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Ball milling of ductile metals
MATERIALS SCIENCE & ENGINEERING ELSEVIER
Materials Science and Engineering A199 (1995) 165-172
A
Ball milling of ductile metals J.Y. Huang, Y.K. Wu, H.Q. Ye Laboratory of' Atomic Imaging of Solids, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110015, People's Republic of China Received 13 June 1994; in revised form 17 October 1994
Abstract
Pure copper powder was employed to study the effects of ball milling on the development of the structure and properties of ductile metals. The results indicate that larger spheres with diameters of about 2-2.5 mm are created after 20 h of ball milling. The formation of such spheres is mainly due to sphere-to-flake or sphere-to-sphere welding. This welding is not complete, leaving large pores and curved voids in the spheres. The average grain size of such spheres is 10-100 nm. The increase in lattice strain is about 0.2%. The microhardness increases from 45 MPa (unmilled) to 220 MPa (milled for 20 h). High-resolution transmission electron microscopy (HRTEM) investigations show the following: (a) the deformation of ball-milled copper proceeds by [112](111) twinning or high-order twinning; (b) the [112](11i) twins are thickened by passage of (a/6)[112] twinning partial dislocations; (c) subgrains tend to form in the twins. In addition to twinning, dislocation slip plays an important role in the deformation process; the mobility of 60 ° dislocations and their pile-up in the crystals can lead to the formation of subgrains. Crystal refinement leads to an increase in the number of grain boundaries; both low-angle and high-angle grain boundaries with local strain and a high density of dislocations are observed. The estimated mean dislocation density is more than 1014 m 2, which is hardly ever reached in plastically deformed metals. The different kinds of structural defects which exist in the grain boundaries and within the crystals may result in increased strength and microhardness, increased free energy and changes in other properties of ball-milled materials.
Keywords: Ball milling; Copper; Ductility
1. Introduction
Mechanical alloying (MA)/ball milling (BM) (often refers to the milling of single component powders) has attracted considerable interest from the materials science community in recent years. This is mainly due to the wide range of materials exhibiting intriguing properties and structures that can be fabricated using this method. Many materials, such as oxide dispersion strengthened (ODS) alloys [1], intermetallics and inorganics [2,3], a m o r p h o u s and nanocrystalline materials [4-6] and a number of non-equilibrium structures [710], have been synthesized by MA. Although the underlying fundamentals of the M A process are increasingly well understood [11], m a n y details during the mechanical treatment of solids remain unclear. For example, it is well known that repeated mechanical deformation and cold welding often lead to the formation of nanocrystals, but the formation mechanisms and the 0921-5093/95/$09.50 © 1995 SSDI 0921-5093(94)09715-1
Elsevier Science S.A. All rights reserved
structural characterizations of the nanocrystals are poorly understood. This is mainly due to the difficulty in preparing specimens from mechanically alloyed powders for direct transmission electron microscopy (TEM) observations. The simplest method involves suspending the powder in a volatile liquid and pipetting the suspension onto a carbon support film. Usually, T E M images can only be obtained from the edges of mechanically alloyed powders. However, the microstructure of the edges may differ from that in the interior of the particles and, for powders with a large size distribution, the fine fraction m a y not be representative of the whole. In this paper, we succeeded in preparing specimens from ball-milled copper. Atomic level information on the deformation behaviour of ductile metals during ball milling and structural characterization of the end product were achieved. To our knowledge, detailed studies of this type on mechanically alloyed materials have not been performed previously.
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Fig. 1. Scanning electron micrographs of copper powder milled for 20 h showing: (a) small spheres; (b) thick flakes; (c) incomplete primary spheres and secondary spheres; (d) regular and smooth spheres with small fragments. Bar = 1 mm.
2. Experimental details An irregular copper powder (purity, better than 99.0%) with a particle size smaller than 70 lam was milled in a vibratory ball mill in an argon atmosphere. A cylindrical steel vial with an external length of 50 mm and a diameter of 65 mm containing 40 g of powder was used. One 35.5 g, 21 mm (diameter), six 13.7 g, 15 mm, seven 5.5 g, 11 m m and four 0.9 g, 7 mm steel balls, with a ball-to-powder weight ratio of 4 : 1, were used. Scanning electron microscopy (SEM) observations were performed using a Cambridge Instrument $360. For optical microscopy, the resultant materials were mounted on epoxy resin, mechanically ground and etched with a reagent consisting of 50 vol.% H2NO 3 and 50 vol.% water. Microhardness was measured by a Shimadzu Micromet Vicker's tester at a load of 25 gf. After 20 h of ball milling, some of the powder had agglomerated to spheres or flat fragments with diameters of about 2-2.5 mm due to cold welding (as shown in Fig. 1, see Section 3). Some of these were mechanically ground to a thickness of about 30 p.m, and then dimpled with a Gatan Model 656 dimple grinder with 5 gm diamond paste. The as-prepared samples were then ion thinned for high-resolution transmission electron microscopy ( H R T E M ) observations carried out with a JEM-2000EXII electron microscope operated at 200 kV.
