science (JSPS) for the award of an Invitation Fellowship for Research in Japan (short term). The authors are grateful to Dr. S. Kishimoto (NIMS) for his support and helpful discussion for SPS. This work was supported in part by the Center for Nanostructured Materials Technology (CNMT) under the 21st Century Frontier R&D Programs of the Ministry of Science and Technology, Korea through the Korea Institute of Science and Technology (KIST) and also by the Grant-in-Aid for Scientific Research on Priority Areas “Giant Straining Process for Advanced Materials Containing Ultra-High Density Lattice Defects” and the World Premier International Research Center Initiative (WPI Initiative) on Materials Nanoarchitronics, MEXT, Japan. References
Fig. 11. Typical compressive stress–strain curves for the samples obtained from (a) Route I, (b) Route II and (c) Route III.
4. Conclusions The present study clearly demonstrates the feasibility of using mechanical alloying for the production of high strength Al–Zr nanocomposites through different processing routes. A high strength of more than 1 GPa with a plastic strain of 10% could be obtained. The strength of these compacts was due to the presence of nanocrystalline Al, coarse grained Al, metastable Al3 Zr (L12 ) particles and to some extent because of contamination from the milling media. The presence of nanocrystalline Al even after sintering at 400 ◦ C suggests that these alloys exhibit good thermal stability.
[1] A. Inoue, H. Kimura, Mater. Sci. Eng. A 286 (2000) 1–10. [2] V. Viswanathan, T. Laha, K. Balani, A. Agarwal, S. Seal, Mater. Sci. Eng. R 54 (2006) 121–285. [3] P.Y. Li, H.J. Yu, S.C. Chai, Y.R. Li, Scripta Mater. 49 (2003) 819–824. [4] V. Neubert, B. Smola, I. Stulikova, A. Bakkar, J. Reuter, Mater. Sci. Eng. A464 (2007) 358–364. [5] Y.R. Mahajan, S.D. Kirchoff, Scripta Metall. 21 (1987) 1125–1130. [6] M.N. Rittner, J.R. Weertman, J.A. Eastman, K.B. Yoder, D.S. Stone, Mater. Sci. Eng. A 237 (1997) 185–190. [7] H. Sasaki, N. Kobayashi, K. Kita, J. Nagahora, A. Inoue, Mater. Trans. 44 (2003) 1948–1954. [8] M.S. Chuang, G.C. Tu, Scripta Metall. Mater. 33 (1995) 1999–2005. [9] B. Forbord, H. Hallem, N. Ryum, K. Marthinsen, Mater. Sci. Eng. A 387–389 (2004) 936–939. [10] C. Suryanarayana, Mechanical Alloying and Milling, Marcel Dekker, New York, 2004. [11] C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1–184. [12] E.J. Lavernia, B.Q. Han, J.M. Schoenung, Mater. Sci. Eng. A 493 (2008) 207– 214. [13] K.M. Youssef, R.O. Scattergood, K.L. Murty, C.C. Koch, Scripta Mater. 54 (2006) 251–256. [14] K.I. Moon, K.S. Lee, J. Alloy. Compd. 291 (1999) 312–321. [15] H. Araki, S. Saji, T. Okabe, Y. Minamino, T. Yamane, Y. Miyamoto, Mater. Trans. JIM 36 (1995) 465–468. [16] K.I. Moon, H.S. Park, K.S. Lee, Mater. Sci. Eng. A 323 (2002) 293–300. [17] K.V. Rajulapati, R.O. Scattergood, K.L. Murty, Z. Horita, T.G. Langdon, C.C. Koch, Metall. Mater. Trans. A 39 (2008) 2528. [18] D. Witkin, Z. Lee, R. Rodriguez, S. Nutt, E. Lavernia, Scripta Mater. 