Hot Rolling Texture Development in CMnCrSi Dual-Phase Steels T. WATERSCHOOT, L. KESTENS, and B.C. De COOMAN The amount of strain below the temperature of nonrecrystallization, Tnr , has an important influence on the phase fractions and the final crystallographic texture of a hot-rolled dual-phase ferrite ⫹ martensite CMnCrSi steel. The final texture is influenced by three main microstructural processes: the recrystallization of the austenite, the austenite deformation, and the austenite-to-ferrite transformation. The amount of strain below Tnr plays a major role in the relative amounts of deformed and recrystallized austenite after rolling. Recrystallized and deformed austenite have clearly different texture components and, due to the specific lattice correspondence relations between the parent austenite phase and its transformation products, the resulting ferrite textures are different as well. In addition, austenite deformation textures result from either dislocation glide or the combination of dislocation glide and mechanical twinning, depending on the stacking fault energy (SFE). The texture components in hot-rolled dual-phase steels were studied by means of X-ray diffraction (XRD) measurements and orientation imaging microscopy (OIM). A clear crystallographic orientation difference was observed between the ferrite phase, transformed at temperatures near Ar3, and the ferritic bainite and martensite phases, formed at lower temperatures. The results suggest that the primary ferrite, nucleated at temperatures close to Ar3, transformed from the deformed austenite. The low-temperature constituents, bainite and martensite, form in the recrystallized austenite.
I. INTRODUCTION
THE microstructure of a dual-phase steel consists typically of a dispersion of a hard second phase in a matrix of ferrite. This second phase is usually martensite, but other low-temperature constituents, such as bainite, can be present as well.[1,2] A minimum amount of martensite (4 to 5 pct) is required to obtain the typical mechanical dual-phase behavior: a continuous yielding, a low yield-to-tensilestrength ratio (YS/TS), a high uniform and total elongation, and high work-hardening rates.[3,4] In recent years, both coldand hot-rolled grades have been developed for automotive applications such as wheel discs. A general deformation and transformation schedule for the hot rolling of dual-phase steels is shown in Figure 1. Primary ferrite formation needs to be promoted, as there is only a very short time for high-temperature transformation, after the hot rolling and before the start of the fast cooling. Deformation at temperatures below the temperature of nonstatic recrystallization Tnr makes strain accumulation in the austenite phase possible. The increased dislocation density in the pancaked austenite decreases the incubation time for the transformation to ferrite in the hot-rolled strip. At the start of the fast cooling, the microstructure is a mixture of a low-C (⬃0.01 wt pct C) ferrite and a higher-C (⬃0.3 to 0.5 wt pct C) austenite. The hardenability of the austenite, i.e., the ability to form martensite at reasonable cooling rates, can be increased through an increase in the C content, through the addition of alloying elements such as Mn, Si, Cr and Mo,[2,5,6] and through a fast cooling prior to coiling. T. WATERSCHOOT, Graduate Student, and L. KESTENS and B.C. DE COOMAN, Professors, are with the Department of Metallurgy and Materials Science, Laboratory for Iron and Steelmaking, Ghent University, B-9052 Ghent, Belgium. Contact e-mail:
[email protected] Manuscript submitted February 26, 2001. METALLURGICAL AND MATERIALS TRANSACTIONS A
The present contribution focuses on the texture formation in hot-rolled dual-phase steels as part of a wider effort to find ways in which to improve the forming properties of dual-phase steels. An important factor influencing the formability of sheet steel is the crystallographic texture of the material. Much research has been carried out on the hot band textures of interstitial free (IF) and extra low carbon (ELC) steels.[7] The transformation texture of the ferrite is influenced by three main micro-structural processes: the austenite recrystallization, the austenite deformation, and the austenite-to-ferrite transformation. In all cases the strongest ferrite texture component can be found on the rolling direction (RD) fiber (具110典␣ //RD). If the austenite recrystallization is not suppressed, the maximum intensity of the final product is found on the ferrite rotated cube component, {001}具110典␣, which is the transformation product of the cube component {001}具010典␥ , the main recrystallization component in austenite.[7,8] If the recrystallization is suppressed, e.g., by the addition of certain alloying elements such as Nb and Ti, the main ferrite texture component is shifted to {112}-{113}具110典␣, and stronger ferrite  -fiber components are obtained.[8,9] The ferrite  fiber is the transformation equivalent of the  rolling texture in austenite, which is composed of the main austenite deformation components: the Cu, Brass, and S components. At the present time, no data are available on the effect of the hot-rolling process parameters on the texture development in dual-phase steels. Hutchinson et al.,[10] however, investigated a CMnSi transformation induced plasticity (TRIP) steel and determined a strong ferrite texture, resulting from the transformation of partially recrystallized austenite. In the present work, the texture of the final, hot-rolled dual-phase steel was investigated. A clear relation between the amount of deformation below Tnr and the texture components specific for each of the different phases was identified. Furthermore, the crystallographic textures of the ferrite, VOLUME 33A, APRIL 2002—1091
Fig. 1—General deformation and transformation characteristics of hot-rolled dual-phase steel. The left figure shows the equilibrium diagram; and the right figure explains the deformation, recrystallization, and transformation behavior during hot rolling of the dual-phase steel.
