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A microstructural study of the origins of γ recrystallization textures in 75% warm rolled IF steel
Acta Materialia 54 (2006) 4337–4350 www.actamat-journals.com
A microstructural study of the origins of c recrystallization textures in 75% warm rolled IF steel M.Z. Quadir *, B.J. Duggan Department of Mechanical Engineering, University of Hong Kong, Pokfulam, Hong Kong Received 16 November 2005; received in revised form 3 April 2006; accepted 15 May 2006 Available online 1 August 2006
Abstract IF steel was warm rolled at 700 C in a single pass. The resulting texture and microstructure were remarkably similar to those of the same steel after cold rolling. A detailed investigation of the microstructure by orientation imaging microscopy and scanning transmission electron microscopy showed microbands to have a mutual misorientation of less than 4 and shear bands to contain material misoriented from the parent matrix by less than 10. Recrystallization did not occur preferentially at high-angle grain boundaries nor in shear bands. Instead the recrystallization nuclei were confined in the original hot band grain envelopes in crystals belonging to the c fiber. These c deformed grains had systematically developed deformation bands which consisted of elements that had rotated by up to 30 about the Æ1 1 1æ parallel to the normal direction. This is essentially the same nucleation process as observed in cold rolled and annealed IF steel. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: IF steel; Warm rolling; Shear bands; Deformation bands; Recrystallization
1. Introduction Steel used for pressing motorcar body panels has changed over the past 20 years from the Al-killed variety to the interstitial-free (IF) kind. The reason is partly economic: the processing takes 50 h for Al-killed steels but only a few minutes for IF steels; and partly technical: the formability is not very sensitive to processing conditions for IF steel, but is very sensitive for Al killed steels [1]. The IF steels are so called because they contain strong carbide formers, such as titanium and niobium, which getter the nitrogen and carbon atoms from solid solution and instead incorporate them into complex precipitates. These IF steels can be processed in continuous rolling and annealing lines which has allowed the successful processing of IF steel to occur for a wide range of parameters, which leads to significant commercial advantage for IF steels over other lowcarbon steels.
*
Corresponding author. Tel.: +852 2857 8615; fax: +852 2858 5415. E-mail address: [email protected] (M.Z. Quadir).
The formation of textures in steel, because of its technological importance, has been continuously investigated over many decades. Early work established that drawability was directly linked to the texture and it was quickly shown that for good drawability the steel should have as high a proportion as possible of grains having a Æ1 1 1æ direction parallel to the sheet normal direction (ND) and as small a proportion as possible with Æ1 0 0æ parallel to the ND [2]. This particular type of texture responsible for good drawability develops during recrystallization annealing from the rolling texture which consists of deformed grains clustered into sets of orientations known as a(Æ1 1 0æ//RD) and c(Æ1 1 1æ//ND) fibers. These are routinely formed at 80–85% cold rolling reduction in industrial practice and can be most conveniently shown in the /2 = 45 section in Bunge’s version of Euler space [3] shown in Fig. 1a. The normal sheet-making route followed for drawable sheets goes through a hot rolling stage in which the sheet is deformed in the austenite state to give a random texture before coiling and cold rolling, but in prize-winning work Barnett and Jonas [4] showed that rolling in the ferrite
1359-6454/$30.00 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2006.05.026
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Fig. 1. /2 = 45 ODF sections of warm rolled (75% at 700 C) and annealed IF steels (intensity levels: 1, 2, 4, 6, 8, 11).
range at temperatures up to 700 C (i.e. warm rolling) also produced well-developed a and c fiber textures. Clearly, annealing of material with these fiber components should produce drawable steel as it does in the cold rolled case, and this proved to be true [5]. However, the scientifically exciting part of their work was that it considerably extended the range of variables over which deformed and recrystallized microstructures could be investigated in ferrite while keeping the textures unchanged. Critical metallurgical factors such as strain rate sensitivity m and dynamic strain aging were determined from room temperature to 700 C and these were related to the microstructures and texture sharpness formed in an ultralow-carbon steel (ULCS) and an IF steel. The work demonstrated that for the ULCS, the value of m changed from negative to positive as temperature increased beyond 320 C and the texture was much sharper at temperatures in excess of 500 C than that found in IF steel at the same temperatures. In contrast, the IF steel showed a slightly positive value of m over a wide temperature range and formed weaker textures, and this was linked to material instability in the form of shear bands, which are narrow, plate-like structures which cut the microstructure at ±35 to the roll-
ing direction (RD) when viewed in the longitudinal section and which carry a relatively large shear strain. These shear bands, in the same way as shown earlier in rolled low stacking fault energy (SFE) face-centered cubic (fcc) materials [6], were claimed to be responsible for generating a weaker texture than formed in ULCS. Furthermore since the IF steel gave a better texture on annealing, the conclusion was drawn, backed by some optical microscopy results, that shear bands in the IF steel were the sites of {1 1 1}Æh k læ nuclei. Now this is not an unreasonable conclusion because Duggan et al. [7] had earlier shown using transmission electron microscopy (TEM) that in Al-killed steel shear bands did indeed provide nuclei of the orientation {1 1 1}Æh k læ. However, in later work, Tse et al. [8] showed that cold rolling produced orientation splitting in a single grain to produce mutually misoriented large volumes of material separated by either a sharp boundary or a more gently curving transition from one orientation to the other. This was shown to occur systematically and they called the process deformation banding (DB). The deformation-induced grain boundary they associated with nucleation of recrystallization, and since DB occurred preferentially in crystals belonging to the c fiber, they proposed that {1 1 1}Æh k læ
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originated from such sites. A link between DB and recrystallization was not established at the time of the work of Barnett and Jonas, i.e. 1997, and so the possible role of DB in the recrystallization texture formation in IF steels was not explored. This is the starting point for the present investigation. It is clearly necessary to investigate whether DB is as important in recrystallization texture formation in warm rolled IF steel as it is in the cold rolled case [8]. The investigation involves orientation imaging microscopy (OIM), TEM and scanning TEM (STEM) for microstructural characterization and uses conventional X-ray diffraction (XRD) global orientation distribution functions (ODFs) to cast new light on the subject of the annealing of warm rolled IF steels. 2. Experimental The material used had the composition 0.0043 wt.% C stabilized by adding 0.057% Ti in the melt. The other impurities were 0.006% P, 0.10% Mn, 0.035% Al, <0.01% Si, 0.002% N, 0.0065% S and 0.0031% O. The IF steel slab was hot rolled to 70% reduction in three passes starting at 1100 C. The exit temperature of the last pass was above 950 C. This resulted in randomly oriented grains with 70 lm average size which is about twice the grain size in commercially produced steels. The sample was then reheated to 750 C and allowed to cool to 700 C before being rolled up to maximum 75% reduction in a single pass, the maximum capacity of the rolling mill used. This high reduction in a single pass was preferred, because if the rolling was done in two or three passes it is inevitable that the waiting period between exit and re-entry to the mill would allow recovery and probably recrystallization to occur. This would affect subsequent processing. The single pass produced a material without any sign of recrystallization, and the hardness of the warm rolled material was 172 VHN, which is to be compared with a value of 178 VHN after cold rolling. High-temperature lubricant was sprayed after each pass onto the rolls to keep the friction effect to a minimum. Recrystallization annealing was done in an aircirculation furnace preheated at 710 C which allowed some minor grain growth to occur at longer annealing times. The global textures were measured using the Schulz reflection technique and Co Ka radiation to produce 1 1 0, 2 0 0 and 2 1 1 pole figures from which ODFs were calculated using the BATE Textan-III software. For the SEM study, longitudinal section (LS) and rolling plane section (RP) were electropolished in an electrolyte having 9:1 volume ratio of acetic and perchloric acids at 30 V DC. Samples from the mid-thickness and away from the edges of the sample were taken for all experiments to avoid depthdependent shear and plate widening effects. TEM foils were produced from the LS by mechanical thinning followed by electrolytic thinning using the standard window method. The sample was observed using a JEOL-STEM FX2000 operated at 200 kV. Slightly convergent beam electron dif-
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fraction (CBED) was used to obtain diffraction patterns for orientation measurements. The foils used for TEM were further observed with SEM to provide links between the microstructures observed using these two techniques. 3. Results 3.1. Rolling and annealing texture Fig. 1a shows the /2 = 45 section of Bunge’s Euler space and the major orientation components along the a and c fibers. The global texture shown in Fig. 1b of 75% warm rolled IF steel consists of these two fibers. Annealing of this material to complete recrystallization produced the texture shown in Fig. 1c, with peak intensity close to {5 5 4}Æ2 2 5æ and the almost complete annihilation of the a fiber components, except at the orientation {1 1 1}Æ1 1 0æ which is common to both a and c fibers. 3.2. Deformed microstructure 3.2.1. Microbands and shear bands Fig. 2 shows a low-magnification channeling contrast (CC) image of the LS showing several grains. Grains C and D are composed of microbands cut through by narrow shear bands in a sense opposite to the habit plane of the microbands. Such grains usually belong to the c fiber. Grains A and B belong to the a fiber and show the characteristic contrast associated with larger subgrains with small mutual misorientations. Grain A has shear banding which is rare in a fiber grains, and grain B has no such features which is much more common. There is evidence, given the width of B, that this grain has undergone much higher deformation than the imposed strain would indicate. Again thinned grains such as B are common. Fig. 3 shows a higher magnification image from a different position in the same specimen and shows three grains, A, B and C. The shear bands in grains B and C are very narrow, 0.5 lm or less. Observation of shear bands in the rolling plane shows that these produce offsets in the transverse direction (TD), proving that there is a shear component parallel to the TD. Another point of interest in this figure is that the left-hand end of grain B has a-type contrast, while from the center to the right-hand end a typical c-type contrast is observed. It is possible that these are, in fact, two different grains, but extensive use of OIM shows that this is unlikely in the majority of cases. Considerable evidence obtained from lower rolling reductions shows how features such as this arise. Presumably, due to neighborhood constraints or energy saving, one part of a grain rotates towards one orientation and forms structures differently to other parts of the same grain which are rotating to another orientation. This is evident in Fig. 4, which shows a single grain after 25% single-pass warm rolling reduction at 700 C. Clearly one part has a different microstructure from the other, and this is undoubtedly the origin of what is visible in Fig. 3. The boundary between grains B and C in
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Fig. 2. SEM CC image of LS of 75% warm rolled IF steel showing several deformed grains with smooth and in-grain shear banded substructures cutting through the microbands.
