Physical parameters related to the developments of recrystallization textures in an ultra low carbon steel

PII:

Acta mater. Vol. 47, No. 1, pp. 55±65, 1999 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00337-1 1359-6454/99 $19.00 + 0.00

PHYSICAL PARAMETERS RELATED TO THE DEVELOPMENTS OF RECRYSTALLIZATION TEXTURES IN AN ULTRA LOW CARBON STEEL I. SAMAJDAR1, B. VERLINDEN1{, L. KESTENS2 and P. VAN HOUTTE1 Department MTM, KU Leuven, de Croylaan 2, 3001 Heverlee, Belgium and 2Department Flat Rolling, CRM, Technologiepark 9, 9052 Gent, Belgium

1

(Received 13 March 1998; accepted 25 September 1998) AbstractÐDevelopments in deformation and recrystallization textures were studied in cold-rolled (50±90% reduction) ultra low carbon (ULC) steel using X-ray texture measurements and orientation imaging microscopy (OIM). During deformation, g-®bre (ND//h111i) increased between 0 and 50% reduction but then did not change signi®cantly, while a-®bre (RD//h110i) increased progressively from 0 to 90% reduction. After complete recrystallization, however, a steady increase in g and almost no changes in a were observed with increasing strain. Developments in recrystallization textures were attributed to two parameters: (1) spacings (li, as measured along the normal direction, ND, where i can be a speci®c component of a/g-®bres) of the a/g deformed bands; and (2) their relative ability to form recrystallized grains. While li was determined by the deformation texture and the thicknesses of the deformed grains/bands and naturally decreased with increasing strain, estimations of parameter (2) were obtained from the so-called nucleation factors (Ni, de®ned as the number of recrystallized i grains per i bandÐas measured/estimated along the ND). At higher strains, noticeable drops in the Nis of a-®bre were observed. Two plausible causes for such drops were increased stored energy advantages for g bands and orientation pinning in some of the a regions. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.

1. BACKGROUND

tion will highlight that the ®nal recrystallization texture is due to availability of certain sources of recrystallized nuclei and the relative potency of such sources. A recent study [7] on Ti-bearing IF steel has shown that recrystallized grains of a- (RD// h110i) and g-®bre nucleate from the deformed bands of these respective orientations. In planestrain deformation (e.g. rolling), a deformed a/g grain/band is expected and observed [7, 13] to be considerably elongated along the RD (rolling direction), while its thickness along the TD (transverse direction) is expected to be independent of reduction. The band may, however, thin down along the ND (normal direction) [7, 13]Ðthe thickness (along the ND) of the deformed bands being dependent on the amount of reduction and developments (if any) of new high-angle grain boundaries [7, 13± 16]. Evidently, availability of recrystallized a/g grains will depend on the spacings (as measured along the ND) of the respective deformed grains/ bands. The relative potency (i.e. ease with which recrystallized grains may be formed) of a/g bands is, however, a more tricky subjectÐas it may involve nucleation as well as growth. A more generalized statistical quanti®cation of the relative potency may be sought in the form of nucleation factorsÐNi. Ni is de®ned [7] as the number of recrystallized grains of an orientation i (where i may be any particular component of the a/g-®bre) per deformed i band, as

Ultra low carbon (ULC) steel and interstitial free (IF) steel are two important classes of deep drawing quality (DDQ) steels [1±9]. In IF steel, carbon is mostly arrested by suitable interstitial elements (e.g. Ti, Nb), while existence of free carbon (albeit at very low levels) is possible in ULC steel. The excellent deep drawability in both steels is due to the existence of a strong g (ND//h111i) ®bre recrystallization texture [1±9]. Typically, the development of g recrystallization texture is attributed to preferred nucleation and/or selective growth of g grains, presumably from g-oriented deformed regions [1±12]. Considerable scienti®c e€orts have been directed to the discussion and clari®cation of these two ``mechanisms''. In general, although the existence of preferred nucleation is widely accepted [1±12], questions still exist on the relative contribution of selective growth [2, 5, 7]. Rather than doing in-depth research to resolve (i.e. if possible) the ``dispute'' between preferred nucleation and/or selective growth, it is possibly far more important, at least technologically, to identify the factors responsible for the developments in g recrystallization texture and to relate them to some physical characteristics of the deformed microstructure/microtexture. Any introduction to recrystalliza{To whom all correspondence should be addressed. 55