Fig. 2—Thermomechanical processing of the different hot-rolled samples.
bainite, and martensite phases were studied separately, using orientation imaging microscopy (OIM). II. EXPERIMENTAL PROCEDURE The composition of the material used in the course of the present work was as follows: 0.075 wt pct C, 1.39 wt pct Mn, 0.69 wt pct Cr, and 0.1 wt pct Si. Plates with an initial thickness of 32 mm were laboratory hot rolled after being reheated for 1 hour at 1270 ⬚C. A final hot band thickness of 2.5 mm was obtained in 6 passes, with reductions of 25, 25, 30, 40, 40, and 45 pct, respectively. A schematic overview of the time-temperature dependence during hot rolling is shown in Figure 2. All sheets 1092—VOLUME 33A, APRIL 2002
were hot rolled in the austenite region, i.e., at temperatures above Ar3, which was determined to be 720 ⬚C by means of dilatometry. The starting temperature of the first rolling pass was decreased from 1050 ⬚C to 900 ⬚C, while the interpass time for this series of samples with names ending in C was kept constant at 5 seconds. A larger deformation was given at temperature below Tnr (880 ⬚C)[12] for lower starting temperatures. The hot rolling of sample 1050L started at a temperature of 1050 ⬚C and longer interpass times of 10 to 20 seconds were used, to evaluate the effect of longer static recrystallization times between the passes. After rolling, all materials were air cooled for 10 seconds, with a cooling rate of 15 ⬚C/s, and water spray cooled at a rate of 60 ⬚C/s to room temperature. The hot-rolled samples were examined by light optical microscopy (LOM), using the Le Pera etching technique,[13] which colors both the martensite and the retained austenite white, the bainite dark brown or black, and the ferrite browngray. The volume fraction of each phase was determined using the digital image processing program LUCIA.* *LUCIA is a trademark of Laboratory Imaging, Prague, Czech Republic.
Local texture measurements were made using the OIMelectron back scattering diffraction (OIM-EBSD) technique. The EBSD attachment was mounted on a PHILIPS* XL30 *PHILIPS is a trademark of Philips Electronic Instruments Corp., Mahwah, NJ.
environmental scanning electron microscope (SEM) equipped with a LaB6 filament and operated at 30 kV. The METALLURGICAL AND MATERIALS TRANSACTIONS A
Table I. Relation between Start Rolling Temperature, the Effective Strain below Tnr , and the Phase Composition Tstart rolling Deformation below Ferrite Martensite Bainite Tnr (Pct) (Pct) (Pct) Sample (⬚C) below Tnr 1050C 1010C 970C 900C 1050L
1050 1010 970 900 1040
50 65 80 90 80
pct pct pct pct pct
0.7 1.3 1.9 2.6 1.9*
69.3 81.7 82.9 84.7 72.5
19.7 16.1 15.6 14.3 13.7
11.0 2.2 1.5 1.0 13.8
*Longer interpass times.
specimen tilt was 70 deg during OIM experiments. The samples were prepared for these measurements using mechanical polishing, electropolishing, and a light etching with 2 pct Nital. The OIM scans were carried out in a rectangular grid, with a step size of 0.5 m. According to conventional standards, it is tacitly understood in these types of measurements that the crystallographic anisotropy is not correlated with a morphological anisotropy of the grain size. Although it might theoretically be possible to design an optimized scanning strategy, previous experience has shown that reasonably selected scanning areas will produce statistically reliable texture data.[11] Further investigation of the phases was done using the image quality (IQ) and the confidence index (CI) of the measurements provided by the commercial TSL OIM* software. The IQ evaluates the contrast
Fig. 3—Phase balance as a function of the effective strain below Tnr .
