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Texture and structure development in cross-rolled α brass
Acra meroll Vol Primed m Great
34. No. 4. pp. 653-660. 1986 Britam. All rights reserved
TEXTURE
Department
CopyrIght
<
AND STRUCTURE DEVELOPMENT CROSS-ROLLED ct BRASS
0001-6160 86 53.00 + 0.00 1986 Pergamon Press Ltd
IN
W. Y. YEUNG and B. J. DUGGAN of Mechanical Engineering. University of Hong Kong. Pokfulam Road. Hong Kong
Abstract-Cross rolling produces a complicated framework structure of lath-like bands after medium strains while deformation twins and shear bands rarely occur. {I 10}(223) texture begins to develop after 50% reduction and becomes well established after 80% rolling. The texture development is rather steady and does not involve the complex changes that occur in straight-rolled brass. Cracking occurs along grain boundaries after 60% reduction at - 196.C and at 70% reduction at room temperature. Changes in mechanical properties are found to correlate with microstructure evolved in cross-rolling. The steady development of { I10)(223) texture is associated with homogeneous deformation on normal slip systems. Dislocation activity and work hardening on eight slip systems necessary lo accomodate the imposed strain leads to loss of plasticity in grain boundary regions and intergranular cracking. R&um&Un laminage croisP produit une structure complexe de bandes semblables i des lattes apres des d&formations moyennes. alors que des macles mecaniques et des bandes de cisaillement ne se produisent que rarement. La texture {I 10}(223) commence B se divelopper apris une reduction de 50% et elle devient bien Ctablie apres un laminage de 80%. Le diveloppement de la texture est assez continu et il ne met pas en jeu les changements complexes qui se produixnt dans le laiton Iamine unidiuctionrellement. La fissuration intergranulaire se produit pour une riduction de 60% d - l96’C et pour une reduction de 70% B la tem$rature ambiante. Les changements de propri&Cs mbcaniques sont lies B la microstructure produit par le laminage croisi. Le dCveloppement continu de la texture (I lO}(223) est associe P une diformation homogene sur des systimes de glissement normaux. L’activite des dislocations et le durcissement sur les huit systemes de glissement nicessaires pour accommoder la d&formation imposee conduisent P une perte de plasticitC dans les rCgions intergranulaires et I une fissuration intergranulaire. Zusammenfassung-Kreuzweises Walzen fiihrt zu einer komplizierten Rahmenstruktur aus latteniihnlichen Btindern bei mittleren Dehnungen; Verformungszwillinge und Scherbander finden sich dagegen selten. Die {110}(223)-Textur entwickelt sich ab einer Reduktion von 50% und ist bei 80% deutlich Die Texturentwicklung verlluft ziemlich stetig und weist nicht die komplexen Verinderungen auf. die in einfach gewalztem Messing auftreten. RiDbildung findet sich nach einer Reduktion von 60% bei - 196-C und von 70% bei Raumtemperatur. Die Anderungen in den mechanischen Eigenschaften korrelieren gut mit der wiihrend des kreuzweisen Walzen entstehenden Mikrostruktur. Die stetige Entwicklung der {ll0)(223)-Textur ist von einer homogenen Verformung auf normalen Gleitsystemen begleitet. Die zur Anpassung an die aufgezwungenen Dehnungen notwendige Versetzungsaktivitlt und Verfestigung auf acht Glcitsystemen fiihrt dazu. daB die PlastizitIt in den Komgrenzbereichen verrnindert wird und intergranulare RiBbildung einsetzt
INTRODUCTION Several workers [l-3] have reported that the major texture component developed in cross-rolled f.c.c. materials is { 1lOj(223). In N&Fe the texture is made up of two symmetrical (lTO)[223] orientations plus a minor cube component [I] and in copper twin components of the major texture are present [2]. Later work showed that cube oriented material was also present in copper. In high-purity dilute aluminiummanganese alloys a very strong { 1lOj(223) texture was formed after cross rolling to 95% reduction in thickness [4]. Information concerning microstructures produced by cross-rolling is also limited and often difficult to interpret because transmission electron micrographs taken through the sheet plane are of limited value in rolled metals (51. Even so. Grtenba and Hu [6] showed that structures developed in a single crystal 653
of iron after unidirectional rolling to 80% reduction were destroyed by rolling a further 15% in the perpendicular direction. In their review of deformation structures Gil Sevillano et al. [7] suggested that changes in microstructure should be closely related to imposed macroscopic shape changes but that particular textures could modify local deformation modes. In the present work, microstructure and texture development are followed in cross-rolled z brass at both room temperature and - 196°C. A computer simulation of cross-rolled texture is presented together with an attempt to evaluate slip activity in relation to grain boundary cracking. EXPERIMENTAL Two batches of material were prepared from a hot rolled commercial
quality
70,‘30 brass slab of 9 mm
654
YEUNG and DL’GGAN:
STRUCTURES
IN CROSS-ROLLED
I BRASS
RESL’LTS
Fig. 1. Sheet geometry of the cross rolled strip
thickness. Batch 1 was cold rolled to 8 mm thick on a Robertson 2 high mill of 127 mm (5 in.) roll diameter and annealed in a salt bath at 850-C for 3 min to obtain material of mean grain size -450 pm. Batch 2 was prepared from the same 9 mm thick slab but cold rolled to 2 mm thick and annealed in salt bath at 850 C for 20 min to produce an average grain size of 500 pm. This thinner material was used for conventional tensile testing on the assumption that the somewhat stronger starting texture arising from heaGer cold rolling would have little effect on mechanical properties. The annealed plate was cut into square sheets and cross rolled. Rolling geometry is shown in Fig. I. and the deformation was done at room temperature (RT) and - 196 C with about 5% reduction per pass up to specimen cracking. For rolling at room temperature the rolled sheets were quenched into cold water after each pass and for low temperature deformation the rolled sheets were immersed in a bath of liquid nitrogen immediately after each run and kept in the bath until the temperature was restored. All roiled materials were stored at -30-C before testing. TEM foils were electropolished in methanol-nitric acid solution (67 ml methanol. 33 ml nitric acid. at 8V d.c. at -30-C) and optical specimens etched with aqueous ferric chloride (IO g FeCl. 20 ml HCI. 80 ml H:O). Scanning electron microscopy was used to study crack initiation and propagation and conventional III pole figures were measured by the Schulz reflection technique with filtered CuK, radiation. In order to obtain representative texture data 10% of total thickness was removed from each surface by chemically polishing with dilute nitric acid. Tensile properties were determined in both RDl and RD2 directions. Fig. I, using the specimen geometr!’ shown in Fig. 2.