56

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL

observed/estimated along the ND. Although originally proposed for f.c.c. texture [15, 16], Nis were quite successful in describing the strong g-®bre recrystallization texture in IF steel [7]. The term nucleation factor is possibly slightly misleading, as Nis are typically [7, 15, 16] statistical estimates of both nucleation and growth. Volume or area fractions of the recrystallized grains (i.e. after full recrystallization) may be obtained as Nidi/ li [7, 15, 16], where di and li are the average recrystallized grain size and the deformed band spacingÐ both for an orientation i and measured/estimated along the ND. Ni and li are two physical parameters responsible for the formation of the ®nal recrystallization texture, which, on the other hand, are related to the deformed state. In this study, an e€ort will be made to investigate the developments/changes in these two parameters at four di€erent cold-rolling reductions in a ULC steel.

2. EXPERIMENTAL METHODS

As-cast ULC steel (chemical composition given in Table 1) was hot-rolled with reheating, ®nishing and coiling temperatures being 12508C, 9208C and 7508C, respectively. The hot-band material had a weak transformation texture, with an average recrystallized grain size of 50 mm. Hot-rolled material was subsequently cold-rolled to 50, 70, 82 and 90% reductions. To study the as-deformed structures by orientation imaging microscopy (OIM), the cold-rolled samples were given a prolonged recovery treatment at 4008C. Other samples were recrystallized in a salt bath at 5508C (for 70±90% rolled samples) and at 6508C (for 50% rolled samples) for di€erent times, so as to obtain di€erent recrystallized fractions. Typically for OIM measurements, mid-thickness (t/2) regions of the long transverse sections (i.e. planes containing rolling and normal directions) were scanned. Details of the OIM sample preparation is given elsewhere [10]. X-ray bulk texture measurements were obtained from the rolling planes (and also from the t/2 sections). Xray ODFs (orientation distribution functions) were measured by the inversion of four incomplete pole ®gures, using the standard series expansion method [17, 18]. For volume fraction measurements, X-ray ODFs were convoluted with suitable model functions (with an integrated ODF value of 1 and which used a 16.58 Gaussian spread).

Table 1. Chemical composition (in mass %) of the ULC steel C 0.0024

N

S

Mn

Al

Si

P

0.0017

0.003

0.079

0.021

0.0045

0.010

3. RESULTS

As in previous studies [7, 10, 13, 19, 20], textures of the cold-rolled and/or recrystallized ULC steel are discussed considering mainly the a (RD//h110i) and g (ND//h111i) ®bre components. While F {111}h112i and E {111}h110i orientations were considered as part of g-®bre, I {112}h110i and H {001}h110i components were considered as part of a. For bulk X-ray texture, volume fractions of individual orientations were obtained using a 16.58 Gaussian spread, while discrete orientation measurements of the OIMs were characterized by a 208 misorientation from the respective ideal components. If an orientation fell within 208 of more than one ideal orientation (i.e. F/E/I/H), then it was considered as part of the component of least misorientation. Orientations misoriented by more than 208 of ideal a/g components were considered as ``random''. Volume fractions with 16.58 Gaussian spread (as obtained by convoluting the X-ray ODFs) were observed to be approximately equivalent to measured average area fractions (from OIM) with the 208 misorientation criteria. Hence a rough comparison of the numerical results (volume fractions vs area fractions) from X-ray and OIM may be justi®ed. 3.1. Global developments in textures 3.1.1. Deformation textures. Figure 1 plots the f2=458 sections of X-ray ODFs for (a) initial hotband texture and (b)±(e) after 50±90% cold rolling. The fully recrystallized hot band contained a weak transformation texture. Note that the initial hotband texture/microstructure is important, as it may a€ect the ®nal (i.e. after cold rolling) recrystallization texture [2, 21]. Figure 2 plots the respective volume fractions of the ®bres and their components at all four reductions. Volume fractions in the starting hot-band texture are also included as reference. As shown in the ®gures, the weak transformation texture of the hot band was converted into relatively strong g- and a-®bres after the lowest reduction of 50% [see Fig. 1(a) and (b)]. With subsequent reductions, g-®bre did not strengthen signi®cantly, although the relative strength of a increased. The changes in g/a were re¯ected in the changes of their respective components (e.g. F/E and I/H). Note that somewhat similar observations were made in channel die-compressed IF steel (7± 75% reduction) [13]. 3.1.2. Recrystallization textures. Figure 3(a)±(d) plots the f2=458 sections of the X-ray ODFs and Fig. 4 shows the volume fractions of the ®bres and their respective components after complete recrystallization. For such complete recrystallization, recrystallization times were selected to fully recrystallize the samples, while avoiding any noticeable post-recrystallization grain growth [20]. The recrystallization texture of the 50% deformed material

