Effect of rolling temperature on the deformation and recrystallization ...

Effect of Rolling Temperature on the Deformation and Recrystallization Textures of Warm-Rolled Steels MOHAMMAD R. TOROGHINEJAD, ALAN O. HUMPHREYS, DONGSHENG LIU, FAKHRADDIN ASHRAFIZADEH, ABBAS NAJAFIZADEH, and JOHN J. JONAS Warm-rolling trials were carried out on three interstitial-free (IF) steels (stabilized with either niobium or titanium), an extralow-carbon (ELC) steel, and an experimental low-carbon chromium (LC Cr) material at temperatures between 440 °C and 850 °C. The influence of rolling temperature on their as-rolled microstructures and deformation and recrystallization textures was investigated. Also, the effect of coiling simulation and degree of rolling reduction on the r values of some of these materials was examined. In-grain shear bands were evident in all as-rolled microstructures, but their sensitivity to deformation temperature varied between steels. Shear bands of moderate intensity were formed in the IF steels across all temperatures. In the ELC material, intense shear bands were formed at low rolling temperatures, but at higher temperatures, this intensity was drastically reduced. The development of shear bands at the higher rolling temperatures was significantly enhanced by alloying with chromium. The deformation textures produced were typical of rolled ferrite materials. The intensity of this texture increased markedly with temperature for the ELC grade. Conversely, the intensity of the recrystallization texture decreased with increasing temperature. The addition of chromium was found to strengthen the {111} component and, hence, the formability. The sharpness of both the deformation and recrystallization textures of the IF steels was relatively unaffected by rolling temperature. These differences are attributed to the intensity and frequency of shear-band formation and the dynamic strain-aging (DSA) behaviors of the various materials.

I. INTRODUCTION

STEELMAKERS have begun to take an interest in warm (ferritic) rolling, as it has the potential to broaden the product range and decrease the cost of hot-rolled strip materials.[1,2] However, rolling in the ferrite phase (for steels containing carbon or nitrogen in solution) can lead to significant dynamic strain aging (DSA), which, in turn, influences their strain-rate sensitivity (m). In-grain shear bands constitute a form of localized flow within a grain, which occurs at a higher strain rate than that of the bulk material. Therefore, their formation is highly influenced by the strain-rate sensitivity. A negative rate sensitivity will enhance the formation of shear bands and, conversely, an unusually high rate sensitivity can suppress their formation within the grain.[3] The formation of shear bands during warm rolling is generally necessary for the development of strong {111} recrystallization textures, which ensure high mean r values and, hence, good formability.[4] It has been shown that interstitial-free (IF) steels have small positive rate sensitivities across a wide range of temperatures.[5] By contrast, low-carbon (LC) steels can exhibit both negative and high-positive rate sensitivities due to DSA effects. Thus, warm-rolled LC steels tend to have weak {111} annealed textures, which are responsible for the poor formabilities of final products.

Previous workers have shown that alloying additions, such as chromium, can modify the behavior of LC steels by extending DSA effects into the warm-rolling temperature region (650 °C to 800 °C).[6,7] This creates a plateau in the flow stress vs temperature plot rather than the sharp peak usually associated with conventional carbon-nitrogen DSA behavior. A small increase in strain rate does not, therefore, result in a large increase in flow stress within this plateau region, i.e., the rate sensitivity of the material is reduced. It has been found that warm-rolled and annealed LC steels containing 1.3 pct chromium have strong normal direction (ND) fiber-texture components, leading to improved formability.[5] However, for further economic benefit, the possibility of using lower levels of chromium for warmrolled LC steels was studied in the present work. Since the deformation microstructure (particularly, the in-grain shear bands) and the deformation texture influence the orientations of the recrystallized grains,[3,4] the effect of rolling temperature on the development of shear bands and of the deformation and recrystallization textures was studied for various grades of IF and LC steels. Furthermore, the effects of rolling reduction and postrolling aging treatment upon the recrystallization texture were also investigated.

II. EXPERIMENTAL PROCEDURE MOHAMMAD R. TOROGHINEJAD, PhD Student, FAKHRADDIN ASHRAFIZADEH, Professor, and ABBAS NAJAFIZADEH, Professor, are with the Department of Materials Engineering, Isfahan University of Technology, Isfahan 84154, Iran. ALAN O. HUMPHREYS, Research Associate, DONGSHENG LIU, Research Associate, and JOHN J. JONAS, Professor, are with the Department of Metallurgical Engineering, McGill University, Montreal, PQ, Canada H3A 2B2. Contact e-mail: [email protected] Manuscript submitted August 26, 2002. METALLURGICAL AND MATERIALS TRANSACTIONS A

The chemistries of the five steels investigated in the current project are listed in Table I. These included three IF steels: two stabilized with niobium and titanium, and one partially stabilized with a limited niobium addition. An extralow-carbon (ELC) steel was included for comparison. These were all regular production alloys received in the VOLUME 34A, MAY 2003—1163