3. Results During the first few minutes of milling, the starting powder was flattened and welded together to form thin flakes, which then tended to weld together to form
thicker flakes. After 2 h of milling, in addition to thick flakes, some spheres with diameters of about 0.1 mm had developed. After 20 h of milling, different kinds of particle morphologies were observed, as shown in Fig. 1. Most particles were small spheres (Fig. l(a)) with an average diameter of about 0.7 mm; some ear-like flakes with a diameter of about 2 mm and a thickness of 0.1 mm were also produced, as shown in Fig. l(b). The edges of such flakes tended to convolute, and some small spheres were welded within the flakes. Some open spheres with diameters of about 2.5 mm were also found, as shown in Fig. l(c). The "primary" spheres were found to contain smaller "secondary" ones. The shell of sphere 1 in Fig. l(c) consists of some thicker flakes, whereas those of spheres 2 and 3 are composed of several layers of thicker fragments. The secondary spheres are also incomplete. In addition to incomplete spheres, some regular, smooth spheres were found, as shown in Fig. 1(d). Small fragments were welded on the surface of these smooth spheres (denoted by arrows). In order to understand the interior structure of the spheres, we examined their cross-section using SEM and optical microscopy. Fig. 2(a) shows that there is one large and several smaller pores and some elongated curved voids within the sphere. From the optical micrograph in Fig. 2(b), it can be seen that a convoluted layered structure was produced in the sphere. The crystallite size obtained from X-ray diffraction using the Scherrer formula for unmilled copper is 340 nm, and it decreases to 34 nm after 20 h of milling. From the bright and dark field images of Figs. 3(a) and 3(b), we can estimate that the average grain size is 10-100 nm. Fig. 3(c) shows the electron diffraction pattern (EDP) corresponding to Fig. 3(a). It shows an
J.Y. Huang et al. / Materials Science and Engineering A 199 (1995) 165-172
f.c.c, pattern, and diffraction rings from oxides were also detected. In order to obtain a precise measurement of the atomic level strains induced by ball milling, the following method was used. The X-ray lineshapes produced by the instrument and the small grain size are represented by the function 1/(1 + kZx2), where k is a constant, and those produced by the lattice strain are represented by 1/(1 +k2x2)2; the following f~rmulae were obtained [12] /~=B-b
(1)
fl = (rn + 2n)2/(m + 4n)
(2)
where fl is the total broadening width caused by the strain and the small grain size, B is the half-height peak width of K~j obtained from the experiment and b, m and n correspond to the broadening width caused by the instrument, the small grain size and the strain respectively, b can be corrected. The average grain size and lattice strain can be obtained from the following formula
Fig, 2. Optical micrographs of the spheres formed after 20 h of milling: (a) large pores and curved voids; (b) convoluted layered structure.
167
L = 0.912/(mcosO)
(3)
n = A ]~Sd/d]tanO
(4)
where L is the average grain size, 2 is the X-ray wavelength, 0 is the diffraction angle, A is a constant depending on the strain distribution and is approximately equal to 4 or 8 for a gaussian or random distribution of strains respectively and 16d/d] is the average strain. If we obtain an accurate value of the grain size from TEM, and solve Eqs. (1)-(4), I6d/d[ can be precisely calculated. If L = 100 nm, the calculated strain is 0.2%, which is comparable with that obtained in ball-milled materials (typically 0.1%-0.5%) [13]. The microhardness of as-milled copper reaches 220 MPa, which is about five times larger than that of unmilled copper (45 MPa). Nevertheless, a higher microhardness value (up to 2.5 GPa) was reported in nanocrystalline copper produced by inert gas condensation [14]. In ball-milled copper powder, ~the dominance of welding events over the fracture of powder does not allow a further refinement of the grain size, thus limiting the microhardness to values typical for nanosized materials. In order to obtain a further comprehension of the process of crystal fragmentation and cold welding of the crystalline structure of the ball-milled materials (grain boundaries, defect structure, etc.), we investigated the structure of the grain boundaries and the structural defects of ball-milled copper at the atomic level. Fig. 4 shows a twin boundary containing a step (indicated by a white arrow). The step is associated with an (a/6)[112] twin dislocation. It was found [15] that the twin boundaries in undeformed specimens are planar and rarely show steps, which is indicative of the thermal nature of these twins. In contrast, twin boundaries in deformed specimens tend to form steps, which are equivalent to the ledge associated with an (a/6)[112] twinning disl__ocation. A stacking fault (SF) is impeded by an (a/6)[112] partial dislocation and terminated in the matrix (indicated by white-edged black arrow). A dislocation dipole is also detected near the SF. The mobility of these dislocations may govern the deformation of ball-milled copper. Fig. 5(a) shows the formation of multiple twins in the whole crystal, and the twin boundaries are not flat. Some steps (marked by arrows) with a height of about 2-3 times the d~1~1)spacing were observed, and are also associated with an (a/6)[112] twin dislocation, indicating that the twins are thickened by passage of individual (a/6)[112] partial dislocations. From the enlarged H R T E M image in Fig. 5(b), some 60 ° dislocations were detected near the twin boundaries. Fig. 6(a) shows high-order twins (marked by T~, T2 and T3) and multiple twins (marked by white-edged black lines) in a crystal. Another interesting phe-
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Fig. 3. Transmission electron micrographs obtained from the spheres: (a) bright field image; (b) dark field image; (c) corresponding EDP nomenon is the operation of an intersect twinning process shown in Fig. 6(a). A [112](111) twin intersects another twin at region O. A stacking fault (marked by SF and an arrow) at the tip of the multiple twins was also detected. A more complex configuration of higher order twins was also observed as shown in Fig. 6(b). The five segments are not arranged in fivefold symmetry; the angles between every two segments are different, i.e. vary from 66 ° to 88 ° . In order to fill the whole space, many internal distortions were produced, as indicated by the mismatch dislocations at the interface between segments 1 and 2, the multiple twins in segment 2 and the edge dislocation in segment 5, etc. Sometimes dislocation dipoles are observed inside grains as shown in Fig. 7(a). Detailed analyses found
Fig. 4. HRTEM image showing a twin step and an SF impeded by an (a/6)[]q2] partial dislocation.