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL

57

Fig. 1. f2=458 sections of the X-ray ODFs for (a) hot-band material and for samples after (b) 50%, (c) 70%, (d) 82% and (e) 90% reductions.

was randomized, while with subsequent reductions the g recrystallization texture strengthened noticeably (see Fig. 3). As in Fig. 4, a gradual increase in g but almost no changes in a was observed with increasing strain. 3.2. As-deformed (but recovered) state Extensive OIM measurements were obtained from deformed but recovered ULC samples. In the OIM scans, deformed grains were typically

observed in the form of pancaked bands of nearly one uniform (i.e. within 208 of an ideal component) orientation (see Fig. 5). The boundaries of such bands were often (but not always) marked by >208 misorientation. In the following sections, three important aspects of the deformed microstructures are discussed. 3.2.1. Deformed band thicknesses (ti). The average band thicknesses (ti, measured along the ND) are listed in Table 2 for all four reductions. Also

58

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL

Fig. 2. Volume fractions (with 16.58 Gaussian spread, as estimated by convoluting respective X-ray ODFs) of g/a®bres and F/E/I/H components in the as-deformed (50± 90% reduction) ULC steel. Volume fractions in the hotband texture are also included.

included are the respective standard deviations for a/g band thicknesses. Although the bands thinned with increased reductions (see Table 2), at any particular reduction, average tis were observed to be more or less independent of orientation (i.e. average thicknesses of a/g/random bands were similar at any given reduction). However, except for 50% reduction, spread or distribution (as represented by the standard deviations) of a band thicknesses were considerably wider than g bands (note that standard deviations for random bands were somewhat similar to g). For example, at 70% reduction observed

Fig. 4. Volume fractions (16.58 Gaussian spread) of g/a®bres and their respective components, as obtained by convoluting X-ray ODFs of 50±90% deformed but fully recrystallized samples.

maximum and minimum band thicknesses were 15 and 1.8 mm for a, but only 9.6 and 3.4 mm for g. Typically, the thick a bands were the relatively wellde®ned pancaked deformed grains (with band edges marked by >208 misorientation), while the thin a bands/regions were often observed as part of a larger pancaked grain with random or g orientation. 3.2.2. Deformed band spacings (li). Deformed band spacings (li, as measured along the ND) were directly measured from the OIM scans and are listed in Table 3 for F/E/I/H/random components. At any given reduction, the spacings of the deformed bands (li) will depend on the average

Fig. 3. f2=458 sections of the X-ray ODFs in the fully recrystallized samples after 50% (a), 70% (b), 82% (c) and 90% (d) reduction.

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL

59

Fig. 5. OIM micrographs of (a) 50% and (b) 90% deformed (but recovered) samples. Orientations within 208 of exact a/g-®bre components are marked in di€erent levels of grey, while grain boundaries are marked for 1±208 (light boundaries) and >208 (dark boundaries) misorientation. Some (though not all) of the deformed bands of F/E/I/H components are also marked. In (b) grain boundaries are drawn for the partial OIM, as the changes in the levels of grey in the other section of the OIM (the part in which no grain boundaries are marked) may give an indication of misoriention/fragmentation in bands of di€erent orientation. Note that the darker grey levels correspond to lesser misorientation from the ideal components.

band thickness (ti) and area fraction (Ai) of the component, as li=(ti/Ai). With increasing reductions, spacings of both a and g bands decreased.

While a drop in g spacings (above 50% reduction) was mainly due to ``thinning'' of the deformed bands (i.e. reduced tis), the drop in a spacings was

60

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL

Table 2. Measured band thicknesses (along the ND) after 50±90% reductions. Values of average band thicknesses (av. ti), which, at any particular reduction, were similar between bands of di€erent orientations, are listed along with the respective standard deviations (SD) for a/g band thicknesses 50% av. ti (mm) 25

