Investigation of Al Pb nanocomposites synthesized by nonequilibrium processes

Transcription

1 Journal of MATERIALS RESEARCH Welcome Comments Help Investigation of Al Pb nanocomposites synthesized by nonequilibrium processes H. W. Sheng, F. Zhou, Z. Q. Hu, and K. Lu State Key Laboratory for RSA, Institute of Metal Research, Chinese Academy of Sciences, Shenyang , People s Republic of China (Received 4 October 1996; accepted 27 March 1997) Two nonequilibrium processes (melt-spinning and ball-milling) were successfully employed to synthesize Al 1 x Pb x (x 5, 10, 20, 30 wt. %) nanocomposites with distinct microstructures. In the melt-spun (MS) Al Pb alloys, the nanometer-sized Pb particles are uniformly distributed in the micrometer-grained Al matrix and have an orientational relationship with the matrix, while in the ball-milled (BM) samples, both Pb and Al components are refined with prolonged milling time, forming nanocomposites with Pb particles homogeneously dispersed into the Al matrix. The minimum particle size of Pb in the milled samples linearly increases with the Pb content. The microhardness of the BM Al Pb samples is much larger than that of the MS samples, which mainly results from strengthening effects of the nanometer scale Al grains following the Hall Petch relationship. The microhardness for both BM and MS Al Pb samples varies with the Pb content, and maximum hardness for both samples exists when Pb content is about 5 wt. %, indicating that small amounts of Pb, in the form of nanoparticles, may strengthen the Al matrix. I. INTRODUCTION The development of immiscible alloys has been largely constrained by the conventional equilibrium processing, which generally results in gross segregation due to the wide miscibility gap, and high disparity in the densities and melting temperatures between the immiscible elements. 1 For example, aluminum alloys with proper distribution of lead have potential as materials for plain bearings, because of the lubrication ability of lead in the aluminum matrix. 2 However, owing to the immiscibility of lead phase in the Al matrix, 3 the grossly segregated lead may sometimes act as a kind of inclusion. Therefore, it is of importance to explore new processing approaches to obtain materials with fine distributions. During recent decades, nonequilibrium processes, such as rapid solidification and ball-milling, have been developed to produce refined controllable microstructures in immiscible systems. By means of a rapid solidification technique, Al-based (Al In, 4 Pb, 5 Sn, 6 Cd 7 ) and Cu-based (Cu Pb, 1,8 Cu Pb Sn 1 ) alloys have been synthesized with uniform distributions of the fine immiscible particles embedded in the matrix, indicating that the microstructure could be considerably refined compared with conventionally processed materials. Similar to rapid solidification, ball-milling is another effective nonequilibrium process for producing novel materials (see, for a review [9]). Besides its original purpose for producing dispersion-strengthened alloys, 10 much of the recent work of ball-milling has been focused on the study of the solid-state reaction. In most cases, finely and homogeneously alloyed products (solid solutions 11,12 or amorphous 13,14 etc.) can be obtained by ball-milling elements even with large positive mixing heats. Nonetheless, in some immiscible alloy systems, it was found that the alloying effect might not take place, and the milled products are nanostructured pure phase mixtures, e.g., Cu Pb, 15 Ge Sn, 16 Al In, 17 indicating the nanocomposites might as well be obtained by milling the immiscible components. Currently there exists a substantial amount of literature regarding the microstructural characterization and properties of binary systems upon milling; however, no detailed study has been reported on the structural evolution of Al Pb alloys, though it is of great potential for applications. Further, the effects of different nonequilibrium processing approaches on the resultant microstructures have scarcely been investigated. In the present work, we report our results on the synthesis of Al Pb nanocomposites via two nonequilibrium processes, rapid-solidification and ball-milling, to cast light on the possibility of forming the controllable twophase mixture of Al Pb. The microstructures as well as mechanical properties of Al Pb samples from different approaches are analyzed and compared in the text. II. EXPERIMENTAL Alloy ingots with compositions of Al 1 x Pb x (x 5, 10, 20, and 30 wt. %) were cast by arc-melting of % pure Al and Pb in water-cooled copper cru- 308 J. Mater. Res., Vol. 13, No. 2, Feb Materials Research Society