that these dislocations belong to two different types (indicated by A and B), and their corresponding Burgers circuits are drawn in Fig. 7(a). The electron beam direction and the dislocation line~ are parallel to [110], the Burgers vector bA = (1/2)[011] or (1/2)[i01] and both have 0 A = 60 ° with the dislocation lines. Therefore the dislocation A is mobile. However, bB= (1/2)[110], 0B = 90 ° and it is immobile. The estimated mean dislocation density is more than 1014 m - z , which is comparable with that of high-strained bulk Cu [16], but is hardly ever reached in plastically deformed metals. The crystal in the upper part of Fig. 7(b) is tilted around the [110] axis by about 4 ° with respect to the grain in the lower part, and a tilted grain boundary is shown by arrows. A series of 60 ° dislocations appear in the grain boundary, which suggests that the low-angle grain boundary is caused by the mobility and pile-up of these dislocations. The two grains may have been one originally. Repeated mechanical deformation may have induced many dislocations. A continuous increase in the dislocation density makes the dislocations difficult to mobilize, and ultimately leads to their pile-up. At a certain dislocation density within these heavily strained regions, the crystal disintegrates into subgrains which are initially separated by low-angle grain boundaries. Multiple twins were also found in the bottom grain of Fig. 7(b), and the twin boundaries stop at the low-angle grain boundary. Fig. 8(a) shows a typical H R T E M image of a highangle grain boundary (at an angle of about 88°). The
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Fig. 5. (a) HRTEM images showing the lenticular morphology and the thickening of multiple twins. The steps at the grain boundary are denoted by arrows. (b) Enlargement of local region in (a), showing the steps and 60 ° dislocations at the twin boundary.
two grains with a common (110) zone axis are heavily deformed; both twins and stacking faults can be seen clearly. The contrast in lattice imaging may result from the orientation changes of local areas caused by internal strain. The grain boundary in Fig. 8(a) is not smooth or flat, but curved (as shown by the arrows). From Fig. 8(b), the local enlargement of Fig. 8(a), we can see that the lattice image near the grain boundary is slightly distorted. Fig. 9(a) shows the triple conjunction of crystals. The angles between crystals 1 and 2, 2 and 3 and 3 and 1 are 165 °, 60 ° and 135 ° respectively. A microtwin is observed in the triple conjunction point, where high internal stress should exist. The formation of the microtwin may thus reduce the shear stress. This is direct evidence for the relaxation of internal stresses through twin formation. A continuous series of grain boundary dislocations was detected in the grain boundary between crystals 1 and 2. From the enlarged H R T E M image of the grain boundary region shown in Fig. 9(b), the dislocations (indicated by marks) are spaced periodically by about 19 X,, which is comparable with the calculated value of 20 A using the formula D = b/ [2sin(0/2)], where b is the Burgers vector and 0 is the angle between the two crystals.
4. Discussion
MA/BM is a very complex process during which the morphology, structures and properties of materials can be significantly changed. In this paper, pure copper was employed to study the effects of ball milling on the development of the morphology, properties and microstructure of ductile metals. The morphology of the powder (Fig. 1) formed after 20 h of milling provides a clue to how the spheres are created. In the first few minutes, ball-to-ball or ball-tovial collisions produce small, thin flakes, and this process continues until small spherical particles are produced. On further milling, the spherical particles increase in size due to sphere-to-flake or sphere-tosphere cold welding. Finallly, larger spheres with diameters of about 2-2.5 mm are developed. In this case, the secondary spheres are small particles welded together to make larger primary spheres, and these spheres are formed by a cold welding mechanism. Obviously, the primary spheres are formed after the secondary spheres. This intriguing phenomenon is very similar to the formation of a "snowball". This type of sphere has also been reported by Harris et al. [17]. They found that, after 222 h of milling, some large hollow
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primary spheres with diameters of about 5.2 mm containing secondary spheres had developed. They later suggested that these spheres were formed by a spalling mechanism: net tensile stresses resolved at or near the inner surface of a shell can create spall fractures, causing fragments to break away from the wall; these fragments then undergo an autogenous milling process to produce smaller, loose secondary spheres. So, in their case, the secondary spheres were formed after the primary spheres, which is contrary to the sphere formation process observed by us. Thus the primary and secondary spheres observed by Harris et al. [17] and in this study are different, and are formed by different mechanisms. The sizes of the spherical particles obtained by Harris et al. [17] (5.2 mm) and in this study (2-2.5 mm) are also different. All these differences are probably caused by the different ball milling conditions, e.g. the ball-to-powder ratio, the weight of the powder, the size and number of the balls used, the mill intensity and the milling time. Sphere-to-flake and sphere-tosphere welding is incomplete, leaving some large pores and voids within the spheres (Fig. 2).
Fig. 7. (a) HRTEM image of 60 ° (denoted by A) and 900 (denoted by B) dislocation dipoles. The Burgers circuits of the two types of dislocations are shown. (b) HRTEM image of a low-angle grain boundary (denoted by arrows) caused by the mobility of 60° dislocations.
Fig. 6. (a) HRTEM image of high-order twins (denoted by T l, T 2 and T3), multiple twins (marked by white-edged black lines) and intersection of twins (the intersection region is denoted by the letter O). A stacking fault is marked by SF and an arrow. (b) HRTEM image of a more complex configuration of higher order twins.
Generally, plastic deformation proceeds by slip and twinning at low and moderate strain rates. MA is a mechanical deformation process during which the powder particles are repeatedly fractured and welded under the ball-to-ball or ball-to-vial collisions. HRTEM investigations show that the deformation process of ballmilled copper occurs via twinning and slip. Two types of twins, i.e. [112](11i) and high-order twins, play a major role in the deformation process, as shown by their high densities. The [112](11]-) twins are thickened by passage of (a/6)[112] partial dislocations, as revealed by the twin steps in Figs. 4 and 5; this suggests that twin thickening is an important mode of strain accommodation in ball-milled copper. The mobility of the twinning dislocations plays an important role in the deformation process during ball milling and depends on the actual spread of the dislocation core. The twins shown in Fig. 5(a) have become quite thick and lenticular. There is a strong tendency to form subgrains.
