Scripta Materialia 57 (2007) 257–260

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Attaining deep drawability and non-earing properties in Ti + Nbinterstitial-free steels through double cold rolling and annealing

Rajib Saha,a,b,* R.K. Raya,b and D. Bhattacharjeeb

aIndian Institute of Technology, Kanpur 208016, IndiabR&D Division, Tata Steel Jamshedpur, Jharkhand 831001, India

Received 13 December 2006; revised 7 February 2007; accepted 30 March 2007Available online 9 May 2007

The intensity and uniformity of the c-fibre improve very significantly in a Ti + Nb interstitial-free steel after the cold rolled andannealed sample is subjected to a second cold rolling and annealing treatment. A uniform grain size as well as a balance between thetwo major components of the c-fibre, namely {111}h112i and {111}h1 10i, are expected to make the material highly deep-drawableand free from earing during processing.� 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Interstitial-free (IF) steels; Texture; Microstructure; Grain boundary character distribution; Double cold rolling and double annealing(DRDA)

Double reduction (second cold rolling) is generallyapplied for the production of thin steel sheets for light-weight packaging applications such as tin plates. Thisdouble reduction enables down-gauging without sub-stantial loss of performance of the steel [1]. This treat-ment imparts additional strength to the steel but at thesame time makes it less ductile as compared with the sin-gle reduced material. This may impose some limitationsto the can manufacturing process. For the production ofcomplicated and intricately shaped cans and for someother applications, the material sometimes needs to havea substantial amount of formability as well as strength.This may be achieved by applying an annealing treat-ment after the second cold rolling. In the present work,double cold rolling and double annealing treatment wereapplied to a Ti + Nb interstitial-free (IF) steel to pro-duce thin sheet for sophisticated packaging applications.The materials after different treatments were character-ized by their microstructure, texture and grain boundarycharacter distribution.

The steel, melted as an industrial heat, has the chem-ical composition: 0.003% C, 0.36% Mn, 0.006% S,0.014% P, 0.012% Si, 0.047% Al, 0.0033% N, 0.011%Nb, 0.05% Ti (in wt.%). A transfer bar of this steelwas controlled hot rolled by �80% in several passes.

1359-6462/$ - see front matter � 2007 Acta Materialia Inc. Published by Eldoi:10.1016/j.scriptamat.2007.03.055

* Corresponding author. Address: R&D Division, Tata Steel Jam-shedpur, Jharkhand 831001, India. Tel.: +91 0657 2147445; e-mail:rajib.saha@tatasteel.com

The finish rolling temperature was kept to within±10 �C of 964 �C (within the austenite recrystallizationrange). The hot bands were then subjected to 90% coldrolling followed by batch annealing at 720 �C for 27 h.The annealed sample was further cold rolled by anamount �65% and subjected to second annealing at650 �C for 90 min. Crystallographic textures were deter-mined using electron backscattered diffraction (EBSD)patterns, in a FEI – Quanta 200 scanning electronmicroscope, from the mid-thickness regions of the coldrolled and annealed sheets. During data collection aminimum number of seven Kikuchi bands were allowedto index a pattern, thus ensuring good reliability ofinformation. It has been reported by Field [2] thatEBSD patterns having a confidence index above 0.1can correctly index an orientation 95% of the time. Inthe present experiments the average confidence indexof the patterns for annealed sample was 0.7, and thisled to more than 98% of the patterns being indexed cor-rectly in every case. The minimum angular resolution ofthe EBSD from alpha iron using W-filament has beenfound to be 1�, as reported by Humphreys [3]. In orderto make accurate measurements, misorientations of lessthan 1.5� were excluded from the data. Orientation dis-tribution functions (ODFs) were measured using TSL-OIM software and U2 = 45� sections (Bunge notation)were determined therefrom. The ODFs were calculatedwith 15� Gaussian spread around the ideal orienta-tions. Using the same software, the grain boundary

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Figure 2. Image quality maps of the (a) first and (b) second annealedsteel.

258 R. Saha et al. / Scripta Materialia 57 (2007) 257–260

distribution character was determined for all the steelsamples. The Brandon criterion was used while report-ing the fractions of low-angle grain boundaries(LAGBs), high-angle grain boundaries (HAGBs) andcoincidence site lattice (CSL) boundaries. Boundarieswith a grain boundary angle of less than 15� wereclubbed together as LAGBs and those with a grainboundary angle greater than 15� were included in theHAGBs. CSL boundaries include those in the range ofR3–R29 as per Brandon’s criterion.

The U2 = 45� sections of the first cold rolled, first an-nealed, second cold rolled and second annealed samplesare presented in Figure 1a–d. The texture of the firstcold rolled steel consists of the c- and a-fibres along withthe rotated cube component {100}h01 1i. The firstannealing treatment leads to a sharpening of the c-fibreand the complete disappearance of the rotated cube.However, the c-fibre is not completely uniform in thiscase and shows maxima at {111}h112i locations. Thesecond cold rolling leads to an increase in the intensityof the c-fibre, but the fibre is quite non-uniform, as be-fore. A drastic change in texture appears after secondcold annealing when a very sharp and uniform c-fibreis found to develop in this material, but practically noa-fibre or any other component. Figure 2a and b showsthe image quality maps of the first and second annealedsamples, respectively. Although both the samples are inan annealed condition and show recrystallized grainsonly, the second annealed sample is made up of largeequiaxed grains of almost uniform size, whereas the firstannealed sample shows a duplex grain structure, consist-ing of both large and small-sized grains.