70% SD a 10.1

SD g 10.9

av. ti (mm) 7

82%

SD a 3.8

SD g 1.9

Table 3. Spacings (as measured along the ND from the OIM scans) of F/E/I/H/random bands after four di€erent reductions Reduction percentages

lF

lE

lI

lH

lrandom

50% 70% 828 90%

125 31 23 14

124 30 21 11

153 37 23 10

442 107 68 30

54 16 13 8

caused by a combination of reduced tis and increased a fractions (see Fig. 2). Evidently the decrease in a spacings was relatively stronger than the drop in g spacings, see Table 3. 3.2.3. Strain localizations. In cold-rolled DDQ steels, plastic instabilities were often reported in the form of micro or shear bands [13, 22, 23]. In typical OIM scans of deformed (but recovered) samples [see Fig. 5(a)], such plastic instabilities appear as ``clusters'' of grain boundaries at an approximate angle of 378 with the RD, and in almost all cases inside the individual pancaked grains. As in a previous study [13], rather than specifying them as micro or shear band, when exact mechanism(s) of their formation is not fully understood, a more generalized nomenclature of ``strain localization'' was used. Indeed, presence of more frequent boundaries of higher than average misorientations justify this name. In plane-strain channel die-compressed IF steel [13], marked appearance of the strain localizations also coincided with stored energy advantage for g bands. An estimation of stored energy may be obtained from y/d (cell size/cell misorientation) values [7, 10, 13, 15, 16], as obtained from the OIM scans. Such estimations may have several limitations: 1. stored energy may scale with y/d when the presence of statistically stored dislocations (e.g. loose cell interior dislocations) is neglected; 2. some of the dislocation cells may fall below the resolution (for OIM, typically 0.5 mm); and

av. ti (mm) 4.4

90%

SD a 2.8

SD g 1.8

av. ti (mm) 2.4

SD a 1.8

SD g 1.2

3. measurements need to be conducted on recovered samples and this may not be truly representative of the as-deformed state. Alternative techniques (e.g. line broadening in neutron di€raction, or y/d measurements in TEM) also have their restrictions. In the latter case the area covered in one TEM foil is extremely small, and a very large number of samples needs to be studied in order to obtain a reliable amount of data. In neutron di€raction the resolution is restricted and the stored energy of minor texture components cannot be measured. Results from the three techniques [10, 12] show that stored energy increases with increasing Taylor factors, except for orientations with very high Taylor factors. For example, E has a slightly higher Taylor factor than F, but a lower stored energy. In spite of the limitations of the OIM, relatively large amounts of OIM data may provide an estimation of stored energy variations among di€erent texture components (see Table 4). Also listed are the spread/distribution (as represented by the respective standard deviations) of the y/d values. Results (as listed in Table 4) were obtained from samples with the same recovery treatments (typically 4008C/48 h annealing in a salt bath) and only one sample was used for each reduction (although OIM data were obtained from several scans on each sample). As given in Table 4, at 50% reduction, y/d values (and the standard deviations) were nearly the same for all four a/g-®bre components (e.g. F/E/I/H). Even at 50% reduction, strain localizations were often seen [see Fig. 5(a)], but their appearance was virtually independent of orientation. For 70% reduction and above, these so-called strain localizations appeared somewhat preferentially (as observed qualitatively from the OIM scans) on the higher Taylor factor [10, 11, 13] orientations of F and E. In general, preferred formation/appearance of the strain localizations had two direct e€ects on the g-oriented deformed bands: (i) it gave them a broken or fragmented appearance

Table 4. Average values of y/d (cell misorientation/cell size) and the estimated standard deviations (SD, providing a measure of the spread/distribution of the measured y/d values) in F/E/I/H components and after 50±90% deformation (plus the recover treatment) Orientation (within 208 of the ideal components) F E I H

50% Av. y/d (8/mm)

SD

70% Av. y/d (8/mm)

SD

82% Av. y/d (8/mm)

SD

90% Av. y/d (8/mm)

SD

1.4 1.3 1.43 1.36

0.38 0.31 0.35 0.26

2.78 2.6 2.12 1.38

0.98 0.9 0.82 0.33

3.1 3.2 2.1 1.36

1.04 1.01 0.87 0.31

3.23 3.16 2.2 1.39

1.24 1.11 0.85 0.36

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL

61

[see Fig. 5(b)] and (ii) as in IF steel [13], a marked stored energy advantage (as both the average y/d and its spread/distribution were higher; see Table 4).