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Another deformation mode in ball-milled copper is the glide of dislocations. Two types of dislocations, i.e. 60 ° and 90 ° , were observed. Since the 90 ° dislocation is immobile, the 60 ° dislocation can play a major role in the glide process. The mobility and pile-up of the 60 ° dislocations in the crystals can also lead to the formation of subgrains. The grain size reduction by ball milling is usually viewed as follows [18,19]. In the early stage of ball milling, deformation is localized in shear bands with a typical width of 0.1-1 I.tm, as confirmed by Hellstern et al. [6] via T E M observations. On further deformation, the crystal lattice breaks into nanosized grains within shear bands, leading ultimately to a random orientation of neighbouring grains separated by high-angle grain boundaries. These shear bands were also observed in high- strained bulk copper by Hasen and Ralph [16]. They reported that the mean cell size in shear bands was 0.36-0.58 lam. However, such large shear bands were not observed in ball-milled copper powder. This is probably caused by the different properties of powders and bulk materials, i.e. the small grain size in the original copper powder (340 nm) may limit the formation of such large shear bands. Alternatively, it may be caused by the different deformation methods, which
Fig. 9. (a) HRTEM image of a triangular grain boundary. (b) Enlargement of the grain boundary between crystals 1 and 2. The mismatch dislocations are marked.
Fig. 8. (a) HRTEM image of a high-angle grain boundary (denoted by arrows). (b) Enlargement of local region in (a). The trace of the unit cell in each crystal is outlined.
suggests that the deformation modes during ball milling are somewhat different from the conventional plastic deformation process. From the H R T E M observations, it is suggested that the fracture of Cu crystals during ball milling may occur in two ways. (l) The whole crystal is deformed by the formation of [112](11i) twins or high-order twins, and the twins are then thickened by passage of (a/6)[112] partial dislocations. Finally, subgrains may be formed in the twins. (2) The whole crystal is deformed directly by the production of 60 ° dislocations. A continuous increase in the dislocation density leads to dislocation pile-up in the crystals. Furthermore, the crystals disintegrate into subgrains which are initially separated by low-angle grain boundaries (Fig. 7(b)). This process is repeated continuously until a final grain size is reached, and many low-angle and high-angle grain boundaries with local stresses and high densities of dislocations are created. All the grain boundaries investigated are not smooth and flat, but curved. The atoms in the high-angle grain
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boundaries are slightly displaced from the equilibrium positions with respect to the neighbouring crystallite lattice. At the same time, m a n y structural defects are retained inside the grains. All of the structural changes result in an alteration of the properties of the materials, such as an increase in microhardness and strength, increase in free energy, etc. Usually the microhardness and free energy of ball-milled materials are increased [6,20], but it is not well understood whether this increase is caused by grain boundaries or strain inside the crystals. This study suggests that the harder grain boundaries and the crystals with high densities of dislocations may contribute to this increase. In summary, the deformation process of ball-milled copper mainly occurs via twinning and dislocation slide, which is similar to the conventional plastic deformation process. However, this does not mean that they are completely the same, as demonstrated by the different structural changes in the two processes. For example, large shear bands are usually observed in high-strained bulk copper [16], but not in ball-milled copper. Lastly, both the strength of the grain boundaries and the grains may result in a strengthening of the materials during ball milling.
5. Conclusions (1) Large spheres with diameters of about 2 - 2 . 5 m m were produced after 20 h of milling. SEM observations and optical micrographs show that large pores and curved voids exist in the spheres. The formation of such spheres is due to sphere-to-flake or sphere-to-sphere welding. (2) The average grain size obtained by T E M is 10-100 nm. The increase in lattice strain is about 0.2%, and the microhardness increases from 45 M P a to 220 M P a after ball milling. The estimated dislocation density is about 1014 m 2, which is hardly ever reached in plastically deformed metals, but is comparable with heavily strained bulk copper. (3) H R T E M investigations show that the deformation of ball-milled copper proceeds by [112](111) twinning or high-order twinning. The [112](111) twins are thickened by the passage of (a/6)[112] twinning partial dislocations. In addition to twinning, dislocation slip plays an important role in the deformation process; the
mobility of 60 ° dislocations and their pile-up in the crystals can lead to the fracture of crystals. (4) All the grain boundaries investigated are not smooth and flat, but curved, and contain local strain and a high density of dislocations. (5) The grain boundaries and interior of the grains, with high local stresses and high densities of structural defects, may contribute to the increase in strength, microhardness and free energy of ball-milled materials.
Acknowledgment This project was supported by the National Natural Science Foundation of China.
