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C. Ghosh et al.: Microstructure, texture, grain boundary characteristics and mechanical properties of a cold rolled and

Chiradeep Ghosh, Arunansu Haldar, Pampa Ghosh, Ranjit Kumar Ray R & D Division, Tata Steel, Jamshedpur, Jharkhand, India

Microstructure, texture, grain boundary characteristics and mechanical properties of a cold rolled and annealed ferrite–bainite dual phase steel The microstructural, textural evolution and changes in grain boundary character distribution during annealing of a prior cold worked (30 %, 50 % and 80 %) ferrite–bainite dual phase steel have been studied and correlated with mechanical properties. It has been shown that submicron sized subgrains can be obtained by selecting the appropriate amount of cold rolling and annealing cycle. Increasing the annealing temperature in all the materials produces the expected results, namely decrease in strength with a simultaneous increase in ductility. Although reasonably sharp c-fibres were obtained in 80 % cold rolled and its 500 8C annealed counterpart, the very low r values (< 1.0) make the steel unsuitable for the purpose of deep drawing. It is envisaged that grain boundary engineering may lead to better strength-ductility combinations in this steel for an enhanced range of applications. Keywords: Microstructure; Texture; Grain boundary character distribution; Cold rolling; Mechanical properties

1. Introduction Conventional dual phase steels are a class of high strength steels that are characterized by a microstructure essentially consisting of a dispersion of 20 to 30 % of a hard phase in a soft ductile ferritic matrix. The dispersed hard phase could be either martensite or bainite [1]. In general dual phase steel can be made in two different ways: (i) during cooling after hot rolling [2] and (ii) after intercritical annealing or holding in the isothermal bainitic transformation temperature range of a previously rolled product [3, 4]. In either case, cooling at an appropriate rate and range essentially controls the transformation of a part of austenite to ferrite and of the rest to martensite or bainite [5]. Although the literature on ferrite–martensite dual phase steels is vast, comparatively less work has been carried out on the ferrite–bainite steels [6 – 13]. In the present investigation an attempt has been made to understand the effect of different amounts of cold deformation followed by annealing on a ferrite–bainite dual phase steel in terms of microstructural and textural changes and grain boundary char-

acter distribution. The relevant mechanical properties have also been evaluated with an eye to possible application of such steels. Finally, suggestions have been made about the ways and means for further improvement of mechanical properties of this class of steels.

2. Experimental procedures Investigations have been carried out on a commercial low carbon ferrite–bainite steel in the form of a hot rolled plate, with chemical composition as given in Table 1. Pieces were cut from the initial plate to the dimensions 5 mm · 25 mm · 200 mm (thickness · width · length). Three different cold rolling reductions were used, namely, 30 %, 50 % and 80 % reduction. The rolling was carried out using a two high laboratory rolling mill with a roll diameter of 250 mm. The rolling was performed at a roll peripheral speed of 27.5 m min – 1 with machine oil as lubricant. The cold rolling was successfully done without any edge cracking. Specimens 15 mm in width and 100 mm in length were cut from the cold rolled sheets along the rolling direction (RD) for subsequent annealing and mechanical testing. The annealing of the cold rolled specimens was carried out at different temperatures ranging from 200 – 700 8C, for 30 min with an interval of 100 8C. This was done to find out if satisfactory strength-ductility combinations may be achieved in this class of steels by recovery annealing and / or recrystallization. From the annealed samples tensile specimens were prepared in accordance with ASTM E8M-04 standard and were subsequently polished. Tensile tests of polished samples were carried out at room temperature in an Instron 4210 machine with a constant cross-head speed of 0.5 mm min – 1. The bulk hardness values (Hv) were calculated from the average of seven separate measurements taken at randomly selected points by using a load of 5 kg for 15 s. A detailed microstructural characterization was performed using optical microscopy, field emission gun scanning electron microscopy (FEG-SEM) and transTable 1. Chemical composition in wt.%, (balance Fe) C