49 (2003) 297–302. [19] B.Q. Han, Z. Lee, D. Witkin, S. Nutt, E.J. Lavernia, Metall. Mater. Trans. A 36 (2005) 957–965. [20] Y.S. Park, K.H. Chung, N.J. Kim, E.J. Lavernia, Mater. Sci. Eng. A 374 (2004) 211–216. [21] L. Shaw, H. Luo, J. Villegas, D. Miracle, Scripta Mater. 50 (2004) 921–925. [22] C.A.D. Rodrigues, D.R. Leiva, K.R. Cardoso, C.S. Kiminami, W.J. Botta Filho, Mater. Sci. Forum 386–388 (2002) 3–38. [23] J. Ye, L. Ajdelsztajn, J.M. Schoenung, Metall. Mater. Trans. A 37 (2006) 2569–2579. [24] J.R. Groza, A. Zavaliangos, Rev. Adv. Mater. Sci. 5 (2003) 24–33. [25] T.T. Sasaki, T. Mukai, K. Hono, Scripta Mater. 57 (2007) 189–192. [26] I.H. Kim, C.S. Kim, K.T. Kim, Y.H. Kim, Key Eng. Mater. 326–328 (2006) 429– 432. [27] R.L. Snyder, H.J. Bunge, J. Fiala (Eds.), Microstructure Analysis from Diffraction, International Union of Crystallography, 1999. [28] F. Tancret, H.K.D.H. Bhadeshia, D.J.C. MacKay, Mater. Sci. Technol. 19 (2003) 283–290. [29] Z.A. Chaudhury, C. Suryanarayana, Metallography 17 (1984) 231–252. [30] D. Nagahama, D.H. Ping, M. Ohnuma, H. Sasaki, K. Kita, K. Hono, Mater. Trans. 44 (2003) 1955–1961. [31] S.H. Lee, K.I. Moon, H.S. Hong, K.S. Lee, Intermetallics 11 (2003) 1039–1045.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical properties of Al–Zr nanocomposite materials B. Srinivasarao a,b , C. Suryanarayana c , K. Oh-ishi b , K. Hono b,a,∗ a
Graduate School of Pure and Applied Sciences, University of Tsukuba, 1-2-1 Sengen, Tsukuba 305-0047, Japan National Institute for Materials Science, 1-2-1 Sengen, Tsukuba 305-0047, Japan c Department of Mechanical, Materials and Aerospace Engineering, University of Central Florida, Orlando, FL 32816-2450, USA b
a r t i c l e
i n f o
Article history: Received 7 February 2009 Received in revised form 11 April 2009 Accepted 17 April 2009 Keywords: Aluminum alloy Nanocrystal Nanocomposites Mechanical alloying Spark plasma sintering
a b s t r a c t Three different processing routes were explored to develop Al–Zr nanocomposite alloys using mechanical alloying and spark plasma sintering methods. Depending on the route of milling adopted, the powder in the as-milled condition consisted of either a solid solution of Zr in Al or a mixture of Al-solid solution and Al3 Zr (L12 ) phases. The alloys after sintering consisted of Al and Al3 Zr (L12 ) with grain sizes of less than 100 nm. These nanocomposite alloys exhibited a high compressive strength of 1 GPa with 10% plasticity. The high strength observed in these alloys was explained on the basis of the retention of nanometer sized grains and also the fine dispersion of the L12 phase. On the other hand, the good amount of plasticity was explained to be due to excellent bonding between the powder particles and the presence of coarse Al grains in the matrix. © 2009 Elsevier B.V. All rights reserved.