2 cibles under Ar atmosphere. The Al Pb thin ribbons 2 3 mm wide, 20 mm thick, and a few meters long were obtained by using the melt-spinning technique. Commercial elemental powder blends of Al and Pb (purity % and particles less than 100 mesh) with compositions of Al 1 x Pb x (x 2, 5, 10, 20, 30, and 40 wt. %) were used as starting materials for the BM samples. Ball-milling was performed in a vibratory ball mill. Mixed powders were sealed under an Ar atmosphere. High hardness, good wear-resistant chrome steel balls were used; the ball-to-powder weight ratio is 30 : 1. The MS and BM Al Pb samples were examined by using x-ray diffraction (XRD), scanning electron microscope (SEM), transmission electron microscopy (TEM), and high resolution transmission electron microscopy (HREM) techniques. XRD experiments were carried out on a Rigaku x-ray diffractometer (D max-ra, 12kW) with Cu K a radiation. The average grain sizes and precise lattice parameters were determined, respectively, by means of quantitative analysis of the XRD spectra. SEM was performed on a Cambridge stereoscan-360 scanning microscope equipped with a Link-AN 1000 energy dispersive x-ray spectroscopy (EDX). TEM and HREM were conducted on a JEOL-2010 microscope. Specimens for TEM and HREM observations were prepared by ion thinning the MS ribbons and the BM samples (which were consolidated from the as-milled powders). Microhardness was measured in a MVK-H3 microhardness tester at a load of 25 gf, for which the BM powders were consolidated with a uniaxial pressure of 2 GPa into pellets (8 mm in diameter and 2 mm in thickness) and then mechanically polished. At least ten measurements for each sample were collected to calculate the average values. III. RESULTS AND DISCUSSION A. Melt spinning Figure 1 shows the low magnified TEM bright images of Al Pb MS samples, which were prepared with similar cooling rate during melt-spinning. It is found that refined Pb particles are uniformly dispersed throughout the Al matrix with a bimodal distribution. With the Al grains, nanometer-sized Pb particles can be discerned; at the grain boundaries of Al, however, are located some larger Pb (several tens of nanometers). With the increase of Pb contents, the embedded Pb nanoparticles become slightly denser, but the size of the Al grains, which are about several mm in size, does not show substantial change. Figure 2 illustrates the detailed characteristics of the embedded Pb particles. Figure 2(a) shows a typical bright-field TEM image of the Al 10 wt. % FIG. 1. Low magnified bright-field transmission electron micrographs showing the distribution of Pb particles in (a) Al 5 wt. % Pb and (b) Al 10 wt. % Pb MS samples. Pb MS ribbon. It can be seen that the fine Pb particles are faceted, in a range of 5 30 nm with a mean diameter of 11 nm, as shown in the inset of Fig. 2(a). In addition, Moirè fringes in the images of the small Pb particles are clearly seen, indicating that the Al Pb interface was highly ordered. Figures 2(b) and 2(c) show typical electron diffraction patterns from Al 110 and 001 zones for the MS Al Pb samples. These results agree with those reported previously 5,18 ; i.e., the Pb particles in the MS Al Pb samples exhibit a cube-cube orientational relationship with the surrounding Al matrix ( 111 Al k 111 Pb and 110 Al k 110 Pb ). Figures 2(d) and 2(e) show HREM images corresponding to Figs. 2(b) and 2(c), respectively. In addition to the orientational relationship between Pb and Al, the Pb particles exhibit hexagonal cross-sectional shapes perpendicular to the Al 110 zone axis, and octagonal cross-sectional shapes perpendicular J. Mater. Res., Vol. 13, No. 2, Feb

3 FIG. 2. (a) A typical bright-field transmission electron micrograph of MS Al 10 wt. % Pb; the inset is a histogram showing the distribution of the Pb particles; (b, c) SAED from MS Al 10 wt. % Pb, with the beam direction parallel to the 110 Al and 001 Al zone axes, respectively. (d, e) HREM images showing the hexagonal and octagonal cross-sectional shapes of Pb particles normal to the 110 Al and 001 Al zone axes, respectively. (f) Schematic diagram of the truncated octahedral Pb particle shape. to the Al 001 zone axis, manifesting that the particle shapes are truncated octahedral bounded by 111 and 100 facets, as schematically shown in Fig. 2(f). B. Ball milling 1. Microstructural evolution Figure 3 shows the XRD results of Al 10 wt. % Pb with different milling times. Up to 10 h, only diffraction peaks for Al and Pb are observed; shifts of peak positions have not been detected. No trace of amorphous or other intermediate phases is found. It is also evident that the intensities of both Al and Pb diffraction lines decrease while the profile of the pattern broadens, suggesting that grains of Al and Pb are refined. The apparent grain size of Al and Pb was calculated using the Scherrer equation. 19 In determining the grain size of Pb, the effect of the internal strain on the line broadening was neglected. Figure 4 displays the average grain sizes of FIG. 3. XRD for Al 10 wt. % Pb milled for different periods of time. 310 J. Mater. Res., Vol. 13, No. 2, Feb 1998

