Shear banding in rolled dispersion hardened AlMg2Si alloys

Scripta METALLURGICA

Vol.

23, pp. 1811-1816, 1989 Printed in the U.S.A.

Pergamon Press plc All rights reserved

SHEAR BANDING IN ROLLED DISPERSION HARDENED A1-Mg2Si ALLOYS John Liu Alcoa Laboratories Alcoa Center, PA 15069, USA Michael Mato and Roger D. Doherty Department of Materials Engineering, Drexel University Philadelphia, PA 19104, USA ( R e c e i v e d March 30, ( R e v i s e d A u g u s t 17,

1989) 1989)

Introduction A crucially important feature of plastic deformation is the homogeneity of the strain distribution. If strain can localize in highly strained "slaear bands", major effects are expected in the mechanical response and particularly in the resulting microstructure of the deformed material. This phenomenon has, in the last decade, received considerable attention in various fields (1-8). In particular, this topic has a long history (9, 10) in the literature of the metallic deformed state and recrystallization. As a result of many microstructural studies of rolled polycrystalline metals, various features have been identified that appear to promote shear banding. These features include prior mechanical twinning (11-13), lamellar dislocation structures (14), high solute (e.g., in aluminum alloys (3, 14-16)), finely dispersed strengthening precipitates (2, 3, 16), and high strain, especially when accompanied by changes of strain state (17). All these microstructural features enhance work hardening that continues out to strains of 2 or higher typical of heavily-rolled metal. Data, up to 1983, on the role of microstructure in the development of shear bands in aluminum, have been reviewed by Chandra-Holm and Embury (18). Although the detailed microstructural mechanisms are not yet clear, the current view is that some type of dynamic recovery mechanism appears to be involved that leads to localized dynamic softening in the shear band (3) even though no macroscopic work softening can usually be detected (17). There is also experimental evidence that deformation-induced heating may play a role, with the extreme case being the adiabatic condition (19, 20). It should be noted that other models for shear banding are available that are purely mechanical (4-6) or based on texture softening (2, 5, 6) and have no direct role for microstructure. One recent major success in recrystaUization modelling has resulted from the detailed studies by Humphreys and his coworkers (21-24) of "deformation zones" produced by plastic strain around large particles (radius > 0.5 I.tm-l.0 I.tm) in a range of alloys. Humphreys (21), Sandstrom (25) and Nes (26-28) have obtained moderate success in understanding and predicting the recrystallization behavior of materials containing coarse particles, based on an understanding of these deformation zones. However, these workers implicitly assumed that even with heavy rolling reduction, the deformation remains homogeneous, at least on a scale greater than that of individual particles and slip planes. The assumption of an absence of shear banding with coarse particles was apparently a reasonable one because such bands had not been reported in any of the studies on recrystallization of heavily-rolled coarse particle containing materials (21-24, 29-33). Moreover, alloys with coarse particles which are not sheared by dislocation slip are expected to show more homogeneous slip, at least at low strains (34, 35). Although coarse particles promote increased work hardening, this saturates at small strains (36). The higher work hardening of solute containing alloys has also been suggested as promoting shear banding (3, 15, 16). A study was recently made on the influence of particle size and spacing on the magnitude of the misorientation of the deformation zones at coarse particles and the consequent recrystailization characteristics in model AI-Mg2Si alloys (37-39). In this work, it was discovered that all two-phase alloys showed intense shear banding after cold rolling to 90% reduction, as did a single-phase alloy with a matrix composition identical to that of the two-phase alloys. It was found that the matrix adjacent to particles in a shear band exhibited very large misorientations, while much smaller misorientations were observed outside these bands (38, 39). The purpose of the present

1811 0036-9748/89 $3.00 + .00 Copyright (c) 1989 Pergamon Press plc

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paper is to describe briefly the direct microstructural evidence for strain localization in coarsely dispersed A1-Mg2Si alloys after 90% reduction and provide surface observations of strain localization produced by incremental rolling reductions in materials given to a wide range of strains prior to roiling.