Table I. Steels Compositions (Weight Percent) Steel IF Nb IF Ti Unstabilized IF Nb ELC LC Cr

C

Mn

P

S

Si

Cr

Al

N

Ti

Nb

0.002 0.004 0.002 0.020 0.037

0.24 0.15 0.47 0.12 0.35

0.060 0.005 0.032 0.004 0.010

0.011 0.010 0.008 0.007 0.010

0.011 0.006 0.009 0.006 0.020

0.020 0.065 0.019 0.071 0.480

0.065 0.041 0.022 0.048 0.036

0.0037 0.0037 0.0039 0.0067 0.0012

0 0.062 0 0 0

0.041 0 0.009 0 0

form of 100-mm-thick transfer bars from the roughing mill. An experimental low-carbon chromium (LC Cr) alloy (cast and forged into a 50-mm-thick slab) was manufactured for the project. This was made from electrolytic iron to minimize impurity levels of the residual elements and, hence, to remove any superfluous dynamic strain-aging (DSA) effects. Rolling trials were carried out using the pilot mill at CANMET (Ottawa, Canada) at an equivalent strain rate of ⬃30 s⫺1. A soluble oil lubricant was applied to the surface of the rolls to minimize friction along the roll gap, as it is well known that friction can produce a texture gradient through the product thickness.[8] In order to prepare the test billets for the warm-rolling trials, each was hot rolled down to a thickness of 11 mm, with the final rolling temperature always exceeding 960 °C to ensure austenite recrystallization, thus ensuring a uniform final texture. Strips of each steel (width of 110 mm) were reheated in a furnace at 1050 °C for 30 minutes. The specimen temperature was monitored using a thermocouple inserted into the center of each strip at its midplane. After austenitization, the steels were air cooled down to the appropriate rolling temperature on a ceramic slab. This heat treatment resulted in the ferrite grain sizes shown in Table II. Single-pass warm-rolling experiments were conducted using a reduction of 65 pct at temperatures of 440 °C, 640 °C, 710 °C, 780 °C, and 850 °C. After rolling, the sheets were immediately water quenched to prevent further static recrystallization and, hence, to preserve the as-rolled microstructure. To investigate the effects of further reduction, the ELC and LC Cr specimens warm rolled at 640 °C (65 pct reduction) were additionally cold rolled by 40 pct (a total 80 pct reduction). Select grades were also warm rolled using two passes, with a 40 pct reduction at 700 °C in the first pass and a 65 pct reduction at 650 °C in the second (a total 80 pct reduction). Specimens for analysis were cut from the midwidth regions of the as-rolled strips so as to avoid possible edge effects from the rolling process. One sample from each condition was completely recrystallized (determined metallographically) using a muffle furnace at 700 °C to 750 °C and annealing times of 1 to 12 hours. In order to investigate the effect of coiling treatment on formability, one sample of each grade (rolled at 640 °C) was aged prior to annealing. The IF steels were aged for 3 hours at 400 °C, and the ELC and LC Cr steels were aged for 6 hours at 450 °C. This aging step was used to simulate the slow cooling of the coiled sheet, as opposed to the rapid quench performed on most of the samples in this study. Textures were determined by measuring three incomplete pole figures ({200}, {110}, and {112}) along the midplanes of the samples using a Siemens D500 goniometer system 1164—VOLUME 34A, MAY 2003

Table II. Ferrite Grain Size (mm) before Rolling Reheat Temperature

IF Nb

IF Ti

1050°C

58

ⱖ300

Unstabilized IF Nb ELC 66

34

LC Cr 54

Fig. 1—␸2 ⫽ 45 deg plot showing the locations of the main rolling (RD or a) and recrystallization (ND or g) texture components.

(Mo tube) and the standard reflection technique. From these pole-figure data, the orientation distribution functions were calculated using Bunge’s series-expansion method,[9] with an expansion degree of 1max ⫽ 22.[10] Because of the cubic crystal and orthorhombic sample symmetries, the textures are presented in the reduced 0 deg ⱕ w1, ⌽, w2 ⱕ 90 deg Euler space. The important components of the rolling texture of LC steels are located along three orientation lines: the a or rolling direction (RD) fiber ( 储具110典), the g or ND fiber ( 储具111典), and the ␧ or transverse direction (TD) fiber (储具110典).[11] All the ideal orientations belonging to the aforementioned fibers can be found in the w2 ⫽ 45 deg section of Euler space, and these are shown in Figure 1. For the skeleton-line plots, there are only two ideal orientations along the ND fiber: {111}具110典 and {111}具112典. As a result, the range w1 ⫽ 60 to 90 deg is sufficient to represent the complete fiber in the latter case.[12] Estimated rm values (mean Lankford values) for these samples were calculated on the basis of the texture data using software developed at McGill University.[13] These r values are presented as functions of the angle to the RD for all the warm-rolled samples. In accordance with previous work,[14] the calculations were made by using the relaxedconstraint method of crystal plasticity, where the critical METALLURGICAL AND MATERIALS TRANSACTIONS A