3.3. Di€erent stages of recrystallization In this study, any strain-free region, i.e. regions which appeared free from >18 (as 0±18 misorientations were considered within the tolerance of the OIM measurements) grain boundaries, above 3 mm (along the ND) in size, was considered as a recrystallized grain. The observations on such recrystallized grains and on the remaining ``nonrecrystallized'' regions are described in the following sections. 3.3.1. Frequency vs size advantage. Formation of recrystallization texture, in general, may be explained from the frequency or size advantage of recrystallized grains [7, 10, 15, 16, 24]. In other words, after recrystallization the strength of a particular texture component i may increase when recrystallized i grains are more in number or larger in size [24]. Such generalized mechanisms are relatively easy to characterize [15, 16, 24] and may provide an overall picture of the g-®bre recrystallization texture [7, 10]. In the present ULC steel and at any given reduction, average grain sizes of g grains were comparable to a and random grains (see Table 5). The maximum grain sizes of g, a and random grains were also comparable (approximately) at all stages of recrystallization. This rules out possibilities of global growth advantage for g [7]. In absence of any size advantage, g recrystallization texture of the ULC steel (irrespective of reduction) was evidently caused by the frequency advantage of the g grains. Figure 6 plots the number fraction of (a) g and (b) a grains after di€erent amounts of recrystallization and at all four reductions. As shown in the ®gure, at 50% reduction the number fraction of g/a grains did not change signi®cantly with the progress of recrystallization. For 70% and larger reductions, however, number fractions of g decreased but a increased (to their respective ®nal values) as the recrystallization progressed to completion (see Fig. 6). Table 5. Average sizes (in mm, and as-measured along the ND) of F/E/I/H/random grains in fully recrystallized 50±90% reduced samples. Average grain sizes irrespective of any orientation are also listed. Recrystallization times were selected to fully recrystallize the samples, but to avoid any post-recrystallization grain growth Orientations F E I H Random Irrespective of orientation

50%

70%

82%

90%

21 22 26 19 24 23

15 19 18 16 15 17

14 12 11 12 15 13

10 11 12 10 11 11

Fig. 6. Number fractions (i.e. number of a/g grains/total number of grains) of (a) gamma (g) and (b) alpha (a) grains at di€erent stages of recrystallization. The captions provide the approximate ranges (e.g. in fraction recrystallized) of recrystallization over which the measurements were obtained.

3.3.2. Recrystallization kinetics. As discussed in Refs [2, 19, 25], recrystallization kinetics in low carbon or IF steel may di€er signi®cantly from the usual sigmoidal behaviour as typically sluggish recrystallization kinetics is observed above 70% recrystallization. The same observations were made in the ULC steel at all four reductions. Considering 0±70% recrystallization (i.e. by avoiding the sluggish recrystallization part [19]), JMAK exponents were measured between 2 and 2.5 for 70±90% reductions, while the JMAK exponent for 50% reduction was signi®cantly higherÐestimated as 3.7. The higher JMAK exponent of 50% reduced material was due to ``delayed'' nucleation [15]. At 6508C, 50% deformed material hardly recrystallized until 50 s annealing (only 3% recrystallization was estimated), while 3±70% recrystallization was estimated between 50 and 80 s. 3.3.3. Above 70% recrystallization. Even above 70% recrystallization, samples of all four reductions showed some remaining bands of nearly one uniform orientation (see Fig. 7). Such bands, as described in detail in the inset of Fig. 7, contained deformed as well as recrystallized regions (considering our convention of existence or non-existence of grain boundaries within a distance of 3 mm). In the 50% deformed material no distinct preferred orientation (i.e. F/E/I/H/random) was noted for these bands, while bands in samples deformed to 70% or

62

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL

Fig. 7. 90% deformed sample with ``remaining'' I band (5508C and 600 s annealed). A magni®ed section of the band is given in the inset. Graphics convention is otherwise the same as in Fig. 5.

more were predominantly of a-®bre. In general, these ``remaining'' bands had several interesting di€erences with the original as-deformed band structure. They were fully recovered (as OIM pattern quality from the band and the regular recrystallized grains were comparable) and contained larger fractions of low-angle boundaries than the original as-deformed bands. Their average thickness was, in general, 2±4 times higher than the corresponding as-deformed band thickness (i.e. for the same reduction). 4. DISCUSSION

4.1. The global picture The overall picture of the developments in global textures (both deformation and recrystallization) is quite clear from Figs 1±4. During deformation, g®bre increased from 0 to 50% reduction, but then did not increase further, while after recrystallization g increased steadily with strain. The a-®bre, on the other hand, increased steadily with strain during deformation, but did not change during recrystallization. Present as well as past [7] OIM results had clearly shown that the g/a grains nucleate from the respective deformed regions. Hence, the apparently strange results on deformation and recrystallization textures (see Figs 2 and 4) indicate that the relative contributions (towards recrystallization) from deformed g regions increased and from deformed a regions decreased with increasing strainÐwhatever might be the reason. 4.2. On the preferred nucleation As discussed earlier in Section 3.3.1., frequency advantage of g grains was mainly responsible for the g recrystallization textures. As shown in Fig. 6, the number fraction of g grains increased with strain, while the number fraction of a grains did