Fig. 3. SEM CC image of LS of 75% warm rolled IF steel showing the in-grain shear bands and their interactions with the grain boundaries.
Fig. 4. SEM CC image of LS of 25% warm rolled IF steel showing different substructures in different regions of the same deformed grain. This grain will probably consist of at least two different orientations as the rolling reduction increases. This orientation splitting is what gives rise to deformation banding.
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Fig. 3 is not flat, but wavelike; again this is common and has been observed by Tse [9] after cold rolling IF steel to 85% reduction. It is also frequently observed in cold rolled copper and a-brass [10] as well as in Al alloys [11] and is therefore probably a general phenomenon. Fig. 5a shows a TEM image of microbands observed in LS. The CBED patterns from five microbands are shown in Figs. 5c–g and the measured orientations are plotted on a 2 0 0 pole figure in Fig. 5h. The orientation of microband 1 is (2.25, 2.81, 1.62)[2,1,1] and the lattices in the contiguous microbands are rotated by a maximum of 6 3.65. The traces of the microbands are along h3:91; 0:26; 0:73i ffi h1 0 0i. In turn, this lies on the (0 1 1) plane, which is a body-centered cubic (bcc) slip plane. This is consistent with the idea that the microbands are formed by slip processes, a conclusion reached by Chen and Duggan in more extensive work [12] and more recently by Halder et al. [13]. To make the TEM results more statistically convincing the foil was removed from the TEM instrument and examined using
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SEM (Fig. 5b). It is clear that the microbands shown in Fig. 5a are part of a large set contained within the same grain and hence it is safe to assume that the TEM results have some general validity, i.e. microbands extend across the whole grain, a distance of 10–20 lm, and have small mutual misorientations. Shear bands similar to those shown in the CC image in Fig. 2 were then located and examined in a TEM thin foil. Fig. 6a shows a typical example of two sets of microbands making ±30 to the RD. Fig. 6b shows higher magnification detail from Fig. 6a showing a shear band. In Fig. 6b microband M1 is sheared to location M2, and this is proved by the fact that the orientations of M1 and M2 are identical (2.58, 2.15, 2.15)[3.02, 1.44, 2.18] (Figs. 6c and e). S is rotated 7.46 away from M1 and M2 to (2.3, 2.38, 2.24)[3.23, 1.23, 2.00] orientation (Fig. 6d). This is a small rotation. Another example is shown in Fig. 7 which has a pair of shear bands cutting through a grain containing a single set of parallel microbands.
Fig. 5. (a) TEM image from the LS showing microbands. (b) SEM micrograph of the foil area shown in (a). (c–h) CBED patterns and the corresponding 2 0 0 pole figure showing the orientations of microbands 1–5 labeled in the TEM image.
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Fig. 5 (continued)
Sheared elements 1, 2, 3 have the orientation (2.14, 2.92, 1.68)[3.15, 2.45, 0.23], and microbands 4, 5, 6 (2.06, 2.82, 1.92)[3.21, 2.37, 0.033] and 7, 8, 9 (2.02, 2.76, 2.06)[3.25, 2.31, 0.08], i.e. the shear band material 1, 2, 3 is rotated 5 from 4, 5, 6 and 8 from
7, 8, 9. This small rotation was confirmed as the most common case from the larger area scans available using OIM. Fig. 8 shows a CC and OIM image with the associated pole figures for the whole area, together with the line scan through fine parallel shear bands. The average misorienta-
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Fig. 6. TEM image of LS of 75% warm rolled IF steel showing the shear band cutting through one set of microbands while being parallel to another set of microbands and the corresponding CBED patterns of sheared and unsheared material.
tion is 8 and is not simply related to any of the sample reference axes RD, TD, ND. 3.2.2. Deformation bands Fig. 9 shows a CC micrograph taken from the RP section at the mid-plane and the grain boundaries are outlined black after observing at higher magnification. Grain X has the characteristic contrast of grains belonging to the a
fiber, which is verified by the pole figure shown in Fig. 9b. It maintained a fairly constant orientation along the line across the width of the grain. Grains Y and Z show the more complex microstructures which are very widely found in rolled IF steel. The line scans show, in each case, the grain orientation {1 1 1}Æu v wæ to be rotated about Æ1 1 1æ by 20–30. In Y it is from Y1 {1 1 1}Æ1 1 2æ to Y2 {1 1 1}Æ1 1 0æ (Fig. 9c) and in Z from Z1
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Fig. 7. TEM image of LS of 75% warm rolled IF steel with the corresponding CBED patterns showing the magnitude of crystallographic rotation produced by shear band material 1, 2, 3.
{1 1 1}Æ1 2 3æ to Z2 {1 1 1}Æ1 1 2æ (Fig. 9d). This kind of orientation rotation across a single grain, called by the present authors deformation bands (DBs) and by others [14] deformation-induced high-angle boundaries (DIHABs), are commonly found in both warm and cold rolled IF steel [8,15] and a model has been proposed for this process [8]. Grain Y has a continuous orientation variation from one side to the other, while Z has a relatively sharp boundary. Such lattice curvature, i.e. either sharp or more gentle across a grain, has also been observed in cold rolled IF steel [8], and in fcc metals and alloys [16]. 3.3. Recrystallization studies The material was recrystallized to the 8% level at 700 C and very large montages were prepared from both the LS and RP sections. Fig. 10 shows a CC image of a represen-
tative area from the LS showing the variety of annealed states in a stack of warm deformed hot band grains observed in the partially recrystallized condition. Grain A contains recrystallized grains of 3–9 lm in diameter formed in a finely divided microstructure which is characteristic of c fiber material. The growth of these newly formed recrystallized grains is observed to be within the deformed hot band grain, i.e. they have not, at this stage, grown through a preexisting hot band grain boundary. Grain B contains obvious shear bands, indicated by arrows and no obvious recrystallization is associated with them. Instead, only minor spheroidization of subgrains has occurred along the shear bands which is also true of grain E. Inspection of Figs. 3, 5, 6 and 7 shows deformed c material to consist of elongated microbands of 0.2–0.5 lm in width, with occasional ‘‘block type’’ cells of 1 · 1 lm. Annealing produces significant changes in the elongated
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Fig. 8. SEM CC image with OIM and 2 0 0 pole figures of LS of 75% warm rolled IF steel showing the magnitude of misorientation created at the shear band and matrix interfaces.