AII Dimensions in mm Fig. 2. Specimen geometr! of tensile specimens
A rolling reduction increased from 10 to !OoO. strain markings increased on RDI and RD2 surfaces. At 400/o strain. grains uere distinguishable b! both angle and number of traces in each grain. Fig. 3(a). After 50% reduction the strain markings \\ere \er! intense. Fig. 3(b). but there were no obvious differences between room and loa temperature deformation. Shear bands in cross rolled brass n’ere rarel! observed and appeared after 70% reduction at room temperature and 60% at - 196 C. Fig. 3(c). Transmission electron micrographs taken from RDI and RD2 showed similar structures. After 50”0 rolling thin lath-like features crossing heali]! dislocated regions and shearing other laths \+ere frequently seen. Fig. 4(a). Clusters of laths separated b! heavily dislocated regions were also common]! observed. Fig. 4(b).
(b)
Fig. 3. Optical microwucwre of I brass cross rolled at room temperature. (a) 40?/0 (b) 50”0 (c) 70”“. x 60 Trace of rolling plane horizontal. RDI sections.
YEL’:SG and DUGG.Q:
STRUCTURES
fK CROSS-ROLLED
x BRASS
655
(4
08 Fig. 4. Electron micro~raphs of 2 brass cross rolled 50% at room temperature. fa) Thin lath-like structure crossing hee\il> dislocated regions and shearing other laths. RDI section. fb) Clusters of laths. RD2 section.
Figure 5 shows that after 60Sb rolling a complicated frametnorh structure is formed. Two sets of laminar structures intersect at 20 in this figure and in some cases a third set appears. Fig. Sfb). ln~reasin~ the rolling reduction to 70% produced no new features. Shear bands were rare and found in two kinds of locations Figure 6(a) shows a shear band cutting through a heavily twinned region which is identical to the n~i~rostru~ture found in unjd~rectio~a~ly rolled brass 181. The shear band makes 35” to the twin boundaries, Shear bands were also formed in grains showing framework structures. Fig, 6(b). In this case and in all others found. shear bands were parallel to one of the lath directions. but these lath directions, showed no common relationship to tolling geometry.
Cracks formed during cross rolling and these lead to the material breaking up at reductions ~80%. Fine cracks appeared at the edges of the sheet after 709~0deformation at room temperature and at 60%
m Fig. 5. (a) Electron mi~ro~raph af !xbrass cross rolled h0” II at room temperature showing microstructure of lath-li!-x bands, RD2 section. Ib) Electron m~~ro~raph after cross rolling of 60% at - 196 C showing a “framet\ork” structure. REX?section.
at - 196.C. Figure 7 shows a room temperature example and Fig. 8 is a scanning electron micrograph showing cracking between areas with strain markings of different orientations. Many examples were examined and it was concluded that cracking ws basieall! intergranular. There were no fundamental digerences between cracks formed at loti and room temperatures and no simple angular relationships between propa gation paths and the principal stress axes.
Textures developed in brass with increasing crossrolling strain at room temperature and - 196 C are shown in Figs 9 and IO. The texture evolution in these samples is simple and does not involve the complex changes reported in straight-rolled 2 brass [9]. In Fig. 9 the 111 pole figures show that a {I 10)(12?) texture begins to form after 50% rolling at room temperature. The strength of this component steadi& increases and a sharp texture becomes uell established at 80% rolling. A weak i I I 1.; sheet component is also observed to exist after 70% reduction. For brass cross-rolled at - I96’C. Fig. IO. the pattern of
656
1’EL’NCi and DUGGAN:
(a)
STRUCTURES
IN CROSS-ROLLED
3 BRASS
(b)
Fig. 6. Electron micropraphs of z brass cross rolled 70% at room temperature. (a) Shear band formed in a heavily twinned region. RD2 section. (b) Shear band developed in the “framework” structure. RDI section.
Fig. 7. Optical micropraph of sheet surface of z brass cross rolled 70% at room temperature showing a fine crack along a grain boundary.
Fig. 8. Scanning electron micrograph of sheet surface of a brass cross rolled 70°% at - 196 C intergranular cracking.
shoumg
YEUNG and DUGGAN:
STRUCTURES IN CROSS-ROLLED a BRASS
657
DlSCUSSION
Fig. 9. {ill} Pole figures of a brass after various cross rolling reductions at room temperature. (a) SO%, (b) 60%, (c) 70%, (d) 80%.
texture development is very similar to that observed at room temperature, i.e. {110}(223) starts to build up around 50% reduction and is well established at 70% rolling. It is apparent that lowering the temperature of deformation brings a small increase in texture
sharpness.