Figure 1. ODF U2 = 45� sections of the: (a) first cold rolled, (b) first anneal

Two major texture components of the c-fibre, namely{111}h11 2i and {111}h110i, were next considered andcrystal orientation maps for these two orientations onlywere obtained for the first cold rolled and first annealedas well as the second cold rolled and second annealedsamples. The orientations were calculated with a devia-tion within 15�. These are represented in Figure 3a–d,respectively. It is at once clear from Figure 3 thatwhereas, in the first annealed sample, the grains with{111}h11 2i orientation occupy 31.4% volume fractionof the total grains, the volume fraction of {11 1}h110ioriented grains is only 18.3%. On the other hand, inthe second annealed sample the volume fractions of{111}h11 2i and {111}h11 0i oriented grains are34.7% and 38.4%, respectively, i.e. they are nearly equal.

ed, (c) second cold rolled and (d) second annealed steel.

Figure 3. Crystal orientation maps of the steels after differenttreatment: (a) first cold rolled, (b) first annealed, (c) second coldrolled and (d) second annealed steel.

R. Saha et al. / Scripta Materialia 57 (2007) 257–260 259

A similar behaviour regarding these two componentswas also observed in the first cold rolled and second coldrolled samples.

The grain boundary character distribution (GBCD)of the first cold rolled, first annealed, second cold rolledand second annealed samples are depicted in Figure4a–d, respectively. As expected, after first cold rollingthe majority (�0.55) of the boundaries are LAGBs,presumably because of the formation of cells and sub-grains [4], and due to the development of a strong tex-ture [5–7]. A substantial HAGB fraction also formsdue to the accumulation of dislocations [8,9]. The

Figure 4. GBCD of the steels after different treatment: (a) first coldrolled, (b) first annealed, (c) second cold rolled and (d) secondannealed steel.

HAGB fraction increases drastically from 0.37 to 0.61after annealing in the first annealed sample, due to theformation of recrystallized grains [10,11]. The CSLboundaries remain very similar in both samples. Aremarkable change in GBCD appears after the first an-nealed sample is second cold rolled (Fig. 4c). Here, inthe second cold rolled sample, the LAGB fraction isnearly 0.87, whereas the HAGB fraction decreases to avery low value of nearly 0.10. The CSL fraction also de-creases to a very low value. Presumably the lower levelof cold deformation (�65%) during the second cold roll-ing does not lead to much dislocation accumulation, dueto which the HAGB fraction decreases [8,9]. However,the much sharper c-fibre in the second cold rolled sam-ple as compared with the first cold rolling is also respon-sible for the much higher LAGB fraction in the former.The EBSD image quality value (�47) for the first coldrolled sample is less than that of the second cold rolledsample (�70). This indicates a larger strain accumula-tion in the former as compared with the latter[12,3,13–16]. Finally, the second annealed sample showsa GBCD which is not much different from that of thefirst annealed sample.

It would be tempting to correlate texture with micro-structure and grain boundary character distribution.Both the first and second annealed samples are charac-terized by nearly similar GBCD, although their texturesare very different. The texture of the second annealedsample is much stronger and shows a more uniformc-fibre as compared with the first annealed sample.The similarity of the corresponding GBCDs are there-fore rather surprising. There is, however, a significantdifference between the GBCDs of the first and secondcold rolled samples. A rather strong c-fibre texture inthe second cold rolled sample has ensured that theLAGB fraction is very high (>0.85) while the HAGBand CSL fractions are insignificant. Recrystallizationannealing of such a material has yielded a very sharpand uniform c-fibre texture. It appears that a cold rolledmaterial showing a very high LAGB fraction is condu-cive to the development of a very sharp and uniformc-fibre after recrystallization. The second annealed sam-ple with a sharp and uniform c-fibre is also character-ized by a nearly uniform grain size, as shown inFigure 2b. The uniformity of the c-fibre in the secondannealed sample (with nearly equal volume fractionsof the {111}h112i and {111}h110i components) alsoensures that the problem of earing during processing willbe minimal as compared with the first annealed sample[17–20]. The strength of the c-fibre in the second an-nealed sample is also expected to impart a remarkabledeep drawing behaviour in this material.

In conclusion, a moderate amount of second coldrolling followed by annealing appears to improve thestrength and uniformity of the annealed steel. The c-fibre becomes stronger and more uniform after a secondcold rolling and annealing treatment. A uniform c-fibreis associated with a uniform grain size distribution. Agrain boundary character distribution consisting of avery high low-angle grain boundary fraction, as pro-duced after second cold rolling, appears to be highlybeneficial for the production of a very sharp anduniform c-fibre after subsequent annealing.

260 R. Saha et al. / Scripta Materialia 57 (2007) 257–260

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