Mn

Si

Cr

V

S

P

0.17

1.6

0.52

0.083

0.13

0.014

0.018

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mission electron microscopy (TEM). Specimens after etching with 2 % Nital, were observed along the cross section in a FEG-SEM operated at 15 kV. Thin foil TEM study was conducted in a PHILIPS CM200 TEM operated at 100 kV. Thin foils were prepared by twin jet electro polishing using a solution of 10 % HClO4 + 90 % CH3COOH. Global texture measurements were carried out on the half thicknesses of the 30 %, 50 % and 80 % cold rolled as well as annealed (at 500 8C and 700 8C) samples using X-ray diffraction (XRD) method in a PANalytical MRD machine. The (110), (200), (112) and (222) pole figures were determined, from which the orientation distribution functions (ODFs) were calculated using the method of Bunge [14]. Grain boundary characterization was carried out in the FEG-SEM operated at 20 kV equipped with an orientation imaging microscopy/electron back scattered diffraction (OIM / EBSD) attachment. Samples after polishing with colloidal silica, were put in the sample holder of the SEM/ EBSD equipment. Depending on the fineness of the structure, step sizes ranging from 0.5 – 1.0 lm were used and a number of different areas were selected for scanning. Two typical grain boundary maps by EBSD are shown here (Fig. 1a and b). All the maps possess high band contrast (BC) and band slope (BS) and low mean angle deviation (MAD), indicating high CI (confidence index) value and image quality. The final scanned data were processed using HKL software to calculate the grain boundary character distribution (GBCD) charts. For the purpose of this paper low angle grain boundaries (LAGBs) have been defined to include those boundaries with less than 158 grain boundary angle. Boundaries with greater than 158 grain boundary angle have been termed as high angle grain boundaries (HAGBs). The Brandon criterion has been followed in identifying the CSL (coincidence site lattice) boundaries and the maximum deviation from the exact CSL values has been taken as p15ﬃﬃR8 [15]. The CSL boundaries considered in this paper range between R3 and R29b.

3. Results 3.1. Microstructural changes The La Pera etched optical micrographs of the hot rolled ferrite–bainite, and the 30 %, 50 % and 80 % cold rolled steels are presented in Fig. 2a – d. No trace of martensite could be found in any of these steels, i. e. these are truly ferrite–bainite steels. After cold rolling a distinct deformed structure is obtained and this shows a gradual alignment along the rolling direction, with increasing amount of cold deformation level. The structure becomes visibly finer as the amount of cold deformation increases. A detailed FEG-SEM study of the microstructural features of the cold rolled and annealed steel samples shows significant changes only after annealing at 500 8C and above, for the 30 % and the 50 % cold rolled steels. For the 80 % cold rolled samples distinct changes in the microstructure appear after annealing at 400 8C. Figure 3a – d depict

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(b) Fig. 1. Grain boundary maps for (a) the 30 % cold rolled steel; (b) the 50 % cold rolled steel which has been annealed at 500 8C; Red: High Angle Boundary (> 158); Green: Low Angle Boundary (< 158); Black: Coincidence Site Lattice (CSL) Boundary.

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(d) Fig. 2. Optical microstructure of (a) hot rolled; (b) 30 %; (c) 50 % and (d) 80 % cold rolled steel.

Int. J. Mat. Res. (formerly Z. Metallkd.) 101 (2010) 10


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the FEG-SEM microstructures of the 30 % cold rolled steel annealed at 500 8C, 600 8C and 700 8C. The structures essentially show a mixture of ferrite and bainite. In this low Si steel there is a likelihood of cementite existing along with bainite. The microstructure (Fig. 3d) in fact indicates the presence of pearlite at a few places. With increasing annealing temperature the ferritic areas grow at the expense of the bainitic regions. Figure 4a – d shows the FEG-SEM microstructures of the 50 % cold rolled steel annealed at 500 8C and 700 8C. Distinct ferritic and bainitic regions can be observed at 500 8C (Fig. 4a). At this temperature the bainitic plates show a dis-

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cernible change to a somewhat globular shape. The carbide particles produced by the breakdown of the bainite, decorate the interiors of the ferritic grains (Fig. 4b). There is not much perceptible change in the microstructure after annealing at 600 8C. At 700 8C the microstructure mainly consists of ferritic areas with some fine precipitates inside, the bulk of the carbides (lumpy in shape) decorate the boundaries of individual ferrite grains (Fig. 4c and Fig. 4d). Figure 5a – f show the FEG-SEM microstructures of the 80 % cold rolled steel annealed at 400 8C, 500 8C, 600 8C and 700 8C. Annealing at 400 8C results in a highly deformed and elongated ferrite–bainite structure aligned

Fig. 3. FEG-SEM micrographs of 30 % cold rolled and (a) 500 8C; (b) 600 8C; (c) and (d) 700 8C annealed steel.

Fig. 4. FEG-SEM micrographs of 50 % cold rolled and (a) and (b) 500 8C; (c) and (d) 700 8C annealed steel.