(d) Mechanical properties
Figure 11 shows how tensile strength and elongation-to-fracture varies with cross rolling strain at both room temperature and - 196°C. Tensile strength and elongations are similar in both RDl and RD2 directions and the material is stronger after deformations at - 196°C.
Fig. 10. { 1I l} Pole figures of a brass after various cross rolling reductions at - 196°C. (a) 50%, (b) 60%, (c) 70%.
The only feature found in electron micrographs which correlated with the strain markings shown in Fig. 3(a) and (b) are lath structures, Fig. 4(a) and (b). It is therefore most likely that lath structures, either in isolation or in clusters, are the deformation feature which etch to give strain markings. At deformations of less than 50% it was rare to find intersecting laths, but at rolling reductions of greater than 50% the electron microstructure was dominated by a framework of strongly diffracting deformation bands, Fig. 5, which were often not as well delineated as in those formed at lower strains, cf. Figs 4 and 5. Their diffraction contrast was always high and this plus their general appearance lead to the conclusion that they were sites of localised shear deformation containing a high density of dislocations. The angles of these bands varied from grain to grain but the set within a particular crystal seemed to be characteristic of that crystal, Figs 4 and 5. In the absence of goniometric facilities and limited by the uncertainties of 100 kV conventional SAD methods [lo] it was not possible to determine the precise nature of the deformation bands. It has been established that in general a laminar structure approximately parallel to the rolling plane is a necessary precursor to shear band formation [I 1, 121. This is readily formed in low stacking fault energy materials during unidirectional rolling by a deformation twinning process [9]. Grains producing this kind of twin/shear band morphology were extremely rare in cross rolled material but those found showed remarkably similar microstructures to that of straight rolled brass, Fig. 6(a). Figure 6(b) shows a shear band in all respects similar to those found in straight rolled brass but it is contained in a framework of deformation laths. It is thus clear that cross rolling constraints are capable of producing typical shear band structures in grains which are not copiously twinned.
Fig. 1 I. Variations of ultimate tensile strength and elongation to fracture along RDI and RD2 directions of a brass after cross rolling at room temperature and - 196°C.
658
YEUNG and DUGGAN:
STRUCTURES IN CROSS-ROLLED
Fig. 12. Texture of f.c.c. materials after various simulated cross rolling reductions, f I I I ) pole figures. (a) O%, (b) 50%, (c) 70%, (d) 80%.
The relationships between dev~Iopi~g structure and texture in unidirectionally rolled brass have been the subject of much recent research and are quite well understood. By 40% reduction a copper-type texture emerges in brass by normal processes of slip. At greater rolling strains large volumes of material deform by twinning and at 60% reduction a texture transition occurs with a build-up of (11 l)(uuw) components. This arises from coupled rotation of twins and matrix until twinning planes are parallel to the sheet surface, producing two twin related f I1 lf( Il2) components [9]. Grains in which the laminar twin structure are close- to this orientation form shear bands. Shear bands increase in near until { 11 I ] components are replaced by the texture of material in shear bands. Homogeneous slip in sheared material produces the chara~te~stic high strain brass texture (llOf(112). In contrast both structure and texture development in cross rolled brass are quite different to that occurring in unidirectional rolling. (llOf(223) is present after 50% cross rolling and steadily increases in intensity up to the maxims strain of 80%, Figs 9 and 10. There is no copper-type texture formed at low reductions and only a very weak (111) component appears at -800/c, a strain at which this com~nent begins to decrease in unidirectionally rolled brass [9]. The mechanism by which the {l I If texture is formed however, might be identical because of the great correspondence shown between Fig. 6(a) and typical straight-rolled structures [8]. Both single crystal and polycrystal studies have established that (I 10)<112) is truly stable up to quite high strains in unidirectional rolling. Rillamore and Roberts f13], assuming homogeneous slip, showed that rotations due to slip on the two slip systems with
a BRASS
highest resolved shear stresses would lead to the development of f 1IO){ 112). Changing the imposed strain to cross rolling favours production of (110){223), an orientation 8’ from fl IO){ 112). An attempt was made to simulate texture development in cross-rolling using a programme developed by DiIlamore [14]. In this a random selection of 100 grain orientations is made giving the weak starting texture shown in Fig. 12(a). Homogeneous slip based on {11l)( 110) systems is assumed in accordance with Taylor [15] and Bishop and Hill [16] and when ambiguity arises it is assumed that the operative slip systems are the set of f’ive with the largest individual shears. Using strain increments of 5% alternately for RDi and RD2 it is clear that orientations close to f 111f are signifi~ntly depleted by 50% reduction, Fig. 12(b). As strain increases to 70%, Fig. 12(c), there is clustering of {1lO}(uvu~) poles and by 80% a (110)(223) texture has formed, Fig. 12(d). In the computer generated pole figures the apparent reduction in the total number of points is due to overlap and so the most densely populated region are actually of higher intensity than is shown. The computer simulations when compared with the measured pole figures allows the conclusion that the development of (110)(223) is due primarily to homo~neous deformation based on normal slip activity. Stability of the texture (1310](223) was investigated by computing crystal rotation with respect to cross rolling strain. Results are plotted in terms of Euler angles [17] in Fig. 13, and these show that a 40% simulated rolling strain imparted in increments of 5% alternately in RDl and RD2 direction to a (I 10]<223) oriented crystal produces minor selfr~to~ng deviations from the original o~entation. Simulated straight rotling of this texture shows that it is not stable, Fig. 14.
ST
‘$6 Rnliing Fig. 13. Computer simulation of crystal rotation (in terms of Roe’s Euler angles) of material in (~~~)~~~3~orie~tatjon subjected to cross rolling.
