MICROSTRUCTURE AND TEXTURE FORMATION OF THIN HOTROLLED STEEL STRIPS

Roumen Petrov*, Leo Kestens*, Rafael Colás** , and Yvan Houbaert*

* Department of Metallurgy and Materials Science, Ghent University, Technologiepark 903,B-9052, Zwijnaarde (Gent), Belgium

** Facultad de Ingeniería Mecánica y Eléctrica, Universidad Autónoma de Nuevo León, A.P.149-F, 66451 San Nicolás de los Garza, N.L., Mexico

SUMMARYPlain carbon steel and an Nb containingsteel were directly hot rolled, after beingcast into slabs of around 50 mm in thick-ness, in an industrial, six-stand continuousrolling mill to a final thickness that variedfrom 2.7 to 1.06 mm. The microstructureand texture of the thin strips was studied bymeans of classical metallography and ori-entation image microscopy (OIM) in themiddle thickness of the strips.

It was found that the average grain diameterof plain carbon hot rolled strips varies from7 to 10 µm and depends on the total amountof reduction, whereas for the Nb steel it de-creases to 3.6 µm.

The middle thickness texture of the Nb steelis typical transformation texture originatingfrom strained austenite, whereas the typeand intensity of the ferrite texture in thestrips depends on the thermo-mechanicalhistory of the steel strip. The increase ofstrain at elevated temperatures leads to in-tensification of the ferrite texture and alltypical transformation components from therecrystallized austenite phase are presented.When the strain at elevated temperature islow the texture is weak and additionaltransformation components emerging fromdeformed austenite are dominant. The dataconfirm that by proper control of the proc-essing parameters it is possible to obtainthin hot rolled strips with fine grains andappropriate texture providing satisfactorylevels of strength and formability.

1. INTRODUCTIONWhen the high surface quality of the thinsteel strips is not of the prime importancefor their application, hot rolled thin stripscould be used as a viable economical alter-native of cold rolled steel strips. This re-placement can be more effective when me-chanical properties of the hot rolled thinstrips, i.e. strength, toughness and form-ability are high enough and are combinedwith low values of anisotropy. Such a com-bination of properties may be achieved bymeans of thermo-mechanical treating thatresults in a fine-grained microstructure withan appropriate texture.

Improvements in thin slab casting, hot di-rect charging together with hot rolling tech-niques and practices have resulted in theproduction at an industrial scale, of hotrolled strips of gauges thinner than 2 mm.Such processing implies that the reductionsof thickness in the originally cast slab ishigher that 96%. Furthermore, hot directcharging of thin cast slab does not permitthe common isotropic transformations thattake place in conventionally processed ma-terial 1).

The goal of the present work is to study theinfluence of processing parameters on themicrostructure and texture of two carbonsteels, one plain and the other with Nb ad-dition, which were cast into thin slabs andhot rolled in a compact strip mill to differ-ent final thickness.

2. EXPERIMENTALTwo low carbon steels, designated as A(plain carbon) and B (with higher Mn andNb contents), were hot rolled in an indus-trial six-stand continuous rolling mill to dif-ferent final thickness directly after beingcast into slabs of around 50 mm in thick-ness. The chemical composition of thesteels is shown in Table 1. Full details ofthe mill can be found elsewhere 1).

Three different rolling schedules were usedwith steel of composition A (identified asA1, A2 and A3), whereas only one sched-ule, similar to that of A2, was applied withsteel B. The samples from either type ofmaterial were cut at room temperature fromthe coiled product. Coiling of the stripstakes place at the end of the run out table.The later was equipped with a cooling sys-tem, which assure coiling temperature be-low 650°C. Some parameters of the rollingschedules – rolling temperature, thicknessand reduction at each stand are shown inFig.1.

Fig.1 Data of the rolling schedules (rollingtemperature, thickness and reduction) forstrips A1, A2 and A3. The Ar3 temperature(860°C) is the dashed line.