resolved shear-stress ratio (a) for glide on the {112}具111典 and {110}具111典 slip systems (t{112}/t{110}) is 0.95. The relaxed-constraint approach was selected, based upon the observation that the grains in the present steels are of a pancake shape. III. RESULTS A. Deformation Microstructures The deformation microstructures were composed of ferrite grains elongated along the RD, some of which contained

in-grain shear bands, as shown for a typical warm-rolling temperature of 640 °C in Figure 2. It can be seen that for the IF Nb and IF Ti steels, moderately intense shear bands were formed (Figures 2(a) and (b), respectively), but in the unstabilized IF Nb material, the shear bands were less-well formed (Figure 2(c)). However, in the ELC steel, low densities and intensities of shear bands were produced (Figure 2(d)). The shear bands in the LC Cr steel were less intense than those formed in the IF steels, but were more densely populated than those of the ELC material (Figure 2(e)). The shear-band content was quantified using a pointcounting technique, and the percentage of material contain-

Fig. 2—Typical deformation microstructures produced by rolling at 640 °C: (a) IF Nb, (b) IF Ti, (c) unstabilized IF Nb, (d) ELC, and (e) LC Cr steel. METALLURGICAL AND MATERIALS TRANSACTIONS A

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Fig. 3—Dependence of shear band content on rolling temperature.

ing in-grain shear bands at the midplane of each strip was determined. The dependence of the fraction of banded material on rolling temperature is summarized in Figure 3. At the higher rolling temperatures of 780 °C and 850 °C, static recrystallization of the deformed ferrite frequently occurred prior to quenching, leading to a final microstructure of strain-free ferrite without any trace of shear bands. It can be seen from Figure 3 that the population of shear bands in all three IF steels decreased slightly as the rolling temperature was increased. However, in the ELC steel, the shear-band content displayed a strong dependence on rolling temperature, with small fractions at temperatures of 640 °C and above. The LC Cr steel displayed a trend similar to that of the ELC material; however, the chromium addition appeared to increase the overall population of shear bands. The frequency of in-grain shear banding in the LC Cr steel after two-pass warm rolling was increased to about 60 pct (in contrast to 50 pct after single-pass rolling at similar temperatures). By contrast, this additional deformation did not significantly change the shear-band content of the ELC steel. However, cold rolling after warm rolling caused more severe shear banding in the ELC and LC Cr steels. B. Deformation Textures Typical deformation textures after rolling at 640 °C are shown in Figure 4. These were all typical of ferrite rolling, characterized by a partial RD fiber (from {001}具110典 to {111}具110典) as well as a complete ND fiber. Although these were qualitatively similar, higher intensities were present in the ELC and LC Cr steels. The RD, ND, and TD fibers of the unstabilized IF Nb and ELC steels after rolling at various temperatures are shown in Figures 5 and 6, respectively, in the form of skeleton lines. Given the similarities between the rolling textures, only two sets of skeleton lines are presented here. It is observed in Figure 5 that the rotated cube ({001}具110典), ND fiber, and {554}具225典 components were dominant at all rolling temperatures in the unstabilized IF 1166—VOLUME 34A, MAY 2003

Fig. 4—Deformation textures of the (a) IF Nb, (b) IF Ti, (c) unstabilized IF Nb, (d ) ELC, and (e) LC Cr steels after warm rolling at 640 °C (intensity levels 2, 3, 4, . . .).

Nb sample and that the texture sharpness remained relatively unchanged with increasing temperature. The highest intensity in this sample was less than 10 times random. By contrast, the rolling temperature had a significant effect on texture intensity in the ELC sample. The maximum texture intensity in this material increased with rolling temperature to about 20 times random (at 710 °C, Figure 6). The dominant component was the rotated cube at all rolling temperatures, and there were secondary maxima at other locations on the RD, ND, and TD fibers. The LC Cr steel had textures similar to those of the ELC material, but the maximum intensities were lower after rolling at 710 °C and 640 °C (14 and 11 times random, respectively). C. Recrystallization Textures To achieve complete recrystallization, the IF Ti, IF Nb, and unstabilized IF Nb steels were annealed at 750 °C for 1, 2, and 5 hours, respectively. The ELC steel was annealed for 1 hour at 700 °C, and the LC Cr grade was annealed for 12 hours at 725 °C. (The Ac1 temperature of the LC Cr METALLURGICAL AND MATERIALS TRANSACTIONS A

(a)

(a)

(b)

(b)

(c)

(c)

Fig. 5—(a) RD, (b) ND, and (c) TD fiber representations of the influence of rolling temperature on the deformation texture of the unstabilized IF Nb steel.