YEUNG and DUGGAN:
STRUCTURES
IN CROSS-ROLLED
659
a BRASS
deform so as to keep continuity and intergranular cracking will ensue. It is worth noting that unidirectionally rolled brass which is capable of extremely large reductions without cracking does so because of the intervention of twinning followed by shear banding. Cross roiling, because of its different constraints, allows only normal slip processes and this leads directly through work hardening to reduced ductility.
CONCLUSIONS
40""' 0
20
10 %
”
30
”
40
Rolling
Fig. 14. Computer simulation of crystal rotation (in terms of Roe’s Euler angles) of material in (T10)[223] orientation after unidirectional rolling.
Cracks first appeared after 60% reduction for cross-rolling at - 196°C and at 70% reduction at room temperature. Figures 7 and 8 show that cracking is associated with localized deformation and develops along grain boundaries. According to Taylor theory [ 151 five independent slip systems are required to satisfy the continuity condition between grains. Honeff and Mecking [ 181 have recently modified this theory with the “relaxed constraints” model such that the operation of three slip systems can readily accomodate the constraints in the interior of rolled-out grains, but five slip systems are still required near grain boundaries. It is possible to compute the shear strain yn for crystals of orientation {110}(223) on each of the five operative slip systems for a 5% increment in grain-strain (n denotes the particular slip system concerned). The direction of rolling is then changed and the shear strains on the new set of five systems recalculated. Some of the slip systems are common to both deformation directions. The total shears on each system for imposed deformations of 30 and 40% are calculated from m = Zyn and using standard nomenclature as in [19] are plotted in Fig. 15. It is clear that eight systems are necessary to accomodate the imposed cross rolling strain for the orientation {1lOj(223), compared with five for straight rolling of this orientation. If work hardening is assumed to be roughly proportional to the total shear on each system then exhaustion of plasticity in the grain boundary region is possible because alternative systems have hardened by deformation in the other direction. This hardening of many slip systems will undoubtedly occur in crystals of other orientations and it is reasonable to suppose that at a critical strain the grains will not be able to
1. Cross-rolling produces a complicated framework of intersecting lath-like structures which accomplish shear deformations and contain a high density of dislocations. The framework structure is identical on both RDl and RD2 faces and mechanical properties measured in these direction are similar. 2. Development of the framework structure coincided with a steady increase in the texture {110}(223). There were no complex transitions in either structure or texture, which is very different to the situation in unidirectionally cold rolled brass. 3. The texture (110}(223) is stable up to 80% reduction. A computer simulation of cross rolling based on {11 I}( 110) slip produced a concentration of poles at {110}(223) and so it is concluded that this texture is formed by homogeneous deformation and normal slip. No account was made taken in the texture simulation of the framework structure and so it is thought that, although this is essential for continued plastic flow, the small volume fraction of
(al
Slip L
Systems
cb)
= 0.15
VI
0.10
d
3 0.05
s
0
d 82
Slip
2=
Systems
Fig. 15. Computer calculation of total shear r, on {II I}( 110) slip systems for a (T10)[223] crystal after (a) 30%, (b) 40% cross rolling reduction.
660
YEUNG and DUGGAN:
STRU~UR~S
2. 3. 4. 5.
a BRASS
6. T. Grzenba and H. Wu, 2. ~efuIik. 60, 944 (1969). 7. 3. Gil Sevillano, P. Van Houtte and E. Aernoudt, Frog. Mater. Sei. U, 69 (1982). , M. Hat~erIy, W. 8. Hut~~ins~n and P. 7. Wakefield, Metal Sci. 12, 343 (1978). 9. W. B. Hutchinsan, B. J. Duggan and M. Hatherly, Metals Technol. 6, 398 (1979). IO. B. J. Duggan and 1. P. Jones, Texture Crystall. Solids 2, 205 (1977). 1I. K. Morii and Y. Nakayama, Proc. &II Int. Co& on Texfures o~~#te~~~s, Vol. 1, p. 327. Iron & see1 Inst. Japan (1981). 12. B. J. Duggan and W. Y. Yeung, Pruc. 7th Int. Conf, on Texrutes of M@teriuls, p. 65. Netherla~s Sot. Mater. Sci. (t984). 13. I. L. Diiiamore and W. T. Roberts, &Q metail.12,281 0. W~~a~n, Texturen Metalii~her W~rkat~ffe, (1964). 14, I. L. Dillamore, private communication. Berlin, p. 96 (1939). IS. G, 1. Taylor, J. Infr. MS 62, 30’7(1938). R. W. Brick, D. L. Martin and R. P. Arq$ier, Trans. Am. 16. J. F. W. Bishop and R. Hill, Phil. Mulg.42,414 (1951). Sot. Met&Is3f, 675 (1943). 17. R. J. Roe, J. appl. Phys. 36, 2024 (1965). A. Merlini and P. A. Beck, Acfa metail. 1, 598 (1953). 18. H, Honeff and H. Mecking, Proc. 5th Int. ConA on R. Rixen, R. Musick, H. Gdker and K. Liicke, 2. Textures of Materials, Vol. 1, p. 265. Springer, Berlin Metaiik. 66, 16 (1975). H. Hu, Textures in research arid Practice, p, ZOO. (I978). 19. J. F. W. Bishop and R. Hill, Phil. Msg. 42, 1298 (1951). Springer, New York (1969).
material contained within these features does not cont~b~te sign~~~ntty to the final texture. 4. Shear banding as it occurs in unidir~tionally rolled material is rare in cross relied brass. 5. Intergranular cracking happens at cross rolling strains of 60% at - 196°C and 70% at room temperature. Cracking is att~buted to the exhaustion of plasticity due to slip activity and work hardening on the eight systems necessary to a~omp~ish the imposed shape change in grain boundary regions.