The value of critical temperaturesA1=718°C and A3=860°C were calculatedwith the equations proposed by Andrews 2).The temperature changes in the strips dur-ing rolling were predicted by computermodeling 3), as this approach allows to havean idea whether the final rolling tempera-ture (FRT) stays within the austenite phasefield or within the two phase, α+γ, region.The values shown in Fig.1 indicate that thefinal rolling temperatures predicted by themodel for the rolling schedules A1, A2 andB were above A3, without entering the in-

tercritical region, but not for the thinnestsample (A3), on which the final pass mayhave been imparted in the intercritical re-gion (T=850°C).

Evaluation of the forming characteristics ofthe hot rolled steels was carried out fol-lowing the standard procedure for determi-nation of the average value of the normalanisotropy R of the sheets, by means of ten-sile tests as it is reported elsewhere 4, 5). Thesame characteristic was calculated on thebase of the texture measurements using thefull constraints model.

The microstructure of the samples was ex-amined in their cross section after standardsample preparation procedure finished with2 % nital etching. The plane for observationwas perpendicular to the transverse direc-tion of the rolled sheet. In the centre of thisplane (middle thickness of the strip) wereperformed local microtexture measurementsby means of orientation imaging micros-copy (OIM). The OIM attachment was in-stalled on a Philips XL30 ESEM with anLaB6 filament and the electron backscat-tering diffraction (EBSD) patterns were ac-quired and analysed by means of the com-

Chemical composition, mass.%SteelC Mn P S Si Al Nb N

A 0.054 0.199 0.01 0.006 0.015 0.033 0.003 0.0043

B 0.064 0.804 0.01 0.006 0.008 0.033 0.023 0.0055

Table 1. Composition of the steels.

mercial TSL OIM* software 6). The orienta-tion data of at least 3 local measurements ofeach sample were summarized and furtherpost processed by means of FHM-MTMsoftware developed by Van Houtte7) in or-der to represent the texture of the strips bymeans of the orientation distribution func-tions (ODFs). Every ODF created in thisway represents statistically reliable texturedata.

3. RESULTS AND DISCUSSION 3.1. MicrostructureThe microstructure observation (Fig.2a-d)displays that the average grain diameterDav, which was used to define the grain sizeof the strips, changes significantly as afunction of the chemical composition of thesteel and the processing parameters. Fig.3ashows the changes in the grain diameter andthe area fraction of the grains in the mid- * TSL OIM is a trade mark of TexSem LaboratoriesINC., Draper, UT, USA..

thickness of the strips. The differences be-tween the microstructure and grain size inthe mid-thickness and in the surface layersof the strips are negligibly small and theyare shown in Fig.3b.

Strong grain refinement is observed in steelB, which contains 0.023%Nb. The averagegrain diameter (Dav) measured in this steel(cf. Fig. 3a) is 3.6 µm that is 2.5 times

smaller than the average grain diametermeasured in strips A1, A2 and A3 (comparethe data in Fig. 3a). The grains smaller than5µm include between 75% and 95% of allnumber of grains but cover only 30 % ofthe area.

The ferrite grains in plain carbon steel arelarger than those in the Nb containing steelB, and their average diameter vary from 7.2µm, for strip A3, to 10 µm, for strip A1.Metallographic observation did not showsignificant differences between the

Fig.2. Microstructure of the middle thickness of the strips; (a) Steel B, t=2.18 mm; (b)Steel A3, t=1.06 mm; (c) Steel A2, t=1.92 mm; (d) Steel A1, t=2.69 mm

microstructure (grain size and shape) in thesurface and in the mid-thickness of thestrips (Fig.3b), but somehow, the grainsclose to the surface are slightly smaller indiameter than those in the mid-thickness.The grain refining effect, which was ob-served in the plain carbon strips, as a func-tion of the rolling reduction, could be asso-ciated with the changes in the dislocationdensity of the parent austenite phase priorto transformation. An indirect proof of thisassumption is the difference between theaustenite grain size predicted by the com-puter model 3) and the grains measured inthe final ferrite.