not change signi®cantly. These clearly show the same trend as global textural developments. Relative changes (with progress of recrystallization) in the number fraction of grains of particular orientation(s) may serve as an index of preferred nucleation [7, 15]. In other words, if the majority of grains with particular orientation nucleate ®rst, then their number fraction is expected to fall from a relatively higher to a lower value (i.e. from the beginning to the completion of recrystallization). As shown in Fig. 6, at 50% reduction the number fraction of g/a grains did not change with recrystallization, while for 70% and larger reductions and with increasing recrystallization, the number fraction of g fell, but a increased. These indicate lack of any preferred nucleation at 50% reduction, while at and above 70% deformation the nucleation of g grains was de®nitely preferred over a. Such an index of preferred nucleation (based on the statistics of number fractions, as in Fig. 6) may include the classical nucleation, but may also include limited or selective growth [4±7, 28]. In an experimental characterization of recrystallized grains (as in OIM), some convention needs to be ®xed to separate the recrystallized grains from the deformed microstructure. Usually in OIM, such a convention can be based on the absence of grain boundaries over a minimum size (taken as 3 mm in the present study) [7, 10, 15]. In gaining such a minimum size, nucleation as well as limited/selective growth may be involved. In the absence of a global growth advantage (see Section 3.3.1.), limited or selective growth at the early nucleation stages may be considered (as in Refs [7, 10]) as part of the overall nucleation process. There are two plausible mechanisms behind the preferred nucleation of g grains: stored energy advantage for deformed g regions [7, 10±12] and/or micro-growth advantage/selection for F-oriented

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL

grains growing into I-oriented deformed regions [2± 4, 7, 10, 23, 26]. The measurements on y/d clearly showed higher values (both the average values and the standard deviations; see Table 4) for 70% and more deformed (plus recovered) g bands. Combination of higher stored energies and wider distributions (as in Table 4) may explain the preferred nucleation behaviour (i.e. Fig. 6) [7, 10]. However, this alone does not eliminate possibilities of micro-growth advantage or selection. The OIM observations at the early stages of nucleation (irrespective of reduction), however, did not show any signi®cant advantage for F-recrystallized grains forming next to (i.e. nucleating next to or selectively growing into) deformed I regions. This observation is quite similar to that in IF steel [7], and possibly indicates the predominant role of stored energy advantage [7, 10±12] in determining the preferred nucleation of g grains. Although a debate may still continue on the relative merits of stored energy advantage or micro-growth advantage/selection, the role of frequency advantage (as caused by the preferred nucleation) in establishing the g recrystallization texture seems quite convincing. 4.3. The physical parametersÐli and Ni As discussed in Section 1, the recrystallization texture components may depend on the two physical parametersÐli and Ni. Obviously this approach assumes that a, g and random grains nucleate from their respective deformed bands, an assumption supported by the present experimental evidence. As given in Table 3, spacings of the deformed bands (li) can be measured directly from the OIM scans. In another approach, spacings may be estimated from Ai/ti, where Ai is the area fraction of the texture component i and ti is the deformed band thickness. Average tis are listed in Table 2. The original hot-band grain size was 50 mm and hence after rolling (or a plane-strain deformation [13]) the expected ti=50 (1 ÿ Fd)/d, where Fd is the fraction deformation and d is the grain-splitting factor (a statistical estimation on the formation of high-angle boundaries [7, 13]). After 50% reduction d was 1, while d was approximately 2 at and above 70% reduction. In other words, statistically no new highangle grain boundaries developed at 50% reduction, while for reductions of 70% and more each grain, on average, split into 2. The appearance (i.e. for reductions of 70% and above) of more intense strain localizations in g-oriented regions may explain more splittings in g bands (and hence the fragmented/broken appearance), while the comparable tis (but wider spread) for a bands may be explained from a di€erent standpoint. Two types of a bands were observed: relatively thicker and thinner bands. While the thicker bands were possibly formed by the pancaking of the original grains, the thinner bands were possibly breakaway parts of an original grain (i.e. after pancaking of an original grain only