I.
IN CROSS-R~~~~D
34. No. 4. pp. 653-660. 1986 Britam. All rights reserved
TEXTURE
Department
CopyrIght
<
AND STRUCTURE DEVELOPMENT CROSS-ROLLED ct BRASS
0001-6160 86 53.00 + 0.00 1986 Pergamon Press Ltd
IN
W. Y. YEUNG and B. J. DUGGAN of Mechanical Engineering. University of Hong Kong. Pokfulam Road. Hong Kong
Abstract-Cross rolling produces a complicated framework structure of lath-like bands after medium strains while deformation twins and shear bands rarely occur. {I 10}(223) texture begins to develop after 50% reduction and becomes well established after 80% rolling. The texture development is rather steady and does not involve the complex changes that occur in straight-rolled brass. Cracking occurs along grain boundaries after 60% reduction at - 196.C and at 70% reduction at room temperature. Changes in mechanical properties are found to correlate with microstructure evolved in cross-rolling. The steady development of { I10)(223) texture is associated with homogeneous deformation on normal slip systems. Dislocation activity and work hardening on eight slip systems necessary lo accomodate the imposed strain leads to loss of plasticity in grain boundary regions and intergranular cracking. R&um&Un laminage croisP produit une structure complexe de bandes semblables i des lattes apres des d&formations moyennes. alors que des macles mecaniques et des bandes de cisaillement ne se produisent que rarement. La texture {I 10}(223) commence B se divelopper apris une reduction de 50% et elle devient bien Ctablie apres un laminage de 80%. Le diveloppement de la texture est assez continu et il ne met pas en jeu les changements complexes qui se produixnt dans le laiton Iamine unidiuctionrellement. La fissuration intergranulaire se produit pour une riduction de 60% d - l96’C et pour une reduction de 70% B la tem$rature ambiante. Les changements de propri&Cs mbcaniques sont lies B la microstructure produit par le laminage croisi. Le dCveloppement continu de la texture (I lO}(223) est associe P une diformation homogene sur des systimes de glissement normaux. L’activite des dislocations et le durcissement sur les huit systemes de glissement nicessaires pour accommoder la d&formation imposee conduisent P une perte de plasticitC dans les rCgions intergranulaires et I une fissuration intergranulaire. Zusammenfassung-Kreuzweises Walzen fiihrt zu einer komplizierten Rahmenstruktur aus latteniihnlichen Btindern bei mittleren Dehnungen; Verformungszwillinge und Scherbander finden sich dagegen selten. Die {110}(223)-Textur entwickelt sich ab einer Reduktion von 50% und ist bei 80% deutlich Die Texturentwicklung verlluft ziemlich stetig und weist nicht die komplexen Verinderungen auf. die in einfach gewalztem Messing auftreten. RiDbildung findet sich nach einer Reduktion von 60% bei - 196-C und von 70% bei Raumtemperatur. Die Anderungen in den mechanischen Eigenschaften korrelieren gut mit der wiihrend des kreuzweisen Walzen entstehenden Mikrostruktur. Die stetige Entwicklung der {ll0)(223)-Textur ist von einer homogenen Verformung auf normalen Gleitsystemen begleitet. Die zur Anpassung an die aufgezwungenen Dehnungen notwendige Versetzungsaktivitlt und Verfestigung auf acht Glcitsystemen fiihrt dazu. daB die PlastizitIt in den Komgrenzbereichen verrnindert wird und intergranulare RiBbildung einsetzt
INTRODUCTION Several workers [l-3] have reported that the major texture component developed in cross-rolled f.c.c. materials is { 1lOj(223). In N&Fe the texture is made up of two symmetrical (lTO)[223] orientations plus a minor cube component [I] and in copper twin components of the major texture are present [2]. Later work showed that cube oriented material was also present in copper. In high-purity dilute aluminiummanganese alloys a very strong { 1lOj(223) texture was formed after cross rolling to 95% reduction in thickness [4]. Information concerning microstructures produced by cross-rolling is also limited and often difficult to interpret because transmission electron micrographs taken through the sheet plane are of limited value in rolled metals (51. Even so. Grtenba and Hu [6] showed that structures developed in a single crystal 653
of iron after unidirectional rolling to 80% reduction were destroyed by rolling a further 15% in the perpendicular direction. In their review of deformation structures Gil Sevillano et al. [7] suggested that changes in microstructure should be closely related to imposed macroscopic shape changes but that particular textures could modify local deformation modes. In the present work, microstructure and texture development are followed in cross-rolled z brass at both room temperature and - 196°C. A computer simulation of cross-rolled texture is presented together with an attempt to evaluate slip activity in relation to grain boundary cracking. EXPERIMENTAL Two batches of material were prepared from a hot rolled commercial
quality
70,‘30 brass slab of 9 mm
654
YEUNG and DL’GGAN:
STRUCTURES
IN CROSS-ROLLED
I BRASS
RESL’LTS
Fig. 1. Sheet geometry of the cross rolled strip
thickness. Batch 1 was cold rolled to 8 mm thick on a Robertson 2 high mill of 127 mm (5 in.) roll diameter and annealed in a salt bath at 850-C for 3 min to obtain material of mean grain size -450 pm. Batch 2 was prepared from the same 9 mm thick slab but cold rolled to 2 mm thick and annealed in salt bath at 850 C for 20 min to produce an average grain size of 500 pm. This thinner material was used for conventional tensile testing on the assumption that the somewhat stronger starting texture arising from heaGer cold rolling would have little effect on mechanical properties. The annealed plate was cut into square sheets and cross rolled. Rolling geometry is shown in Fig. I. and the deformation was done at room temperature (RT) and - 196 C with about 5% reduction per pass up to specimen cracking. For rolling at room temperature the rolled sheets were quenched into cold water after each pass and for low temperature deformation the rolled sheets were immersed in a bath of liquid nitrogen immediately after each run and kept in the bath until the temperature was restored. All roiled materials were stored at -30-C before testing. TEM foils were electropolished in methanol-nitric acid solution (67 ml methanol. 33 ml nitric acid. at 8V d.c. at -30-C) and optical specimens etched with aqueous ferric chloride (IO g FeCl. 20 ml HCI. 80 ml H:O). Scanning electron microscopy was used to study crack initiation and propagation and conventional III pole figures were measured by the Schulz reflection technique with filtered CuK, radiation. In order to obtain representative texture data 10% of total thickness was removed from each surface by chemically polishing with dilute nitric acid. Tensile properties were determined in both RDl and RD2 directions. Fig. I, using the specimen geometr!’ shown in Fig. 2.