References [1] J.S. Benjamin, Metall. Trans., 1 (1970) 2943. [2] J.S.C. Jang, G. Donelly, P. Godavaati and C.C. Koch, Int. J. Powder Metall., 24 (1988) 315. [3] M.S. Kim and C.C. Koch, J. Appl. Phys., 62 (1987) 3450. [4] C.C. Koch, O.B. Cavin, C.G. McKamey and J.O. Scarbrough, Appl. Phys. Lett., 43 (1983) 1017. [5] R.B. Schwarz, Mater. Sci. Eng., 97 (1988) 71. [6] E. Hellstern, H.J. Fecht, Z. Fu and W.L. Johnson, J. Appl. Phys., 65 (1989) 305. [7] A.O. Aning, Z. Wang and T.H. Courtney, Acta Metall. Mater., 41 (1993) 165. [8] E. Hellstern, L. Schulz, R. Bormann and D. Lewc, Appl. Phys. Lett., 53 (1988) 1399. [9] J.Y. Huang, Y.K. Wu, X.M. Meng and K.Y. Hu, Acta Metall. Sinica B (English Edition), 6 (1993) 286. [10] J.Y. Huang, A.Q. He and Y.K. Wu, Nanostructured Mater., 4 (1994) 1. [11] J.S. Benjamin, Mater. Sci. Forum, 88-90 (1992) 1. [12] X. Fan, X-Ray Diffraction, Tianjian University, Tiangjian, 1981. [13] J. Eckert, J.C. Holzer and W.L. Johnson, J. Appl. Phys., 73 (1991) 131. [14] G.W. Nieman, J.R. Weertman and R.W. Siegel, J. Mater. Res., 6 (1991) 1012. [15] S.R. Singh and J.M. Howe, Philos. Mag. Lett., 65 (1992) 233. [16] N. Hasen and B. Ralph, Acta Metall., 30 (1982) 411. [17] A.M. Harris, G.B. Schaffer and N.W. Page, J. Mater. Sci. Lett., 12 (1993) 1103. [18] H.J. Fecht, E. Hellstern and W.L. Johnson, Metall. Trans. A, 21 (1990) 2333. [19] J. Eckert, J.C. Holzer and W.L. Johnson, Scr. Metall., 27 (1992) 1105. [20] P. LeBrun, E. Gaffet, L. Frogen and L. Delaey, Scr. Metall., 26 (1992) 1743.
Materials Science and Engineering A199 (1995) 165-172
A
Ball milling of ductile metals J.Y. Huang, Y.K. Wu, H.Q. Ye Laboratory of' Atomic Imaging of Solids, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110015, People's Republic of China Received 13 June 1994; in revised form 17 October 1994
Abstract
Pure copper powder was employed to study the effects of ball milling on the development of the structure and properties of ductile metals. The results indicate that larger spheres with diameters of about 2-2.5 mm are created after 20 h of ball milling. The formation of such spheres is mainly due to sphere-to-flake or sphere-to-sphere welding. This welding is not complete, leaving large pores and curved voids in the spheres. The average grain size of such spheres is 10-100 nm. The increase in lattice strain is about 0.2%. The microhardness increases from 45 MPa (unmilled) to 220 MPa (milled for 20 h). High-resolution transmission electron microscopy (HRTEM) investigations show the following: (a) the deformation of ball-milled copper proceeds by [112](111) twinning or high-order twinning; (b) the [112](11i) twins are thickened by passage of (a/6)[112] twinning partial dislocations; (c) subgrains tend to form in the twins. In addition to twinning, dislocation slip plays an important role in the deformation process; the mobility of 60 ° dislocations and their pile-up in the crystals can lead to the formation of subgrains. Crystal refinement leads to an increase in the number of grain boundaries; both low-angle and high-angle grain boundaries with local strain and a high density of dislocations are observed. The estimated mean dislocation density is more than 1014 m 2, which is hardly ever reached in plastically deformed metals. The different kinds of structural defects which exist in the grain boundaries and within the crystals may result in increased strength and microhardness, increased free energy and changes in other properties of ball-milled materials.
Keywords: Ball milling; Copper; Ductility
1. Introduction
Mechanical alloying (MA)/ball milling (BM) (often refers to the milling of single component powders) has attracted considerable interest from the materials science community in recent years. This is mainly due to the wide range of materials exhibiting intriguing properties and structures that can be fabricated using this method. Many materials, such as oxide dispersion strengthened (ODS) alloys [1], intermetallics and inorganics [2,3], a m o r p h o u s and nanocrystalline materials [4-6] and a number of non-equilibrium structures [710], have been synthesized by MA. Although the underlying fundamentals of the M A process are increasingly well understood [11], m a n y details during the mechanical treatment of solids remain unclear. For example, it is well known that repeated mechanical deformation and cold welding often lead to the formation of nanocrystals, but the formation mechanisms and the 0921-5093/95/$09.50 © 1995 SSDI 0921-5093(94)09715-1
Elsevier Science S.A. All rights reserved
structural characterizations of the nanocrystals are poorly understood. This is mainly due to the difficulty in preparing specimens from mechanically alloyed powders for direct transmission electron microscopy (TEM) observations. The simplest method involves suspending the powder in a volatile liquid and pipetting the suspension onto a carbon support film. Usually, T E M images can only be obtained from the edges of mechanically alloyed powders. However, the microstructure of the edges may differ from that in the interior of the particles and, for powders with a large size distribution, the fine fraction m a y not be representative of the whole. In this paper, we succeeded in preparing specimens from ball-milled copper. Atomic level information on the deformation behaviour of ductile metals during ball milling and structural characterization of the end product were achieved. To our knowledge, detailed studies of this type on mechanically alloyed materials have not been performed previously.
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J.Y. Huang et al./ Materials Science and Engineering A 199 (1995) 165-172
Fig. 1. Scanning electron micrographs of copper powder milled for 20 h showing: (a) small spheres; (b) thick flakes; (c) incomplete primary spheres and secondary spheres; (d) regular and smooth spheres with small fragments. Bar = 1 mm.
2. Experimental details An irregular copper powder (purity, better than 99.0%) with a particle size smaller than 70 lam was milled in a vibratory ball mill in an argon atmosphere. A cylindrical steel vial with an external length of 50 mm and a diameter of 65 mm containing 40 g of powder was used. One 35.5 g, 21 mm (diameter), six 13.7 g, 15 mm, seven 5.5 g, 11 m m and four 0.9 g, 7 mm steel balls, with a ball-to-powder weight ratio of 4 : 1, were used. Scanning electron microscopy (SEM) observations were performed using a Cambridge Instrument $360. For optical microscopy, the resultant materials were mounted on epoxy resin, mechanically ground and etched with a reagent consisting of 50 vol.% H2NO 3 and 50 vol.% water. Microhardness was measured by a Shimadzu Micromet Vicker's tester at a load of 25 gf. After 20 h of ball milling, some of the powder had agglomerated to spheres or flat fragments with diameters of about 2-2.5 mm due to cold welding (as shown in Fig. 1, see Section 3). Some of these were mechanically ground to a thickness of about 30 p.m, and then dimpled with a Gatan Model 656 dimple grinder with 5 gm diamond paste. The as-prepared samples were then ion thinned for high-resolution transmission electron microscopy ( H R T E M ) observations carried out with a JEM-2000EXII electron microscope operated at 200 kV.