1. Introduction The development of high strength heat resistant Al-based alloys (both monolithic and composite) has gained considerable interest recently [1–4]. For this purpose, transition metals such as Fe, Cr, Mo, Sc and Zr were added since these elements have a low solid solubility in Al, low diffusivity and additionally they form high melting temperature intermetallic phases with Al. Low solid solubility and diffusivity makes the growth of these intermetallic phases very difficult, making these alloys suitable for high temperature applications. Although the solid solubility of these transition elements in Al is limited, it can be significantly increased by rapid solidification or mechanical alloying methods and this supersaturated solid solution is stable without decomposition up to 320 ◦ C. Further heat treatment can lead to precipitation hardening with a large volume fraction of the precipitate, which enhances the strength of the alloy. Among these transition elements, Zr has the lowest diffusivity and also forms a metastable Al3 Zr (L12 ) phase which displays a very small lattice mismatch with the Al-matrix and therefore has high resistance for coarsening [5–9]. Rittner et al. [6] have synthesized nanocrystalline Al–Zr alloys by the inert gas condensation technique and observed a maximum hardness of 3 GPa. Subsequently Sasaki et al. [7] reported a high tensile strength of 800 MPa for ultrafine grained (260 nm) Al95.3 Zr4 Fe0.7 alloy prepared by physi-
∗ Corresponding author. E-mail address: [email protected] (K. Hono). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.04.032
cal vapor deposition technique. This alloy had also exhibited good high temperature mechanical properties. However, the samples processed by this method were only plates of less than 0.6 mm in thickness. Mechanical alloying (MA) is a process that can produce supersaturated solid solutions even in immiscible systems. The process involves repeated cold welding, fracturing, and rewelding of powder particles in a high-energy ball mill. This technique is capable of synthesizing supersaturated solid solutions, nanocrystalline alloys, amorphous alloys, and their composites with a variety of microstructural modifications [10–14]. Since this technique can be used to synthesize a large amount of powder at a time, it is an industrially viable process. The powders of supersaturated solid solutions or amorphous phase can be consolidated to bulk materials by subsequent consolidation processes like hot pressing [15–17], extrusion [18–22], and spark plasma sintering (SPS) [23–25]. Recently, Sasaki et al. [25] reported an ultrahigh strength bulk nanocrystalline Al–Fe alloy processed by MA and subsequent SPS process, which exhibited yield strength of higher than 1 GPa with excellent heat resistance. In line with the work, we have explored further possibility to process high strength heat resistant aluminum alloys using the Al–Zr system. 2. Experimental procedure Pure Al (99.99%), Pure Zr (99.9%) and home made (arc melted and crushed) Al3 Zr (D023 ) powders were used in this study. MA was conducted in a Fritsch P-4 planetary ball mill using 10 mm diameter
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440C stainless steel balls in a 440C stainless steel vial under argon atmosphere at a speed of 300 rpm. All of the powder handling was carried out in a glove box having Ar atmosphere. Three different methods were employed to make the nanocrystalline composite powders. Route I: Pure elemental Al and Zr powders (resulting total alloy compositions of 0.5, 1.5 and 3 at.% Zr) were milled with a ball to powder weight ratio of 10:1 for a milling time of 300 h to obtain a solid solution of Zr in Al. Route II: Pure elemental Al powder and D023 -Al3 Zr were used (total alloy composition were 2.1, 3, 4 and 4.5 at.% Zr) with a ball to powder weight ratio of 10:1 for a milling time of 120 h to synthesize a solid solution of Zr in Al. Route III: Prealloyed L12 -Al3 Zr was first prepared by mixing stoichiometric amounts of pure Al and Zr powders and milling them for 10 h. Subsequently, different amounts of this resultant powder (10, 25 and 40 wt.%) were mixed with pure Al powder and milled with a ball to powder weight ratio of 20:1 for 5 h to make a nanocomposite. (The total alloy compositions after converting into atomic percentages are 1.63, 4.33 and 7.37 at.% Zr, respectively.) Ethanol (1.5–4 ml or 6–15 wt.%) was used as a process control agent (PCA) in all the experiments. Following milling, the powder was packed in a tungsten carbide die (coated with liquid BN spray) with an inner diameter of 10 mm and compacted by an SPS machine (Sumitomo Coal Mining Company, Model 1050) under a vacuum of <10−3 Pa. Sintering was carried out as follows: at first a stress of 445 MPa was applied, and then the material was heated to desired temperature (i.e., 400–490 ◦ C) at a heating rate of 80 ◦ C/min and held for 10 min. After finishing sintering the load was removed and the sample was water cooled. The final sintered compact had dimensions of 5 mm thickness and 10 mm diameter. The surface of the compact was cleaned by grinding on emery paper to mirror finish and then the density was measured using the Archimedes method with water as a liquid medium. The density of all the compacts was found to be >97% of the theoretical value. Compression tests were performed at a quasi-static strain rate of 1 × 10−4 s−1 in an Instron machine using rectangular samples of 2 mm width, 2 mm breadth and 4 mm height. X-ray diffraction (XRD) analysis was carried out with a Rigaku RINT-2500 X-ray diffractometer using Cu K␣ radiation. Both scanning electron microscopy (1540EsB CrossBeam) and transmission electron microscopy (TEM) (Philips CM200 and TECNAI G2 F30) techniques were used to characterize the microstructural features of the alloys. 3. Results and discussion Figs. 1(a) and 2(a) show X-ray diffraction profile of ball milled powders in Route I and Route II, respectively. In both routes, either the elemental zirconium (in Route I) or intermetallic D023 -Al3 Zr (in Route II) is gradually dissolved into the Al-matrix with increase in milling time to form a solid solution. Note that the amount of the D023 -Al3 Zr added in Route II was such that the total Zr content in the alloy was 3 at.%. The solid solution was formed after 300 h of milling in Route I, while it took only 120 h in Route II. This suggests that the time required for the elemental Zr to completely dissolve was much longer than that for the intermetallic phase. This is because interdiffusion between different phases predominantly occurs only when the grain size of the two phases reaches a critical value during the MA process [11]. Since Al3 Zr is a brittle material compared to elemental Zr, it is easy for the intermetallic phase to reach the minimum critical grain size in a very short time, and hence diffusion occurs quickly. It is also possible that the initial formation of ZrH2 (formed due to reaction between PCA and Zr) phase in the
Fig. 1. Typical XRD profiles of powders made by Route I. (a) Evolution of Al-solid solution during milling of Al–0.5 at.% Zr and (b) summary of the various alloys after 300 h of milling.