Fig. 6—(a) RD, (b) ND, and (c) TD fiber representations of the influence of rolling temperature on the deformation texture of the ELC steel.

deformed steel was determined by dilatometry to be around 740 °C). The recrystallization textures of the five steels after warm rolling at 640 °C are illustrated in Figure 7. The IF Nb and IF Ti steel textures (Figures 7(a) and (b)) contained TD fibers with maximum intensities of 7 and 8 times random,

respectively, at the {554}具225典 component. The recrystallization textures of the unstabilized IF Nb and ELC steels are shown in Figures 7(c) and (d), respectively. The former exhibits its maximum intensity at {554}具225典, while that of the latter is located near {112}具110典. By contrast, the recrystallization texture of the LC Cr sample

METALLURGICAL AND MATERIALS TRANSACTIONS A


(a)

(b) Fig. 7—Recrystallization textures of the (a) IF Nb, (b) IF Ti, (c) unstabilized IF Nb, (d ) ELC, and (e) LC Cr steels after warm rolling at 640 °C (intensity levels 2, 3, 4, . . .).

contained a partial RD fiber and a discontinuous ND fiber (Figure 7(e)). The skeleton lines pertaining to the RD, ND, and TD fibers of the recrystallized steels after rolling at various temperatures are shown in Figures 8 through 12. It can be seen from Figure 8 that the IF Nb steel rolled at 850 °C and annealed at 750 °C exhibited peaks at the {554}具225典, {111}具110典, and rotated-cube ({001}具110典) positions. It also shows that the intensity of the rotated-cube component increased with rolling temperature, although the intensities of the other components remained approximately constant. The recrystallization texture of the IF Ti steel after rolling at 850 °C was weak and, in part, close to random (Figure 9). It can be seen that as the rolling temperature was increased, the intensities of the {554}具225典, {111}具112典, and {111}具110典 components decreased. The Goss component ({110]具001典) was present at all temperatures and, after rolling at 640 °C, the rotated Goss component ({110}具011典) was also observed. The texture intensities associated with the unstabilized IF Nb steel are illustrated in Figure 10. These are similar to 1168—VOLUME 34A, MAY 2003

(c) Fig. 8—(a) RD, (b) ND, and (c) TD fiber representations of the influence of rolling temperature on the recrystallization texture of the IF Nb steel.

those observed in the IF Nb steel (Figure 8), but with lower intensities. The Goss component decreased in strength with increasing rolling temperature. The influence of rolling temperature was greater for the ELC than for the IF steels. At a rolling temperature of 850 °C, METALLURGICAL AND MATERIALS TRANSACTIONS A

(a)

(a)

(b)

(b)

(c)

(c)

Fig. 9—(a) RD, (b) ND, and (c) TD fiber representations of the influence of rolling temperature on the recrystallization texture of the IF Ti steel.

Fig. 10—(a) RD, (b) ND, and (c) TD fiber representations of the influence of rolling temperature on the recrystallization texture of the unstabilized IF Nb steel.

the rotated-cube component was dominant (Figure 11) and its intensity decreased with decreasing temperature. After rolling at 640 °C, the {112}具110典 component of the RD fiber was dominant. Finally, the highest intensity after rolling at 440 °C was associated with the Goss component.

The texture intensities of the LC Cr material are illustrated in Figure 12. After rolling at 640 °C, the near-{113}具110典 and near-{111}具110典 components dominated. The highestintensity component after rolling at 440 °C was the Goss, with the ND fiber components of secondary importance.



(a)

(a)

(b)

(b)

(c)

(c)

Fig. 11—(a) RD, (b) ND, and (c) TD fiber representations of the influence of rolling temperature on the recrystallization texture of the ELC steel.

Fig. 12—(a) RD, (b) ND, and (c) TD fiber representations of the influence of rolling temperature on the recrystallization texture of the LC Cr steel.

D. Coiling Simulation

dom, Figure 13(a)) than its intensity without aging (7 times random, Figure 7(a)). The LC Cr specimen (Figure 13(e)) contained a partial RD fiber (13 times random) and a continuous ND fiber (rather than the incomplete ND fiber formed without the aging treatment, Figure 7(e)). By contrast, the coiling simulation had a negligible effect on the IF Ti,

By comparison with Figure 7, it is shown in Figure 13 that the aging treatment had a considerable influence on the recrystallization textures of the IF Nb and LC Cr samples. That of the IF Nb steel was sharper (11 times ran1170—VOLUME 34A, MAY 2003


Fig. 14 —Recrystallization textures of the (a) ELC and (b) LC Cr steels after two-pass warm rolling and aging at 450 °C (intensity levels 2, 3, 4, . . .).