63

part of it may rotate to a). Note that such thin a bands (especially for the I component and for 80 and 90% reductions) were often closely spaced (i.e. only separated by 1/2 bands of other orientation) and the continued increase in a texture with strain, coincided with an increasing appearance of thin a bands. Ni values were estimated (see Table 6) for di€erent texture components (and for all the reductions) by using Ai=Nidi/li (the formulation is a purely geometrical one; see Refs [7, 15, 16]) where Ai is the area fraction (as estimated from the OIM scans). Note that the area and volume fractions (measured from the X-ray ODFs) were nearly similar. di is the average recrystallized grain size (measured along the ND) of the component i. Estimated Nis (see Table 6), do provide an index of nucleation as well as growth. Considering only nucleation, and that too from a purely geometric point of view, maximum Nis were generallized as 2 (for a thick deformed band, e.g. more than 4 subgrains thick) or 1 (for a thin deformed band, e.g. 1±4 subgrains thick) [15, 16]. Extending the same logic, expected maximum Ni for 90% reduction is 1, while for 70 and 80% reductions it is 2. However, a maximum Ni of >2 can even be possible if di>ti (average band thickness), as in 50% reduction (using the same geometric approach as in Refs [15, 16]). The relative roles of nucleation and growth on the changes in the numerical values of Nis possibly need to be made with respect to such maximum possible Ni values. At 50% reduction, Nis of a/g components were comparable, while Ni for random orientation was distinctly higher. Recrystallization behaviour at 50% reduction was quite di€erent from the rest of the reductions. Deformed bands were considerably thick (average ti=25 mm) and new grains were almost always observed to form/nucleate (i.e. at the early stages of recrystallization) from band interior strain localization [i.e. cluster of grain boundaries at an approximate angle of 378 with the RD; see Fig. 5(a)]. The nucleation from band interior strain localizations also coincided with a delay or incubation time and no preferred nucleation for g grains. After 70% reduction, strain localizations appeared somewhat preferentially on the g bands and increased stored energies for g-deformed regions were noted (see Table 4), along with preferred nucleation of g over a (see Fig. 6). Nis for F and E were higher than for I and H (see Table 6). Table 6. Estimated nucleation factors (Ni) for F/E/I/H/random deformed bands Reduction 50% 70% 82% 90%

F

E

I

H

Random

0.69 0.28 0.31 0.28

0.71 0.27 0.29 0.181

0.68 0.21 0.07 0.095

0.56 0.09 0.024 0.03

1.43 0.55 0.54 0.355

64

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL

Interestingly, however, Ni for I dropped signi®cantly above 70% reduction, which may partly be explained from increased stored energy advantage for F/E. There was another plausible cause for such a drop. Thin and closely spaced a bands (especially for the I component) were often observed after 80 and 90% reductions. Some of these bands were de®nitely consumed during recrystallizationÐfor example, average spacing of I bands after 90% reduction increased from 10 to 50 mm as recrystallization progressed to 70%. But the remaining (i.e. above 70% recrystallization) I bands were de®nitely coarser (2±4 times on average) and contained deformed as well as recrystallized regions (see Section 3.3.3). One plausible mechanism of the formation of such remaining bands is variant inhibition [16] or orientation pinning [27]. In other words, growth of recrystallized grains from closely spaced I bands may become restricted by the presence of recrystallized grains/deformed regions of similar orientations. This will de®nitely reduce the Ni (e.g. Ni for I)Ðat least in the regions with closely spaced I bands, an e€ect which may also re¯ect on the overall Ni value. Orientation pinning may also explain the coarsened/thickened a band structures and the changes in a texture with strain [7]. An interesting di€erence between IF and ULC steel is the di€erences in the numerical values of Nis. As given in Ref. [7], after 90% cold rolling, estimated values of NF and NE were 0.74 and 0.39, respectively, which are signi®cantly higher than the values of the present ULC with the same reduction (note that the NI and NH are rather similar). There are two possible explanations for such di€erences. Firstly, the presence of free carbon in ULC and the expected solute drag on the grain boundaries may explain the di€erences in recrystallization behaviour [2]. But another possiblity also exists in the framework of relative stored energy advantage. Relative stored energy advantage (both average y/d and its spread/distribution) for g against a is de®nitely larger for IF [7, 10] than for ULC steel (estimated for the same reduction). In IF steel strain localizations appear more preferentially on g bands than in ULC (strictly on the basis of semi-quantitative comparison between the present results and Ref. [13]). If appearance of strain localizations are responsible for the stored energy advantage of g, then these sets of results may also explain the higher NF and NE in IF steel [7]. Evidently, a more detailed understanding of the formation of these strain localizations is needed. The role of alloy chemistry, deformation conditions and microtexture on the formation of strain localizations need to be clari®ed. Physical data on microstructure and texture, as provided in the present study, can be ®tted in any recrystallization texture evolution model (e.g. Ref. [29]) to test the quantitative or semi-quantitative predictability of the ®nal recrystallization textures.