3. Results During the first few minutes of milling, the starting powder was flattened and welded together to form thin flakes, which then tended to weld together to form
thicker flakes. After 2 h of milling, in addition to thick flakes, some spheres with diameters of about 0.1 mm had developed. After 20 h of milling, different kinds of particle morphologies were observed, as shown in Fig. 1. Most particles were small spheres (Fig. l(a)) with an average diameter of about 0.7 mm; some ear-like flakes with a diameter of about 2 mm and a thickness of 0.1 mm were also produced, as shown in Fig. l(b). The edges of such flakes tended to convolute, and some small spheres were welded within the flakes. Some open spheres with diameters of about 2.5 mm were also found, as shown in Fig. l(c). The "primary" spheres were found to contain smaller "secondary" ones. The shell of sphere 1 in Fig. l(c) consists of some thicker flakes, whereas those of spheres 2 and 3 are composed of several layers of thicker fragments. The secondary spheres are also incomplete. In addition to incomplete spheres, some regular, smooth spheres were found, as shown in Fig. 1(d). Small fragments were welded on the surface of these smooth spheres (denoted by arrows). In order to understand the interior structure of the spheres, we examined their cross-section using SEM and optical microscopy. Fig. 2(a) shows that there is one large and several smaller pores and some elongated curved voids within the sphere. From the optical micrograph in Fig. 2(b), it can be seen that a convoluted layered structure was produced in the sphere. The crystallite size obtained from X-ray diffraction using the Scherrer formula for unmilled copper is 340 nm, and it decreases to 34 nm after 20 h of milling. From the bright and dark field images of Figs. 3(a) and 3(b), we can estimate that the average grain size is 10-100 nm. Fig. 3(c) shows the electron diffraction pattern (EDP) corresponding to Fig. 3(a). It shows an
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f.c.c, pattern, and diffraction rings from oxides were also detected. In order to obtain a precise measurement of the atomic level strains induced by ball milling, the following method was used. The X-ray lineshapes produced by the instrument and the small grain size are represented by the function 1/(1 + kZx2), where k is a constant, and those produced by the lattice strain are represented by 1/(1 +k2x2)2; the following f~rmulae were obtained [12] /~=B-b
(1)
fl = (rn + 2n)2/(m + 4n)
(2)
where fl is the total broadening width caused by the strain and the small grain size, B is the half-height peak width of K~j obtained from the experiment and b, m and n correspond to the broadening width caused by the instrument, the small grain size and the strain respectively, b can be corrected. The average grain size and lattice strain can be obtained from the following formula
Fig, 2. Optical micrographs of the spheres formed after 20 h of milling: (a) large pores and curved voids; (b) convoluted layered structure.
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L = 0.912/(mcosO)
(3)
n = A ]~Sd/d]tanO
(4)
where L is the average grain size, 2 is the X-ray wavelength, 0 is the diffraction angle, A is a constant depending on the strain distribution and is approximately equal to 4 or 8 for a gaussian or random distribution of strains respectively and 16d/d] is the average strain. If we obtain an accurate value of the grain size from TEM, and solve Eqs. (1)-(4), I6d/d[ can be precisely calculated. If L = 100 nm, the calculated strain is 0.2%, which is comparable with that obtained in ball-milled materials (typically 0.1%-0.5%) [13]. The microhardness of as-milled copper reaches 220 MPa, which is about five times larger than that of unmilled copper (45 MPa). Nevertheless, a higher microhardness value (up to 2.5 GPa) was reported in nanocrystalline copper produced by inert gas condensation [14]. In ball-milled copper powder, ~the dominance of welding events over the fracture of powder does not allow a further refinement of the grain size, thus limiting the microhardness to values typical for nanosized materials. In order to obtain a further comprehension of the process of crystal fragmentation and cold welding of the crystalline structure of the ball-milled materials (grain boundaries, defect structure, etc.), we investigated the structure of the grain boundaries and the structural defects of ball-milled copper at the atomic level. Fig. 4 shows a twin boundary containing a step (indicated by a white arrow). The step is associated with an (a/6)[112] twin dislocation. It was found [15] that the twin boundaries in undeformed specimens are planar and rarely show steps, which is indicative of the thermal nature of these twins. In contrast, twin boundaries in deformed specimens tend to form steps, which are equivalent to the ledge associated with an (a/6)[112] twinning disl__ocation. A stacking fault (SF) is impeded by an (a/6)[112] partial dislocation and terminated in the matrix (indicated by white-edged black arrow). A dislocation dipole is also detected near the SF. The mobility of these dislocations may govern the deformation of ball-milled copper. Fig. 5(a) shows the formation of multiple twins in the whole crystal, and the twin boundaries are not flat. Some steps (marked by arrows) with a height of about 2-3 times the d~1~1)spacing were observed, and are also associated with an (a/6)[112] twin dislocation, indicating that the twins are thickened by passage of individual (a/6)[112] partial dislocations. From the enlarged H R T E M image in Fig. 5(b), some 60 ° dislocations were detected near the twin boundaries. Fig. 6(a) shows high-order twins (marked by T~, T2 and T3) and multiple twins (marked by white-edged black lines) in a crystal. Another interesting phe-
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Fig. 3. Transmission electron micrographs obtained from the spheres: (a) bright field image; (b) dark field image; (c) corresponding EDP nomenon is the operation of an intersect twinning process shown in Fig. 6(a). A [112](111) twin intersects another twin at region O. A stacking fault (marked by SF and an arrow) at the tip of the multiple twins was also detected. A more complex configuration of higher order twins was also observed as shown in Fig. 6(b). The five segments are not arranged in fivefold symmetry; the angles between every two segments are different, i.e. vary from 66 ° to 88 ° . In order to fill the whole space, many internal distortions were produced, as indicated by the mismatch dislocations at the interface between segments 1 and 2, the multiple twins in segment 2 and the edge dislocation in segment 5, etc. Sometimes dislocation dipoles are observed inside grains as shown in Fig. 7(a). Detailed analyses found
Fig. 4. HRTEM image showing a twin step and an SF impeded by an (a/6)[]q2] partial dislocation.