*TSL OIM is a trademark of TexSem Laboratories Inc., Draper, UT.
of the Kikuchi lines, whereas the CI quantifies the reliability of the pattern indexing. In addition, texture measurements were carried out in the conventional way by measuring pole figures on polished samples using X-ray diffraction (XRD) and by calculating orientation distribution functions (ODFs). All samples for texture measurements were taken at 50 pct of the thickness. The pole figures were measured by the Schultz reflection method using an Euler cradle goniometer mounted on a Siemens D5000 diffractometer (Siemans, Erlangen, Germany) using Mo K␣ radiation. III. RESULTS The effective strain at temperatures below Tnr , 880 ⬚C, based on the formula of Barrato et al.,[12] is shown as a function of the start-rolling temperature in Table I. The effective strain below Tnr increases with a decreasing startrolling temperature. Samples 970C and 1050L, with startrolling temperatures of 970 ⬚C and 1050 ⬚C, respectively, underwent the same effective strain at a temperature below Tnr , but the interpass times for the latter were between 10 and 20 seconds. The interpass times were typically 5 seconds for the 970C sample. In the following sections, the influence of the amount of deformation at temperatures below Tnr on the phase balance as well as on the texture development of the steels is discussed in detail. A. Microstructural Analysis A summary of the microstructural evolution as a function of the amount of deformation at a temperature below Tnr , METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 4—Microstructures of specimens 1050C, 1010C, and 900C.
for constant interpass times except for sample 1050L, is shown in Table I and in Figure 3. The corresponding microstructures of samples 1050C, 1010C, and 900C are shown in Figure 4. All materials showed the typical hot-rolled dualphase microstructures consisting of large amounts of ferrite with martensite and bainite islands. It is worth noticing that VOLUME 33A, APRIL 2002—1093
Fig. 6— 2 ⫽ 45 deg section of the Euler space, showing the most important rolling and recrystallization texture components for the ferritic phase.
Fig. 5—Schematic C-concentration profile in the martensite-bainite islands and the adjacent ferrite grain.
in some cases the martensite island had a bainitic core. This can be explained by the C-concentration profile in the transformed austenite, as shown in Figure 5. Due to the low diffusivity of C in austenite at lower temperatures, the Cconcentration on the ␥ side of the ␣ -␥ phase boundary is higher than in the center, leading to a difference in hardenability within the transforming austenite grain. As a result, the bainitic transformation is not always suppressed in the center of the grains, and the bainitic core is obtained in the martensite islands. The deformation at temperatures below Tnr resulted in a pancaked austenite microstructure, which enhanced the primary ferrite transformation due to the increased number of nucleation sites and the larger driving force for nucleation. A strong increase in the amount of ferrite (from 60 to 82 pct) was noticed for effective strains below Tnr up to 1.3. For larger deformations at low temperatures, the amount of ferrite shows no further increase, as the amount of ferrite is approaching the equilibrium ferrite content. By means of Thermocalc calculations, the equilibrium amount of ferrite was determined to be 86 pct. The amount of bainite steadily decreased with the increasing ferrite content. This can be understood if one considers that when more ferrite is formed, the austenite is enriched in C due to the very low solubility of C in the primary ferrite (less than 0.02 wt pct C), and the hardenability of the untransformed austenite increases. The martensite content first increased for effective strains up to 1 and then decreased slightly. The initial increase is related to the better hardenability, i.e., the higher C content of the untransformed austenite. The subsequent slight decrease in the martensite phase fraction is due to the overall decreasing amount of the second phase. The sample 1050L, with the longer interpass time, contained less ferrite than sample 970C, although both were 1094—VOLUME 33A, APRIL 2002
deformed the same amount below Tnr . Long interpass times, which increase the amount of static recrystallized austenite, are therefore detrimental for the promotion of austenite-toferrite transformation. B. Texture Analysis The orientations of interest to the processing of ferritic steels can be found in the 2 ⫽ 45 deg section of the Euler space (Figure 6). (1) The ␣ or RD fiber, on the left vertical axis and representing all orientations with a common 具110典␣ direction parallel to the RD. (2) The ␥ or normal direction (ND) fiber, represented by a horizontal line at ⌽ ⫽ 54.7 deg and representing the {111}具uvw典␣ orientations, where the {111}␣ planes are parallel to the rolling plane. (3) The cube {001}具010典␣ and rotated cube {001}具110典␣ component. (4) The Goss {110}具001典␣ and rotated Goss {110}具110典␣ component. (5) The ferrite  fiber, which is the transformation product of the austenite  fiber, is not sharply defined, but forms a continuous series of orientations from {114}具110典␣ on the RD fiber toward {111}具112典␣ on the ND fiber. All samples showed comparable textures, which were relatively weak, as is usually the case for transformationtype textures (Figure 7). Sample 1010C, for example, had a maximum of 3.7 ⫻ random at the upper part of the ␣ fiber, from the rotated cube component ({001}具110典␣) down to the {114}具110典␣ component. Furthermore, an increased intensity was noticed on the  fiber running from {114}具110典 at the left vertical axis over {111}具112典␣ to {554}具225典␣ at the right axis. Local maxima on this fiber were noticed for 1 ⫽ 0 deg (3.4 ⫻ random), 1 ⫽ 26 deg (2.7 ⫻ random), and 1 ⫽ 90 deg (2.8 ⫻ random), and a local minimum for 1 ⫽ 60 deg. A small intensity increase was also observed for the rotated Goss component {110}具110典␣. It is worth METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 7— 2 ⫽ 45 deg section of the ODF of samples 1050C, 1010C, 970C, 900C, and 1050L. A 2 ⫽ 45 deg with the most important ferrite texture components is shown as a reference.