Experimental Methods and Observations A series of quasi-binary AI-Mg2Si alloys from 0.31 to 1.85 wt.% Mg2Si were cast, processed to produce a series of different dispersions, all solute equilibrated at 350°C, at which temperature the solubility of Mg2Si is close to that of the leanest alloy - 0.15 wt.% Mg2Si. One sample, with a coarse dispersion, was given a supersaturation treatment at 480°C followed by a quench and a low temperature age to give a secondary fine dispersion before rolling. Full details of the processing of the alloys, the quantitative analysis of the dispersion and methods of assessment of the deformed microstructures are presented elsewhere (38, 39). Table 1 shows the alloys used for shear band studies. Included in the table are corresponding dispersion details: volume fraction, particle size (as major and minor semiaxes of oblate spheroids) and nearest neighbor spacing A3. In the first study, all the samples were cold rolled in 15 passes in a 20 cm diameter roiling mill to a 90% reduction to 0.48 mm. The material as-deformed and after short annealing at 350"C was examined by optical and transmission electron microscopy. All alloys are described by a numeric/alphabetic combination - digits denote wt.% Mg2Si and letters represent dispersion types (details in Reference (39) and Table 1). Figures 1-3 (alloys 0.31A, 0.85B and 1.85C, see Table I) show the longitudinal section of three alloys as-rolled in which the polarized light micrographs show clear evidence of both in-grain shear bands and those that cross grain boundaries, both seen at close to the usual angle of =, 35" to the rolling direction (3-6). The severity of the banding, at a nearly constant precipitate size, appears to increase with the increased volume fraction of precipitate, Figures 1-3, and at a constant volume fraction with increased precipitate size and spacing. Figure 4 shows the weaker shear banding seen in the precipitate-free alloy (0.15A). Despite the thinness of the samples it was possible to obtain, by a technique developed by Smith (40), TEM foils from the plane normal to the transverse direction, Figure 5 (Alloy 1.85E). This is the same section as that of the optical micrographs. As distinct from the regions outside the shear band, where the subgrains are more equiaxed, within the shear band rite subgrains are elongated in the shear direction, a feature previously reported in other studies of shear banding in aluminum alloys (3, 17). As discussed in more detail elsewhere (38, 39), the subgrains in the shear band are very highly misoriented both adjacent to and away from large particles. Surface observations of strain localization were made by the application of an incremental rolling strain on previously rolled material after polishing the longitudinal section and marking with vertical scratches from the diamond tip of a surface profilometer. This technique for demonstrating shear banding has been previously used in other studies (3, 12). Figure 6 (Alloy 1.85E) shows intense shear bands which were produced by a very small additional rolling reduction (0.08 mm) in material previously rolled with a reduction of 90% from a starting thickness of 4.8 mm in a 0.2 m diameter roiling mill. The surface shear bands are seen producing offsets on the surface scratch and at grain boundaries. Subsequently, seven of the alloys initially used (37, 38) were further investigated by the surface method after rolling 40% to 80% reductions, at 10% intervals, using a 15 cm diameter rolling mill. The sample widths were 25 mm. Here the additional reductions on polished and scratched samples were 3% to 5% of the original thicknesses. Table 1 gives the microstructural details (38) of the alloys used and for various prior reductions, the maximum shear band offset on the vertical scratches resolved into the rolling direction and the angle between that shear band and the rolling direction. For smaller prior reductions, no detectable scratch offsets could be seen on any of six scratches. After 40% prior strain, only the solid solution alloy, 0.15A, showed the offsets. However, only two offsets were seen in six scratches. After 50% reduction, two of the alloys, 0.15A again and the leanest two-phase alloy, 0.31B, showed very few offsets. After 60% reduction, the solid solution alloy showed abundant offsets, and alloy 0.31B showed offsets fast crossing grain boundaries, but offsets were still not seen on every scratch. Alloy 0.54A showed several offsets, while 0.85A and 1.85E had occasional offsets. All seven alloys studied showed offsets after 70% reduction, although for the alloy with highest volume fraction of small particles, 1.85B, only a few offsets were seen. Shear band offsets were seen crossing the grain boundaries that were parallel to the rolling direction in alloys 0.31B, 0.54A, 0.85A and 1.85E, but not in the solid solution alloy 0.15A. After 80% reduction, extensive shear banding was seen at scratches in all samples, and in all samples except 1.85B offsets crossing grain boundaries were seen.