case of Route I, has slowed down the diffusion of Zr atoms into the Al lattice to form the solid solution [26]. In case of Route I, the elemental Zr was reacting with the PCA and forming as ZrH2 and also because of the long milling times the amount of PCA needed is very high (i.e., 4 ml). Whereas in the case of Route II, the formation of ZrH2 is not observed and also the milling time was less (i.e., 120 instead of 300 h) due to which the amount of PCA needed to avoid the agglomeration of powders is less (i.e., minimum of 1.5 ml). Fig. 1(b) shows the XRD patterns of the blended elemental powders with different Zr contents processed for 300 h in Route I. It is to be noted that up to 3 at.% Zr could be dissolved in Al to form the solid solution after 300 h of milling. Note that the alloys with higher Zr contents exhibit broader diffraction peaks, suggesting finer grain size with Zr content. The average grain size of the solid solution phase obtained after 300 h of milling in Route I was calculated from the XRD profiles by Voigt function [27] as 67, 62 and 42 nm for 0.5, 1.5 and 3 at.% Zr, respectively. The structural evolution in the powders with different amounts of Al3 Zr (leading to different total Zr contents in the alloy) milled for 120 h following the Route II process is presented in Fig. 2(b) for the overall Zr content of 2.1, 3, 4.0, and 4.7 at.%. Note that the solid solution formation after 120 h is observed up to 4 at.% Zr. Beyond this composition, the metastable L12 -Al3 Zr phase appeared after 120 h of milling. Because of the presence of metastable phase and overlap of the peak positions it is difficult to calculate the grain size from the XRD plot, instead it was calculated from the dark field electron microscopy images of the powder particles. Typical TEM
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Fig. 2. Typical XRD profiles of powders made by Route II. (a) Evolution of Al-solid solution during milling of Al–3 at.% Zr and (b) summary of the various alloys after 120 h of milling.
micrographs of BF and DF was shown in Fig. 3 for 2.1 at.% Zr alloy powder after 120 h of milling and the figure shows a nanocrystalline solid solution which is consistent with Fig. 2(b). The average grain sizes were 48, 32, 35 and 30 nm respectively for 2.1, 3, 4 and 4.7 at.% Zr alloys. This result is in consistent with the previous reports that the grain size of milled powders is smaller for the higher the solute content in the alloys [10,11]. This is because the alloy with the higher solute content becomes stronger and harder due to solid solution strengthening. But the smaller grain size for the 3 at.% Zr was due to the use of a higher amount of PCA, viz., 3 ml ethanol whereas only 1.5 ml PCA was used for all other powders in this route. In Route III, at first L12 -Al3 Zr alloy powder was prepared by mixing pure elemental Zr and Al with 1.5 ml PCA and Fig. 4(a) shows the XRD pattern of the powder after 10 h of milling. From the figure, it can be seen that Al, L12 -Al3 Zr and ZrH2 phases were present and no Zr peaks were observed. The average grain size of the L12 Al3 Zr phase was estimated to be about 9 nm from the dark field TEM images. Secondly, this powder was mixed with Al in different amounts (10, 25 and 40 wt.%) and milled together for extra 5 h with an extra 0.5 ml of PCA and the XRD profiles were shown in Fig. 4(b). It may be noted from the figure that the powder contains Al, L12 -Al3 Zr and ZrH2 phases. The absence of elemental Zr in these powders suggests that all the Zr was completely used up to form the L12 -Al3 Zr and ZrH2 phases. Alternately, it is possible that its volume fraction is so low that it can be hardly detected by the XRD method. Fig. 5 shows the TEM bright field and corresponding dark field image of Al–25 wt.% Al3 Zr nanocomposite powder. The dark
Fig. 3. Typical TEM micrograph of 2.1 at.% Zr alloy powder after 120 h of milling (Route II). (a) Bright field micrograph and (b) dark field micrograph.