4 FIG. 4. Variation of the grain size of Pb and Al in the milled powders (Al 10 wt. % Pb) versus the milling time, which was determined from Pb (111), (200), and Al(200) of the XRD profiles. Al and Pb, which were determined, respectively, from the half-maximum width of the Pb (111), (200), and Al (111) diffraction peaks, as a function of milling time. The grain size of Pb decreases with an increasing milling time, tends to a steady value (about 7.0 nm) after 10 h of milling, showing the formation of the nanostructured Pb component. Prolonging the milling time, even to 100 h, we find that the grain sizes of Al and Pb are almost unchanged. Meanwhile, lattice parameters for Al and Pb in the BM sample have been quantitatively determined in order to examine the possible formation of the supersaturated Al/Pb solutions. Lattice parameters for Al and Pb in the BM Al Pb sample are nm and nm, respectively. Compared with the tabulated values for Al ( nm) and Pb ( nm), 20 one may find that the lattice parameters for both Al and Pb were unchanged after ball milling, suggesting that no alloying effect between Al and Pb had taken place. SEM observations reveal the microstructural evolution of Al Pb with the milling time, as shown in Fig. 5 (white particles are Pb). At the early stage of ball-milling, the existence of large Pb particles indicates that Pb is not pulverized completely. Since Al and Pb are malleable, kneading and cold welding dominate over fracturing at the early stage of milling, and the lamellar Pb components are clearly observed [Fig. 5(b)]. As ballmilling proceeds, due to the effects of fracturing, Al and Pb are comminuted, while Pb particles are gradually dispersed into the Al matrix, forming individual fine particles. In the SEM photograph for Al Pb milled 10 h [see Fig. 5(c)], although some Pb particles of few hundred nanometers are visible, the fine microstructure is beyond the resolution of SEM. The composition of the milled dispersion, examined by EDX, was nearly FIG. 5. SEM micrographs backscattered electron mode of Al 10 wt. % Pb milled for different times: (a) 1 h, (b) 2 h, and (c) 10 h. the same with the unmilled mixture. The iron content is only 0.25 wt. % after 10 h of milling, indicating that the contamination from the milling media is rather small. J. Mater. Res., Vol. 13, No. 2, Feb

5 TEM observations provide more details as shown in Fig. 6. It can be seen that Pb particles, most of which are in the form of irregular shape, are nanoparticles surrounded by the Al matrix, demonstrating the formation of the Al Pb nanocomposites. The inset in Fig. 6(a) is a selected area electron diffraction pattern for the Al Pb sample, from which we found that both the Pb and the Al phases are random in orientations. No simple orientational relationship between Pb and Al can be identified. Figure 6(b) is the size histogram for Pb particles in the Al Pb sample milled for 10 h. The statistical mean diameter of Pb particles, as determined from the dark-field images, is about 6.5 nm, which coincides well with the XRD result (7.0 nm). During annealing the milled powders, growth of Pb particles was observed. Calculated from the XRD profile width, the Pb particle sizes in the BM sample annealed at 200 ± C, 300 ± C, and 350 ± C (for 10 min) were determined as 13 nm, 15 nm, and 20 nm, respectively. Figure 7(a) shows a typical TEM bright image FIG. 6. (a) A bright TEM image of the Al 10 wt. % Pb milled for 10 h; the inset is the corresponding SAED. (b) A dark TEM image of Pb in the Al 10 wt. % Pb milled for 10 h. (c) A histogram showing the distribution of the Pb particles in the Al 10 wt. % Pb milled for 10 h. FIG. 7. (a) A bright TEM image of Al 10 wt. % Pb sample (milled for 10 h) after annealed at 300 ± C. (b) A histogram showing the distribution of the Pb particle. 312 J. Mater. Res., Vol. 13, No. 2, Feb 1998

