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Deformation induced grain boundaries in commercially pure aluminium
Acta mater. 48 (2000) 3377±3385 www.elsevier.com/locate/actamat
DEFORMATION INDUCED GRAIN BOUNDARIES IN COMMERCIALLY PURE ALUMINIUM C. P. CHANG{, P. L. SUN and P. W. KAO Institute of Materials Science and Engineering, Sun Yat-Sen University, Kaohsiung 80424, Taiwan (Received 7 February 2000; received in revised form 9 May 2000; accepted 9 May 2000) AbstractÐThe structure of commercially pure aluminium processed by equal channel angular extrusion at room temperature, to an equivalent strain up to 8, was studied by transmission electron microscopy. Attention was paid to the structure evolution of the generated boundaries. Three types of boundaries were found, namely polygonized dislocation wall, partially transformed boundary, and grain boundary. By increasing strain, polygonized dislocation walls transform ®rstly into partially transformed boundaries, then into grain boundaries. Dissociation of lattice dislocations is suggested to cause the transformation to occur. At equivalent strain equal to 8, most of the generated boundaries transform into grain boundaries. 7 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Aluminium; Grain boundaries; Plastic deformation
1. INTRODUCTION
Recently, there have been many reports on the microstructure and properties of ultrahigh-strained metals [1±5]. The extremely high amount of deformation can be introduced by several dierent ways, e.g. compressed-torsion [3], equal-channel angular extrusion (ECAE) (also called equal-channel angular pressing) [6±8], cyclic-extrusion-compression [9], and accumulative roll-banding process [10]. Many unusual characters were found in these materials [1±5], including the generation of a submicrongrained structure. Normally, grain re®nement in metals is achieved by either static or dynamic recrystallization at a certain elevated temperature; it is therefore surprising to ®nd that grain re®nement can be achieved by plastic deformation at room temperature. Despite the fact that the generation of a submicron-grained structure has been reported by many authors, there appears to be a lack of detailed and consistent descriptions of the generated boundaries in the ®ne-grained structure. The term ``non-equilibrium grain boundary'' has been intensely used by the Ufa group [1±3] to describe the boundaries that are generated by compressed-torsion and ECAE in
{ To whom all correspondence should be addressed. Fax: +886-7-5254099. E-mail address: [email protected] (C.P. Chang).
several materials, e.g. Al±4%Cu±0.5%Zr, Mg± 1.5%Mn±0.3%Ce, Zn±22%Al, and Cu. They reported extinction contours in these grains, and no ordinary grain boundary fringes were found on these boundaries. The grain boundary fringes could only be seen after the specimens were subjected to annealing. These authors believed that the extinction contours were caused by internal stress, resulting from the attachment of dislocations on grain boundaries. Selected area diraction (SAD) patterns obtained from these materials showed incomplete ring pattern, indicating the boundaries were not low angle ones. They also reported virtually no dislocations within the grains in aluminium alloys, but very high dislocation density in copper. The size of the grains generated varied from 01 mm to 00.2 mm, depending on the material and the strain introduced. The nature of non-equilibrium grain boundary was studied by Horita et al. [11, 12] in Al±Mg alloys. The alloys were subjected to a strain of about 7 by using a compressed-torsion technique. By using lattice imaging in TEM, they found distortion or bending of lattice fringes near non-equilibrium boundaries. They also found some lattice fringes were missing indicating the possible presence of dislocation near the boundaries. All these results suggest that the boundaries generated by torsion were under a high energy state. Ferrasse et al. [5] have studied the microstructure of copper and 3003 aluminium alloy, processed by ECAE up to an equivalent strain of 9.2. They used
1359-6454/00/$20.00 7 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 0 ) 0 0 1 3 8 - 5
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CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
the term ``subgrain'' to describe the ®ne-grained structure. They found that the subgrain interiors appeared to be free of dislocations, and most of the dislocations resided near subgrain boundaries. Extinction contours within subgrain boundaries were found. They observed reasonably complete SAD ring patterns from the ®ne-grained structure. A quite dierent result of ECAE processed copper at an equivalent strain of about 11 was given by Agnew and Weertman [13], they found the structure was composed of a combination of ®ne equiaxed cells and lamellar shear bands. From their transmission electron micrographs, it can be seen that the structure resembles a conventional deformation structure. Systematic study of the eect of processing parameters on the microstructure of pure aluminium (99.99% purity) deformed by ECAE has been done by Nemoto, Langdon, and their coworkers [14±18]. Iwahashi et al. [14] studied the microstructural evolution of pure aluminium at dierent strain levels. They found elongated subgrains were formed at low strains, these elongated subgrains transformed into ®ne equiaxed grains of about 1 mm at an equivalent strain of 010. By measuring the spreading of diffraction spots in SAD patterns, they concluded that generated boundaries had high misorientation. They gave no statement about whether these boundaries were grain boundaries. In the other published papers [15±18], they described the generated boundaries as high angle grain boundaries. Unfortunately, both the magni®cation and the imaging conditions of most of their TEM micrograghs have prevented the characteristic of these boundaries to be seen clearly. From their TEM micrographs, virtually no dislocation can be seen within these ®ne-grained interiors. They did not report any abnormal extinction contour in pure aluminium specimens. Microstructure of pure aluminium (99.99% purity) deformed by cyclic-extrusion-compression up to a strain of 60 was studied by Richert et al. [9]. They reported that equiaxed cells/subgrains with the size of about 1.2±1.4 mm were formed. They used electron microdiraction to measure the misorientation of the generated boundaries, and found that about two out of three of the boundaries had misorientation larger than 158. In ®g. 10 of their paper, some generated high angle boundaries had typical grain boundary fringes, suggesting that these boundaries may be grain boundaries. From the above discussion, it appears that by introducing very large amounts of plastic deformation into a material, ultra®ne-grained structure can be produced. However there remains a question as to whether or not the boundaries generated by deformation are grain boundaries or dense dislocation boundaries. A grain boundary is a twodimensional planar defect, while a dislocation boundary consists of one-dimensional line defects. The aim of the present research is to clarify the
question as to whether or not grain boundaries can be generated by plastic deformation at room temperature using ECAE. We choose ECAE as the process route, since by this novel technique, a very large amount of strain can be introduced into a work piece without altering its size. 2. METHODS
For ECAE, the tool (Fig. 1) is a die with two intersecting channels of identical cross-section. A specimen of almost the same cross-section is placed into one of the channels, and a punch then extrudes it into the second channel. Under these conditions, deformation is achieved by simple shear. Deformed by ECAE, the sample retains the same cross-sectional area so that it is able to repeat the process to several cycles. Therefore, very large plastic strains could be accumulated in the sample. The material used in this work was commercially pure 1050 Al (99.5% pure), the material was homogenized at 873 K for 12 h, and air cooled to room temperature. The initial grain size of the material is about 500 mm. Specimens of 12 12 80 mm were subjected to ECAE for various passes at room temperature. The die for ECAE consists of two channels at an angle of F=908. The angle, C, de®ning the arc of the two channels is 208. The equivalent strain on pressing through the die was calculated to be 1 [19]. The strain rate applied during ECAE is estimated to be about 0.1 sÿ1. Samples were passed through the die repetitively up to eight times. The sample was rotated counterclockwise about the exit extrusion axis by 908 between each pass. This process route was designated as route Bc by Furukawa et al. [20]. Thin foils for TEM were ®rst sliced perpendicular
Fig. 1. Schematic con®guration of ECAE die and the specimen. Surface-1 faces to the front, bottom, rear, and upper side of the exit channel at extrusion pass 1, 2, 3 and 4, respectively.
CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
to the extrusion axis, then mechanically thinned down to about 200 mm thick and ®nally polished by a standard twin-jet polishing method using a electrolyte of 25% nitric acid and 75% methanol at 58C, 10 V. TEM was carried out using a Philips CM200 microscope operated at 200 kV. Boundary misorientation, ym, was measured by microdiraction. A reference grain was ®rstly tilted to a beam direction, say [110], the convergent beam electron diraction (CBED) pattern of this grain was taken, then the neighbouring grain was tilted to a h110i direction, the CBED pattern of this grain was also taken. By recording the tilt readings from the microscope and measuring the rotation angle between the two CBED patterns, the tilt component, yt, and the rotation component, yr, of the misorientation between these two grains can be measured. ym is obtained by taking the minimum value of the square root of the sum of the squares of yt and yr [21]. Caution needs to be taken while
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tilting the neighbouring grain, since sometimes despite tilt to the nearest h110i orientation will get minimum yt, but may not have the minimum value of ym. 3. RESULTS
Figure 2 shows the development of a ®ne-grained structure from low strain to large strain. The equivalent strain, e, of Fig. 2(a)±(c) is 2, 4 and 8, respectively. The shape of grains was found to have developed from an elongated shape into an equiaxed shape, meanwhile the size of the grains decreased with increasing strain. By the linear interception method without any correction factor, the apparent grain size for e=2, 4 and 8 is 0.61, 0.57 and 0.35 mm, respectively. The dislocation density in the grain interiors decreased with increasing strain. At e=8 most of the grain interiors were free of dislocations. In some grain interiors, dislocation
Fig. 2. TEM micrographs showing the structural evolution at an equivalent strain of (a) 2, (b) 4 and (c) 8. Note that by the increase of strain, grain size decreases, grain boundary fringes show up, dislocation density in the grain interiors drops, and the grains become equiaxed.
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CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
Fig. 3. TEM micrograph showing dislocation loop debris was found in some grain interiors.
loop debris could be seen (Fig. 3), indicating that massive dislocation annihilations have occurred. It needs to be noted that these loops were not caused by electron damage at 200 kV, since they were seen right from the start of the observation. After studying many boundaries generated at dierent strains, it was found that they could be classi®ed into three types, namely (a) polygonized dislocation wall (PDW), (b) partially transformed boundary (PTB), and (c) grain boundary. Figure 4 shows a grain, which has all three types of boundaries. A PDW is a conventional subgrain boundary, which consists of perfect lattice dislocations, has very low misorientation and is typically less than 18. The misorientation of PTBs is in the range of about 1±58, thickness fringe can be seen on these
Fig. 4. TEM micrograph showing a grain, which is surrounded by a polygonized dislocation wall (PDW), a partially transformed boundary (PTB), and grain boundaries (GB). The equivalent strain was 8.
boundaries, but the fringe contrast is interfered by dislocations. A PTB also consists of dislocations, but the dislocation spacing is so small that the interference of diracted beams which arise from the periodic displacement ®elds of the dislocations preventing individual dislocation to be resolved under bright ®eld imaging condition [22] [Figs 4 and 5(a)]. Although the dislocations of a PTB could not be resolved under bright ®eld imaging conditions, they could sometimes be resolved under weak beam imaging conditions, as shown in Fig. 5. After carefully examining the contrast of these dislocations under dierent diracted beams, it was found that these dislocations were not perfect dislocations. Since no perfect dislocation in fcc aluminium, could give in and 020, and out-contrast at contrast at g 1 11 g=220, as shown in Fig. 5(b), (c) and (e). Attempt has been made to identify the dislocations by considering Burgers vector b=1/6h211i, but without success, since it is dicult to decide the contrast characteristic of the dislocations in Fig. 5(d). The same kind of problem arises when examining the dislocations in other PTBs. The only conclusion can be made from our analysis is that the dislocations in a PTB had dissociated. Another important characteristic of PTBs is the appearance of strain contrast beside the boundaries, as indicated by the arrows in Fig. 5(a), indicating that the strain ®eld of a PTB may extend a few nanometers from the boundary plane. The misorientation of the PTB in Fig. 5 is 1.78 Perhaps the most striking result was the generation of grain boundary fringes in these specimens, indicating that the transformation of dislocation walls into grain boundaries had occurred. In order to con®rm this, a series of diraction conditions were taken from a grain to ensure that any lattice dislocation existing on the boundaries would be seen. Figure 6 shows the result of a set of analysis. From Fig. 6, it can be seen that (a) no extrinsic grain boundary dislocation can be imaged, (b) no strain contrast is associated with the boundaries, and (c) the grain interior is free of dislocation. Weak beam images of Fig. 6(a)±(d) have also been taken, and no dislocation was found. Misorientation of the grain boundaries in Fig. 6 has also been measured, and is shown in the ®gure, these grain boundaries are high angle boundaries. The misorientation of 79 well-developed grain boundaries have been measured, for e=4 and 8, and the result is given in Fig. 7. At e=4 and 8, 85% and 95% of the grain boundaries, respectively, have misorientations larger than 108. With increasing strain, the distribution of misorientation is shifted towards higher angles. The proportion of well-developed grain boundary was also found to increase with increasing strain. At e=8, more than 090% of the boundaries have transformed into grain boundaries.
CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
Fig. 5. A partially transformed boundary was imaged under bright ®eld, and weak beam imaging conditions under several g vectors. The arrows in (a) indicate the strain ®eld beside the partially transformed boundary. The misorientation of this boundary is 1.78. See text for detailed discussion. The equivalent strain was 8.
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CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES 4. DISCUSSION
From the results presented above, the formation of submicron-grained structure with grain boundaries in pure aluminium by ECAE has been veri®ed. Since the size of the grains decreases with increasing strain, and the deformation was conducted at room temperature, the possibility that the formation of grain boundaries caused by the migration of PDWs or PTBs is excluded. Further supporting evidence on the above conclusion is shown in Fig. 8(a). The grain boundaries of an L-shape grain cannot possibly be formed by a migration mechanism. This picture is not dicult to understand, when the idea that grain boundaries were transformed from PDWs is accepted. Figure 8(b) shows the schematic diagram of the boundary positions in Fig. 8(a) ignoring the structure of the boundary, the morphology of these boundaries is the same as that of dislocation cells generated by conventional defor-
mation methods [23]. Imagine that all the boundaries were PDWs at low strain, with increasing strain, boundaries 1±10, except 11, transformed into grain boundaries, resulting in the structure of Fig. 8(a). Whether or not a boundary transforms, depends on the activity of dislocations on both sides of the boundary, and the stress state of the boundary. According to the above idea, the whole structure evolution is basically a transformation process from dislocation walls into grain boundaries, and can be described as follows. At low plastic strain, say at e 22, the structure is similar to a conventional deformation structure with some degree of recovery. Most of the boundaries generated are PDWs, together with a few PTBs and grain boundaries. The PDWs generated have misorientation less than 018, while grain boundaries generated, at this strain level, have misorientation less than several degrees.
Fig. 6. TEM micrographs showing no extrinsic grain boundary dislocation can be imaged under four dierent g vectors. Misorientation angles of the grain boundaries are also shown. The equivalent strain was 8.
CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
Fig. 7. Histogram showing the relatively frequency of misorientation angle of generated grain boundaries at equivalent strain (a) 4 and (b) 8.
When the strain is increased, say at e 2 4, some PDWs transform into PTBs and some PTBs transform into grain boundaries, so that the proportion
3383
of PDWs decreases, and the proportion of PTBs and grain boundaries increases. When the strain is further increased, say at e 2 8, the proportion of PDWs and PTBs decreases, most of the boundaries now transform into grain boundaries, and most of the grain boundaries now have misorientation larger than 108. Although the detailed mechanism of the transformation from a PDW into a grain boundary is still not clear at this stage, from the fact that PDWs ®rstly transform into PTBs, then transform into grain boundaries, the transformation process is expected to be caused by an increase of dislocation density in these boundaries. For each extrusion pass, the strain imposed is very large, a very large number of dislocations travel between boundaries, massive annihilation between travelling dislocations is expected and the debris loops found in Fig. 3 are the result of this annihilation process. On the other hand, many glide dislocations are forced to be deposited onto the boundaries. Continuously depositing dislocations on the boundaries will reduce dislocation spacing, when the dislocation spacing becomes smaller than some distance, the interaction energy between lattice dislocations becomes so high that the dislocations start to dissociate. As dislocations continue to be deposited onto the boundaries, dislocation spacing is further reduced, and these dislocations eventually lose the identity of individual dislocation, and become grain boundary structure. Once grain boundary structure has developed, glide dislocations that come afterwards, deposit onto the grain boundaries, and dissociation of these extrinsic grain boundary dislocations occurs [24, 25]. The fact that no extrinsic grain boundary dislocation could be imaged in most of the grain boundaries indicates that the dissociation process is very eective. As deformation continues,
Fig. 8. (a) TEM micrograph showing an L-shape grain contained a PDW. (b) Schematic diagram of the boundary positions in (a). See text for detailed discussion. The equivalent strain was 8.
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CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
the rotation of the grains causes the misorientation of grain boundary to increase to a higher value, as shown in Fig. 7. In the literature [26], the generation of a ®negrained structure induced by large strain, either using conventional methods or using ECAE has been regarded as a kind of geometric dynamic recrystallization [27, 28]. Geometric dynamic recrystallization is a process in which the original grains are ¯attened at large strain, and the original grain boundaries become serrations due to dynamic recovery, when the grain boundary spacing is reduced to a certain distance, i.e. about twice the subgrain size, these grain boundaries will pinch-o, and ®ne equiaxed grains are formed. When the deformation is induced by conventional methods, strong evidences for the formation of ®ne-grained structure caused by geometric dynamic recrystallization at large strain, at elevated temperature, have been provided [27, 28]. However, when the deformation is introduced by ECAE, as in the present case, there is no experimental evidence to support
Fig. 9. Photographs showing the shearing of the engraved line marks before extrusion, and after extrusion pass 1±4. N indicates the number of extrusion passes. Note that despite the line marks on surface-1 being restored to its vertical position after only three passes, it needs four passes to restore the line marks on every surface of the specimen back to their original position.