field image clearly shows two regions where clusters of nano-Al3 Zr clustered grains (also shown by broken lines in Fig. 5(a)) in one region and the second region contains Al-matrix grains. The matrix grain sizes were found to be 90, 62 and 48 nm for 10, 25 and 40 wt.% Al3 Zr additions. Fig. 6 represents the XRD patterns of the bulk samples consolidated by SPS at different temperatures. To achieve good high temperature mechanical properties it is important that the alloy contains the metastable L12 -Al3 Zr phase than the equilibrium D023 -Al3 Zr from the solid solution in the sintered compact because high creep resistance can be obtained if the matrix/precipitate interface has smaller lattice mismatch [28]. Fig. 6(a) and (b) shows XRD profiles of the sintered samples containing 3 at.% Zr prepared by Route I (300 h) and 2.1 at.% Zr prepared by Route II (120 h), respectively. From these two figures it is clear that L12 -Al3 Zr phase precipitated from the supersaturated solid solution during sintering up to a critical temperature, above which the equilibrium D023 -Al3 Zr phase had formed. This critical temperature was 490 ◦ C for the Al–3 at.% Zr alloy processed by Route I and 450 ◦ C for the
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Fig. 4. Typical XRD profiles of powders made by Route III. (a) L12 -Al3 Zr phase after 10 h of milling and (b) summary of the various alloys after 15 h of milling.
Al–2.1 at.% Zr alloy processed by Route II. A similar observation was reported in the case of rapidly solidified Al–Zr alloys [5,29] and vapor deposited Al–Ti–Cr alloys [30], which can be understood in terms of the lower activation energy required to form the metastable L12 phase due to coherent precipitation with a small interfacial energy than the incoherent D023 phase. Moon and Lee [14] reported that the thermal stability of the L12 -Al3 Zr phase was
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increased when a ternary element such as Cu, Ni and Mn is added. In case of Route I process, because of large milling time and because the milling media is 440C stainless steel, there is a possibility that Ni and Mn from the milling media could be present in the milled powder. These additions might have made the L12 -Al3 Zr phase stable up to higher temperatures (viz., 490 ◦ C). Fig. 6(c) and (d) shows the phases present in the sintered alloys with different Zr concentrations made by Route I and Route II, respectively. The XRD patterns show the presence of metastable L12 phase after sintering at 470 and 400 ◦ C, respectively. These results should be compared with that of Fig. 6(e), which shows the typical XRD pattern of the Al–25 wt.% Al3 Zr alloy sintered at 450 ◦ C; this figure reveals a number of phases, viz., Al, Al3 Zr (L12 ) ZrC and iron aluminides. ZrC was formed by the reaction between the carbon present in the PCA and Zr, whereas the iron aluminides are due to contamination of iron from the milling media. Nanocrystalline L12 -Al3 Zr has very high Vickers hardness value approaching 10 GPa [31] whereas Al–Zr solid solution was less than 400 Hv (Route I and Route II). Because of the high hardness of L12 phase in Route III the iron contamination will be higher even though the milling time is smaller (15 h) than Route I and Route II processes (120 and 300 h). Fig. 7 shows the SEM backscattered micrographs of sintered compacts prepared by Route I and Route II. The black regions in Fig. 7(a) and (b) correspond to pure Al formed during SPS and the grey region corresponds to the milled powder particles bonded together. All the compacts made by Route I and subsequent sintering (up to 490 ◦ C) show poor bonding between the particles and lot of processing defects such as porosity, cracks and voids indicated by white arrows in Fig. 7(a) whereas for Route II processed and sintered samples the bonding between the particles has greatly improved and very few isolated porosity and voids were present (Fig. 7(b)). The major cause for the poor particle bonding in Route I could be due to the high powder particle strength; a result of contamination from the milling media because of long milling times (300 h) and the use of large amount of PCA (4 ml ethanol). Both these result in incorporation of oxide, carbide and other ceramic particles into the metallic matrix and consequently increase the strength of the powder. And because of the high particle strength, the particles are not sufficiently compressed during the sintering process to establish full contact between them. The presence of pure Al at the particle boundaries is considered to be due to the partial melting because of the local joule heating at the particle interfaces and is commonly observed in spark plasma sintered aluminum alloys [23–25]. During the solidification, Zr is rejected from Al due to the limited solubility, and hence coarse grains of pure Al are formed at the particle inter-
Fig. 5. TEM micrograph of Al–25 wt.% (L12 )Al3 Zr powder particle after milling. (a) Bright field image and (b) dark field image.