Fig. 15—Recrystallization textures of the (a) ELC and (b) LC Cr steels after warm and cold rolling (intensity levels 2, 3, 4, . . .).

Fig. 13—Recrystallization textures of the (a) IF Nb, (b) IF Ti, (c) unstabilized IF Nb, (d ) ELC, and (e) LC Cr steels after warm rolling at 640 °C and aging (intensity levels 2, 3, 4, . . .).

unstabilized IF Nb, and ELC steel textures (Figures 13(b) through (d), respectively), as will be discussed in more detail subsequently. E. Effect of Strain The effect of increased reduction on the recrystallization texture is shown in Figure 14. By comparison with Figure 11, it can be seen that for the ELC grade, the two-pass texture was similar to that produced by single-pass warm rolling above 700 °C; i.e., the texture contains partial RD and ND fibers with a maximum intensity of 10 times random (Figure 14(a)). By contrast, the recrystallization texture of the LC Cr steel after two-pass warm rolling contains a partial RD fiber and a complete ND fiber (Figure 14(b)); this differs completely from that produced by single-pass rolling above 640 °C (Figure 12). The textures of the ELC and LC Cr samples after warm followed by cold rolling are shown in Figures 15(a) and (b), respectively. These are characterized by partial RD and comMETALLURGICAL AND MATERIALS TRANSACTIONS A

plete ND fibers, but of low intensity. Furthermore, the Goss component was observed in the LC Cr grade, an observation that is discussed in more detail subsequently. F. Calculated r Values Mean r (rm) and planar anisotropy (⌬r) values calculated from the recrystallization textures are summarized in Table III. The IF Nb and the IF Ti steels rolled between 440 °C and 710 °C had rm values in the range of 1.2 to 1.4. The ⌬r values for the IF Nb and IF Ti steels lay between 0.55 to 0.60 and 0.75 to 1.0, respectively, at similar temperatures. By contrast, the mean r values for the other grades fell below 1. However, alloying with chromium caused a slight increase in rm from 0.51 to 0.86 after rolling at 640 °C. IV. DISCUSSION A. Deformation Microstructure It is evident from Figure 2 that at the lowest rolling temperature of 440 °C, the degree of shear banding was greater in the ELC and LC Cr steels than in the IF materials. By contrast, in the warm-rolling range, i.e., above 600 °C, the amount of banding in the ELC and LC Cr steels dropped, VOLUME 34A, MAY 2003—1171

Table III. Calculated Mean r and ⌬r Values IF Nb Steel

IF Ti Steel

Unstabilized IF Nb Steel

ELC Steel

LC Cr Steel

Rolling Temperature (°C)

rm

⌬r

rm

⌬r

rm

⌬r

rm

⌬r

rm

⌬r

850 710 640 440 Aged and annealed rolled Two-pass warm rolling Warm and cold rolled

1.0 1.2 1.3 1.3 1.4 — —

0.1 0.6 0.6 0.6 0.3 — —

0.6 1.3 1.3 1.2 1.4 — —

0.5 1.0 0.8 0.8 0.8 — —

0.8 0.7 0.7 1.0 0.9 — —

⫺0.2 0.3 0.3 0.6 0.5 — —

0.3 0.4 0.5 0.9 0.5 0.3 1.0

⫺0.1 0.0 0.1 0.6 0.1 ⫺0.1 0.2

— 0.7 0.9 0.8 0.6 0.5 0.9

— 0.1 ⫺0.1 ⫺0.5 ⫺0.1 0.2 0.5

whereas the proportion of grains containing bands remained nearly constant in the IF materials. It is of interest that the nature and morphology of the bands was also different in the two types of steels. The shear-band morphology of these steels has been studied in detail.[15] The present investigation has shown that moderately intense shear bands form in the IF Nb- and Ti-stabilized steels, regardless of rolling temperature. Furthermore, some unexpectedly heavy shear bands were formed in the IF Ti steel; this can be attributed to the rather large initial grain size (more than 300 mm), as coarse initial grain sizes have been shown to promote heavy shear-band formation in cold-rolled steels.[16,17,18] In the ELC and LC Cr steels, shear bands of low intensity were produced at intermediate rolling temperatures of 640 °C and 710 °C. After rolling at 780 °C, these steels recrystallized prior to quenching and, so, had lost their shear bands. Intense shear bands were always produced in both steels at a rolling temperature of 440 °C. This dependence of shear-banding behavior on rolling temperature can be linked directly to the different DSA behaviors of these materials, as has been discussed extensively by Barnett and Jonas[2] and Jonas.[19] They stressed that DSA alters the rate sensitivity of the material, which, in turn, affects the ease with which rapid flow localization, i.e., shear banding, can occur. The DSA characteristics of the steels in this investigation have been studied previously[20] and are summarized in Figure 16.