6 for Al 10 wt. % Pb annealed at 300 ± C. Except for the grain growth of the Pb particles, the morphologies of the Pb particles do not show substantial change. Figure 7(b) gives the histogram for the distribution of the Pb particles, with a statistic mean diameter of 15 nm which corresponds to the XRD results very well. It is obvious that coarsening of the Pb particles takes place upon annealing, and the higher the annealing temperature, the larger the Pb particles obtained. However, the drastic growth of the Pb particles in the milled sample was not observed, which may result from the isolation of the Al matrix. From the above analysis, one may find the microstructures in the MS and BM Al Pb samples are different in two aspects, i.e., the form of the contained Pb particles and the grain size of the containing Al matrix, which originates from the distinct processing approaches. Strictly speaking, the terminology of nanocomposites is more appropriate for the BM Al Pb samples due to the nanometer scale Al matrix. In fact, the bulk, full density Al Pb nanocomposites have been prepared by warm consolidation of the BM Al Pb powders Minimum size of Pb particles and Al grains Similar to Al 10 wt. % Pb, Al 1 x Pb x alloys with different contents of Pb became nanostructured materials after ball milling for 10 h, as shown by the XRD patterns in Fig. 8. Only diffraction peaks from Pb and Al can be identified, suggesting the final milled products are twophase mixtures of Al and Pb pure components. However, the minimum particle size for Pb in the BM Al 1 x Pb x alloys varies with the content of Pb. Figure 9 illustrates the minimum Pb particle size evaluated from the XRD profiles as a function of the Pb content. It is clear that the minimum value for Pb increases linearly with the Pb content. By milling the pure elements, Eckert et al. 22 found that the ultimate grain size achievable by ball-milling is determined by the competition between the heavy mechanical deformation introduced during milling and the recovery behavior of the metal. The final grain size scales with the melting point and bulk modulus of the corresponding metal: the higher the melting point and bulk modulus, the smaller the minimum grain size of the powders. However, it seems not the case for Pb in the Al Pb alloy, because the recovery and growth of Pb particles might be constrained by the isolation of the Al matrix, which will lead to a further refinement of the Pb component. On the other hand, the refined Pb particles will coalesce because of the impingement during further milling. When the two processes, i.e., fracturing and coalescence, reach a dynamic equilibrium, Pb particles will attain a steady grain size. Since the possibility for fine Pb particles to collide is related to their volume fraction in the alloys, the minimum grain size of Pb particles therefore depends on their volume fraction in Al matrix. The less the volume fraction, the fewer the possibility for Pb particles to impinge, the smaller the minimum particle size Pb particles can be obtained, which was exactly what we have observed from Figs. 8 and 9. On the contrary, we noticed that the minimum grain size of Al is almost the same in each BM Al Pb alloy. It is because the volume fraction for Pb in the Al Pb alloys is too small to affect the refinement of the Al phase such that the attainable minimum grain size of the Al in the BM Al Pb alloys is comparable to that by milling pure Al elements, i.e., about 23 nm. 22 FIG. 8. XRD profiles of Al Pb milled powders (10 h) with different levels of Pb. FIG. 9. Particle size of the Pb component in Al Pb powders milled for 10 h as a function of the content of Pb. J. Mater. Res., Vol. 13, No. 2, Feb

7 3. Microhardness measurements Figure 10 shows the microhardness of the BM and MS Al Pb alloys as a function of Pb content. Several features can be noticed from Fig. 10. Firstly, the microhardness of the BM Al samples is much larger than that of the MS samples, which might result from the nanometer-scaled Al matrix in the BM Al Pb samples. In fact, the microhardness of the Al Pb samples increased gradually with the prolonged milling times or the refinement of both Al and Pb components. At the early stage of ball milling, the hardness rapidly increased, but with further milling, it gradually reached a constant value, being several times larger than that for the pure bulk Al. It has been experimentally observed that, in conventional materials, the hardness varies with the grain size following the empirical Hall Petch relation: H n s 0 1 K 0 d 2 1 2, where H n is the hardness, s 0 is the intrinsic stress resisting dislocation motion, K 0 is a constant, and d is the grain size. The experimental results for the microhardness testing of pure nanophase metals show these materials are considerably harder, by factors of from 2 to 7, than their coarse-grained counterparts. 23 In the present work, although BM Al Pb alloys are two-phase mixtures, their considerably increased hardness might be attributed to the refinement of the main Al component in the BM sample according to the Hall Petch relation. Figure 11 displays the microhardness of the Al Pb sample milled for 10 h as a function of inverse particle size. The best fit of the experimental data is roughly a straight line, which implies the validity of the Hall Petch relationship in the BM Al Pb samples. FIG. 10. Microhardness of the Al Pb powders for MS Al 1 x Pb x (open circle) and Al 1 x Pb x samples milled for 10 h (closed circle) as a function of the content of Pb. FIG. 11. Variation of the Vickers hardness as a function of inverse grain size of Al in the sample of Al 10 wt. % Pb, showing the Hall Petch relation. Secondly, one may find from Fig. 10 that there is a maximum microhardness value corresponding to a Pb content of 5 wt. % for both samples. Since the grain size of the Al matrix in either the BM or the MS Al 1 x Pb x samples are nearly the same, variation of the Vickers hardness might be caused by the addition of Pb. Because of their fine-scale dimension, Pb nanoparticles might act as a kind of strengthening particles rather than coarse inclusions. The dispersed Pb particles are weaker than the Al matrix; they will be sheared by the gliding dislocations, which impedes the dislocation motion, and thus strengthens the Pb dispersed composites, causing the microhardness increases. One can argue that the strengthening effect from Pb nanoparticles depends on their volume fractions as well as particle sizes. A small amount of the fine Pb particles will effectively strengthen the BM Al Pb; the microhardness will increase with the addition of the Pb. When the content of Pb is greater than such a critical value (5 wt. %), which actually corresponds to its critical volume fraction or particle size, the softer large Pb particles will play a role in weakening the Al Pb alloys, as shown in Fig. 10 that the Vickers hardness decreases when the content of Pb is greater than 5 wt. %. IV. CONCLUSIONS The nonequilibrium processes, i.e., rapid solidification and ball-milling, are successfully used to prepare the Al Pb nanocomposites with Pb nanoparticles uniformly embedded in the Al matrix. However, the microstructures of the Al Pb products vary with different processing approaches. In the MS sample, the Pb particles (5 30 nm) have regular shape and an epitaxial orientation relationship with the Al matrix; while in the 314 J. Mater. Res., Vol. 13, No. 2, Feb 1998