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Based on the results of the scratch experiments, there is a weak tendency for offsets to become larger with larger strain. This is not a strong effect - with most offsets being about 2 ~tm. A more important trend is that, judging by the strain required to induce shear banding at surface scratches, the solid solution alloy actually had the strongest tendency to strain localization. Increased volume fraction of precipitates and, at a constant volume fraction of precipitates, smaller and thus more closely spaced precipitates delayed shear banding to slightly larger strains, but clearly did not prevent its onset. However, based on the observation made earlier from metallographic etching effects solely at 90% reduction a different conclusion could be drawn; that is, the solid solution alloy appeared to be less prone to shear banding. This discrepancy in conclusions can be rationalized in light of the following considerations. While the metallographic technique provides some qualitative information about the severity of shear banding, it is not capable of quantifying dimensional displacements at the shear bands. The mid-width location for the metallographic sample, however, substantially satisfies the plane strain condition (there was no lateral spread due to rolling, and hence, there were lateral stresses). On the other hand, the scratch experiments allow shear band offsets to be measured. However, because there are no lateral stresses at a free surface, plane swain conditions are not met. Therefore, the discrepancy is probably the result of a combination of the qualitative nature of the metallographic technique, and the difference in stress state. We point out these issues in order that such information may be of use to those performing similar experiments. They do not, however, influence our major conclusions.

Conclusions The observations based on the results of the scratch experiments show that although a high density of non-shearable precipitates 0.4 Ixm to 2 I.tm in size, measured as the major semiaxis of oblate spheroids (38, 39), slightly delays the onset of shear banding to higher swains, nevertheless strong strain localization can and does occur in particle-dispersed aluminum alloys containing a small amount of solute. The shear bands produce large local misorientations, both at the intersection of shear bands and particles and in the shear bands away from particles. These large misorientations were shown by Liu and Doherty (37-39) to give rise to prolific nucleation of recrystailized grains. It appears, therefore, that it is unlikely to be valid to try to model the recrystallization process and the resulting grain size (and texture) in particle dispersed roiled alloys without recognition of the possibility of shear banding. Given the widespread incidence of strain localization in rolled alloys, it is not perhaps surprising that shear banding can readily occur even in alloys containing coarse particle dispersions. However, it is striking that the phenomenon has not been previously reported despite the detailed work carried out in the recrystallization of two-phase alloys (21-33). It should be noted that Wood and McCormick (41) have recently reported what they describe as a third type of strain localization in addition to those caused by precipitate shear and by dynamic swain aging. This third type of swain localization was found in a strongly deformed and heavily overaged (and thus particle containing) A1-Zn-Mg alloy. This was associated by Wood and McCormick with the shear bands seen in heavily rolled materials. The increased work hardening at small strains, of particle-containing alloys, which leads to slip homogenization in lightly deformed two-phase alloys (34, 35) appears to have little significance for the onset of strain localization seen on the scale of macroscopic shear bands in heavily-rolled alloys.

Acknowledeements The authors would like to thank Dr. S. F. Baumann of Alcoa Laboratories for valuable discussions, and Alcoa Laboratories for permission to publish the results. The authors are also grateful for the helpful comments of our reviewer - these included bringing Reference 41 to our attention.