that these dislocations belong to two different types (indicated by A and B), and their corresponding Burgers circuits are drawn in Fig. 7(a). The electron beam direction and the dislocation line~ are parallel to [110], the Burgers vector bA = (1/2)[011] or (1/2)[i01] and both have 0 A = 60 ° with the dislocation lines. Therefore the dislocation A is mobile. However, bB= (1/2)[110], 0B = 90 ° and it is immobile. The estimated mean dislocation density is more than 1014 m - z , which is comparable with that of high-strained bulk Cu [16], but is hardly ever reached in plastically deformed metals. The crystal in the upper part of Fig. 7(b) is tilted around the [110] axis by about 4 ° with respect to the grain in the lower part, and a tilted grain boundary is shown by arrows. A series of 60 ° dislocations appear in the grain boundary, which suggests that the low-angle grain boundary is caused by the mobility and pile-up of these dislocations. The two grains may have been one originally. Repeated mechanical deformation may have induced many dislocations. A continuous increase in the dislocation density makes the dislocations difficult to mobilize, and ultimately leads to their pile-up. At a certain dislocation density within these heavily strained regions, the crystal disintegrates into subgrains which are initially separated by low-angle grain boundaries. Multiple twins were also found in the bottom grain of Fig. 7(b), and the twin boundaries stop at the low-angle grain boundary. Fig. 8(a) shows a typical H R T E M image of a highangle grain boundary (at an angle of about 88°). The
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Fig. 5. (a) HRTEM images showing the lenticular morphology and the thickening of multiple twins. The steps at the grain boundary are denoted by arrows. (b) Enlargement of local region in (a), showing the steps and 60 ° dislocations at the twin boundary.
two grains with a common (110) zone axis are heavily deformed; both twins and stacking faults can be seen clearly. The contrast in lattice imaging may result from the orientation changes of local areas caused by internal strain. The grain boundary in Fig. 8(a) is not smooth or flat, but curved (as shown by the arrows). From Fig. 8(b), the local enlargement of Fig. 8(a), we can see that the lattice image near the grain boundary is slightly distorted. Fig. 9(a) shows the triple conjunction of crystals. The angles between crystals 1 and 2, 2 and 3 and 3 and 1 are 165 °, 60 ° and 135 ° respectively. A microtwin is observed in the triple conjunction point, where high internal stress should exist. The formation of the microtwin may thus reduce the shear stress. This is direct evidence for the relaxation of internal stresses through twin formation. A continuous series of grain boundary dislocations was detected in the grain boundary between crystals 1 and 2. From the enlarged H R T E M image of the grain boundary region shown in Fig. 9(b), the dislocations (indicated by marks) are spaced periodically by about 19 X,, which is comparable with the calculated value of 20 A using the formula D = b/ [2sin(0/2)], where b is the Burgers vector and 0 is the angle between the two crystals.
4. Discussion
MA/BM is a very complex process during which the morphology, structures and properties of materials can be significantly changed. In this paper, pure copper was employed to study the effects of ball milling on the development of the morphology, properties and microstructure of ductile metals. The morphology of the powder (Fig. 1) formed after 20 h of milling provides a clue to how the spheres are created. In the first few minutes, ball-to-ball or ball-tovial collisions produce small, thin flakes, and this process continues until small spherical particles are produced. On further milling, the spherical particles increase in size due to sphere-to-flake or sphere-tosphere cold welding. Finallly, larger spheres with diameters of about 2-2.5 mm are developed. In this case, the secondary spheres are small particles welded together to make larger primary spheres, and these spheres are formed by a cold welding mechanism. Obviously, the primary spheres are formed after the secondary spheres. This intriguing phenomenon is very similar to the formation of a "snowball". This type of sphere has also been reported by Harris et al. [17]. They found that, after 222 h of milling, some large hollow
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primary spheres with diameters of about 5.2 mm containing secondary spheres had developed. They later suggested that these spheres were formed by a spalling mechanism: net tensile stresses resolved at or near the inner surface of a shell can create spall fractures, causing fragments to break away from the wall; these fragments then undergo an autogenous milling process to produce smaller, loose secondary spheres. So, in their case, the secondary spheres were formed after the primary spheres, which is contrary to the sphere formation process observed by us. Thus the primary and secondary spheres observed by Harris et al. [17] and in this study are different, and are formed by different mechanisms. The sizes of the spherical particles obtained by Harris et al. [17] (5.2 mm) and in this study (2-2.5 mm) are also different. All these differences are probably caused by the different ball milling conditions, e.g. the ball-to-powder ratio, the weight of the powder, the size and number of the balls used, the mill intensity and the milling time. Sphere-to-flake and sphere-tosphere welding is incomplete, leaving some large pores and voids within the spheres (Fig. 2).
Fig. 7. (a) HRTEM image of 60 ° (denoted by A) and 900 (denoted by B) dislocation dipoles. The Burgers circuits of the two types of dislocations are shown. (b) HRTEM image of a low-angle grain boundary (denoted by arrows) caused by the mobility of 60° dislocations.