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Fig. 6. Typical XRD patterns of sintered samples (a) and (c) for Route I, and, (b) and (d) for Route II and (e) for Route III process.
faces. The partial melting of Al results in the density values of >97% of the theoretical value, because the melt fills the spacing between the particles. Fig. 7(c) and (d) shows the SEM backscattered high magnification images taken from Fig. 7(a) and (b), respectively. The white contrast in Fig. 7(c) and (d) indicates the presence of Al3 Zr (L12 ) phase and Al13 Fe4 phase whereas the black contrast is due to oxide particles i.e., Al2 O3 mostly came from polishing. These Al3 Zr particles are in the size range of 10–50 nm. Fig. 8 shows the SEM micrographs of the samples containing different amounts of Al3 Zr (L12 ) made by Route III and after sintering at 450 ◦ C. The figure shows a uniform microstructure and no pro-
cessing defects such as cracks, debonding and porosity are seen. The white contrast in these figures corresponds to Al3 Zr (L12 ) and Al13 Fe4 which are uniformly dispersed throughout the matrix and most of them are well below 2 m in size. Fig. 9(a) shows the bright field electron micrograph of Al–3 at.% Zr alloy made by Route II with 3 ml of ethanol as PCA. The selected area diffraction pattern (SAED) (shown as an inset) indicates the presence of Al3 Zr (L12 ) and Al-phases, confirming the results obtained by XRD, Fig. 6(d). Mostly the second phase particles were of equiaxed grain size forming at the grain boundaries of the matrix Al. Fig. 9(b) gives an idea of the matrix grain size
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Fig. 7. Typical SEM backscattered micrographs of sintered samples. (a) 3 at.% Zr (Route I) sintered at 470 ◦ C, (b) 4 at.% Zr (Route II) sintered at 400 ◦ C (c and d) High magnification back scatter images taken from (a) and (b), respectively.
which is well below 100 nm. The average matrix grain sizes calculated from the dark field images are 65, 58 and 45 nm for 0.5, 1.5 and 3 at.% Zr alloys, respectively for Route I and 85, 45, 48 and 35 nm for 2.1, 3, 4 and 4.5 at.% Zr alloys made by Route II, respectively. The average grain size of the composites made by Route I and Route II are well below 100 nm even after sintering and this small grain size is expected to confer high strength on these materials. Fig. 10(a) shows the microstructure of sintered compact of
Al–25 wt.% Al3 Zr alloy in which the dark regions correspond to cluster of Al3 Zr grains. While the clusters have sizes of typically ≤1 m, the average grain size in the Al3 Zr cluster was found to be about 10 nm, clearly seen from Fig. 10(b). The average matrix grain sizes were 200, 120 and 50 nm for 10, 25 and 40 wt.% Al3 Zr alloys, respectively. Fig. 11 shows the typical stress–strain curves obtained by compression test for the samples prepared by Routes I, II and III. For
Fig. 8. Typical SEM micrographs of sintered samples sintered at 450 ◦ C (Route III process). (a) 10 wt.%, (b) 25 wt.% and (c) 40 wt.% Al3 Zr.