structure, i.e. the microbands are replaced by subgrains ranging in size from 0.5 to 1.5 lm in diameter (Fig. 10). It is arbitrarily assumed that subgrains of 3 lm diameter or larger are either successful or potential nuclei. From the montage of several hundred square micrometers, of which Fig. 10 is a small part, a total of 428 such nuclei/ grains were found, their orientations being determined
and plotted on a 2 0 0 pole figure in Fig. 11. Some 84% of these nuclei/grains had orientations belonging to the c fiber, with the highest concentration near {1 1 1}Æ1 1 2æ and {1 1 1}Æ1 2 3æ. These orientations are clearly part of the spread of the fully recrystallized texture shown in Fig. 1c. Interestingly, at this stage of recrystallization, i.e. 8%, grains such as B with clear shear bands, provided
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Fig. 9. SEM CC image with corresponding 2 0 0 pole figures of RP section of 75% warm rolled IF steel showing orientation gradients across the grains along line scans.
almost nothing to these data. Clearly, there is recovery in grains such as B, but no recrystallization nuclei/grains were found at this stage using the 3 lm criterion. The LS montages gave very little systematic information regarding nucleation of recrystallization because the nucleation events were difficult to correlate with the structures in which they had occurred. This was not the case when the material was examined in the RP. Figs. 12a and b show a CC image and the corresponding OIM image of a deformed c grain neighbored by two a grains, captured in the RP section of the same 8% partially recrystallized material. The c grain at the middle of the micrograph shows well-developed recrystallization nuclei A, B and C growing into the surrounding material. The overall texture is shown in Fig. 12c in the form of a 2 0 0 pole figure. The orientations of the deformed material 1–4 and the nuclei A–C are labeled in the pole figure corresponding to the labeling in the CC and OIM
images. Grain 1 is part of the a fiber which has a spread from near {4 1 1}Æ1 1 0æ to {1 1 2}Æ1 1 0æ and grain 4 is also a, but part of a complementary set of orientations spreading from {1 0 0}Æ1 1 0æ to {1 1 3}Æ1 1 0æ. The c microstructures 2 and 3 have two different orientations: 2 is close to {1 1 1}Æ1 1 0æ and 3 is close to {1 1 1}Æ1 2 3æ. From the OIM image it is very clear that the nucleation events occurred in the deformation gradient formed in the c-oriented material. The orientation of the recrystallized grains A, B, C belongs to the orientation spread of region 3, i.e. {1 1 1}Æ1 2 3æ. This kind of nucleation in a c grain showing a severe orientation gradient was systematically observed in many cases and is identical to the situation in cold rolled IF steel reported by Tse et al. [8] and described there as nucleation in deformation banded grains of the c family. During successive recrystallization stages recrystallized grains grew within c deformed parent grains and completed the recrystallization process
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Fig. 10. SEM CC image of LS at the 8% partially recrystallized condition of 75% warm rolled IF steel.
Fig. 11. Orientations in 2 0 0 pole figures of partially recrystallized grains at the condition shown in Fig 10 (1: {1 1 1}Æ1 2 3æ; 2: {1 1 1}Æ1 1 2æ). A total of 428 new grains of minimum diameter 3 lm are included in this pole figure.
by growing out of the hot band grain boundaries and consuming the a deformed grains. 4. Discussion 4.1. Warm rolled microstructure and texture The texture of 75% single-pass warm rolled material contains a and c fibers shown in the /2 = 45 ODF section in Fig. 1b. This is almost identical to the cold rolled tex-
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tures measured after the same magnitude of reduction as reported elsewhere [4,8]. Since the same textures are formed in both cold rolling and warm rolling, and the strain rates are 5 and 25 s1 [17], respectively, i.e. they differ by a factor of 5, it is plain that rolling texture formation is insensitive to both of these variables. The general microstructure formed in warm rolling is very similar to that formed in cold rolled IF steel [9]. There is a correlation between orientation and the deformed microstructures, i.e. the a fiber grains have characteristic smeared contrast compared to the sharp and finely fragmented contrast of c grains (Figs. 2, 3, 8 and 9). The observation that a grains can be thinner than their starting grain sizes was also reported in cold rolled IF steel. Chen et al. [18] explained this phenomenon by reference to the number of highly stressed slip systems, which is as many as 7 in a orientations. From this they predicted a relatively homogeneous flow to occur in a grains during deformation. This could also be the reason for the rather featureless appearance of material belonging to the a fiber. This featureless appearance in SEM is actually a large cell structure when examined using TEM in both warm [13] and cold rolled material [18]. Sometimes the large cell structure is revealed by a change in contrast in the a grain envelope, as in Figs. 2, 3 and 9. Microbands are the characteristic feature of c grains (and a grains, if they are close to {1 1 0}Æ1 1 2æ [19]) and are associated with less than five slip systems, the number known to be essential for homogeneous deformation. They are volumes of material, relatively free of dislocations enclosed by pairs of dislocation walls between 0.2 and 0.5 lm apart, which appear as parallel lines inclined 25–35 with RD in the LS (Figs. 3, 5 and 6). There are two basic models in the literature for microband formation, that due to Jackson [20] and Chen et al. [18], and based on this work it is not possible to choose between them. It is clear that the trace of the dislocation walls shown in Fig. 5a is along [1 0 0], which is a direction in a {1 1 0} plane and is therefore most probably the operating slip plane. Data such as are shown in Fig. 5 have been obtained in cold rolled IF steel and rigorously analyzed by Chen et al. [18] and Halder et al. [13] who concluded that microbands are a crystallographic phenomenon. This view has recently been questioned by Hurley et al. [21], in careful OIM work on Al–0.13% Mg using two contiguous surfaces, who concluded that microbanding is not a crystallographic phenomenon. Be that as it may, the important role of the microbands is that they constitute obstacles against which dislocations pile up prior to the strain burst which forms the shear bands [14]. The study of shear bands is emphasized in this investigation since they are obviously an important microstructural element and play a vital role in the deformation process, in both cold and warm rolled IF steels [4]. In this investigation shear bands are 0.2–0.5 lm thick, are generally confined within a single grain boundary envelope and carry a large amount of shear strain (i.e. a natural shear strain
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Fig. 12. SEM CC image with corresponding OIM and 2 0 0 pole figure showing the formation of recrystallized grains in a deformation banded grain.
4.6 in Fig. 6) and appear to be crystallographic because they are often formed parallel to a second set of microbands where these occur (Fig. 6). The large strain they carry is evident from the significant offsets created by shear bands in the preexisting microbands and at grain boundaries where they terminate (Figs. 2, 3, 6, 7 and 8). Rare examples of shear bands continuing across grain boundaries are between grains C/D in Fig. 2 and B/C in Fig. 3. Shear bands normally cut one set of microbands in the opposite sense to the RD (25–35) and are usually absent
in the a oriented microstructure. There are significant similarities between microbands and shear bands at the rolling reduction used in the present investigation, i.e. both 1. appear to be crystallographic and are confined within a single grain envelope (Fig. 8a), 2. have a grain orientation dependency, i.e. they are found in c grains and almost absent in a grains, 3. are similar in width to microbands, 4. appear to have a sheet-like structure, and
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5. are parallel to each other and are parallel to a second set of microbands when they are present (Fig. 6a). Therefore presumably the principal difference between them is the stage at which they were formed, i.e. microbands were formed in the early stages of rolling and shear bands were formed later when a c oriented grain is significantly work hardened by a series of parallel sets of dislocation walls which comprise the microbands. The orientation relationship between the shear band material and the matrix in which they have been formed is particularly important in regard to recrystallization. Generally shear band materials were found to be 8–10 misoriented with the matrix material as measured by EBSP of SEM and CBED of TEM (Figs. 6–8). In the TEM micrograph shown in Fig. 6b, the microbands have rotated 35–45 into the shear bands and are therefore expected to have a crystallographic rotation of the same magnitude if a rigid body rotation alone is involved. However the measured 8–10 misorientation implies that there is significant amount of slip involved in the deformation process and a model for how this is achieved has been proposed [22]. This small misorientation is effectively a low-angle boundary between the shear band material and the parent matrix. 4.2. Recrystallization behavior Two factors are essential for successful nucleation: formation of a mobile interface, which implies a critical misorientation with the matrix being consumed; and a driving force, which is usually related to the difference in dislocation density between the nucleus and its surroundings. Boundaries lose their dislocation character when the misorientation between two lattices reaches 15–20, and this is also the misorientation required for mobile interfaces. The shear bands have misorientations of about 50% of these values and therefore are not expected to be mobile. Of course these intermediate values of misorientation can be increased by sub-boundary climb during recovery, but this has to involve migration of the interface as well. Representative grains containing shear bands, such as is shown in Fig. 10, show evidence of recovery, but there are no new grains, unlike the situation in the grain above the shear banded grain. The fact that setting a lower limit of 3 lm to capture nucleation events meant that shear banded grains made almost no contribution to the data plotted in Fig. 11, indicates that they are not effective nucleation sites. The reason why Barnett and Jonas [4] concluded that shear bands gave rise to their final texture is that they used etched surfaces and optical microscopy for their metallographic investigation. It seems likely that the etch marks which covered c grains are traces of microbands and shear bands, because the intensity of the marks is far higher than would be expected for shear banded grains [4]. The optical technique is also of lower resolution than the electron beam