Fig.4: Changes in the average austenitegrain diameter DA

3), average ferrite graindiameter DF and the DA/DF ratio as a func-tion of the rolling reduction in the plaincarbon steel strips. The white dots displaythe average ferrite diameter measured bythe linear interception method

Fig 4 displays the ratio DA/DF for the stripsof plain carbon steel A1, A2 and A3, whereDA is the average austenite grain diameterbefore the transformation predicted by themodel 3) and DF is the size of the ferrite. Itcan be considered that with increase of thisratio more ferrite grains (with a small di-ameter) are nucleated from a single aus-tenite grain, assuming that there is no graingrowth during coiling. Because coiling wasexecuted at temperatures below 650°C,such assumption seems reasonable. There-fore, the number of ferrite grains emergingfrom a single austenite grain will be a func-tion only of the deformed substructure ofthe austenite, and can be used as a qualita-tive characteristic of the accumulated plas-tic deformation in the parent austenite priorto transformation. The data in Fig. 4 showthat austenite grains of a given size produceferrite grains of smaller size as the total ac-cumulated strain increases, i.e. the DA/DF

ratio increases with an increase of the roll-ing reduction.

3.2. Texture.The texture of strip B (Fig. 5a), which wascalculated on the basis of 5 different localmeasurements, displays two importantcomponents of ferrite that were generatedfrom deformed austenite. The {332}�113�ferrite texture component, which is consid-ered as a favorable orientation from theviewpoint of drawability8,9) (high R-value)

a b

Fig.3: (a) Area fraction of the grains plotted vs. grain diameter for steels B(Dav=3.6µm), A3 (Dav=7.2µm) and A1 (Dav=10µm); (b) Grain diameter measured inthe center of the strip ad in the edge zone (strip A2, Davcenter=9µm, Davedge=7µm).

is presented with an intensity of 2.5X. The{112} �131� component, which also repre-sents a transformation product from de-formed austenite, appears with a maximumintensity of 3.2X. This component, togetherwith the rotated cube that was observedwith an intensity of 2 X, are considered tohave a detrimental effect on the R-value 9) .

Figure 5: Mid-thickness texture of the strips(a) B, t= 2.18mm; (b) A3, t=1.06mm; (c)A2, t=1.92mm; (d) A1, t=2.69mm

The analysis of the distribution of the tex-ture components among grains of differentsize was done by calculating the ODF fromgrains smaller than 3 µm and larger than 10µm, as they can be clearly discriminated insteel B. The ODF plots shown in Fig. 6cand d correspond to the smaller and largergrains shown in Fig. 6a and b, respectively.The small and large ferritic grains both dis-play the transformation texture originatingfrom deformed parent γ phase. The trans-formation texture components of the de-formed austenite {332}�113� (shifted to{554} �225�), {111}�112� displaying an in-tensity of 6X and {001}�010� with an inten-sity of 4 X are the strongest components inthe ODF of the grains larger than 10 µm,whereas the {001}�110� and {110}�110�components appear with intensities of 3X,respectively (Fig.6 d). According to someauthors 9,10,11,) these texture componentscould emerge from either deformed or re-crystallized austenite but the later is less

probable, taking into account the Nb con-tent of steel B and the rolling parameters.

Fig.6: Image quality maps showing thegrains smaller than 3 µm (a) and larger than10 µm (b) together with the correspondingODFs gathered from small (c) and large (d)grains

The ODF of the grains smaller then 3µm(Fig. 6c) displays almost equal intensity ofboth ferrite texture components that couldemerge from deformed austenite({223} �110� and {332}�113�) with intensityof 2~3X and from recrystallized austenite({001} �110� and {110}�110�) with intensityof 2~3X.

Considering the data of the grain size di-ameter and how the texture components aredistributed among them it can be concludedthat the strong deformed grains are amongthe first to transform into ferrite and hence,they have a growth advantage with respectto later transformed grains that might haveemerged from the partially recrystallizedaustenite.

Fig.7: Modeled ferrite transformation tex-tures obtained after γ/α transformation of γphase with (a) Brass, Copper and (b) Brass,Copper and Goss orientation

A simple modeling of ferrite transformationtextures produced after γ/α phase transfor-mation regarding the K-S crystallographicorientation relationship (i.e. rotation of 90°about �112� axis) is shown in Fig.7. The γparent phase with Brass {110}�112�, Cop-per {112}�111� and Goss {110}�001�,which are characteristic texture componentsof strained austenite orientations, gives riseafter the transformation to the same trans-formation texture components that wereobserved in the industrially hot rolled thinstrip B. It should be mentioned that almostexact correspondence between measuredand modeled texture were found in the largeferrite grains. (Compare the textures inFig.6d and Fig.7b.).