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Fig. 9. Typical TEM micrograph of 3 at.% Zr sintered compact (Route II). (a) Bright field image, (b) dark field image showing nanocrystalline grain size of Al.
the case of Route I a high compressive strength of 1 GPa could be achieved in all of the compositions but they broke into pieces immediately after yielding and sometimes in the elastic regime itself (to remove the effect of brittle phase Al3 Zr-D023 on compression properties sintered samples of 480 ◦ C was used because D023 form starting from 490 ◦ C as observed from Fig. 6). This type of failure can be attributed to the improper bonding (many powder particles are barely touching each other) between the particles and also due to the presence of lot of defects (Fig. 7(a)). The defects oriented perpendicular to the compression axis will be closed during the test, but at other angles they will initiate cracks and these cracks will grow faster because of the high stress levels and therefore lead to fracture. For the Route II processed samples, the compression testing (Fig. 11(b)) results show some plasticity (∼10%). This improvement in plasticity is due to the improvement in bonding between the particles and a significant reduction of process defects (Fig. 7(b)). A high strength of more than 1 GPa was observed for 3 at.% Zr and this is because of the use of a large amount of PCA (i.e., 3 ml whereas only 1.5 ml was used in other cases) which will increase the amount of carbon, oxygen and also will affect the contamination level from the milling media. From the XRD patterns as
well as low magnification TEM studies, these carbides and oxides were not observed suggesting that these were present either in small sizes or still dissolved in the matrix. This could be due to rapid sintering by the SPS technique (the total sintering time was less than 20 min). Fig. 11(c) indicates that Route III processed nanocomposites exhibit good plasticity and this plasticity had decreased when the amount of Al3 Zr (L12 ) phase increased. All of the compression data in this study show strain softening behavior which could be due to either small size (i.e., <100 nm) or because of the use pure Al-matrix (Route III process). The high strength found in the Route III processed alloys is due to the high hardness of L12 -Al3 Zr clusters [31]. From the above results Route II and Route III processes are the most suitable methods for scaling up, because of the minimized processing time, reasonable compressive strength and plasticity. Currently, scaling up of these alloys by extrusion and their high temperature properties are under investigation. So far, we obtained the plastic strain of 10% for the strength level of 1 GPa only in the compression test. Further optimization of the alloy compositions, processing routes and microstructures is necessary to obtain tensile ductility.
Fig. 10. Typical TEM micrograph of Al–25 wt.% Al3 Zr sintered compact (Route III). (a) Bright field micrograph showing L12 clusters and matrix and (b) dark field image showing L12 grains.
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Acknowledgments One of the authors (BSR) acknowledges the National Institute for Materials Science for the provision of a NIMS Graduate Research Assistantship. C. S. is grateful to the Japan Society for Promotion of Science (JSPS) for the award of an Invitation Fellowship for Research in Japan (short term). The authors are grateful to Dr. S. Kishimoto (NIMS) for his support and helpful discussion for SPS. This work was supported in part by the Center for Nanostructured Materials Technology (CNMT) under the 21st Century Frontier R&D Programs of the Ministry of Science and Technology, Korea through the Korea Institute of Science and Technology (KIST) and also by the Grant-in-Aid for Scientific Research on Priority Areas “Giant Straining Process for Advanced Materials Containing Ultra-High Density Lattice Defects” and the World Premier International Research Center Initiative (WPI Initiative) on Materials Nanoarchitronics, MEXT, Japan. References
Fig. 11. Typical compressive stress–strain curves for the samples obtained from (a) Route I, (b) Route II and (c) Route III.
4. Conclusions The present study clearly demonstrates the feasibility of using mechanical alloying for the production of high strength Al–Zr nanocomposites through different processing routes. A high strength of more than 1 GPa with a plastic strain of 10% could be obtained. The strength of these compacts was due to the presence of nanocrystalline Al, coarse grained Al, metastable Al3 Zr (L12 ) particles and to some extent because of contamination from the milling media. The presence of nanocrystalline Al even after sintering at 400 ◦ C suggests that these alloys exhibit good thermal stability.