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technique used here, where very fine details are revealed. The conclusion is that shear bands in warm rolled IF steel do not significantly contribute to the formation of the recrystallization texture. Careful observation of large CC montages covering many hundreds of square micrometers at the 8% recrystallized stage showed almost no evidence at all of straininduced boundary migration (SIBM). The most likely candidate orientation for this process is grains belonging to the a orientation, for these are naturally low-energy blocks. They do not provide nucleation sites at this stage, but only recovered microstructures (Fig. 12). Since the final texture (Fig. 1c) shows only very low intensity along the a fiber, it is plain that a components are mostly consumed by other orientations. Another factor that must be borne in mind when considering the SIBM mechanism is the fact that in IF steels there is good evidence that the hot band boundaries are preferred sites for the nucleation of FeTiP and MnS particles [23,24]. These, plus any solute segregation provides pinning or drag on such a boundary, and SIBM has to overcome this problem. In both this warm rolled and also in cold rolled IF steel [9,15] there is abundant evidence that nucleation occurs inside the original hot band grain enveloped first, and growth out of these does not occur until later, typically at the 50% recrystallized volume fraction stage. The two factors essential for successful nucleation, formation of a mobile interface and a driving force between the nucleus and its surroundings, are obviously met in the deformation banded grains. The orientation gradients measured across the rolling plane surface of c grains which have DBs provide the necessary lattice curvature as is clearly shown in Fig. 9. This is a result typical for both warm rolling and cold rolling [8,15]. Further examples are shown in recovered c grain material in Fig. 12 where the rotation is from {1 1 1}Æ1 2 3æ to {1 1 1}Æ1 1 0æ. Thus the lattice curvature condition is met by the deformation banding process, which is systematic and is always found in the RP section in c grains. A simple model for deformation banding in bcc metals is given by Liu and Duggan [25]. Information regarding the driving force, however, is lacking in this investigation. However, early work on drawable steels using XRD [26] and in later work in IF steel using neutron diffraction [27] showed that recovery processes in both steels changed the as-deformed stored energy Def: Recov: relationships: EDef: f1 1 1gh1 1 0i hE f1 1 1gh1 1 2i changed to E f1 1 1gh1 1 0i i Def: Def: Recov: Recov: ERecov: f111gh112i and E f1 1 1gh1 1 0i hEf11 1gh1 23i to E f1 11gh1 1 0i iEf1 1 1gh12 3i .
If these inequalities hold true at the single-grain level, there is therefore an internal driving force inside a single recovered c grain which has deformation banded. It is difficult in principle to know where a nucleation event has taken place inside a three-dimensional object using only two-dimensional sections, because what is observed is a cross section of a developing crystal. Such a crystal can grow in any direction until it is stopped, and there is good reason to suppose that these new grains are
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stopped by grain boundaries. This fact has led to the notion that the regions adjacent to grain boundaries are actually the nucleation sites. This may well be so, but the evidence for this must always be ambiguous when twodimensional sections are being examined. For example, the new grains in Fig. 12 could be nucleated away from a grain boundary or at a grain boundary above or below the section being examined. Likewise the new grains visible in Fig. 10 could be nucleated adjacent to the grain boundary, or this could be a section along a line near the center of the deformed grain in which the new grain grew upwards or downwards to impinge the grain boundary. Therefore it is concluded that DBs provide the necessary conditions for nucleation of c oriented grains in c oriented material but it is impossible to rule out ‘‘near boundary’’ nucleation in the manner described by Senuma et al. [28] and Inagaki [29]. At a later stage of recrystallization, new grains grow out of the hot band envelopes to give the equiaxed structure and consume the remaining a deformed hot band. A mechanism is outlined by Duggan and Tse for cold rolled IF steel [30]. This happens when the new grains attain a critical length L, given by the SIBM formula [31] L P 2c/DE, where c is the surface energy of the original grain boundary and DE is the difference in energy across the hot band grain boundary. Of course, this is the simplest formulation, a more realistic formula involving both solute drag and pinning force has been given by Tse and Duggan [9] but this is not discussed further here, since this work adds nothing to the analysis. However, these terms have the effect of increasing L to the order of 10 lm or more, and hence the growing out of the nucleated grains occurs relatively late, on average, in the recrystallization process. 5. Conclusions 1. The rolling and annealing textures of warm rolled IF steel are very similar in terms of orientation and sharpness to those of cold rolled IF steel at the same rolling reduction. 2. Single sets of shear bands are commonly found but are not effective in recrystallization because the shear band material is misoriented 8–10 from its parent matrix, hence forming a low-angle interface with the matrix. 3. The in-grain shear bands, which are similar to the microbands in terms of geometry, alignment with sample axes and thickness, are predominantly formed in deformed c grains. 4. DBs produce high lattice curvature which can be either sharp or cumulative over several subgrain diameters and are best observed in the RP section. The banded seg-
ments are of c orientation having a general rotation relationship around Æ1 1 1æ//ND and are effective sites for nucleation of c recrystallized grains.
Acknowledgements It is a pleasure to acknowledge the support of this work by Grants Nos. CERG/HKU 7067/97E, 7323/98E and 7316/99E given by the Hong Kong Special Administration Region of China. References [1] Hutchinson WB. Mater Sci Forum 1994;157–162:1917. [2] Held JF. In: Proceedings of the conference on mechanical working and steel processing IV. New York (NY): Metallurgical Society of AIME; 1965. [3] Bunge HJ. Mathematische Methoden der Texturanalyse. Berlin: Akademie-Verlag; 1969. [4] Barnett MR, Jonas JJ. ISIJ Int 1997;37:697. [5] Barnett MR. ISIJ Int 1998;38:78. [6] Morii K, Mecking H, Nakyama Y. Acta Metall 1985;33:379. [7] Duggan BJ, Liu GL, Ning H, Zhang LX. Thermomech Process Theory Model Practice 1997:384. [8] Tse YY, Liu GL, Duggan BJ. Scr Mater 2000;42:25. [9] Tse YY, Duggan BJ. Mater Metall Trans A 2006;37A:1055. [10] Sindel M, Ko¨hlhoff GD, Lu¨cke K, Duggan BJ. Texture Microstruct 1990;12:37. [11] Hansen N. Metall Trans A 1985;16A:2167. [12] Chen QZ, Duggan BJ. Metall Mater Trans A 2004;35:3423. [13] Halder A, Huang X, Leffers T, Hansen N, Ray RK. Acta Mater 2004;52:5405. [14] Hughes DA, Hansen N. Metall Trans A 1993;24:2021. [15] Quadir MZ. Ph.D. thesis. University of Hong Kong; 2003. [16] Lee CS, Duggan BJ. Acta Metall 1993;41:2691. [17] Lam KT, Quadir MZ, Duggan BJ. Key Eng Mater 2003;233–236:437. [18] Chen QZ, Ngan AHW, Duggan BJ. Proc Roy Soc Lond A 2003;459:1661. [19] Quadir MZ, Duggan BJ. Acta Mater 2004;52:4011. [20] Jackson PA. Scr Mater 1983;17:199. [21] Hurley PJ, Bate PS, Humphreys FJ. Acta Mater 2003;51:4737. [22] Chen QZ, Quadir MZ, Duggan BJ. Philos Mag 2006;86:3633. [23] Deardo AJ. In: Takechi H, editor. Proceeding of the international forum properties and applications of IF steel. Tokyo: Arcadia Ichigaya; 2003. p. 240. [24] De Cooman BC, de Vyt A. In: Takechi H, editor. Proceeding of the international forum properties and applications of IF steel. Tokyo: Arcadia Ichigaya; 2003. p. 249. [25] Liu G, Duggan BJ. Mater Metall Trans A 2001;32:125. [26] Takechi H, Katoh H, Nagishima S. Trans AIME 1968;242:56. [27] Rajmohan N, Hayakawa Y, Szpunar JA, Root JH. Acta Mater 1997;45:2485. [28] Senuma T, Yada H, Shimizu R, Harase J. Acta Metall Mater 1990;38:2673. [29] Inagaki H. Trans ISIJ 1984;28:266. [30] Duggan BJ, Tse YY. Acta Mater 2004;52:387. [31] Bailey JE, Hirsch PB. Proc Roy Soc Lond A 1962;267:11.