AII Dimensions in mm Fig. 2. Specimen geometr! of tensile specimens
A rolling reduction increased from 10 to !OoO. strain markings increased on RDI and RD2 surfaces. At 400/o strain. grains uere distinguishable b! both angle and number of traces in each grain. Fig. 3(a). After 50% reduction the strain markings \\ere \er! intense. Fig. 3(b). but there were no obvious differences between room and loa temperature deformation. Shear bands in cross rolled brass n’ere rarel! observed and appeared after 70% reduction at room temperature and 60% at - 196 C. Fig. 3(c). Transmission electron micrographs taken from RDI and RD2 showed similar structures. After 50”0 rolling thin lath-like features crossing heali]! dislocated regions and shearing other laths \+ere frequently seen. Fig. 4(a). Clusters of laths separated b! heavily dislocated regions were also common]! observed. Fig. 4(b).
(b)
Fig. 3. Optical microwucwre of I brass cross rolled at room temperature. (a) 40?/0 (b) 50”0 (c) 70”“. x 60 Trace of rolling plane horizontal. RDI sections.
YEL’:SG and DUGG.Q:
STRUCTURES
fK CROSS-ROLLED
x BRASS
655
(4
08 Fig. 4. Electron micro~raphs of 2 brass cross rolled 50% at room temperature. fa) Thin lath-like structure crossing hee\il> dislocated regions and shearing other laths. RDI section. fb) Clusters of laths. RD2 section.
Figure 5 shows that after 60Sb rolling a complicated frametnorh structure is formed. Two sets of laminar structures intersect at 20 in this figure and in some cases a third set appears. Fig. Sfb). ln~reasin~ the rolling reduction to 70% produced no new features. Shear bands were rare and found in two kinds of locations Figure 6(a) shows a shear band cutting through a heavily twinned region which is identical to the n~i~rostru~ture found in unjd~rectio~a~ly rolled brass 181. The shear band makes 35” to the twin boundaries, Shear bands were also formed in grains showing framework structures. Fig, 6(b). In this case and in all others found. shear bands were parallel to one of the lath directions. but these lath directions, showed no common relationship to tolling geometry.
Cracks formed during cross rolling and these lead to the material breaking up at reductions ~80%. Fine cracks appeared at the edges of the sheet after 709~0deformation at room temperature and at 60%
m Fig. 5. (a) Electron mi~ro~raph af !xbrass cross rolled h0” II at room temperature showing microstructure of lath-li!-x bands, RD2 section. Ib) Electron m~~ro~raph after cross rolling of 60% at - 196 C showing a “framet\ork” structure. REX?section.
at - 196.C. Figure 7 shows a room temperature example and Fig. 8 is a scanning electron micrograph showing cracking between areas with strain markings of different orientations. Many examples were examined and it was concluded that cracking ws basieall! intergranular. There were no fundamental digerences between cracks formed at loti and room temperatures and no simple angular relationships between propa gation paths and the principal stress axes.
Textures developed in brass with increasing crossrolling strain at room temperature and - 196 C are shown in Figs 9 and IO. The texture evolution in these samples is simple and does not involve the complex changes reported in straight-rolled 2 brass [9]. In Fig. 9 the 111 pole figures show that a {I 10)(12?) texture begins to form after 50% rolling at room temperature. The strength of this component steadi& increases and a sharp texture becomes uell established at 80% rolling. A weak i I I 1.; sheet component is also observed to exist after 70% reduction. For brass cross-rolled at - I96’C. Fig. IO. the pattern of
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(b)
Fig. 6. Electron micropraphs of z brass cross rolled 70% at room temperature. (a) Shear band formed in a heavily twinned region. RD2 section. (b) Shear band developed in the “framework” structure. RDI section.
Fig. 7. Optical micropraph of sheet surface of z brass cross rolled 70% at room temperature showing a fine crack along a grain boundary.
Fig. 8. Scanning electron micrograph of sheet surface of a brass cross rolled 70°% at - 196 C intergranular cracking.
shoumg
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DlSCUSSION
Fig. 9. {ill} Pole figures of a brass after various cross rolling reductions at room temperature. (a) SO%, (b) 60%, (c) 70%, (d) 80%.
texture development is very similar to that observed at room temperature, i.e. {110}(223) starts to build up around 50% reduction and is well established at 70% rolling. It is apparent that lowering the temperature of deformation brings a small increase in texture
sharpness.