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References 1) 2) 3) 4) 5) 6) 7) 8) 9) 10) 11) 12) 13) 14) 15) 16) 17) 18) 19) 20) 21) 22) 23) 24) 25) 26) 27) 28) 29) 30) 31) 32) 33) 34) 35) 36) 37) 38) 39) 40) 41)

J. Grewen, T. Noda and D. Sauer, Z. Metallkde., 68 (1977) 260. I.L. Dillamore, J. G. Roberts and A. C. Bush, Metal Science, 13 (1979) 73. A. Korbel et al., Acta Metall., 34 (1986) 1999. L. Anand and W. A. Spitzig, Acta Metall., 30 (1982) 553. J. Gil-Sevillano, P. van Houtte and E. Aernoudt, Prog. in Materials Sci., 25 (1980). J.W. Hutchinson et al., Viewpoint Set #6, Scripta Metall., 18 (1984) 421-458. H.C. Rogers, in "Deformation, Processing and Structure," ASM (1984) 425. N. Thompson, N. Wadsworth and N. Louat, Phil Mag., 1 (1956) 113. F. Adcock, J. Inst. Met., 27 (1922) 73. M. Cook and T. L. I. Richards, J. Inst. Met., 78 (1951) 463. M. Blicharski and S. Gorczyca, Metal Science, 13 (1978) 303. B.J. Duggan, M. Hatherley, W. B. Hutchinson and P. T. Wakefield, Metal Science, 12 (1978) 343. M. Blicharski, S. Nourbakhsh and J. Nutting, Metal Science, 13 (1979) 516. K. Morii, Y. Nakayama and H. Mecking, in "Textures of Materials - ICOTOM 7," Ed. C. M. Brakman, P. Jongenberger and E. J. Mittemeijer, Netherlands Society for Materials Science (1984) 117. D.J. Lloyd, E. F. Butryn and M. Ryvola, Microstructural Science, 10 (1982) 373. D.J. Lloyd and M. Ryvola, Microstructural Science, 12 (1984) 577. A. Korbel and M. Richert, Acta Metall., 33 (1985) 1971. H. Chandra-Holm and J. D. Embury, in "Yield, Flow and Fracture of Polycrystals," Ed. T. N. Baker, Applied Science Publishers, U.K. (1983) 275. H. Rogers, Ann. Rev. Mater. Sci., 9 (1979) 283. M.R. Lin and R. H. Wagoner, Metall. Trans. A, 18A (1987) 1035. F.J. Humphreys, Acta MetaU., 25 (1977) 1323. J.R. Porter and F. J. Humphreys, Metal Science, 13 (1979) 83. F.J. Humphreys, Metal Science, 13 (1979) 136. H° M. Chan and F. J. Humphreys, Acta Metall., 32 (1984) 235. R. Sandstrom, in "Recrystallization and Grain Growth of Multi-Phase and Particle Containing Materials," Ed. N. Hansen et al., Riso National Lab., Denmark (1980) 45. E. Nes, in "Recrystallization and Grain Growth of Multi-Phase and Particle Containing Materials," Ed. N. Hansen et al., Riso National Lab., Denmark (1980) 85. E. Nes, in "Light Metals Congress," Vienna (1981) 154. E. Nes and J. A. Welt, Scripta Metall., 18 (1984) 1433. R.D. Doherty and J. W. Martin, J. Inst. Metals, 91 (1962/3) 33. P.R. Mould and P. Cotterill, J. Mater. Sci., 2 (1967) 241. R.W. Benedek, D. Phil Thesis, University of Sussex, U.K. (1974). L.R. Morris, H. Sang and D. M. Moore in "4th International Conf. on the Strength of Metals and Alloys," Nancy, France (1976) 131. J.A. Weft, N. E. Paton, C. H. Hamilton and M. W. Mahoney, Met. Trans., 12A (1981) 1267. E. Hombogen and K. Zum Gahr, Metallography, $ (1975) 181. J.W. Martin, Micromechanisms in Particle Hardened Alloys, C.U.P. Cambridge, U.K. (1980) 99-110. K. Campbell, 1. Dover, T. R. Ramacharandran and J. D. Embury, Metals Forum, 2 (1972) 229. J. Liu and R. D. Doherty, "Aluminium Technology '86," Institute of Metals, London, U.K. (1986) 347. J. Liu and R. D. Doherty, Submitted to Acta MetaU. (1989). J. Liu, Ph.D. Thesis, Drexel University, 1987. S.D. Smith, Alcoa Technical Center, Private Communication, described in Reference 38. J. Wood and P. G. McCormick, Acta Metall., 35 (1987) 247.