Fig. 16—Dependence of measured strain rate sensitivity (m) on temperature[19] (measured at a mean strain rate of 10⫺2 s⫺2).

near-{111}具112典 annealing textures.[21] This may arise because grains with {112–001}具110典 orientations promote the formation of in-grain shear bands in adjacent {111} (ND fiber) grains. These in-grain shear bands, in turn, lead to the enhanced nucleation of {111}具112典-type orientations.[21]

B. Deformation Textures In addition to the favorable {111} component, the deformation textures also included a large number of grains of the detrimental {001} orientation. Another important feature of the current texture results is the marked increase in sharpness of the rolling components with temperature in the ELC and LC Cr steels. This effect can be rationalized in terms of the characteristics of in-grain shear-band formation.[3] The latter weaken the deformation texture by (1) causing grain fragmentation and (2) transferring the strain away from the matrix, such that it does not rotate as rapidly as under conditions of homogeneous deformation. In the absence of ingrain shear bands, therefore, the rolling texture is expected to be sharper, as observed in the ELC grade rolled at higher temperatures.[3] Furthermore, the lower intensity of the deformation texture in the LC Cr steel, compared to the ELC grade, after rolling at 640 °C and above is due to the presence of more in-grain shear bands in this steel. It has long been known that high RD fiber intensities in the deformation textures of rolled steels correlate with sharp 1172—VOLUME 34A, MAY 2003

C. Recrystallization Textures It is evident from Figures 8 through 12 that the recrystallization textures of samples after rolling at 850 °C and annealing were weak, except for that of the IF Nb steel. At this relatively high temperature, rolling occurs in the twophase region and, after processing, the austenite deformation texture components transform into their equivalent ferrite texture components. These are known to be detrimental to the deep-drawing properties of the material.[22] However, the IF Nb steel has a higher transformation temperature than the other alloys; hence, a greater fraction of ferrite will be present during rolling. Furthermore, in this steel, ferrite recrystallization is retarded by both Nb in solid solution and the presence of a fine dispersion of NbC precipitates.[23] Thus, in this alloy, there will be greater strain accumulation during rolling and, hence, more shear bands present to nucleate a strong {111} texture. At intermediate rolling temperatures of 710 °C and 640 °C, higher intensities of the ND fiber were present in METALLURGICAL AND MATERIALS TRANSACTIONS A

all samples, as compared with the case of 850 °C rolling. After rolling at 440 °C, the {111}具112典 and {554}具225典 intensities of the ELC steel were increased slightly compared to the trends observed at other temperatures. As indicated previously, the influence of rolling temperature on annealing texture is far greater for the ELC grade than for the IF materials. This can be attributed to the dependence of the population and morphology of the shear bands on rolling temperature and the presence of alloying elements. Barnett[4] showed that during the annealing of as-rolled IF steels, the presence of shear bands gives rise to recrystallization nuclei with 具111典LND-type orientations. These components are further intensified when the RD fiber components are consumed during the later stages of recrystallization.[4,24] Shear bands frequently provide preferential sites for recrystallization; however, the orientations of these nuclei depend on the nature of the shear bands. When these are of moderate intensity, as in warm-rolled IF steels, recrystallized grains with 具111典LND-type orientations are nucleated.[4] On the other hand, when the bands are intense, as in coldrolled LC steels, the Goss component ({110}具001典) has been observed to form in their vicinity.[19] It can, therefore, be concluded that a proportion of the ND fiber is attributable to nucleation in the vicinity of shear bands of moderate intensity and in the positive ratesensitivity (m) region (Figure 16). The strong Goss component observed in the present ELC and LC Cr annealing textures after rolling at 440 °C, on the other hand, is caused by the nucleation of this component in the vicinity of severe or heavy shear bands; this, in turn, is associated with negative rate sensitivity. Also, the near absence of shear bands evident in the structure of the warm-rolled ELC steel is responsible for the weak Goss and ND fiber in this material after annealing (this material possesses a high positive rate sensitivity at this temperature).[20] Furthermore, because of the link between rate sensitivity and dynamic strain aging, the high solute carbon levels during rolling and annealing of the ELC and LC Cr steels promote nucleation of the Goss texture.[18] In the LC Cr grade, the addition of chromium resulted in a significant improvement in the {111} intensity after rolling at 710 °C and 640 °C followed by recrystallization. This was consistent with the increased frequency of shearband formation at higher rolling temperatures due to the decrease in rate sensitivity. In an earlier work,[5] it was found that the intensity of the {111} components was increased by the addition of 1.3 pct chromium to a 0.03 pct carbon steel. The present results indicate that a chromium content of 0.5 pct is insufficient to produce a sharp {111} recrystallization texture, at least under the current processing conditions. Furthermore, the high manganese content of the steel studied here probably made a further contribution to reducing the intensity of the {111} component. Based on measurements of recrystallization kinetics, it has been suggested that carbon-manganese dipoles may encourage strain-induced boundary migration and, therefore, the nucleation of unfavorable {hkl}具110典 texture components.[17] For this reason, it is expected that lowering the manganese content of the Cr-modified material will further increase the sharpness of the ND fiber component of the annealed product.[25] METALLURGICAL AND MATERIALS TRANSACTIONS A