noticing that samples 900C and 970C show a maximum intensity of 4.5 to 5 ⫻ random close to the {114}具110典␣ component, whereas the intensity along the ␣ fiber was METALLURGICAL AND MATERIALS TRANSACTIONS A
almost constant from the {001}具110典␣ rotated cube component to the {114}具110典␣ component for the higher-temperature samples 1010C and 1050C. Sample 1050L, with the VOLUME 33A, APRIL 2002—1095
(a)
(a)
(b) Fig. 9—Image and distribution chart of the CI and IQ of the OIM measurement on the 900C sample: backscattering pattern (a) IQ and (b) CI.
(b) Fig. 8—Intensities of components along the (a) ␣ fiber and the (b)  of all the samples.
observed for all the samples at 1 ⫽ 0 deg, 1 ⫽ 20 to 30 deg, and 1 ⫽ 90 deg, and a minimum in the 1 ⫽ 50 to 60 deg range. The maximum at 1 ⫽ 20 to 30 deg, was relatively more intense for the low-rolling-temperature samples 900C and 970C. These maxima correspond to the following texture components: {114}具110典␣ (1 ⫽ 0 deg), a component varying from {112}具131典␣ to {111}具112典␣ (1 ⫽ 20 to 30 deg), and {554}具225典␣ (1 ⫽ 90 deg). The minimum is positioned on the {111}具110典␣ component. C. OIM Measurements
longer interpass times, was less textured, with the overall maximum less than 3 ⫻ random. The intensities along the ferrite ␣ and  fibers are shown in Figure 8, in order to compare the fiber intensities quantitatively for the samples. The intensity of the ␣ fiber increased with increasing amounts of deformation at a temperature below Tnr . Almost no intensity difference was noticed between the 970C and 900C samples with 80 and 90 pct, respectively, of deformation below Tnr . It was observed that, for the lower deformation temperatures, there was an intensity peak at the {114}具110典␣ component, whereas an almost constant value from {001}具110典␣ to {114}具110典␣ was noticed for the samples with higher deformation temperatures. Although the amount of deformation below Tnr was the same for samples 970C and 1050L, the latter was less textured and had no maximum around {114}具110典␣, but had an almost constant intensity along the partial RD fiber 0 deg ⬍ ⌽ ⬍ 40 deg. The same general conclusions can be made for the  fiber: the intensity decreased for an increasing rolling-start temperature, i.e., a decreasing amount of deformation below Tnr . Sample 970C even showed somewhat higher intensities than the 900C specimen. Local maxima on the  fiber were 1096—VOLUME 33A, APRIL 2002
In order to obtain a detailed analysis of the texture of each phase present in the microstructure, OIM measurements were carried out on samples 900C and 1050C and the results are shown in Figures 9 through 12. The CI and the IQ maps, together with the frequency distribution charts of sample 900C, are shown in Figure 9. The lightest points in the maps indicate the highest values of CI and IQ; the darkest points designate the lowest values. Dark arrays of points, usually with a 0 CI and a low IQ, represent grain and phase boundaries. Figure 10(a) shows the macroscopic texture of the 900C sample. Despite small differences, e.g., the intensity on the rotated cube component {001}具110典␣, there was a very good correspondence between the XRD and the OIM measured textures. Different phases can be separated, based on the IQ and the CI factors of the diffraction patterns, which were matched to the bcc phase. According to the CI map and the associated histogram of Figure 9(b), three different types of grains were observed: grains with a high CI (⬎0.6:75 pct), with an intermediate CI (between 0.2 and 0.6:10 pct) and with a low CI (⬍0.2:15 pct). The observed morphology of the grains based on their CI or IQ is very similar to that of the light optical microscopy (LOM) micrographs in Figure 4. It is METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 10—OIM texture measurement on the 900C sample. (a) The general texture. (b) The texture of the ferrite, based on CI selection (CI ⬎ 0.6). (c) The texture of the bainitic ferrite with (0.2 ⬍ IC ⬍ 0.5).