A microstructural study of the origins of c recrystallization textures in 75% warm rolled IF steel M.Z. Quadir *, B.J. Duggan Department of Mechanical Engineering, University of Hong Kong, Pokfulam, Hong Kong Received 16 November 2005; received in revised form 3 April 2006; accepted 15 May 2006 Available online 1 August 2006
Abstract IF steel was warm rolled at 700 C in a single pass. The resulting texture and microstructure were remarkably similar to those of the same steel after cold rolling. A detailed investigation of the microstructure by orientation imaging microscopy and scanning transmission electron microscopy showed microbands to have a mutual misorientation of less than 4 and shear bands to contain material misoriented from the parent matrix by less than 10. Recrystallization did not occur preferentially at high-angle grain boundaries nor in shear bands. Instead the recrystallization nuclei were confined in the original hot band grain envelopes in crystals belonging to the c fiber. These c deformed grains had systematically developed deformation bands which consisted of elements that had rotated by up to 30 about the Æ1 1 1æ parallel to the normal direction. This is essentially the same nucleation process as observed in cold rolled and annealed IF steel. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: IF steel; Warm rolling; Shear bands; Deformation bands; Recrystallization
1. Introduction Steel used for pressing motorcar body panels has changed over the past 20 years from the Al-killed variety to the interstitial-free (IF) kind. The reason is partly economic: the processing takes 50 h for Al-killed steels but only a few minutes for IF steels; and partly technical: the formability is not very sensitive to processing conditions for IF steel, but is very sensitive for Al killed steels [1]. The IF steels are so called because they contain strong carbide formers, such as titanium and niobium, which getter the nitrogen and carbon atoms from solid solution and instead incorporate them into complex precipitates. These IF steels can be processed in continuous rolling and annealing lines which has allowed the successful processing of IF steel to occur for a wide range of parameters, which leads to significant commercial advantage for IF steels over other lowcarbon steels.
*
Corresponding author. Tel.: +852 2857 8615; fax: +852 2858 5415. E-mail address: [email protected] (M.Z. Quadir).
The formation of textures in steel, because of its technological importance, has been continuously investigated over many decades. Early work established that drawability was directly linked to the texture and it was quickly shown that for good drawability the steel should have as high a proportion as possible of grains having a Æ1 1 1æ direction parallel to the sheet normal direction (ND) and as small a proportion as possible with Æ1 0 0æ parallel to the ND [2]. This particular type of texture responsible for good drawability develops during recrystallization annealing from the rolling texture which consists of deformed grains clustered into sets of orientations known as a(Æ1 1 0æ//RD) and c(Æ1 1 1æ//ND) fibers. These are routinely formed at 80–85% cold rolling reduction in industrial practice and can be most conveniently shown in the /2 = 45 section in Bunge’s version of Euler space [3] shown in Fig. 1a. The normal sheet-making route followed for drawable sheets goes through a hot rolling stage in which the sheet is deformed in the austenite state to give a random texture before coiling and cold rolling, but in prize-winning work Barnett and Jonas [4] showed that rolling in the ferrite
1359-6454/$30.00 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2006.05.026
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Fig. 1. /2 = 45 ODF sections of warm rolled (75% at 700 C) and annealed IF steels (intensity levels: 1, 2, 4, 6, 8, 11).
range at temperatures up to 700 C (i.e. warm rolling) also produced well-developed a and c fiber textures. Clearly, annealing of material with these fiber components should produce drawable steel as it does in the cold rolled case, and this proved to be true [5]. However, the scientifically exciting part of their work was that it considerably extended the range of variables over which deformed and recrystallized microstructures could be investigated in ferrite while keeping the textures unchanged. Critical metallurgical factors such as strain rate sensitivity m and dynamic strain aging were determined from room temperature to 700 C and these were related to the microstructures and texture sharpness formed in an ultralow-carbon steel (ULCS) and an IF steel. The work demonstrated that for the ULCS, the value of m changed from negative to positive as temperature increased beyond 320 C and the texture was much sharper at temperatures in excess of 500 C than that found in IF steel at the same temperatures. In contrast, the IF steel showed a slightly positive value of m over a wide temperature range and formed weaker textures, and this was linked to material instability in the form of shear bands, which are narrow, plate-like structures which cut the microstructure at ±35 to the roll-
ing direction (RD) when viewed in the longitudinal section and which carry a relatively large shear strain. These shear bands, in the same way as shown earlier in rolled low stacking fault energy (SFE) face-centered cubic (fcc) materials [6], were claimed to be responsible for generating a weaker texture than formed in ULCS. Furthermore since the IF steel gave a better texture on annealing, the conclusion was drawn, backed by some optical microscopy results, that shear bands in the IF steel were the sites of {1 1 1}Æh k læ nuclei. Now this is not an unreasonable conclusion because Duggan et al. [7] had earlier shown using transmission electron microscopy (TEM) that in Al-killed steel shear bands did indeed provide nuclei of the orientation {1 1 1}Æh k læ. However, in later work, Tse et al. [8] showed that cold rolling produced orientation splitting in a single grain to produce mutually misoriented large volumes of material separated by either a sharp boundary or a more gently curving transition from one orientation to the other. This was shown to occur systematically and they called the process deformation banding (DB). The deformation-induced grain boundary they associated with nucleation of recrystallization, and since DB occurred preferentially in crystals belonging to the c fiber, they proposed that {1 1 1}Æh k læ
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originated from such sites. A link between DB and recrystallization was not established at the time of the work of Barnett and Jonas, i.e. 1997, and so the possible role of DB in the recrystallization texture formation in IF steels was not explored. This is the starting point for the present investigation. It is clearly necessary to investigate whether DB is as important in recrystallization texture formation in warm rolled IF steel as it is in the cold rolled case [8]. The investigation involves orientation imaging microscopy (OIM), TEM and scanning TEM (STEM) for microstructural characterization and uses conventional X-ray diffraction (XRD) global orientation distribution functions (ODFs) to cast new light on the subject of the annealing of warm rolled IF steels. 2. Experimental The material used had the composition 0.0043 wt.% C stabilized by adding 0.057% Ti in the melt. The other impurities were 0.006% P, 0.10% Mn, 0.035% Al, <0.01% Si, 0.002% N, 0.0065% S and 0.0031% O. The IF steel slab was hot rolled to 70% reduction in three passes starting at 1100 C. The exit temperature of the last pass was above 950 C. This resulted in randomly oriented grains with 70 lm average size which is about twice the grain size in commercially produced steels. The sample was then reheated to 750 C and allowed to cool to 700 C before being rolled up to maximum 75% reduction in a single pass, the maximum capacity of the rolling mill used. This high reduction in a single pass was preferred, because if the rolling was done in two or three passes it is inevitable that the waiting period between exit and re-entry to the mill would allow recovery and probably recrystallization to occur. This would affect subsequent processing. The single pass produced a material without any sign of recrystallization, and the hardness of the warm rolled material was 172 VHN, which is to be compared with a value of 178 VHN after cold rolling. High-temperature lubricant was sprayed after each pass onto the rolls to keep the friction effect to a minimum. Recrystallization annealing was done in an aircirculation furnace preheated at 710 C which allowed some minor grain growth to occur at longer annealing times. The global textures were measured using the Schulz reflection technique and Co Ka radiation to produce 1 1 0, 2 0 0 and 2 1 1 pole figures from which ODFs were calculated using the BATE Textan-III software. For the SEM study, longitudinal section (LS) and rolling plane section (RP) were electropolished in an electrolyte having 9:1 volume ratio of acetic and perchloric acids at 30 V DC. Samples from the mid-thickness and away from the edges of the sample were taken for all experiments to avoid depthdependent shear and plate widening effects. TEM foils were produced from the LS by mechanical thinning followed by electrolytic thinning using the standard window method. The sample was observed using a JEOL-STEM FX2000 operated at 200 kV. Slightly convergent beam electron dif-
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fraction (CBED) was used to obtain diffraction patterns for orientation measurements. The foils used for TEM were further observed with SEM to provide links between the microstructures observed using these two techniques. 3. Results 3.1. Rolling and annealing texture Fig. 1a shows the /2 = 45 section of Bunge’s Euler space and the major orientation components along the a and c fibers. The global texture shown in Fig. 1b of 75% warm rolled IF steel consists of these two fibers. Annealing of this material to complete recrystallization produced the texture shown in Fig. 1c, with peak intensity close to {5 5 4}Æ2 2 5æ and the almost complete annihilation of the a fiber components, except at the orientation {1 1 1}Æ1 1 0æ which is common to both a and c fibers. 3.2. Deformed microstructure 3.2.1. Microbands and shear bands Fig. 2 shows a low-magnification channeling contrast (CC) image of the LS showing several grains. Grains C and D are composed of microbands cut through by narrow shear bands in a sense opposite to the habit plane of the microbands. Such grains usually belong to the c fiber. Grains A and B belong to the a fiber and show the characteristic contrast associated with larger subgrains with small mutual misorientations. Grain A has shear banding which is rare in a fiber grains, and grain B has no such features which is much more common. There is evidence, given the width of B, that this grain has undergone much higher deformation than the imposed strain would indicate. Again thinned grains such as B are common. Fig. 3 shows a higher magnification image from a different position in the same specimen and shows three grains, A, B and C. The shear bands in grains B and C are very narrow, 0.5 lm or less. Observation of shear bands in the rolling plane shows that these produce offsets in the transverse direction (TD), proving that there is a shear component parallel to the TD. Another point of interest in this figure is that the left-hand end of grain B has a-type contrast, while from the center to the right-hand end a typical c-type contrast is observed. It is possible that these are, in fact, two different grains, but extensive use of OIM shows that this is unlikely in the majority of cases. Considerable evidence obtained from lower rolling reductions shows how features such as this arise. Presumably, due to neighborhood constraints or energy saving, one part of a grain rotates towards one orientation and forms structures differently to other parts of the same grain which are rotating to another orientation. This is evident in Fig. 4, which shows a single grain after 25% single-pass warm rolling reduction at 700 C. Clearly one part has a different microstructure from the other, and this is undoubtedly the origin of what is visible in Fig. 3. The boundary between grains B and C in
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Fig. 2. SEM CC image of LS of 75% warm rolled IF steel showing several deformed grains with smooth and in-grain shear banded substructures cutting through the microbands.