(d) Mechanical properties
Figure 11 shows how tensile strength and elongation-to-fracture varies with cross rolling strain at both room temperature and - 196°C. Tensile strength and elongations are similar in both RDl and RD2 directions and the material is stronger after deformations at - 196°C.
Fig. 10. { 1I l} Pole figures of a brass after various cross rolling reductions at - 196°C. (a) 50%, (b) 60%, (c) 70%.
The only feature found in electron micrographs which correlated with the strain markings shown in Fig. 3(a) and (b) are lath structures, Fig. 4(a) and (b). It is therefore most likely that lath structures, either in isolation or in clusters, are the deformation feature which etch to give strain markings. At deformations of less than 50% it was rare to find intersecting laths, but at rolling reductions of greater than 50% the electron microstructure was dominated by a framework of strongly diffracting deformation bands, Fig. 5, which were often not as well delineated as in those formed at lower strains, cf. Figs 4 and 5. Their diffraction contrast was always high and this plus their general appearance lead to the conclusion that they were sites of localised shear deformation containing a high density of dislocations. The angles of these bands varied from grain to grain but the set within a particular crystal seemed to be characteristic of that crystal, Figs 4 and 5. In the absence of goniometric facilities and limited by the uncertainties of 100 kV conventional SAD methods [lo] it was not possible to determine the precise nature of the deformation bands. It has been established that in general a laminar structure approximately parallel to the rolling plane is a necessary precursor to shear band formation [I 1, 121. This is readily formed in low stacking fault energy materials during unidirectional rolling by a deformation twinning process [9]. Grains producing this kind of twin/shear band morphology were extremely rare in cross rolled material but those found showed remarkably similar microstructures to that of straight rolled brass, Fig. 6(a). Figure 6(b) shows a shear band in all respects similar to those found in straight rolled brass but it is contained in a framework of deformation laths. It is thus clear that cross rolling constraints are capable of producing typical shear band structures in grains which are not copiously twinned.
Fig. 1 I. Variations of ultimate tensile strength and elongation to fracture along RDI and RD2 directions of a brass after cross rolling at room temperature and - 196°C.
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Fig. 12. Texture of f.c.c. materials after various simulated cross rolling reductions, f I I I ) pole figures. (a) O%, (b) 50%, (c) 70%, (d) 80%.
The relationships between dev~Iopi~g structure and texture in unidirectionally rolled brass have been the subject of much recent research and are quite well understood. By 40% reduction a copper-type texture emerges in brass by normal processes of slip. At greater rolling strains large volumes of material deform by twinning and at 60% reduction a texture transition occurs with a build-up of (11 l)(uuw) components. This arises from coupled rotation of twins and matrix until twinning planes are parallel to the sheet surface, producing two twin related f I1 lf( Il2) components [9]. Grains in which the laminar twin structure are close- to this orientation form shear bands. Shear bands increase in near until { 11 I ] components are replaced by the texture of material in shear bands. Homogeneous slip in sheared material produces the chara~te~stic high strain brass texture (llOf(112). In contrast both structure and texture development in cross rolled brass are quite different to that occurring in unidirectional rolling. (llOf(223) is present after 50% cross rolling and steadily increases in intensity up to the maxims strain of 80%, Figs 9 and 10. There is no copper-type texture formed at low reductions and only a very weak (111) component appears at -800/c, a strain at which this com~nent begins to decrease in unidirectionally rolled brass [9]. The mechanism by which the {l I If texture is formed however, might be identical because of the great correspondence shown between Fig. 6(a) and typical straight-rolled structures [8]. Both single crystal and polycrystal studies have established that (I 10)<112) is truly stable up to quite high strains in unidirectional rolling. Rillamore and Roberts f13], assuming homogeneous slip, showed that rotations due to slip on the two slip systems with
a BRASS
highest resolved shear stresses would lead to the development of f 1IO){ 112). Changing the imposed strain to cross rolling favours production of (110){223), an orientation 8’ from fl IO){ 112). An attempt was made to simulate texture development in cross-rolling using a programme developed by DiIlamore [14]. In this a random selection of 100 grain orientations is made giving the weak starting texture shown in Fig. 12(a). Homogeneous slip based on {11l)( 110) systems is assumed in accordance with Taylor [15] and Bishop and Hill [16] and when ambiguity arises it is assumed that the operative slip systems are the set of f’ive with the largest individual shears. Using strain increments of 5% alternately for RDi and RD2 it is clear that orientations close to f 111f are signifi~ntly depleted by 50% reduction, Fig. 12(b). As strain increases to 70%, Fig. 12(c), there is clustering of {1lO}(uvu~) poles and by 80% a (110)(223) texture has formed, Fig. 12(d). In the computer generated pole figures the apparent reduction in the total number of points is due to overlap and so the most densely populated region are actually of higher intensity than is shown. The computer simulations when compared with the measured pole figures allows the conclusion that the development of (110)(223) is due primarily to homo~neous deformation based on normal slip activity. Stability of the texture (1310](223) was investigated by computing crystal rotation with respect to cross rolling strain. Results are plotted in terms of Euler angles [17] in Fig. 13, and these show that a 40% simulated rolling strain imparted in increments of 5% alternately in RDl and RD2 direction to a (I 10]<223) oriented crystal produces minor selfr~to~ng deviations from the original o~entation. Simulated straight rotling of this texture shows that it is not stable, Fig. 14.
ST
‘$6 Rnliing Fig. 13. Computer simulation of crystal rotation (in terms of Roe’s Euler angles) of material in (~~~)~~~3~orie~tatjon subjected to cross rolling.