5. SUMMARY

During deformation, g-®bre increased from 0 to 50% reduction, but then did not change, while a®bre increased steadily with strain. During recrystallization, however, a steady increase in g with strain coincided with no signi®cant changes in a. Strain localizations, as characterized by more frequent grain boundaries of relatively higher misorientations, were observed at an approximate angle of 378 with the RD. Although the presence of such strain localizations were noted even at 50% reduction, at and above 70% reduction they appeared somewhat more preferentially on the g-oriented deformed regions. This coincided with (i) stored energy advantage for deformed g over deformed a and (ii) grain splitting (until 50% reduction, and as estimated statistically, no new high-angle boundaries were created, while at and above 70% reduction each grain, on average, was split into two). After 50% reduction, recrystallization was observed in band interior strain localizations. A delay or incubation time was involved and no preferred nucleation of g over a was observed. Above 50% reduction, the estimated stored energy advantages in deformed g bands also coincided with the preferred nucleation of g grains. Grains of the three main components (i.e. g, a and random) of the recrystallization texture nucleated from the respective deformed bands. Two physical parameters, li and Ni, are identi®ed to establish such recrystallization textures, where li is the spacing and Ni is the nucleation factor (de®ned as the number of i grains per i band)Ðboth measured/estimated along the ND. For all components li decreased with reduction, the drop being relatively stronger in a than in g. Nis for g and a were comparable at 50% reduction. For reductions of 70% and above, Nis for a dropped quite signi®cantly. Two probable causes of the drop were increased stored energy advantage for g and orientation pinning for some of the a grains. AcknowledgementsÐFinancial support from IUAP contract no. P4/33 and supply of the rolled materials from CRM are greatly appreciated. REFERENCES 1. Ito, K., International Forum on Physics of Metals in IF Steel. Iron and Steel Institute of Japan, Tokyo, 1994, p. 99. 2. Hutchinson, W. B. and Ryde, L., in Proc. 16th Risù Symposium on Materials Science, ed. N. Hansen, D. Juul Jensen, Y. L. Liu and B. Ralph. Risù National Laboratory, Roskilde, Denmark, 1995, p. 105. 3. Ray, R. K., Jonas, J. J. and Hook, R. E., Int. Mater. Rev., 1994, 39(4), 129. 4. Emren, F., von Schlippenbach, U. and LuÈcke, K., Acta metall., 1986, 34, 2105. 5. Hutchinson, W. B., Int. Mater. Rev., 1984, 29(1), 25. 6. Kestens, L. and Jonas, J. J., Metall. Mater. Trans., 1996, 27A, 155.

SAMAJDAR et al.: RECRYSTALLIZATION TEXTURES IN ULC STEEL 7. Samajdar, I., Verlinden, B. and Van Houtte, P., Acta mater., 1998, 46(8), 2751. 8. Hayakawa, Y. and Szpunar, J. A., Acta mater., 1997, 45(6), 2425. 9. Park, Y. B., Lee, D. N. and Gottstein, G., in Proc. 16th Risù Symposium on Materials Science, ed. N. Hansen, D. Juul Jensen, Y. L. Liu and B. Ralph. Risù National Laboratory, Roskilde, Denmark, 1995, p. 479. 10. Samajdar, I., Verlinden, B., Van Houtte, P. and Vanderschueren, D., Mater. Sci. Engng, 1997, A238, 343. 11. Dillamore, I. L., Smith, C. J. E. and Watson, T. W., Metal Sci. J., 1967, 1, 49. 12. Rajmohan, N., Hayakawa, Y., Szpunar, J. A. and Root, J. H., Acta mater., 1997, 45(6), 2485. 13. Samajdar, I., Verlinden, B. and Van Houtte, P., Iron Steel Inst. Japan Int., 1998, 38(7), 759. 14. Hughes, D. A. and Hansen, N., Acta mater., 1997, 45(9), 387. 15. Samajdar, I. and Doherty, R. D., Acta mater., 1998, 46(9), 3145. 16. Doherty, R. D., Samajdar, I., Necker, C. T., Vatne H. E. and Nes, E., in Proc. 16th Risù Symposium on Materials Science, ed. N. Hansen, D. Juul Jensen, Y. L. Liu and B. Ralph. Risù National Laboratory, Roskilde, Denmark, 1995, p. 1.

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