The textures of plain carbon steel strips areshown in Fig. 5b, c and d. In general thetexture of the strips A3 and A2 differ sig-nificantly from the texture of the Nb steel,whereas the texture of strip A1 looks simi-lar to it. The intensity of the texture in-creases with the increase of the total rollingreduction and the thinnest strip A3 displaysthe strongest texture. Two main groups oftexture components are presented in theODFs of the plain carbon steel. The first

group corresponds mainly to transformationproducts emerging from recrystallized aus-tenite (i.e. rotated cube, rotated Goss andGoss). The intensity of the rotated cubecomponent is stronger in the thinner stripsA3 and A2 displaying a value of 8 and 6.4X respectively and only 2X in the thickeststrip A1. The rotated Goss and Goss com-ponents are concurrently present only in thetexture of strip A2, whereas the rotatedGoss component with a certain deviationfrom the exact position is observed also inthe texture of the thinnest strip A3. Thetransformation components from deformedaustenite are weak in the textures of thestrips A3 and A2 but they are the strongestcomponents in the texture of the strip A1with a thickness of 2.69 mm, which dis-plays the highest R-value.

Hence, the improvement of the R-value inthe thickest strip is due to the two main rea-sons: (i) a general decrease of the intensityof unfavorable components and (ii) the de-velopment of the {111}�112� texture com-ponent which has a positive effect on thedeep drawability 9). The calculation of theR-values based only on the texture meas-urements supports this assumption(cf.Fig.9).

Fig.8: True strain per pass for the differentstrips a function of the rolling temperature

Fig.9: Measured and modeled data for theaverage normal anisotropy (R-value) of thestrips

For better understanding the texture forma-tion in hot rolling of thin strips, it is neces-sary to know also the strain distributionamong the different passes. Fig. 8 displaysthe evolution in the reduction of the stripsthickness presented by the true strain forevery rolling pass. Strips A2 and A3 weresubjected to higher reductions at elevatedtemperatures than strip A1 but in the finalpass all strips received an almost identicalrolling strain of 0.25. If the reduction atelevated temperatures is higher the intensityof the ferrite transformation texture fromrecrystallized austenite is stronger (Cf.Fig.8 and 5b, c, d), because the high tem-perature deformation intensifies the recrys-tallization of the austenite. Hence, takinginto account the texture data in Fig. 5b,c,dand the data in Fig.8 it is possible to assumethat recrystallization in the thickest strip A3develops weaker than in the thinnest strips,as far as it was subjected to the lower strain.In a later stage it will be possible to employthe present data to extend the γ-recrystallization model 3) to a fully quanti-tative texture dependant description

4. CONCLUSIONSThe microstructure and texture of thin stripsproduced by hot rolling of hot charged-thincast slabs is studied in this work. Two typesof steels were studied, plain carbon steel,subjected to total hot rolling reduction of 94to 98%, and Nb containing steel finish

rolled to a total reduction of 94%. It wasfound that the final grain size on hot rolledplain carbon steels ranged from 7 to 10 µm,was dependant on the total amount of re-duction during hot rolling. Significant grainrefining was observed to occur in the Nbcontaining steel in which the average grainsize was 3.6 µm.

A low intensity transformation texture fromdeformed austenite was observed in the Nbcontaining steel, whereas the type and in-tensity of ferrite texture in the plain carbonsteels depends on the thermo-mechanicalhistory of the steel strip. When the reduc-tion at elevated temperatures is higher theferrite texture intensifies and possesses alltypical transformation components from therecrystallized austenite phase. When thestrain at elevated temperature is low thetexture is weak and additional transforma-tion components emerging from deformedaustenite are dominant. The observed tex-ture data of plain carbon steel are in goodcorrelation with the results from R-valuemeasurements.

The data confirm that by proper control ofthe processing parameters it is possible toobtain thin hot rolled strips with fine grainsand appropriate texture providing satisfac-tory deep drawability.

Acknowledgements:The authors gratefully acknowledge the fi-nancial support of this study granted by theNetherlands Institute for Materials Re-search (NIMR) and Hylsa, S.A. de C.V. isacknowledged for providing the material.

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