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deform so as to keep continuity and intergranular cracking will ensue. It is worth noting that unidirectionally rolled brass which is capable of extremely large reductions without cracking does so because of the intervention of twinning followed by shear banding. Cross roiling, because of its different constraints, allows only normal slip processes and this leads directly through work hardening to reduced ductility.
CONCLUSIONS
40""' 0
20
10 %
”
30
”
40
Rolling
Fig. 14. Computer simulation of crystal rotation (in terms of Roe’s Euler angles) of material in (T10)[223] orientation after unidirectional rolling.
Cracks first appeared after 60% reduction for cross-rolling at - 196°C and at 70% reduction at room temperature. Figures 7 and 8 show that cracking is associated with localized deformation and develops along grain boundaries. According to Taylor theory [ 151 five independent slip systems are required to satisfy the continuity condition between grains. Honeff and Mecking [ 181 have recently modified this theory with the “relaxed constraints” model such that the operation of three slip systems can readily accomodate the constraints in the interior of rolled-out grains, but five slip systems are still required near grain boundaries. It is possible to compute the shear strain yn for crystals of orientation {110}(223) on each of the five operative slip systems for a 5% increment in grain-strain (n denotes the particular slip system concerned). The direction of rolling is then changed and the shear strains on the new set of five systems recalculated. Some of the slip systems are common to both deformation directions. The total shears on each system for imposed deformations of 30 and 40% are calculated from m = Zyn and using standard nomenclature as in [19] are plotted in Fig. 15. It is clear that eight systems are necessary to accomodate the imposed cross rolling strain for the orientation {1lOj(223), compared with five for straight rolling of this orientation. If work hardening is assumed to be roughly proportional to the total shear on each system then exhaustion of plasticity in the grain boundary region is possible because alternative systems have hardened by deformation in the other direction. This hardening of many slip systems will undoubtedly occur in crystals of other orientations and it is reasonable to suppose that at a critical strain the grains will not be able to
1. Cross-rolling produces a complicated framework of intersecting lath-like structures which accomplish shear deformations and contain a high density of dislocations. The framework structure is identical on both RDl and RD2 faces and mechanical properties measured in these direction are similar. 2. Development of the framework structure coincided with a steady increase in the texture {110}(223). There were no complex transitions in either structure or texture, which is very different to the situation in unidirectionally cold rolled brass. 3. The texture (110}(223) is stable up to 80% reduction. A computer simulation of cross rolling based on {11 I}( 110) slip produced a concentration of poles at {110}(223) and so it is concluded that this texture is formed by homogeneous deformation and normal slip. No account was made taken in the texture simulation of the framework structure and so it is thought that, although this is essential for continued plastic flow, the small volume fraction of
(al
Slip L
Systems
cb)
= 0.15
VI
0.10
d
3 0.05
s
0
d 82
Slip
2=
Systems
Fig. 15. Computer calculation of total shear r, on {II I}( 110) slip systems for a (T10)[223] crystal after (a) 30%, (b) 40% cross rolling reduction.
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2. 3. 4. 5.
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6. T. Grzenba and H. Wu, 2. ~efuIik. 60, 944 (1969). 7. 3. Gil Sevillano, P. Van Houtte and E. Aernoudt, Frog. Mater. Sei. U, 69 (1982). , M. Hat~erIy, W. 8. Hut~~ins~n and P. 7. Wakefield, Metal Sci. 12, 343 (1978). 9. W. B. Hutchinsan, B. J. Duggan and M. Hatherly, Metals Technol. 6, 398 (1979). IO. B. J. Duggan and 1. P. Jones, Texture Crystall. Solids 2, 205 (1977). 1I. K. Morii and Y. Nakayama, Proc. &II Int. Co& on Texfures o~~#te~~~s, Vol. 1, p. 327. Iron & see1 Inst. Japan (1981). 12. B. J. Duggan and W. Y. Yeung, Pruc. 7th Int. Conf, on Texrutes of M@teriuls, p. 65. Netherla~s Sot. Mater. Sci. (t984). 13. I. L. Diiiamore and W. T. Roberts, &Q metail.12,281 0. W~~a~n, Texturen Metalii~her W~rkat~ffe, (1964). 14, I. L. Dillamore, private communication. Berlin, p. 96 (1939). IS. G, 1. Taylor, J. Infr. MS 62, 30’7(1938). R. W. Brick, D. L. Martin and R. P. Arq$ier, Trans. Am. 16. J. F. W. Bishop and R. Hill, Phil. Mulg.42,414 (1951). Sot. Met&Is3f, 675 (1943). 17. R. J. Roe, J. appl. Phys. 36, 2024 (1965). A. Merlini and P. A. Beck, Acfa metail. 1, 598 (1953). 18. H, Honeff and H. Mecking, Proc. 5th Int. ConA on R. Rixen, R. Musick, H. Gdker and K. Liicke, 2. Textures of Materials, Vol. 1, p. 265. Springer, Berlin Metaiik. 66, 16 (1975). H. Hu, Textures in research arid Practice, p, ZOO. (I978). 19. J. F. W. Bishop and R. Hill, Phil. Msg. 42, 1298 (1951). Springer, New York (1969).
material contained within these features does not cont~b~te sign~~~ntty to the final texture. 4. Shear banding as it occurs in unidir~tionally rolled material is rare in cross relied brass. 5. Intergranular cracking happens at cross rolling strains of 60% at - 196°C and 70% at room temperature. Cracking is att~buted to the exhaustion of plasticity due to slip activity and work hardening on the eight systems necessary to a~omp~ish the imposed shape change in grain boundary regions.
I.
IN CROSS-R~~~~D