D. Effect of Coiling on Recrystallization Texture The present work showed that aging had a significant effect on the textures of the IF Nb, unstabilized IF Nb, and LC Cr steels, but had less influence on those of the IF Ti and ELC samples. This may be because, according to the solubility product,[26] niobium carbonitride was dissolved during preheating to 1050 °C. Then, after air cooling to the warm-rolling temperature, solute niobium, carbon, and nitrogen continued to be present (in addition to some Nb(C,N) precipitates. The carbon and nitrogen in solution can be reduced by further Nb(C,N) formation during aging. On removal of the carbon and nitrogen before annealing, the intensity of the {111} recrystallization texture can be expected to increase.[25] By contrast, the temperature of dissolution for titanium carbonitride in the IF Ti steel is higher than that of niobium carbonitride in the IF Nb material.[27,28] Hence, the former is not dissolved during preheating to 1050 °C. In the unstabilized IF Nb steel, there is insufficient Nb to remove solute carbon from ferrite. In the ELC sample, there are no carbide-forming elements to tie up the carbon during the coiling simulation. In the LC Cr steel, it is expected that both Fe(Cr)xCy precipitates as well as Cr-C(N) complexes could form during coiling after warm rolling.[29] However, further study is needed to verify this possibility. E. Effect of Reduction on Recrystallization Texture It was demonstrated earlier that the effect of two-pass warm rolling on the LC Cr sample was greater than on the ELC specimen. This was probably because fewer intense shear bands were formed in the LC Cr steel when the strain was increased, an effect that can be attributed, via DSA and the rate sensitivity, to the chromium addition. The observation that the sharpness of the {111} component in the ELC steel was increased by cold rolling after warm rolling can be explained by the increased intensity of shear bands. The formation of the Goss component in the LC Cr steel can be attributed to the presence of intense shear bands in this material.[19] F. Characteristics of the rm and ⌬r Values From the previous results, it appears that acceptable rm values in warm-rolled and annealed IF steels are a direct result of the presence of the shear bands after rolling (and, therefore, of the presence of shear-band nucleation during annealing). Conversely, the low rm values associated with the warm-rolled and annealed ELC steel are a direct result of the absence of shear bands in this material after rolling. The data of Table III reveal that the ⌬r values for the IF Nb steel are lower than those for the IF Ti steel processed under the same conditions. This is in agreement with the findings of other workers.[30] The results of this research indicate that niobium is more effective than titanium for lowering the planar anisotropy of the mechanical properties. By comparing the rm values of the ELC and LC Cr steels, it can be concluded that the addition of chromium leads to a significant improvement. However, the level of formability is still below that required for deep-drawing applications VOLUME 34A, MAY 2003—1173

(rm ⬎ 1.0). It is expected that further increases can be achieved by (1) increasing the chromium level,[29] (2) decreasing the manganese content,[25] (3) increasing the rolling reduction,[25] and (4) decreasing the initial grain size.[25] Such experiments are currently underway, and the results obtained will be reported when they become available. V. CONCLUSIONS 1. Shear-band formation in LC steels is sensitive to rolling temperature, because solute carbon affects the rate sensitivity of the material. Chromium additions modify the DSA behavior in such a way as to increase the shearband content at warm-rolling temperatures. 2. The deformation textures of the ELC and LC Cr steels increase in sharpness with increased rolling temperature, while remaining relatively unchanged in the IF specimens. This is due to the change in shear-band content, as heavy banding at the lower temperatures in the former materials fragments the bulk texture, thus reducing its intensity. 3. The recrystallization textures of the IF grades are relatively unaffected by rolling temperature and contain a continuous ND fiber. However, the ELC steel contains a partial RD fiber and an incomplete ND fiber at high rolling temperatures. The {111}-oriented grains appear to nucleate on moderate shear bands, whereas the Goss component is associated with severe shear bands. 4. A chromium addition of 0.5 pct to LC steel strengthens the {111} recrystallization component and, hence, improves the formability. Although this improvement was small, it is expected that higher levels of alloying will be of greater practical benefit. 5. The use of a coiling simulation strengthens the recrystallization texture of warm-rolled IF Nb and LC Cr steels as it removes solute carbon by carbonitride precipitation. ACKNOWLEDGMENTS The authors are grateful to Stelco Inc. and Dofasco Inc. for the provision of materials for this project and to Farid Hassani (Stelco) and Maria-Lynn Turi (Dofasco) for useful discussions. Rolling was conducted at the Materials Technology Laboratory (CANMET), Natural Resources Canada (Ottawa), with the assistance of Claude Galvani. They also acknowledge with gratitude the financial support received from the Natural Sciences and Engineering Research Council of Canada (NSERC). One of the authors (MRT) also expresses his thanks to the Isfahan University of Technology (Isfahan, Iran) for providing a visiting scholarship during which