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TABLE 1

Microstructural Details of Alloys for Studying Shear Banding

Material( 1)

Precipitate Vol.%

a x b(2)

0.15A

0.0

0.31B

0.21

1.9 x 0.6 lam

0.54A

0.51

0.85A

Nearest Neighbor Spacing. A3

(No Precipitates)

Reduction

Offset

An~le(3)

40%* 50%* 60% 70% 80%

1.9 I~m 1.9 I~m 1.9 ~tm 2.2 !Ltm 1.9 IXm

30" 24" 28" 33" 30"

7.8 ~tm

50%* 60% 70% 80%

1.9 Ixm 2.5 lam 1.6 I.tm+ 2.5 I.tm

30" 32" 30* 35"

1.8 x 0.6 gm

4.6 ktm

60% 70% 80%

1.3 I.tm 1.6 ~tm 2.2 }am

31" 29" 31 °

0.92

0.4 x 0.1 pan

1.5 ktm

60%* 70% 80%#

1.6 ~tm 1.9 ktm 3.8 ~tm

38 ° 33 ° 29 °

1.85B

2.27

0.8 x 0.3 }.tm

1.5 Ism

70%* 80%*

1.0 ~m 1.6 ~tm

33" 39"

1.85D

2.27

1.7 x 0.7 ~m

3.6 llm

70% 80%

1.9 ~m 2.2 ktm

29" 34"

1.85E

2.27

2.3 x0.7 ktm

5.0 ktm

60%* 70% 80%

1.6 gm 1.3 ~tm 1.9 ~tm

35 ° 36° 38 °

Notes: (1) (2) (3) * + #

All A1-Mg2Si alloys equilibrated at 350"C; digits denote wt.% Mg2Si, letters represent dispersion types; details in reference (39). Precipitates approximated as oblate spheroids; a and b are major and minor semiaxes, respectively. Angle between shear and rolling directions. Only very few offsets seen. Larger offsets were visible, but could not be recorded. Accidentally given a larger than usual reduction.

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Figure 1. Alloy 0.31A (Vol.%ffi0.21, a x b=l.3 x 0.5 gin, A3=6.8 p m - Polarized light micrograph of section normal to Iransvvrse direclion.

Figure 2. Alloy 0.85B (Vol.%ffi0.92, a x b=l.2 x 0.4 ;tin, A3=2.7 gin) - Polarized light micrograph of section normal to lransvea~ direction.

Figure 3. Alloy 1.85C (Vol.%ffi2.27, a x b=l.3 x 2.6 gm, A3=2.9 pro) - Polarized light micrograph of section normal to wamvea'sedin~on.

Figure 4. Alloy 0.15A (single-phase) - Polarized light micmgraph of section normal to transvezse direction.

Figure 5. Alloy 1.85E (Vol.%ffi2.27, a x b=2.3 x 0.7 pro, A3=5.0 vm) - TEM micrograph of section normal to transverse direction showing a she,~ band. This band makes an angle of 27" with the rolling direction.

Figure 6. 1.85E (Vol.%ffi2.27, a x bffi2.3 x 0.7 pro, A3=5.0 lun) - Differential interference contrast light micrograph of section normal to transverse direction showing offsets after incremental roiling of sheet initially rolled with a 90% reduction.