Fig. 3. SEM CC image of LS of 75% warm rolled IF steel showing the in-grain shear bands and their interactions with the grain boundaries.
Fig. 4. SEM CC image of LS of 25% warm rolled IF steel showing different substructures in different regions of the same deformed grain. This grain will probably consist of at least two different orientations as the rolling reduction increases. This orientation splitting is what gives rise to deformation banding.
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Fig. 3 is not flat, but wavelike; again this is common and has been observed by Tse [9] after cold rolling IF steel to 85% reduction. It is also frequently observed in cold rolled copper and a-brass [10] as well as in Al alloys [11] and is therefore probably a general phenomenon. Fig. 5a shows a TEM image of microbands observed in LS. The CBED patterns from five microbands are shown in Figs. 5c–g and the measured orientations are plotted on a 2 0 0 pole figure in Fig. 5h. The orientation of microband 1 is (2.25, 2.81, 1.62)[2,1,1] and the lattices in the contiguous microbands are rotated by a maximum of 6 3.65. The traces of the microbands are along h3:91; 0:26; 0:73i ffi h1 0 0i. In turn, this lies on the (0 1 1) plane, which is a body-centered cubic (bcc) slip plane. This is consistent with the idea that the microbands are formed by slip processes, a conclusion reached by Chen and Duggan in more extensive work [12] and more recently by Halder et al. [13]. To make the TEM results more statistically convincing the foil was removed from the TEM instrument and examined using
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SEM (Fig. 5b). It is clear that the microbands shown in Fig. 5a are part of a large set contained within the same grain and hence it is safe to assume that the TEM results have some general validity, i.e. microbands extend across the whole grain, a distance of 10–20 lm, and have small mutual misorientations. Shear bands similar to those shown in the CC image in Fig. 2 were then located and examined in a TEM thin foil. Fig. 6a shows a typical example of two sets of microbands making ±30 to the RD. Fig. 6b shows higher magnification detail from Fig. 6a showing a shear band. In Fig. 6b microband M1 is sheared to location M2, and this is proved by the fact that the orientations of M1 and M2 are identical (2.58, 2.15, 2.15)[3.02, 1.44, 2.18] (Figs. 6c and e). S is rotated 7.46 away from M1 and M2 to (2.3, 2.38, 2.24)[3.23, 1.23, 2.00] orientation (Fig. 6d). This is a small rotation. Another example is shown in Fig. 7 which has a pair of shear bands cutting through a grain containing a single set of parallel microbands.
Fig. 5. (a) TEM image from the LS showing microbands. (b) SEM micrograph of the foil area shown in (a). (c–h) CBED patterns and the corresponding 2 0 0 pole figure showing the orientations of microbands 1–5 labeled in the TEM image.
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Fig. 5 (continued)
Sheared elements 1, 2, 3 have the orientation (2.14, 2.92, 1.68)[3.15, 2.45, 0.23], and microbands 4, 5, 6 (2.06, 2.82, 1.92)[3.21, 2.37, 0.033] and 7, 8, 9 (2.02, 2.76, 2.06)[3.25, 2.31, 0.08], i.e. the shear band material 1, 2, 3 is rotated 5 from 4, 5, 6 and 8 from
7, 8, 9. This small rotation was confirmed as the most common case from the larger area scans available using OIM. Fig. 8 shows a CC and OIM image with the associated pole figures for the whole area, together with the line scan through fine parallel shear bands. The average misorienta-
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Fig. 6. TEM image of LS of 75% warm rolled IF steel showing the shear band cutting through one set of microbands while being parallel to another set of microbands and the corresponding CBED patterns of sheared and unsheared material.
tion is 8 and is not simply related to any of the sample reference axes RD, TD, ND. 3.2.2. Deformation bands Fig. 9 shows a CC micrograph taken from the RP section at the mid-plane and the grain boundaries are outlined black after observing at higher magnification. Grain X has the characteristic contrast of grains belonging to the a
fiber, which is verified by the pole figure shown in Fig. 9b. It maintained a fairly constant orientation along the line across the width of the grain. Grains Y and Z show the more complex microstructures which are very widely found in rolled IF steel. The line scans show, in each case, the grain orientation {1 1 1}Æu v wæ to be rotated about Æ1 1 1æ by 20–30. In Y it is from Y1 {1 1 1}Æ1 1 2æ to Y2 {1 1 1}Æ1 1 0æ (Fig. 9c) and in Z from Z1
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Fig. 7. TEM image of LS of 75% warm rolled IF steel with the corresponding CBED patterns showing the magnitude of crystallographic rotation produced by shear band material 1, 2, 3.
{1 1 1}Æ1 2 3æ to Z2 {1 1 1}Æ1 1 2æ (Fig. 9d). This kind of orientation rotation across a single grain, called by the present authors deformation bands (DBs) and by others [14] deformation-induced high-angle boundaries (DIHABs), are commonly found in both warm and cold rolled IF steel [8,15] and a model has been proposed for this process [8]. Grain Y has a continuous orientation variation from one side to the other, while Z has a relatively sharp boundary. Such lattice curvature, i.e. either sharp or more gentle across a grain, has also been observed in cold rolled IF steel [8], and in fcc metals and alloys [16]. 3.3. Recrystallization studies The material was recrystallized to the 8% level at 700 C and very large montages were prepared from both the LS and RP sections. Fig. 10 shows a CC image of a represen-
tative area from the LS showing the variety of annealed states in a stack of warm deformed hot band grains observed in the partially recrystallized condition. Grain A contains recrystallized grains of 3–9 lm in diameter formed in a finely divided microstructure which is characteristic of c fiber material. The growth of these newly formed recrystallized grains is observed to be within the deformed hot band grain, i.e. they have not, at this stage, grown through a preexisting hot band grain boundary. Grain B contains obvious shear bands, indicated by arrows and no obvious recrystallization is associated with them. Instead, only minor spheroidization of subgrains has occurred along the shear bands which is also true of grain E. Inspection of Figs. 3, 5, 6 and 7 shows deformed c material to consist of elongated microbands of 0.2–0.5 lm in width, with occasional ‘‘block type’’ cells of 1 · 1 lm. Annealing produces significant changes in the elongated
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Fig. 8. SEM CC image with OIM and 2 0 0 pole figures of LS of 75% warm rolled IF steel showing the magnitude of misorientation created at the shear band and matrix interfaces.