Fig. 6. (a) HRTEM image of high-order twins (denoted by T l, T 2 and T3), multiple twins (marked by white-edged black lines) and intersection of twins (the intersection region is denoted by the letter O). A stacking fault is marked by SF and an arrow. (b) HRTEM image of a more complex configuration of higher order twins.
Generally, plastic deformation proceeds by slip and twinning at low and moderate strain rates. MA is a mechanical deformation process during which the powder particles are repeatedly fractured and welded under the ball-to-ball or ball-to-vial collisions. HRTEM investigations show that the deformation process of ballmilled copper occurs via twinning and slip. Two types of twins, i.e. [112](11i) and high-order twins, play a major role in the deformation process, as shown by their high densities. The [112](11]-) twins are thickened by passage of (a/6)[112] partial dislocations, as revealed by the twin steps in Figs. 4 and 5; this suggests that twin thickening is an important mode of strain accommodation in ball-milled copper. The mobility of the twinning dislocations plays an important role in the deformation process during ball milling and depends on the actual spread of the dislocation core. The twins shown in Fig. 5(a) have become quite thick and lenticular. There is a strong tendency to form subgrains.
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Another deformation mode in ball-milled copper is the glide of dislocations. Two types of dislocations, i.e. 60 ° and 90 ° , were observed. Since the 90 ° dislocation is immobile, the 60 ° dislocation can play a major role in the glide process. The mobility and pile-up of the 60 ° dislocations in the crystals can also lead to the formation of subgrains. The grain size reduction by ball milling is usually viewed as follows [18,19]. In the early stage of ball milling, deformation is localized in shear bands with a typical width of 0.1-1 I.tm, as confirmed by Hellstern et al. [6] via T E M observations. On further deformation, the crystal lattice breaks into nanosized grains within shear bands, leading ultimately to a random orientation of neighbouring grains separated by high-angle grain boundaries. These shear bands were also observed in high- strained bulk copper by Hasen and Ralph [16]. They reported that the mean cell size in shear bands was 0.36-0.58 lam. However, such large shear bands were not observed in ball-milled copper powder. This is probably caused by the different properties of powders and bulk materials, i.e. the small grain size in the original copper powder (340 nm) may limit the formation of such large shear bands. Alternatively, it may be caused by the different deformation methods, which
Fig. 9. (a) HRTEM image of a triangular grain boundary. (b) Enlargement of the grain boundary between crystals 1 and 2. The mismatch dislocations are marked.
Fig. 8. (a) HRTEM image of a high-angle grain boundary (denoted by arrows). (b) Enlargement of local region in (a). The trace of the unit cell in each crystal is outlined.
suggests that the deformation modes during ball milling are somewhat different from the conventional plastic deformation process. From the H R T E M observations, it is suggested that the fracture of Cu crystals during ball milling may occur in two ways. (l) The whole crystal is deformed by the formation of [112](11i) twins or high-order twins, and the twins are then thickened by passage of (a/6)[112] partial dislocations. Finally, subgrains may be formed in the twins. (2) The whole crystal is deformed directly by the production of 60 ° dislocations. A continuous increase in the dislocation density leads to dislocation pile-up in the crystals. Furthermore, the crystals disintegrate into subgrains which are initially separated by low-angle grain boundaries (Fig. 7(b)). This process is repeated continuously until a final grain size is reached, and many low-angle and high-angle grain boundaries with local stresses and high densities of dislocations are created. All the grain boundaries investigated are not smooth and flat, but curved. The atoms in the high-angle grain
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boundaries are slightly displaced from the equilibrium positions with respect to the neighbouring crystallite lattice. At the same time, m a n y structural defects are retained inside the grains. All of the structural changes result in an alteration of the properties of the materials, such as an increase in microhardness and strength, increase in free energy, etc. Usually the microhardness and free energy of ball-milled materials are increased [6,20], but it is not well understood whether this increase is caused by grain boundaries or strain inside the crystals. This study suggests that the harder grain boundaries and the crystals with high densities of dislocations may contribute to this increase. In summary, the deformation process of ball-milled copper mainly occurs via twinning and dislocation slide, which is similar to the conventional plastic deformation process. However, this does not mean that they are completely the same, as demonstrated by the different structural changes in the two processes. For example, large shear bands are usually observed in high-strained bulk copper [16], but not in ball-milled copper. Lastly, both the strength of the grain boundaries and the grains may result in a strengthening of the materials during ball milling.
5. Conclusions (1) Large spheres with diameters of about 2 - 2 . 5 m m were produced after 20 h of milling. SEM observations and optical micrographs show that large pores and curved voids exist in the spheres. The formation of such spheres is due to sphere-to-flake or sphere-to-sphere welding. (2) The average grain size obtained by T E M is 10-100 nm. The increase in lattice strain is about 0.2%, and the microhardness increases from 45 M P a to 220 M P a after ball milling. The estimated dislocation density is about 1014 m 2, which is hardly ever reached in plastically deformed metals, but is comparable with heavily strained bulk copper. (3) H R T E M investigations show that the deformation of ball-milled copper proceeds by [112](111) twinning or high-order twinning. The [112](111) twins are thickened by the passage of (a/6)[112] twinning partial dislocations. In addition to twinning, dislocation slip plays an important role in the deformation process; the
mobility of 60 ° dislocations and their pile-up in the crystals can lead to the fracture of crystals. (4) All the grain boundaries investigated are not smooth and flat, but curved, and contain local strain and a high density of dislocations. (5) The grain boundaries and interior of the grains, with high local stresses and high densities of structural defects, may contribute to the increase in strength, microhardness and free energy of ball-milled materials.
Acknowledgment This project was supported by the National Natural Science Foundation of China.
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