the view that the formation of ®ne-grained structure is caused by geometric dynamic recrystallization. Intuitively, since very large strain is induced by ECAE, the ¯attening of original grains seems inevitable, therefore the generation of a ®ne-grained structure was said to be caused by geometric dynamic recrystallization [26]. In the literature, four dierent ECAE processing routes have been adopted, namely route A, BA, BC, and C [5, 7, 20]. For route A, the specimen is not rotated between each extrusion pass; for route BA, the specimen is rotated 908 in clockwise and counterclockwise directions consecutively about the extrusion axis between each extrusion pass; for route BC, the specimen is rotated 908, either clockwise or counterclockwise, about the extrusion axis between each extrusion pass; and for route C, the specimen is rotated 1808 about the extrusion axis between each extrusion pass. Furukawa et al. [20] have analyzed the shearing characteristics associated with each route. By considering the shape of a cube grain after each ECAE route, Furukawa et al. indicated that the grain is restored to its cube shape after every 2N and 4N passes for route C and BC respectively, where N is number of extrusion passes. This means that when a specimen is deformed to a very large strain by route BC or C, the original grain shape will not necessarily be ¯attened. Experimental evidence for the restoration of grain shape by route C has been given by Segal [7]. In order to check the shape restoration in route BC, line marks were engraved on the surfaces of a specimen. After each extrusion pass, photograph of the marks on each surface was taken. Figure 9 shows the photographs taken from surface-1, indicated in Fig. 1, before extrusion and after pass 1±4. Although aected by surface friction, the restoration was not perfect, but it is clearly that ¯attening eect did not occur. In general, the shearing characteristics and restoration behaviour that were found in ECAE route BC followed the analysis of Furukawa et al. [20]. The generation of a ®negrained structure in this study is therefore not caused by geometric dynamic recrystallization. The formation of a ®ne-grained structure appears to be aected by many factors; such as strain level [14], material [3, 17], processing route [5, 15], strain rate [29], and angular angle between two channels [16]. Strain level is not the only factor that causes the formation of grain boundaries. Further study on the detailed transformation mechanism is needed in order to understand this phenomenon fully. 5. CONCLUSIONS
By using equal channel angular extrusion, submicron-grained structure can be obtained in commercially pure aluminium at room temperature. During this process, three types of boundaries were found, namely polygonized dislocation wall, partially trans-
CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
formed boundary, and grain boundary. By increasing strain, polygonized dislocation walls transform ®rstly into partially transformed boundaries, then into grain boundaries. Dissociation of lattice dislocations is believed to cause the transformation to occur. At an equivalent strain equal to 8, most of the generated boundaries transform into grain boundaries, and 95% of these grain boundaries had a misorientation larger than 108. According to the present study, genuine submicron grains can be produced in commercially pure aluminum by equal channel angular extrusion. AcknowledgementsÐThe authors wish to acknowledge the ®nancial support of the National Science Council with contracts NSC 88-2216-E-110-006 and 89-2216-E-110-004. REFERENCES 1. Valiev, R. Z., Krasilnikov, N. A. and Tsenev, N. K., Mater. Sci. Engng, 1991, A137, 35. 2. Valiev, R. Z., Korznikov, A. V. and Mulyukov, R. R., Mater. Sci. Engng, 1993, A168, 141. 3. Valiev, R. Z., Mater. Sci. Engng, 1997, A234-6, 59. 4. Valiev, R. Z., Kozlov, E. V., Ivanov, Y. F., Lian, J., Nazarov, A. A. and Baudelet, B., Acta metall. mater., 1994, 42, 2467. 5. Ferrasse, S., Segal, V. M., Hartwig, K. T. and Goforth, R. E., Metall. Mater. Trans., 1997, 28A, 1047. 6. Segal, V. M., Patent of the USSR, no. 575892, 1977. 7. Segal, V. M., Mater. Sci. Engng, 1995, A197, 157. 8. Segal, V. M., Goforth RE, Hartwig KT, US Patent no. 5,400,633, 1995. 9. Richert, M., Liu, Q. and Hansen, N., Mater. Sci. Engng, 1999, A260, 275. 10. Tsuji, N., Saito, Y., Utsunomiya, H. and Tanigawa, S., Scripta mater., 1999, 40, 795. 11. Horita, Z., Smith, D. J., Furukawa, M., Nemoto, M.,
12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29.
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Valiev, R. Z. and Langdon, T. G., Mater. Characterization, 1996, 37, 285. Horita, Z., Smith, D. J., Furukawa, M., Nemoto, M., Valiev, R. Z. and Langdon, T. G., J. Mater. Res., 1996, 11, 1880. Agnew, S. R. and Weertman, J. R., Mater. Sci. Engng, 1998, A244, 145. Iwahashi, Y., Horita, Z., Nemoto, M. and Langdon, T. G., Acta mater., 1997, 45, 4733. Iwahashi, Y., Horita, Z., Nemoto, M. and Langdon, T. G., Acta mater., 1998, 46, 3317. Nakashima, K., Horita, Z., Nemoto, M. and Langdon, T. G., Acta mater., 1998, 46, 1589. Iwahashi, Y., Horita, Z., Nemoto, M. and Langdon, T. G., Metall. Mater. Trans., 1998, 29A, 2503. Iwahashi, Y., Furukawa, M., Horita, Z., Nemoto, M. and Langdon, T. G., Metall. Mater. Trans., 1998, 29A, 2245. Iwahashi, Y., Wang, J., Horita, Z., Nemoto, M. and Langdon, T. G., Scripta Mater., 1996, 35, 143. Furukawa, M., Iwahashi, Y., Horita, Z., Nemoto, M. and Langdon, T. G., Mater. Sci. Engng, 1998, A257, 328. Barker, I., Hansen, N. and Ralph, B., Mater. Sci. Engng, 1989, A113, 449. Forwood, C. T. and Clarebrough, L. M., Electron Microscopy of Interfaces in Metals and Alloys. Adam Hilger, Bristol, 1991, chapter 4. Bay, B., Hansen, N., Hughes, D. A. and KuhlmannWilsdorf, D., Acta metall. mater., 1992, 40, 205. Darby, T. P., Schindler, R. and Ballu, R. W., Phil. Mag., 1978, 37, 245. Pumphrey, P. H. and Gleiter, H., Phil. Mag., 1974, 30, 593. Humphreys, F. J., Acta mater., 1997, 45, 4231. Drury, M. R. and Humphreys, F. J., Acta metall., 1986, 34, 2259. Solberg, J. K., McQueen, H. J., Ryum, N. and Nes, E., Phil. Mag., 1989, A60, 447. Berbon, P. B., Furukawa, M., Horita, Z., Nemoto, M. and Langdon, T. G., Metall. Mater. Trans., 1999, 30A, 1989.