empirically determined that a CI larger than 0.1 yields absolutely reliable data. In Figure 10(b), the texture of the scanned surface with a high CI is shown, revealing the same patterns as for the total scan. Figure 10(c) shows the texture of the scanned surface with an intermediate CI (10 pct). The texture was different from the total texture, showing a maximum in the rotated cube component {001}具110典␣ of 4 ⫻ random, which is much stronger than in the total texture, and a weakened ferrite  fiber with a maximum of 2.5 times random on {113}具110典␣, compared to more than 5 times random for the total texture. For the lowest CI, an increased intensity (2 times random) in the rotated cube components was noticed as well, but one has to take into account the low reliability of these data points. Similar conclusions can be drawn for sample 1050C. In Figure 11, it is shown that 65 pct of the data points had a CI higher than 0.6, 15 pct had an intermediate CI, and for 20 pct of the data points, the CI was lower than 0.2. The METALLURGICAL AND MATERIALS TRANSACTIONS A
overall textures determined by OIM (Figure 12(a)) and XRD (Figure 7) for sample 1050C showed the same tendencies, again with an increase on the {114}具110典␣ component and a decrease on the rotated cube component for the OIM measurement. No major differences were observed between the ODFs, based on the data points with a CI higher than 0.6 (Figure 12(b)) and the overall ODF, but for the textures derived from data points with an intermediate CI (Figure 12(c)), a higher intensity for rotated cube and a weakened  fiber were observed. A sensitivity analysis showed that the type of texture, revealed by selecting the data points from a certain range of CI, is not sensitive to small changes in the choice of this limit. IV. DISCUSSION The OIM measurements revealed three different types of grains with high, medium, and low CIs. There is a linear VOLUME 33A, APRIL 2002—1097
Based on these presumptions, a reasonable relation is found between the fractions of phases derived from microscopy and from the OIM data. In case of sample 1050C, 69 pct of ferrite is found based on LOM and 65 pct of ferrite is found based on OIM. For sample 900C, the amounts of ferrite based on LOM and OIM determinations are 85 and 75 pct, respectively. A. Texture Evolution during Thermomechanical Processing The final texture of the material is the result of the deformation and recrystallization processes in the austenite phase, followed by the austenite-to-ferrite transformation. (a)
(b) Fig. 11—Distribution chart of the CI and IQ of the OIM measurement on the 1050C sample: (a) image quality (IQ) and (b) confidence index (CI)
relation between the IQ and CI values, which can be observed in Figure 9. For example, the dark grains with the low IQ of Figure 9(a) correspond to the dark, low-CI grains of Figure 9(b). The majority of the grains had both a high CI and a high IQ value, which is characteristic for sharp, clear, and well-defined diffraction patterns. This strongly suggests that these grains are low-dislocation-density grains, which were formed at a high temperature (⬎600 ⬚C). These grains can thus be considered high-temperature primary-ferrite grains. The detailed texture of these ferrite grains is shown in Figures 10(b) and 12(b) for the 900C and 1050C samples, respectively. Grains with an intermediate CI value had also lower IQ values, which is characteristic of a more diffuse diffraction pattern presumably caused by the presence of a larger amount of dislocations. It is reasonable to assume that these dislocations were introduced by a transformation that occurred at lower temperature and that produced a bainitic transformation product. The resulting bainitic textures are shown in Figures 10(c) and 12(c). The very low CI and IQ factors of some grains correspond to diffraction patterns, which were very blurred due to an extremely high density of crystal defects, such as the dislocations and twin structures of martensite. The very low CI for martensite is further enhanced by misfitting, since the orientation of the bct crystal of the martensite, with a c/a value close to but different from 1, is compared to a analogous but slightly different bcc crystal. 1098—VOLUME 33A, APRIL 2002
1. Textures produced by austenite recrystallization Recrystallization of hot-rolled austenite leads to the formation of a cube texture component {100}具001典␥ . The intensity of this component increases with the higher amounts of strain accumulated prior to recrystallization. High finishrolling temperatures and lean chemistries favor the austenite recrystallization and thus the predominance of the cube component. Based on the Bain orientation relationships (Figure 13) between the parent austenite phase and its transformation product, the austenite cube component {100}具001典␥ can transform into 3 ferrite texture components: rotated cube {100}具011典␣, Goss {110}具001典␣, and rotated Goss {110}具110典␣ (Figure 6).[7] Other crystallographic orientation relations such as the Kurdjumov–Sachs (KS) and the Nishiyama–Wasserman (NW) orientation relations (Figure 13) result in nearly the same bcc orientations. Based on the OIM measurements, an obvious intensity increase in the rotated cube orientation was noticed for the bainite phase (Figures 10(c) and 12(c)). It can thus be assumed that most of the austenite, undercooled down to low temperatures and leading to low-temperature constituents such as bainite and martensite, was recrystallized. This proves indirectly that austenite-to-ferrite transformation is promoted by deformation accumulation in the austenite prior to transformation. Similar observations were made by Hutchinson et al.