[1] A. Inoue, H. Kimura, Mater. Sci. Eng. A 286 (2000) 1–10. [2] V. Viswanathan, T. Laha, K. Balani, A. Agarwal, S. Seal, Mater. Sci. Eng. R 54 (2006) 121–285. [3] P.Y. Li, H.J. Yu, S.C. Chai, Y.R. Li, Scripta Mater. 49 (2003) 819–824. [4] V. Neubert, B. Smola, I. Stulikova, A. Bakkar, J. Reuter, Mater. Sci. Eng. A464 (2007) 358–364. [5] Y.R. Mahajan, S.D. Kirchoff, Scripta Metall. 21 (1987) 1125–1130. [6] M.N. Rittner, J.R. Weertman, J.A. Eastman, K.B. Yoder, D.S. Stone, Mater. Sci. Eng. A 237 (1997) 185–190. [7] H. Sasaki, N. Kobayashi, K. Kita, J. Nagahora, A. Inoue, Mater. Trans. 44 (2003) 1948–1954. [8] M.S. Chuang, G.C. Tu, Scripta Metall. Mater. 33 (1995) 1999–2005. [9] B. Forbord, H. Hallem, N. Ryum, K. Marthinsen, Mater. Sci. Eng. A 387–389 (2004) 936–939. [10] C. Suryanarayana, Mechanical Alloying and Milling, Marcel Dekker, New York, 2004. [11] C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1–184. [12] E.J. Lavernia, B.Q. Han, J.M. Schoenung, Mater. Sci. Eng. A 493 (2008) 207– 214. [13] K.M. Youssef, R.O. Scattergood, K.L. Murty, C.C. Koch, Scripta Mater. 54 (2006) 251–256. [14] K.I. Moon, K.S. Lee, J. Alloy. Compd. 291 (1999) 312–321. [15] H. Araki, S. Saji, T. Okabe, Y. Minamino, T. Yamane, Y. Miyamoto, Mater. Trans. JIM 36 (1995) 465–468. [16] K.I. Moon, H.S. Park, K.S. Lee, Mater. Sci. Eng. A 323 (2002) 293–300. [17] K.V. Rajulapati, R.O. Scattergood, K.L. Murty, Z. Horita, T.G. Langdon, C.C. Koch, Metall. Mater. Trans. A 39 (2008) 2528. [18] D. Witkin, Z. Lee, R. Rodriguez, S. Nutt, E. Lavernia, Scripta Mater. 49 (2003) 297–302. [19] B.Q. Han, Z. Lee, D. Witkin, S. Nutt, E.J. Lavernia, Metall. Mater. Trans. A 36 (2005) 957–965. [20] Y.S. Park, K.H. Chung, N.J. Kim, E.J. Lavernia, Mater. Sci. Eng. A 374 (2004) 211–216. [21] L. Shaw, H. Luo, J. Villegas, D. Miracle, Scripta Mater. 50 (2004) 921–925. [22] C.A.D. Rodrigues, D.R. Leiva, K.R. Cardoso, C.S. Kiminami, W.J. Botta Filho, Mater. Sci. Forum 386–388 (2002) 3–38. [23] J. Ye, L. Ajdelsztajn, J.M. Schoenung, Metall. Mater. Trans. A 37 (2006) 2569–2579. [24] J.R. Groza, A. Zavaliangos, Rev. Adv. Mater. Sci. 5 (2003) 24–33. [25] T.T. Sasaki, T. Mukai, K. Hono, Scripta Mater. 57 (2007) 189–192. [26] I.H. Kim, C.S. Kim, K.T. Kim, Y.H. Kim, Key Eng. Mater. 326–328 (2006) 429– 432. [27] R.L. Snyder, H.J. Bunge, J. Fiala (Eds.), Microstructure Analysis from Diffraction, International Union of Crystallography, 1999. [28] F. Tancret, H.K.D.H. Bhadeshia, D.J.C. MacKay, Mater. Sci. Technol. 19 (2003) 283–290. [29] Z.A. Chaudhury, C. Suryanarayana, Metallography 17 (1984) 231–252. [30] D. Nagahama, D.H. Ping, M. Ohnuma, H. Sasaki, K. Kita, K. Hono, Mater. Trans. 44 (2003) 1955–1961. [31] S.H. Lee, K.I. Moon, H.S. Hong, K.S. Lee, Intermetallics 11 (2003) 1039–1045.