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this work was carried out. The help of Mr. E. Fernandez with testing is warmly appreciated. REFERENCES 1. J.J. Jonas: Proc. Processing & Manufacturing of Advanced Materials, THERMEC ‘2000, T. Chandra, K. Higashi, C. Suryanarayana, and C. Tome, eds., Elsevier Science, U.K., 2000. 2. M.R. Barnett and J.J. Jonas: Iron Steel Inst. Jpn. Int., 1999, vol. 39, pp. 856-73. 3. M.R. Barnett and J.J. Jonas: Iron Steel Inst. Jpn. Int., 1997, vol. 37, pp. 697-705. 4. M.R. Barnett: Iron Steel Inst. Jpn. Int., 1998, vol. 38, pp. 78-85. 5. M.R. Barnett: Proc. Materials ‘98, M. Ferry, ed., IMEA, Wollongong, Australia, 1998, pp. 167-72. 6. J. Glen: J. Iron Steel Inst., 1957, pp. 21-48. 7. M.R. Barnett: in Modern LC and ULC Sheet Steels for Cold Forming: Processing and Properties, W. Bleck, ed., Verlag Mainz, Aachen, 1998, pp. 61-72. 8. T. Sakai, Y. Saito, and K. Kato: Trans. Iron Steel Inst. Jpn., 1988, vol. 28, pp. 1036-42. 9. H.J. Bunge: Texture Analysis in Materials Science, Butterworth and Co., London, 1982. 10. H.J. Bunge: Z. Metallkd., 1965, vol. 56, pp. 872-74. 11. L. Seidal, M. Holscher, and K. Lücke: Text. Microstr., 1989, vol. 11, pp. 171-85. 12. U. von Schlippenbach, F. Emren, and K. Lücke: Acta Metall., 1986, vol. 34, pp. 1289-1301. 13. J. Savoie: Texture Menu Software, McGill University, Montreal, Internal Report, 1994. 14. D. Daniel and J.J. Jonas: Metall. Trans. A, 1990, vol. 21A, pp. 33143. 15. D. Liu, A.O. Humphreys, M.R. Toroghinezhad, and J.J. Jonas: Iron Steel Inst. Jpn. 751-59. 16. W.B. Hutchinson, K.-I. Nilson, and J. Hirsch: Metallurgy of VacuumDegassed Steel Products, R. Pradhan, ed., TMS, Warrendale, PA, 1990, pp. 109-25. 17. T. Haratani, W.B. Hutchinson, I.L. Dillamore, and P. Bate: Met. Sci., 1984, vol. 18, pp. 57-65. 18. M.R. Barnett and L. Kestens: Text. Microstr., 2000, vol. 34, pp. 1-22. 19. J.J. Jonas: in Modern LC and ULC Sheet Steels for Cold Forming: Processing and Properties, W. Bleck, ed., Verlag Mainz, Aachen, 1998, pp. 73-84. 20. A.O. Humphreys, D. Liu, M.R. Toroghinezhad, E. Essadiqi, and J.J. Jonas: Mater. Sci. Technol., in press. 21. M.R. Barnett: in Thermomechanical Processing of Steel; COM 2000, S. Yue and E. Essadiqi, eds., Met. Soc. Montreal, 2000, pp. 265-77. 22. J.J. Jonas: in IF Steels 2000 Proc., ISS, Pittsburgh, PA, 2000, pp. 233-45. 23. R.E. Hook and H. Nyo: Metall. Trans. A., 1975, vol. 6A, pp. 1443-51. 24. D. Vanderschueren, N. Yoshinaga, and K. Koyama: Iron Steel Inst. Jpn. Int., 1996, vol. 36, 1046-54. 25. R.K. Ray, J.J. Jonas, and R.E. Hook: Int. Mater. Rev., 1994, vol. 39, pp. 129-72. 26. K. Narita: Trans. Iron Steel Inst. Jpn., 1975, vol. 15, pp. 145-52. 27. K.J. Irvine, F.B. Pickering, and T. Gladman: J. Iron Steel Inst., 1967, vol. 205, pp. 161-82. 28. S. Matsuda and N. Okumura: Trans. Iron Steel Inst. Jpn., 1998, vol. 18, pp. 198-205. 29. M.R. Barnett: Steel Res., 2000, vol. 71, pp. 295-302. 30. S. Satoh, T. Obara, and M. Nishida: Trans. Iron Steel Inst. Jpn., 1984, vol. 24, pp. 838-46.