DEFORMATION INDUCED GRAIN BOUNDARIES IN COMMERCIALLY PURE ALUMINIUM C. P. CHANG{, P. L. SUN and P. W. KAO Institute of Materials Science and Engineering, Sun Yat-Sen University, Kaohsiung 80424, Taiwan (Received 7 February 2000; received in revised form 9 May 2000; accepted 9 May 2000) AbstractÐThe structure of commercially pure aluminium processed by equal channel angular extrusion at room temperature, to an equivalent strain up to 8, was studied by transmission electron microscopy. Attention was paid to the structure evolution of the generated boundaries. Three types of boundaries were found, namely polygonized dislocation wall, partially transformed boundary, and grain boundary. By increasing strain, polygonized dislocation walls transform ®rstly into partially transformed boundaries, then into grain boundaries. Dissociation of lattice dislocations is suggested to cause the transformation to occur. At equivalent strain equal to 8, most of the generated boundaries transform into grain boundaries. 7 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Aluminium; Grain boundaries; Plastic deformation
1. INTRODUCTION
Recently, there have been many reports on the microstructure and properties of ultrahigh-strained metals [1±5]. The extremely high amount of deformation can be introduced by several dierent ways, e.g. compressed-torsion [3], equal-channel angular extrusion (ECAE) (also called equal-channel angular pressing) [6±8], cyclic-extrusion-compression [9], and accumulative roll-banding process [10]. Many unusual characters were found in these materials [1±5], including the generation of a submicrongrained structure. Normally, grain re®nement in metals is achieved by either static or dynamic recrystallization at a certain elevated temperature; it is therefore surprising to ®nd that grain re®nement can be achieved by plastic deformation at room temperature. Despite the fact that the generation of a submicron-grained structure has been reported by many authors, there appears to be a lack of detailed and consistent descriptions of the generated boundaries in the ®ne-grained structure. The term ``non-equilibrium grain boundary'' has been intensely used by the Ufa group [1±3] to describe the boundaries that are generated by compressed-torsion and ECAE in
{ To whom all correspondence should be addressed. Fax: +886-7-5254099. E-mail address: [email protected] (C.P. Chang).
several materials, e.g. Al±4%Cu±0.5%Zr, Mg± 1.5%Mn±0.3%Ce, Zn±22%Al, and Cu. They reported extinction contours in these grains, and no ordinary grain boundary fringes were found on these boundaries. The grain boundary fringes could only be seen after the specimens were subjected to annealing. These authors believed that the extinction contours were caused by internal stress, resulting from the attachment of dislocations on grain boundaries. Selected area diraction (SAD) patterns obtained from these materials showed incomplete ring pattern, indicating the boundaries were not low angle ones. They also reported virtually no dislocations within the grains in aluminium alloys, but very high dislocation density in copper. The size of the grains generated varied from 01 mm to 00.2 mm, depending on the material and the strain introduced. The nature of non-equilibrium grain boundary was studied by Horita et al. [11, 12] in Al±Mg alloys. The alloys were subjected to a strain of about 7 by using a compressed-torsion technique. By using lattice imaging in TEM, they found distortion or bending of lattice fringes near non-equilibrium boundaries. They also found some lattice fringes were missing indicating the possible presence of dislocation near the boundaries. All these results suggest that the boundaries generated by torsion were under a high energy state. Ferrasse et al. [5] have studied the microstructure of copper and 3003 aluminium alloy, processed by ECAE up to an equivalent strain of 9.2. They used
1359-6454/00/$20.00 7 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 0 ) 0 0 1 3 8 - 5
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CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
the term ``subgrain'' to describe the ®ne-grained structure. They found that the subgrain interiors appeared to be free of dislocations, and most of the dislocations resided near subgrain boundaries. Extinction contours within subgrain boundaries were found. They observed reasonably complete SAD ring patterns from the ®ne-grained structure. A quite dierent result of ECAE processed copper at an equivalent strain of about 11 was given by Agnew and Weertman [13], they found the structure was composed of a combination of ®ne equiaxed cells and lamellar shear bands. From their transmission electron micrographs, it can be seen that the structure resembles a conventional deformation structure. Systematic study of the eect of processing parameters on the microstructure of pure aluminium (99.99% purity) deformed by ECAE has been done by Nemoto, Langdon, and their coworkers [14±18]. Iwahashi et al. [14] studied the microstructural evolution of pure aluminium at dierent strain levels. They found elongated subgrains were formed at low strains, these elongated subgrains transformed into ®ne equiaxed grains of about 1 mm at an equivalent strain of 010. By measuring the spreading of diffraction spots in SAD patterns, they concluded that generated boundaries had high misorientation. They gave no statement about whether these boundaries were grain boundaries. In the other published papers [15±18], they described the generated boundaries as high angle grain boundaries. Unfortunately, both the magni®cation and the imaging conditions of most of their TEM micrograghs have prevented the characteristic of these boundaries to be seen clearly. From their TEM micrographs, virtually no dislocation can be seen within these ®ne-grained interiors. They did not report any abnormal extinction contour in pure aluminium specimens. Microstructure of pure aluminium (99.99% purity) deformed by cyclic-extrusion-compression up to a strain of 60 was studied by Richert et al. [9]. They reported that equiaxed cells/subgrains with the size of about 1.2±1.4 mm were formed. They used electron microdiraction to measure the misorientation of the generated boundaries, and found that about two out of three of the boundaries had misorientation larger than 158. In ®g. 10 of their paper, some generated high angle boundaries had typical grain boundary fringes, suggesting that these boundaries may be grain boundaries. From the above discussion, it appears that by introducing very large amounts of plastic deformation into a material, ultra®ne-grained structure can be produced. However there remains a question as to whether or not the boundaries generated by deformation are grain boundaries or dense dislocation boundaries. A grain boundary is a twodimensional planar defect, while a dislocation boundary consists of one-dimensional line defects. The aim of the present research is to clarify the
question as to whether or not grain boundaries can be generated by plastic deformation at room temperature using ECAE. We choose ECAE as the process route, since by this novel technique, a very large amount of strain can be introduced into a work piece without altering its size. 2. METHODS
For ECAE, the tool (Fig. 1) is a die with two intersecting channels of identical cross-section. A specimen of almost the same cross-section is placed into one of the channels, and a punch then extrudes it into the second channel. Under these conditions, deformation is achieved by simple shear. Deformed by ECAE, the sample retains the same cross-sectional area so that it is able to repeat the process to several cycles. Therefore, very large plastic strains could be accumulated in the sample. The material used in this work was commercially pure 1050 Al (99.5% pure), the material was homogenized at 873 K for 12 h, and air cooled to room temperature. The initial grain size of the material is about 500 mm. Specimens of 12 12 80 mm were subjected to ECAE for various passes at room temperature. The die for ECAE consists of two channels at an angle of F=908. The angle, C, de®ning the arc of the two channels is 208. The equivalent strain on pressing through the die was calculated to be 1 [19]. The strain rate applied during ECAE is estimated to be about 0.1 sÿ1. Samples were passed through the die repetitively up to eight times. The sample was rotated counterclockwise about the exit extrusion axis by 908 between each pass. This process route was designated as route Bc by Furukawa et al. [20]. Thin foils for TEM were ®rst sliced perpendicular
Fig. 1. Schematic con®guration of ECAE die and the specimen. Surface-1 faces to the front, bottom, rear, and upper side of the exit channel at extrusion pass 1, 2, 3 and 4, respectively.
CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
to the extrusion axis, then mechanically thinned down to about 200 mm thick and ®nally polished by a standard twin-jet polishing method using a electrolyte of 25% nitric acid and 75% methanol at 58C, 10 V. TEM was carried out using a Philips CM200 microscope operated at 200 kV. Boundary misorientation, ym, was measured by microdiraction. A reference grain was ®rstly tilted to a beam direction, say [110], the convergent beam electron diraction (CBED) pattern of this grain was taken, then the neighbouring grain was tilted to a h110i direction, the CBED pattern of this grain was also taken. By recording the tilt readings from the microscope and measuring the rotation angle between the two CBED patterns, the tilt component, yt, and the rotation component, yr, of the misorientation between these two grains can be measured. ym is obtained by taking the minimum value of the square root of the sum of the squares of yt and yr [21]. Caution needs to be taken while
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tilting the neighbouring grain, since sometimes despite tilt to the nearest h110i orientation will get minimum yt, but may not have the minimum value of ym. 3. RESULTS
Figure 2 shows the development of a ®ne-grained structure from low strain to large strain. The equivalent strain, e, of Fig. 2(a)±(c) is 2, 4 and 8, respectively. The shape of grains was found to have developed from an elongated shape into an equiaxed shape, meanwhile the size of the grains decreased with increasing strain. By the linear interception method without any correction factor, the apparent grain size for e=2, 4 and 8 is 0.61, 0.57 and 0.35 mm, respectively. The dislocation density in the grain interiors decreased with increasing strain. At e=8 most of the grain interiors were free of dislocations. In some grain interiors, dislocation
Fig. 2. TEM micrographs showing the structural evolution at an equivalent strain of (a) 2, (b) 4 and (c) 8. Note that by the increase of strain, grain size decreases, grain boundary fringes show up, dislocation density in the grain interiors drops, and the grains become equiaxed.
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Fig. 3. TEM micrograph showing dislocation loop debris was found in some grain interiors.
loop debris could be seen (Fig. 3), indicating that massive dislocation annihilations have occurred. It needs to be noted that these loops were not caused by electron damage at 200 kV, since they were seen right from the start of the observation. After studying many boundaries generated at dierent strains, it was found that they could be classi®ed into three types, namely (a) polygonized dislocation wall (PDW), (b) partially transformed boundary (PTB), and (c) grain boundary. Figure 4 shows a grain, which has all three types of boundaries. A PDW is a conventional subgrain boundary, which consists of perfect lattice dislocations, has very low misorientation and is typically less than 18. The misorientation of PTBs is in the range of about 1±58, thickness fringe can be seen on these
Fig. 4. TEM micrograph showing a grain, which is surrounded by a polygonized dislocation wall (PDW), a partially transformed boundary (PTB), and grain boundaries (GB). The equivalent strain was 8.
boundaries, but the fringe contrast is interfered by dislocations. A PTB also consists of dislocations, but the dislocation spacing is so small that the interference of diracted beams which arise from the periodic displacement ®elds of the dislocations preventing individual dislocation to be resolved under bright ®eld imaging condition [22] [Figs 4 and 5(a)]. Although the dislocations of a PTB could not be resolved under bright ®eld imaging conditions, they could sometimes be resolved under weak beam imaging conditions, as shown in Fig. 5. After carefully examining the contrast of these dislocations under dierent diracted beams, it was found that these dislocations were not perfect dislocations. Since no perfect dislocation in fcc aluminium, could give in and 020, and out-contrast at contrast at g 1 11 g=220, as shown in Fig. 5(b), (c) and (e). Attempt has been made to identify the dislocations by considering Burgers vector b=1/6h211i, but without success, since it is dicult to decide the contrast characteristic of the dislocations in Fig. 5(d). The same kind of problem arises when examining the dislocations in other PTBs. The only conclusion can be made from our analysis is that the dislocations in a PTB had dissociated. Another important characteristic of PTBs is the appearance of strain contrast beside the boundaries, as indicated by the arrows in Fig. 5(a), indicating that the strain ®eld of a PTB may extend a few nanometers from the boundary plane. The misorientation of the PTB in Fig. 5 is 1.78 Perhaps the most striking result was the generation of grain boundary fringes in these specimens, indicating that the transformation of dislocation walls into grain boundaries had occurred. In order to con®rm this, a series of diraction conditions were taken from a grain to ensure that any lattice dislocation existing on the boundaries would be seen. Figure 6 shows the result of a set of analysis. From Fig. 6, it can be seen that (a) no extrinsic grain boundary dislocation can be imaged, (b) no strain contrast is associated with the boundaries, and (c) the grain interior is free of dislocation. Weak beam images of Fig. 6(a)±(d) have also been taken, and no dislocation was found. Misorientation of the grain boundaries in Fig. 6 has also been measured, and is shown in the ®gure, these grain boundaries are high angle boundaries. The misorientation of 79 well-developed grain boundaries have been measured, for e=4 and 8, and the result is given in Fig. 7. At e=4 and 8, 85% and 95% of the grain boundaries, respectively, have misorientations larger than 108. With increasing strain, the distribution of misorientation is shifted towards higher angles. The proportion of well-developed grain boundary was also found to increase with increasing strain. At e=8, more than 090% of the boundaries have transformed into grain boundaries.
CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
Fig. 5. A partially transformed boundary was imaged under bright ®eld, and weak beam imaging conditions under several g vectors. The arrows in (a) indicate the strain ®eld beside the partially transformed boundary. The misorientation of this boundary is 1.78. See text for detailed discussion. The equivalent strain was 8.
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CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES 4. DISCUSSION
From the results presented above, the formation of submicron-grained structure with grain boundaries in pure aluminium by ECAE has been veri®ed. Since the size of the grains decreases with increasing strain, and the deformation was conducted at room temperature, the possibility that the formation of grain boundaries caused by the migration of PDWs or PTBs is excluded. Further supporting evidence on the above conclusion is shown in Fig. 8(a). The grain boundaries of an L-shape grain cannot possibly be formed by a migration mechanism. This picture is not dicult to understand, when the idea that grain boundaries were transformed from PDWs is accepted. Figure 8(b) shows the schematic diagram of the boundary positions in Fig. 8(a) ignoring the structure of the boundary, the morphology of these boundaries is the same as that of dislocation cells generated by conventional defor-
mation methods [23]. Imagine that all the boundaries were PDWs at low strain, with increasing strain, boundaries 1±10, except 11, transformed into grain boundaries, resulting in the structure of Fig. 8(a). Whether or not a boundary transforms, depends on the activity of dislocations on both sides of the boundary, and the stress state of the boundary. According to the above idea, the whole structure evolution is basically a transformation process from dislocation walls into grain boundaries, and can be described as follows. At low plastic strain, say at e 22, the structure is similar to a conventional deformation structure with some degree of recovery. Most of the boundaries generated are PDWs, together with a few PTBs and grain boundaries. The PDWs generated have misorientation less than 018, while grain boundaries generated, at this strain level, have misorientation less than several degrees.
Fig. 6. TEM micrographs showing no extrinsic grain boundary dislocation can be imaged under four dierent g vectors. Misorientation angles of the grain boundaries are also shown. The equivalent strain was 8.
CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
Fig. 7. Histogram showing the relatively frequency of misorientation angle of generated grain boundaries at equivalent strain (a) 4 and (b) 8.
When the strain is increased, say at e 2 4, some PDWs transform into PTBs and some PTBs transform into grain boundaries, so that the proportion
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of PDWs decreases, and the proportion of PTBs and grain boundaries increases. When the strain is further increased, say at e 2 8, the proportion of PDWs and PTBs decreases, most of the boundaries now transform into grain boundaries, and most of the grain boundaries now have misorientation larger than 108. Although the detailed mechanism of the transformation from a PDW into a grain boundary is still not clear at this stage, from the fact that PDWs ®rstly transform into PTBs, then transform into grain boundaries, the transformation process is expected to be caused by an increase of dislocation density in these boundaries. For each extrusion pass, the strain imposed is very large, a very large number of dislocations travel between boundaries, massive annihilation between travelling dislocations is expected and the debris loops found in Fig. 3 are the result of this annihilation process. On the other hand, many glide dislocations are forced to be deposited onto the boundaries. Continuously depositing dislocations on the boundaries will reduce dislocation spacing, when the dislocation spacing becomes smaller than some distance, the interaction energy between lattice dislocations becomes so high that the dislocations start to dissociate. As dislocations continue to be deposited onto the boundaries, dislocation spacing is further reduced, and these dislocations eventually lose the identity of individual dislocation, and become grain boundary structure. Once grain boundary structure has developed, glide dislocations that come afterwards, deposit onto the grain boundaries, and dissociation of these extrinsic grain boundary dislocations occurs [24, 25]. The fact that no extrinsic grain boundary dislocation could be imaged in most of the grain boundaries indicates that the dissociation process is very eective. As deformation continues,
Fig. 8. (a) TEM micrograph showing an L-shape grain contained a PDW. (b) Schematic diagram of the boundary positions in (a). See text for detailed discussion. The equivalent strain was 8.
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the rotation of the grains causes the misorientation of grain boundary to increase to a higher value, as shown in Fig. 7. In the literature [26], the generation of a ®negrained structure induced by large strain, either using conventional methods or using ECAE has been regarded as a kind of geometric dynamic recrystallization [27, 28]. Geometric dynamic recrystallization is a process in which the original grains are ¯attened at large strain, and the original grain boundaries become serrations due to dynamic recovery, when the grain boundary spacing is reduced to a certain distance, i.e. about twice the subgrain size, these grain boundaries will pinch-o, and ®ne equiaxed grains are formed. When the deformation is induced by conventional methods, strong evidences for the formation of ®ne-grained structure caused by geometric dynamic recrystallization at large strain, at elevated temperature, have been provided [27, 28]. However, when the deformation is introduced by ECAE, as in the present case, there is no experimental evidence to support
Fig. 9. Photographs showing the shearing of the engraved line marks before extrusion, and after extrusion pass 1±4. N indicates the number of extrusion passes. Note that despite the line marks on surface-1 being restored to its vertical position after only three passes, it needs four passes to restore the line marks on every surface of the specimen back to their original position.