[10] for TRIP steels. Furthermore, an obvious variant selection in favor of the rotated cube component {100}具011典␣ compared to Goss {110}具001典␣ and rotated Goss {110}具110典␣ was noticed, such as is generally observed for this type of transformation.[7,8,10] Although the physical explanation for this variant selection is as yet unclear, it can be related to the micro stress-strain equilibrium at the interface of the transforming nuclei.[14] 2. Textures produced by austenite deformation Austenite has an fcc crystal structure, and the main texture components of deformed austenite produced by hot rolling are located on the austenite  fiber. These texture components are the copper (C) {112}具111典␥ , S {123}具634典␥ , brass (B) {110}具112典␥ , and Goss (G) {110}具001典␥ orientations (Figure 14).[15] The intensities of those components depend on the strain that can be accumulated without recrystallization and on the stacking fault energy (SFE). The SFE is considered to be the most important factor influencing the presence of specific deformation texture components in the deformed fcc lattice for a given strain. As shown by Humphreys.[16] for large amounts of rolling METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 12—OIM texture measurement on the 1050C sample. (a) The general texture. (b) The texture of the ferrite, based on CI selection (CI ⬎ 0.6). (c) The texture of the bainitic ferrite with (0.2 ⬍ IC ⬍ 0.5).
reduction, brass (low SFE: 25 mJ/m2) develops a crystallographic texture with the highest intensity on the brass component {110}具112典␥ , a high-intensity Goss component {110}具001典␥ , and low-intensity copper component {112}具111典␥ . For copper (high SFE: 80 mJ/m2), on the other hand, the intensities for all the components are similar, though they show a small increase in intensity for the copper {112}具111典␥ and S {123}具634典␥ components. Obviously, in the case of steels, which undergo an allotropic transformation, this phenomenon has important consequences for the texture of the final transformation product. The  -fiber components of the austenite transform into the  fiber of the ferrite. Based on the KS transformation relations, the texture components in a 2 ⫽ 45 deg section that result from the transformation of the copper, S, brass, and Goss components are shown in Figure 15. In Table II, an overview of the resulting possible texture components for each of the deformation components is shown, together with an estimated intensity in case of a KS transformation without any variant selection. METALLURGICAL AND MATERIALS TRANSACTIONS A
By means of the OIM data, it is possible to study the transformation products of the deformed austenite separately. Figures 10(b) and 12(b) show the ferrite textures of samples 900C and 1050C, respectively. Although the possibility that some of these ferrite orientations are due to the transformation of recrystallized austenite cannot be ruled out, most of the orientations certainly result from the transformation of deformed austenite. The maximum near the {113}具110典␣ component is the result of the transformation of the partial  fiber between the Cu and S texture components. Because the maximum intensity is exactly at {113}具110典␣, it strongly suggests that this peak originates from Cu-oriented grains in the deformed austenite phase. As the transformation product of S is situated 10 deg lower (i.e., near {114}具110典␣ instead of {113}具110典␣), a larger amount of S in the deformed austenite would have led to a shift of the maximum towards the {114}具110典 component. No difference between the two samples is noticed, as both maxima are positioned on {113}具110典␣. The results show that for the transformation VOLUME 33A, APRIL 2002—1099
(a)
(b)
(c) Fig. 13—Schematic of orientation relationships between the ␥ (fcc) and the ␣ (bcc) lattice. The differences in lattice parameters between ␣ and ␥ phases are not taken into account for reasons of clarity. (a) Bain orientation relations: {001}␥ //{001}␣ and 具001典␥ //具110典␣. (b) KS orientation relations: {111}␥ //{011}␣ and 具011典␥ //具111典␣. (c)NW orientation relations: {111}␥ // {011}␣ and 具211典␥ //具011典␣.
of the copper component, a variant selection must be operating. The {110}具110典␣ component is expected to have an intensity comparable to the {113}具110典␣ component in the absence of a variant selection. The experimental results show that the former orientation has an intensity 5 times lower than the latter. The maxima at 1 ⫽ 30 deg and 90 deg, in the ⫽ 45 deg section of the euler space for the ferrite phase, result from the transformation of the components between the Brass and the S component. Also in this case, the exact position of the maximum gives a clear indication about the relative proportion of these two  -fiber components (B and S) in the deformed austenite before transformation. The component resulting from the transformation of the Brass component is positioned exactly in the {554}具225典␣ orientation (⌽ ⫽ 60 deg), whereas the transformation component of S is at 1100—VOLUME 33A, APRIL 2002
Fig. 14—Position of some main components of the ␥ phase in the Euler space. The position of the  fiber with the copper (C), S, and brass (B) component, leading to the Goss (G) component.[14].