8 BM Al Pb sample, both Al and Pb components are nanometer-sized, and Pb particles are randomly oriented with the Al matrix. The Pb particle size decreases with the increasing milling times; the minimum particle size of the Pb linearly increases with the Pb content. Microhardness measurements show that a small amount of Pb, in the form of nanoparticles, may strengthen the Al matrix. The microhardness of the BM Al Pb samples increases with increasing milling time, following the Hall Petch relation. ACKNOWLEDGMENTS The authors thank Dr. G. Ren and Dr. D. H. Ping for TEM and HREM observations. Financial support from the Chinese Academy of Sciences and the National Sciences Foundation of China is acknowledged. REFERENCES 1. A. N. Patel and S. Diamond, Mater. Sci. Eng. 98, 329 (1988). 2. J. P. Pathak, V. Singh, and S. N. Tiwari, J. Mater. Sci. Lett. 11, 639 (1992). 3. T. B. Massalki, Binary Alloy Phase Diagrams (American Society for Metals, Metals Park, OH, 1986). 4. H. Saka, Y. Nishikawa, and T. Imura, Philos. Mag. A 57, 895 (1988). 5. D. L. Zhang and B. Cantor, Acta Metall. Mater. 39, 1595 (1991). 6. W. T. Kim and B. Cantor, J. Mater. Sci. 26, 2869 (1991). 7. D. L. Zhang, J. L. Hutchinson, and B. Cantor, J. Mater. Sci. 29, 2147 (1994). 8. S. N. Tiwari, J. Mater. Sci. Lett. 8, 1098 (1989). 9. C. C. Koch, Mater. Tran. JIM 36, 85 (1995). 10. J. S. Benjamin, Metall. Trans. 1, 2943 (1970). 11. J. Eckert, J. C. Holzer, C. E. Krill III, and W. L. Johnson, J. Appl. Phys. 73, 2794 (1993). 12. C. Gente, M. Oehring, and R. Bormann, Phys. Rev. B 48, (1993). 13. K. Sakurai, Y. Yamada, M. Ito, C. H. Lee, T. Fukunaga, and U. Mizutani, Appl. Phys. Lett. 57, 2660 (1990). 14. T. D. Shen, K. Y. Wang, M. X. Quan, and J. T. Wang, J. Appl. Phys (1992). 15. T. Ohashi and Y. Tanaka, Mater. Tran. JIM 32, 387 (1991). 16. J. S. C. Jang and C. C. Koch, J. Mater. Res. 5, 325 (1990). 17. K. Uenishi, H. Kawaguchi, and K. F. Kobayashi, J. Mater. Sci. 29, 4860 (1994). 18. K. I. Moore, K. Chattopadhyay, and B. Cantor, Proc. R. Soc. A 414, 499 (1987). 19. H. P. Klug and L. E. Alexander, X-ray Diffraction Procedures for Polycrystalline and Amorphous Materials, 2nd ed. (Wiley, New York, 1974). 20. R. W. G. Wyckoff, Crystal Structures (Wiley, New York, 1963), Vol. I. 21. F. Zhou, H. W. Sheng, and K. Lu, J. Mater. Res., Vol J. Eckert, J. C. Holzer, C. E. Krill III, and W. L. Johnson, J. Mater. Res. 7, 1751 (1992). 23. R. W. Siegle and G. E. Fougere, Nanostruc. Mater. 6, 205 (1995). J. Mater. Res., Vol. 13, No. 2, Feb