structure, i.e. the microbands are replaced by subgrains ranging in size from 0.5 to 1.5 lm in diameter (Fig. 10). It is arbitrarily assumed that subgrains of 3 lm diameter or larger are either successful or potential nuclei. From the montage of several hundred square micrometers, of which Fig. 10 is a small part, a total of 428 such nuclei/ grains were found, their orientations being determined
and plotted on a 2 0 0 pole figure in Fig. 11. Some 84% of these nuclei/grains had orientations belonging to the c fiber, with the highest concentration near {1 1 1}Æ1 1 2æ and {1 1 1}Æ1 2 3æ. These orientations are clearly part of the spread of the fully recrystallized texture shown in Fig. 1c. Interestingly, at this stage of recrystallization, i.e. 8%, grains such as B with clear shear bands, provided
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Fig. 9. SEM CC image with corresponding 2 0 0 pole figures of RP section of 75% warm rolled IF steel showing orientation gradients across the grains along line scans.
almost nothing to these data. Clearly, there is recovery in grains such as B, but no recrystallization nuclei/grains were found at this stage using the 3 lm criterion. The LS montages gave very little systematic information regarding nucleation of recrystallization because the nucleation events were difficult to correlate with the structures in which they had occurred. This was not the case when the material was examined in the RP. Figs. 12a and b show a CC image and the corresponding OIM image of a deformed c grain neighbored by two a grains, captured in the RP section of the same 8% partially recrystallized material. The c grain at the middle of the micrograph shows well-developed recrystallization nuclei A, B and C growing into the surrounding material. The overall texture is shown in Fig. 12c in the form of a 2 0 0 pole figure. The orientations of the deformed material 1–4 and the nuclei A–C are labeled in the pole figure corresponding to the labeling in the CC and OIM
images. Grain 1 is part of the a fiber which has a spread from near {4 1 1}Æ1 1 0æ to {1 1 2}Æ1 1 0æ and grain 4 is also a, but part of a complementary set of orientations spreading from {1 0 0}Æ1 1 0æ to {1 1 3}Æ1 1 0æ. The c microstructures 2 and 3 have two different orientations: 2 is close to {1 1 1}Æ1 1 0æ and 3 is close to {1 1 1}Æ1 2 3æ. From the OIM image it is very clear that the nucleation events occurred in the deformation gradient formed in the c-oriented material. The orientation of the recrystallized grains A, B, C belongs to the orientation spread of region 3, i.e. {1 1 1}Æ1 2 3æ. This kind of nucleation in a c grain showing a severe orientation gradient was systematically observed in many cases and is identical to the situation in cold rolled IF steel reported by Tse et al. [8] and described there as nucleation in deformation banded grains of the c family. During successive recrystallization stages recrystallized grains grew within c deformed parent grains and completed the recrystallization process
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Fig. 10. SEM CC image of LS at the 8% partially recrystallized condition of 75% warm rolled IF steel.
Fig. 11. Orientations in 2 0 0 pole figures of partially recrystallized grains at the condition shown in Fig 10 (1: {1 1 1}Æ1 2 3æ; 2: {1 1 1}Æ1 1 2æ). A total of 428 new grains of minimum diameter 3 lm are included in this pole figure.
by growing out of the hot band grain boundaries and consuming the a deformed grains. 4. Discussion 4.1. Warm rolled microstructure and texture The texture of 75% single-pass warm rolled material contains a and c fibers shown in the /2 = 45 ODF section in Fig. 1b. This is almost identical to the cold rolled tex-
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tures measured after the same magnitude of reduction as reported elsewhere [4,8]. Since the same textures are formed in both cold rolling and warm rolling, and the strain rates are 5 and 25 s1 [17], respectively, i.e. they differ by a factor of 5, it is plain that rolling texture formation is insensitive to both of these variables. The general microstructure formed in warm rolling is very similar to that formed in cold rolled IF steel [9]. There is a correlation between orientation and the deformed microstructures, i.e. the a fiber grains have characteristic smeared contrast compared to the sharp and finely fragmented contrast of c grains (Figs. 2, 3, 8 and 9). The observation that a grains can be thinner than their starting grain sizes was also reported in cold rolled IF steel. Chen et al. [18] explained this phenomenon by reference to the number of highly stressed slip systems, which is as many as 7 in a orientations. From this they predicted a relatively homogeneous flow to occur in a grains during deformation. This could also be the reason for the rather featureless appearance of material belonging to the a fiber. This featureless appearance in SEM is actually a large cell structure when examined using TEM in both warm [13] and cold rolled material [18]. Sometimes the large cell structure is revealed by a change in contrast in the a grain envelope, as in Figs. 2, 3 and 9. Microbands are the characteristic feature of c grains (and a grains, if they are close to {1 1 0}Æ1 1 2æ [19]) and are associated with less than five slip systems, the number known to be essential for homogeneous deformation. They are volumes of material, relatively free of dislocations enclosed by pairs of dislocation walls between 0.2 and 0.5 lm apart, which appear as parallel lines inclined 25–35 with RD in the LS (Figs. 3, 5 and 6). There are two basic models in the literature for microband formation, that due to Jackson [20] and Chen et al. [18], and based on this work it is not possible to choose between them. It is clear that the trace of the dislocation walls shown in Fig. 5a is along [1 0 0], which is a direction in a {1 1 0} plane and is therefore most probably the operating slip plane. Data such as are shown in Fig. 5 have been obtained in cold rolled IF steel and rigorously analyzed by Chen et al. [18] and Halder et al. [13] who concluded that microbands are a crystallographic phenomenon. This view has recently been questioned by Hurley et al. [21], in careful OIM work on Al–0.13% Mg using two contiguous surfaces, who concluded that microbanding is not a crystallographic phenomenon. Be that as it may, the important role of the microbands is that they constitute obstacles against which dislocations pile up prior to the strain burst which forms the shear bands [14]. The study of shear bands is emphasized in this investigation since they are obviously an important microstructural element and play a vital role in the deformation process, in both cold and warm rolled IF steels [4]. In this investigation shear bands are 0.2–0.5 lm thick, are generally confined within a single grain boundary envelope and carry a large amount of shear strain (i.e. a natural shear strain
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Fig. 12. SEM CC image with corresponding OIM and 2 0 0 pole figure showing the formation of recrystallized grains in a deformation banded grain.
4.6 in Fig. 6) and appear to be crystallographic because they are often formed parallel to a second set of microbands where these occur (Fig. 6). The large strain they carry is evident from the significant offsets created by shear bands in the preexisting microbands and at grain boundaries where they terminate (Figs. 2, 3, 6, 7 and 8). Rare examples of shear bands continuing across grain boundaries are between grains C/D in Fig. 2 and B/C in Fig. 3. Shear bands normally cut one set of microbands in the opposite sense to the RD (25–35) and are usually absent
in the a oriented microstructure. There are significant similarities between microbands and shear bands at the rolling reduction used in the present investigation, i.e. both 1. appear to be crystallographic and are confined within a single grain envelope (Fig. 8a), 2. have a grain orientation dependency, i.e. they are found in c grains and almost absent in a grains, 3. are similar in width to microbands, 4. appear to have a sheet-like structure, and
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5. are parallel to each other and are parallel to a second set of microbands when they are present (Fig. 6a). Therefore presumably the principal difference between them is the stage at which they were formed, i.e. microbands were formed in the early stages of rolling and shear bands were formed later when a c oriented grain is significantly work hardened by a series of parallel sets of dislocation walls which comprise the microbands. The orientation relationship between the shear band material and the matrix in which they have been formed is particularly important in regard to recrystallization. Generally shear band materials were found to be 8–10 misoriented with the matrix material as measured by EBSP of SEM and CBED of TEM (Figs. 6–8). In the TEM micrograph shown in Fig. 6b, the microbands have rotated 35–45 into the shear bands and are therefore expected to have a crystallographic rotation of the same magnitude if a rigid body rotation alone is involved. However the measured 8–10 misorientation implies that there is significant amount of slip involved in the deformation process and a model for how this is achieved has been proposed [22]. This small misorientation is effectively a low-angle boundary between the shear band material and the parent matrix. 4.2. Recrystallization behavior Two factors are essential for successful nucleation: formation of a mobile interface, which implies a critical misorientation with the matrix being consumed; and a driving force, which is usually related to the difference in dislocation density between the nucleus and its surroundings. Boundaries lose their dislocation character when the misorientation between two lattices reaches 15–20, and this is also the misorientation required for mobile interfaces. The shear bands have misorientations of about 50% of these values and therefore are not expected to be mobile. Of course these intermediate values of misorientation can be increased by sub-boundary climb during recovery, but this has to involve migration of the interface as well. Representative grains containing shear bands, such as is shown in Fig. 10, show evidence of recovery, but there are no new grains, unlike the situation in the grain above the shear banded grain. The fact that setting a lower limit of 3 lm to capture nucleation events meant that shear banded grains made almost no contribution to the data plotted in Fig. 11, indicates that they are not effective nucleation sites. The reason why Barnett and Jonas [4] concluded that shear bands gave rise to their final texture is that they used etched surfaces and optical microscopy for their metallographic investigation. It seems likely that the etch marks which covered c grains are traces of microbands and shear bands, because the intensity of the marks is far higher than would be expected for shear banded grains [4]. The optical technique is also of lower resolution than the electron beam