⌽ ⫽ 73 deg, corresponding to a {332}具113典␣ component. The more the measured maximum is shifted to higher values of ⌽, the larger the number of grains with an S orientation, as opposed to the Brass orientation. For the 1050C sample, the maximum is very near to the {554}具225典␣ component, whereas it is shifted toward the {332}具113典␣ component for the 900C sample. This observation suggests that, in the case of lower deformation temperatures, more S components and fewer Brass components are present. For both samples, an increased intensity for the rotated cube component is noticed as well, which suggests that there is no variant selection for the transformation of the Brass component. For the OIM measurement of sample 1050C, a small local maximum is observed on the {111}具110典␣ component, which is the transformation product of an austenite Goss component. In all other cases, including the XRD measurements, a minimum is observed for this component. Based on the measurements of the bcc textures of the ferrite, a maximum in intensity is suggested for the Cu component of the austenite texture. Furthermore, no marked differences were observed between the intensities of the different texture components of the deformed austenite transformation products. According to the above results, the deformed austenite state must have had a deformation behavior leading to Cu-texture components. This is suggestive of a high SFE. Within the range of investigated deformation temperatures, no influence of these temperatures on the general type of deformation behavior was noticed. A small but significant intensity increase near the {112}具131典-{111}具112典 component is noticed for the 970C and 900C samples, based on the XRD measurements. A possible origin for this component is deformation twinning in the parent austenite phase, which can be induced by latent hardening of the {111} slip systems. This can be caused by large amounts of deformation accumulation, due to high METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 15—Resulting bcc transformation components of the copper, S, brass, and Goss component in fcc, based on the KS orientation relations for transfo rmation
Table II. Resulting Transformation Texture Components from the Austenite  Fiber Texture Components Copper, S, Brass, and Goss Austenite  Fiber Texture Component Copper {112}具111典 S {123}具634典 Brass {110}具112典 Goss {110}具001典
Resulting Ferrite Texture Components
Intensity
{112}具110典 to {113}具110典 {110}具110典 {201}具102典 {114}具110典 to {115}具110典 near {110}具110典 near {332}具113典 (⌽ ⫹ 10 deg) {112}具131典 {001}具110典 {111}具112典 {554}具225典 {001}具100典 {111}具110典 {223}具110典
strong strong very weak very strong weak normal normal strong strong strong strong strong strong
amounts of rolling reduction at a temperature below Tnr , which was the case for samples 970C and 900C. The critical resolved shear stress required to activate the slip systems thus can be drastically increased to make it more likely that other deformation mechanisms, such as mechanical twinning, will occur. However, because the ␥ textures strongly suggest a high SFE, mechanical twinning is not likely to play a major role during deformation. In terms of hot band texture control for this type of material, it is important to realize that an ND fiber texture with the associated favorable deep drawing properties cannot only be obtained by the conventional recrystallization mechanism, as has been commonly supposed. The present data strongly suggest that, by the appropriate control of the parent phase texture, a favorable deep-drawing texture can be obtained after an austenite-to-ferrite transformation, which is characterized by KS crystallographic orientation relationships. The deformation, recrystallization, and transformation behavior of dual-phase steel, together with its influence on the texture development, is summarized in Figure 16. The METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 16—General deformation and transformation characteristics of hotrolled dual-phase steel explaining the deformation, recrystallization, and transformation behavior during hot rolling of dual-phase steel and the development of the final texture.