the view that the formation of ®ne-grained structure is caused by geometric dynamic recrystallization. Intuitively, since very large strain is induced by ECAE, the ¯attening of original grains seems inevitable, therefore the generation of a ®ne-grained structure was said to be caused by geometric dynamic recrystallization [26]. In the literature, four dierent ECAE processing routes have been adopted, namely route A, BA, BC, and C [5, 7, 20]. For route A, the specimen is not rotated between each extrusion pass; for route BA, the specimen is rotated 908 in clockwise and counterclockwise directions consecutively about the extrusion axis between each extrusion pass; for route BC, the specimen is rotated 908, either clockwise or counterclockwise, about the extrusion axis between each extrusion pass; and for route C, the specimen is rotated 1808 about the extrusion axis between each extrusion pass. Furukawa et al. [20] have analyzed the shearing characteristics associated with each route. By considering the shape of a cube grain after each ECAE route, Furukawa et al. indicated that the grain is restored to its cube shape after every 2N and 4N passes for route C and BC respectively, where N is number of extrusion passes. This means that when a specimen is deformed to a very large strain by route BC or C, the original grain shape will not necessarily be ¯attened. Experimental evidence for the restoration of grain shape by route C has been given by Segal [7]. In order to check the shape restoration in route BC, line marks were engraved on the surfaces of a specimen. After each extrusion pass, photograph of the marks on each surface was taken. Figure 9 shows the photographs taken from surface-1, indicated in Fig. 1, before extrusion and after pass 1±4. Although aected by surface friction, the restoration was not perfect, but it is clearly that ¯attening eect did not occur. In general, the shearing characteristics and restoration behaviour that were found in ECAE route BC followed the analysis of Furukawa et al. [20]. The generation of a ®negrained structure in this study is therefore not caused by geometric dynamic recrystallization. The formation of a ®ne-grained structure appears to be aected by many factors; such as strain level [14], material [3, 17], processing route [5, 15], strain rate [29], and angular angle between two channels [16]. Strain level is not the only factor that causes the formation of grain boundaries. Further study on the detailed transformation mechanism is needed in order to understand this phenomenon fully. 5. CONCLUSIONS
By using equal channel angular extrusion, submicron-grained structure can be obtained in commercially pure aluminium at room temperature. During this process, three types of boundaries were found, namely polygonized dislocation wall, partially trans-
CHANG et al.: DEFORMATION INDUCED GRAIN BOUNDARIES
formed boundary, and grain boundary. By increasing strain, polygonized dislocation walls transform ®rstly into partially transformed boundaries, then into grain boundaries. Dissociation of lattice dislocations is believed to cause the transformation to occur. At an equivalent strain equal to 8, most of the generated boundaries transform into grain boundaries, and 95% of these grain boundaries had a misorientation larger than 108. According to the present study, genuine submicron grains can be produced in commercially pure aluminum by equal channel angular extrusion. AcknowledgementsÐThe authors wish to acknowledge the ®nancial support of the National Science Council with contracts NSC 88-2216-E-110-006 and 89-2216-E-110-004. REFERENCES 1. Valiev, R. Z., Krasilnikov, N. A. and Tsenev, N. K., Mater. Sci. Engng, 1991, A137, 35. 2. Valiev, R. Z., Korznikov, A. V. and Mulyukov, R. R., Mater. Sci. Engng, 1993, A168, 141. 3. Valiev, R. Z., Mater. Sci. Engng, 1997, A234-6, 59. 4. Valiev, R. Z., Kozlov, E. V., Ivanov, Y. F., Lian, J., Nazarov, A. A. and Baudelet, B., Acta metall. mater., 1994, 42, 2467. 5. Ferrasse, S., Segal, V. M., Hartwig, K. T. and Goforth, R. E., Metall. Mater. Trans., 1997, 28A, 1047. 6. Segal, V. M., Patent of the USSR, no. 575892, 1977. 7. Segal, V. M., Mater. Sci. Engng, 1995, A197, 157. 8. Segal, V. M., Goforth RE, Hartwig KT, US Patent no. 5,400,633, 1995. 9. Richert, M., Liu, Q. and Hansen, N., Mater. Sci. Engng, 1999, A260, 275. 10. Tsuji, N., Saito, Y., Utsunomiya, H. and Tanigawa, S., Scripta mater., 1999, 40, 795. 11. Horita, Z., Smith, D. J., Furukawa, M., Nemoto, M.,
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Valiev, R. Z. and Langdon, T. G., Mater. Characterization, 1996, 37, 285. Horita, Z., Smith, D. J., Furukawa, M., Nemoto, M., Valiev, R. Z. and Langdon, T. G., J. Mater. Res., 1996, 11, 1880. Agnew, S. R. and Weertman, J. R., Mater. Sci. Engng, 1998, A244, 145. Iwahashi, Y., Horita, Z., Nemoto, M. and Langdon, T. G., Acta mater., 1997, 45, 4733. Iwahashi, Y., Horita, Z., Nemoto, M. and Langdon, T. G., Acta mater., 1998, 46, 3317. Nakashima, K., Horita, Z., Nemoto, M. and Langdon, T. G., Acta mater., 1998, 46, 1589. Iwahashi, Y., Horita, Z., Nemoto, M. and Langdon, T. G., Metall. Mater. Trans., 1998, 29A, 2503. Iwahashi, Y., Furukawa, M., Horita, Z., Nemoto, M. and Langdon, T. G., Metall. Mater. Trans., 1998, 29A, 2245. Iwahashi, Y., Wang, J., Horita, Z., Nemoto, M. and Langdon, T. G., Scripta Mater., 1996, 35, 143. Furukawa, M., Iwahashi, Y., Horita, Z., Nemoto, M. and Langdon, T. G., Mater. Sci. Engng, 1998, A257, 328. Barker, I., Hansen, N. and Ralph, B., Mater. Sci. Engng, 1989, A113, 449. Forwood, C. T. and Clarebrough, L. M., Electron Microscopy of Interfaces in Metals and Alloys. Adam Hilger, Bristol, 1991, chapter 4. Bay, B., Hansen, N., Hughes, D. A. and KuhlmannWilsdorf, D., Acta metall. mater., 1992, 40, 205. Darby, T. P., Schindler, R. and Ballu, R. W., Phil. Mag., 1978, 37, 245. Pumphrey, P. H. and Gleiter, H., Phil. Mag., 1974, 30, 593. Humphreys, F. J., Acta mater., 1997, 45, 4231. Drury, M. R. and Humphreys, F. J., Acta metall., 1986, 34, 2259. Solberg, J. K., McQueen, H. J., Ryum, N. and Nes, E., Phil. Mag., 1989, A60, 447. Berbon, P. B., Furukawa, M., Horita, Z., Nemoto, M. and Langdon, T. G., Metall. Mater. Trans., 1999, 30A, 1989.