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technique used here, where very fine details are revealed. The conclusion is that shear bands in warm rolled IF steel do not significantly contribute to the formation of the recrystallization texture. Careful observation of large CC montages covering many hundreds of square micrometers at the 8% recrystallized stage showed almost no evidence at all of straininduced boundary migration (SIBM). The most likely candidate orientation for this process is grains belonging to the a orientation, for these are naturally low-energy blocks. They do not provide nucleation sites at this stage, but only recovered microstructures (Fig. 12). Since the final texture (Fig. 1c) shows only very low intensity along the a fiber, it is plain that a components are mostly consumed by other orientations. Another factor that must be borne in mind when considering the SIBM mechanism is the fact that in IF steels there is good evidence that the hot band boundaries are preferred sites for the nucleation of FeTiP and MnS particles [23,24]. These, plus any solute segregation provides pinning or drag on such a boundary, and SIBM has to overcome this problem. In both this warm rolled and also in cold rolled IF steel [9,15] there is abundant evidence that nucleation occurs inside the original hot band grain enveloped first, and growth out of these does not occur until later, typically at the 50% recrystallized volume fraction stage. The two factors essential for successful nucleation, formation of a mobile interface and a driving force between the nucleus and its surroundings, are obviously met in the deformation banded grains. The orientation gradients measured across the rolling plane surface of c grains which have DBs provide the necessary lattice curvature as is clearly shown in Fig. 9. This is a result typical for both warm rolling and cold rolling [8,15]. Further examples are shown in recovered c grain material in Fig. 12 where the rotation is from {1 1 1}Æ1 2 3æ to {1 1 1}Æ1 1 0æ. Thus the lattice curvature condition is met by the deformation banding process, which is systematic and is always found in the RP section in c grains. A simple model for deformation banding in bcc metals is given by Liu and Duggan [25]. Information regarding the driving force, however, is lacking in this investigation. However, early work on drawable steels using XRD [26] and in later work in IF steel using neutron diffraction [27] showed that recovery processes in both steels changed the as-deformed stored energy Def: Recov: relationships: EDef: f1 1 1gh1 1 0i hE f1 1 1gh1 1 2i changed to E f1 1 1gh1 1 0i i Def: Def: Recov: Recov: ERecov: f111gh112i and E f1 1 1gh1 1 0i hEf11 1gh1 23i to E f1 11gh1 1 0i iEf1 1 1gh12 3i .
If these inequalities hold true at the single-grain level, there is therefore an internal driving force inside a single recovered c grain which has deformation banded. It is difficult in principle to know where a nucleation event has taken place inside a three-dimensional object using only two-dimensional sections, because what is observed is a cross section of a developing crystal. Such a crystal can grow in any direction until it is stopped, and there is good reason to suppose that these new grains are
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stopped by grain boundaries. This fact has led to the notion that the regions adjacent to grain boundaries are actually the nucleation sites. This may well be so, but the evidence for this must always be ambiguous when twodimensional sections are being examined. For example, the new grains in Fig. 12 could be nucleated away from a grain boundary or at a grain boundary above or below the section being examined. Likewise the new grains visible in Fig. 10 could be nucleated adjacent to the grain boundary, or this could be a section along a line near the center of the deformed grain in which the new grain grew upwards or downwards to impinge the grain boundary. Therefore it is concluded that DBs provide the necessary conditions for nucleation of c oriented grains in c oriented material but it is impossible to rule out ‘‘near boundary’’ nucleation in the manner described by Senuma et al. [28] and Inagaki [29]. At a later stage of recrystallization, new grains grow out of the hot band envelopes to give the equiaxed structure and consume the remaining a deformed hot band. A mechanism is outlined by Duggan and Tse for cold rolled IF steel [30]. This happens when the new grains attain a critical length L, given by the SIBM formula [31] L P 2c/DE, where c is the surface energy of the original grain boundary and DE is the difference in energy across the hot band grain boundary. Of course, this is the simplest formulation, a more realistic formula involving both solute drag and pinning force has been given by Tse and Duggan [9] but this is not discussed further here, since this work adds nothing to the analysis. However, these terms have the effect of increasing L to the order of 10 lm or more, and hence the growing out of the nucleated grains occurs relatively late, on average, in the recrystallization process. 5. Conclusions 1. The rolling and annealing textures of warm rolled IF steel are very similar in terms of orientation and sharpness to those of cold rolled IF steel at the same rolling reduction. 2. Single sets of shear bands are commonly found but are not effective in recrystallization because the shear band material is misoriented 8–10 from its parent matrix, hence forming a low-angle interface with the matrix. 3. The in-grain shear bands, which are similar to the microbands in terms of geometry, alignment with sample axes and thickness, are predominantly formed in deformed c grains. 4. DBs produce high lattice curvature which can be either sharp or cumulative over several subgrain diameters and are best observed in the RP section. The banded seg-
ments are of c orientation having a general rotation relationship around Æ1 1 1æ//ND and are effective sites for nucleation of c recrystallized grains.
Acknowledgements It is a pleasure to acknowledge the support of this work by Grants Nos. CERG/HKU 7067/97E, 7323/98E and 7316/99E given by the Hong Kong Special Administration Region of China. References [1] Hutchinson WB. Mater Sci Forum 1994;157–162:1917. [2] Held JF. In: Proceedings of the conference on mechanical working and steel processing IV. New York (NY): Metallurgical Society of AIME; 1965. [3] Bunge HJ. Mathematische Methoden der Texturanalyse. Berlin: Akademie-Verlag; 1969. [4] Barnett MR, Jonas JJ. ISIJ Int 1997;37:697. [5] Barnett MR. ISIJ Int 1998;38:78. [6] Morii K, Mecking H, Nakyama Y. Acta Metall 1985;33:379. [7] Duggan BJ, Liu GL, Ning H, Zhang LX. Thermomech Process Theory Model Practice 1997:384. [8] Tse YY, Liu GL, Duggan BJ. Scr Mater 2000;42:25. [9] Tse YY, Duggan BJ. Mater Metall Trans A 2006;37A:1055. [10] Sindel M, Ko¨hlhoff GD, Lu¨cke K, Duggan BJ. Texture Microstruct 1990;12:37. [11] Hansen N. Metall Trans A 1985;16A:2167. [12] Chen QZ, Duggan BJ. Metall Mater Trans A 2004;35:3423. [13] Halder A, Huang X, Leffers T, Hansen N, Ray RK. Acta Mater 2004;52:5405. [14] Hughes DA, Hansen N. Metall Trans A 1993;24:2021. [15] Quadir MZ. Ph.D. thesis. University of Hong Kong; 2003. [16] Lee CS, Duggan BJ. Acta Metall 1993;41:2691. [17] Lam KT, Quadir MZ, Duggan BJ. Key Eng Mater 2003;233–236:437. [18] Chen QZ, Ngan AHW, Duggan BJ. Proc Roy Soc Lond A 2003;459:1661. [19] Quadir MZ, Duggan BJ. Acta Mater 2004;52:4011. [20] Jackson PA. Scr Mater 1983;17:199. [21] Hurley PJ, Bate PS, Humphreys FJ. Acta Mater 2003;51:4737. [22] Chen QZ, Quadir MZ, Duggan BJ. Philos Mag 2006;86:3633. [23] Deardo AJ. In: Takechi H, editor. Proceeding of the international forum properties and applications of IF steel. Tokyo: Arcadia Ichigaya; 2003. p. 240. [24] De Cooman BC, de Vyt A. In: Takechi H, editor. Proceeding of the international forum properties and applications of IF steel. Tokyo: Arcadia Ichigaya; 2003. p. 249. [25] Liu G, Duggan BJ. Mater Metall Trans A 2001;32:125. [26] Takechi H, Katoh H, Nagishima S. Trans AIME 1968;242:56. [27] Rajmohan N, Hayakawa Y, Szpunar JA, Root JH. Acta Mater 1997;45:2485. [28] Senuma T, Yada H, Shimizu R, Harase J. Acta Metall Mater 1990;38:2673. [29] Inagaki H. Trans ISIJ 1984;28:266. [30] Duggan BJ, Tse YY. Acta Mater 2004;52:387. [31] Bailey JE, Hirsch PB. Proc Roy Soc Lond A 1962;267:11.