deformation at temperatures above Tnr is followed by recrystallization, which results in a strong cube {001}具010典␥ annealing component. Most of the strain applied at temperatures below Tnr will be accumulated, which leads to pancaked austenite grains with the specific  -fiber texture components. However, depending on the specific circumstances, some static or dynamic recrystallization occurs as well. The deformed austenite transforms into primary ferrite, which inherits its  fiber from the deformed austenite as a result of the lattice correspondence between the parent VOLUME 33A, APRIL 2002—1101
and the product phase. The recrystallized austenite is enriched in C due to the low solubility of C in the ferrite, and during the fast cooling it remains present in the microstructure at low temperatures. Therefore, the texture of the lowtemperature transformation products, bainite and martensite, is related to the texture of the recrystallized austenite phase. In agreement with the well-known KS orientation relationships, an increased intensity on the rotated cube component of the bainite and the martensite is obtained. V. CONCLUSIONS It has been shown that the amount of rolling deformation at temperatures below Tnr plays an important role in the development of the microstructure and the crystallographic texture of the different phases present in a MnCrSi-alloyed dual-phase steel. Large amounts of rolling deformation at temperatures below Tnr result in a pancaked austenite microstructure and stimulate the transformation to a primary ferrite phase with a deformation texture. The XRD and OIM measurements showed comparable textures for all the samples. The textures were relatively weak, as is usually the case for the transformation-type textures. An increased intensity on both the ferrite ␣ and  fibers of the bcc structure and a maximum between the {114}具110典␣ and the {113}具110典␣ component were observed for the samples with the largest amount of low-temperature austenite deformation. Longer interpass times resulted in a weakening of the overall texture. Based on the OIM measurements, it was possible to distinguish polygonal ferrite, formed at high temperatures, and low-temperature constituents such as bainite and martensite. No significant difference between the ferrite texture and the overall texture was noticed, but the bainitic and martensitic textures differed strongly from the global texture. The ferrite had a texture originating from deformed austenite. The textures of the low-temperature constituents had a maximum intensity on the {001}具110典␣, rotated cube component. This strongly suggests these phases originated from recrystallized austenite grains. Based on the OIM measurements, the deformation of the austenite was found to favor a copper type of deformation
1102—VOLUME 33A, APRIL 2002
behavior. Within the range of investigated deformation temperatures, no influence of these temperatures on this general type of deformation behavior was noticed. Comparing calculations of ODF transformations based on the KS orientation relations with the measured OIM textures, an obvious variant selection for the transformation products of the austenitic copper and Goss components is noticed. ACKNOWLEDGMENTS The authors gratefully acknowledge the financial support of OCAS NV, the Corporate Research Center of Sidmar, the Flat Rolled Products Division of the ARBED Group. REFERENCES 1. D.T. Llewellyn and D.J. Hillis: Ironmaking and Steelmaking, 1996, vol. 23 (6), pp. 471-78. 2. T. Waterschoot, B.C. De Cooman, D. Vanderschueren, and B. Staelens: Proc. Thermomechanical Processing of Steels Conf., May 24–26, 2000, IOM Communications, London, pp. 407-17. 3. A.R. Marder: Metall. Trans. A, 1981, vol. 12A, pp. 1569-79. 4. R.G. Davies: Metall. Trans. A, 1978, vol. 9A, pp. 41-52. 5. W.W. Cias: Climax Molybdenum Company, Greenwich, CT. 6. C. Mesplont, T. Waterschoot, B.C. De Cooman, S. Vandeputte, and D. Vanderschueren: Proc. Thermomechanical Processing of Steels Conf., May 24–26, 2000, IOM Communications, London, pp. 495-504. 7. J.J. Jonas: IF Steels 2000 Proc., Pittsburgh, PA, June 5–7, 2000, ISS, Warrendale, PA, pp. 233-46. 8. R.K. Ray, M.P. Butron-Guillen, J.J. Jonas, and G.E. Ruddle: Iron Steel Inst. Jpn. Int., 1992, vol. 32 (2), pp. 203-12. 9. M.P. Butron-Guillen, C.S. Da Costa Viana, and J.J. Jonas: Metall. Mater. Trans. A, 1997, vol. 28A, pp. 1755-68. 10. B. Hutchinson, L. Ryde, and E. Lindh: Mater. Sci. Eng., 1998, vol. A257, pp. 9-17. 11. L. Kestens and J.J. Jonas: Met. Mater., 1999, vol. 5, pp. 419-27. 12. F. Barrato, R. Barbosa, S. Yue, and J.J. Jonas: Proc. Int. Conf. Physical Metallurgy of Thermomechanical Processing of Steel and Other Metals, I. Tamura, eds., ISIJ, Tokyo, 1988, pp. 383-89. 13. F.S. LePera: JOM, 1980, Mar. pp. 38-39. 14. L. Kestens, N. Yoshinaga, D. Vanderschueren, and B.C. De Cooman: IF Steels 2000 Proc., Pittsburgh, PA, June 5–7, 2000, ISS, Warrendale, PA, pp. 271-78. 15. L. Duprez, B.C. De Cooman, and N. Akdut: 6th World DUPLEX Stainless Steel Conf., Venezia, Oct. 17–20, 2000, Associazione Italiana Metallurgia, Milan, Italy, pp. 355-66. 16. F.J. Humphreys and M. Hatherly: Recrystallization and Related Annealing Phenomena, Elsevier Science Ltd., New York, NY, 1995, pp. 44-52.
METALLURGICAL AND MATERIALS TRANSACTIONS A