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along the rolling direction, showing signs of recovery at many places (Fig. 5a). The elongated nature of the microstructure still persists even after annealing at 500 8C (Fig. 5b). Breakdown of some of the bainitic plates with signs of globularization can be observed at this stage (Fig. 5c). Mostly fine precipitates of carbides within the ferrite grains as well as on the grain boundaries can be observed after annealing at 600 8C (Fig. 5d). Figure 5e illustrates the microstructure of the 80 % cold rolled and annealed sample at 700 8C. Here it is clearly evident that the bainitic areas have not totally broken down although lots of carbides have formed and grown to bigger sizes. The microstructure is basically an aggregate of ferrite grains with precipitation of small as well as of large sizes residing mostly at the grain boundaries. All these features are shown on a magnified scale in Fig. 5f. In order to examine the microstructures in greater detail, TEM of thin foils from a selected number of samples was

undertaken. Although there are minor variations in the microstructure over extended areas in any one sample, the images reported here represent the predominant features only. Figure 6a – c shows TEM micrographs of selected positions of 30 %, 50 % and 80 % cold deformed steel. The 30 % cold rolled steel shows a lightly deformed duplex ferrite–bainite structure (Fig. 6a). Sheaves of bainitic plates in several physical orientations can be seen in the foil from 50 % cold rolled steel. In the 80 % cold rolled steel the ferritic areas are highly dislocated and show cell formation. At a few places micron size grains (or subgrains) without any dislocation inside can also be seen. Some typical microstructural features obtained after annealing the cold rolled steels are illustrated in Fig. 7a – d. Figure 7a shows the TEM micrograph of the 30 % cold rolled material annealed at 400 8C. Typical cell formation in the 80 % cold rolled material, even after annealing at a low temperature of 200 8C, is depicted in Fig. 7b. In a

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Fig. 5. FEG-SEM micrographs of 80 % cold rolled and (a) 400 8C; (b) and (c) 500 8C; (d) 600 8C (e) and (f) 700 8C annealed steel.

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Fig. 6. Transmission Electron Micrographs (TEM) of steel at (a) 30 %; (b) 50 % and (c) 80 % cold reduction.

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mixed ferrite–bainite area, extensive recovery leading to the formation of large number of subgrains in the deformed ferrite can easily be noticed (Fig. 7c). Formation of submicron sized subgrains can be easily seen in the 80 % cold rolled steel after annealing at 400 8C. The TEM micrographs clearly show the formation of substructures in the ferrite after low temperature annealing of the cold rolled steels, which could not be observed in the FEG-SEM micrographs taken at these stages. 3.2. Textural changes The u2 = 458 sections of the ODFs of 30 %, 50 % and 80 % cold rolled samples are shown in Fig. 8a – c. All the three materials show a fibre running from u1 = 08 to u1 = 908 at u = 08 in the u2 = 458 section. In addition, the 80 % cold rolled sample shows a well developed c-fibre. The u2 = 458 sections of the ODFs of 30 %, 50 % and 80 % cold rolled materials after annealing at 500 8C and 700 8C are presented in Fig. 8d – i. It is evident from these figures that c-fibre formation takes place in all the cold rolled samples after annealing at 500 8C. This fibre is most intense in the 80 % cold rolled material. Some a-fibre along with the rotated cube component is also present in the 500 8C annealed materials. In general, the textural intensity decreases after annealing at 700 8C. Int. J. Mat. Res. (formerly Z. Metallkd.) 101 (2010) 10

Fig. 7. Transmission Electron Micrographs (TEM) of steel at (a) 30 % cold rolled and annealed at 400 8C; (b) and (c) 80 % cold rolled and annealed at 200 8C and (d) 80 % cold rolled and annealed at 400 8C.

The 30 % cold rolled material at this stage does not show any c-fibre at all. All these features connected with the cold rolled as well as the annealed samples are quite apparent from the c-fibre plots for the different cold rolled and annealed steels, as shown in Fig. 9. The strongest c-fibre is obtained for the 80 % cold rolled material followed by its 500 8C annealed counterpart. These are followed by the 30 % and 50 % cold rolled materials annealed at the same temperature. Over all the c-fibres are not very intense, the maximum intensity *5.5 times random for the 80 % cold rolled steel and *4 times random for the 80 % cold rolled and 500 8C annealed material. The c-fibres of the 700 8C annealed steels are nearly of random intensity. The calculated r values for all the cold rolled and annealed samples have been found to be less than 1.0. The procedure for calculating the r value from texture is given in reference [16]. 3.3. Changes in grain boundary character distribution (GBCD) Figure 10a – c represents the grain boundary character distributions of the initially hot rolled ferrite–bainite steel as well as the 30 %, 50 % and 80 % cold rolled materials. From Fig. 10a it is evident that the low angle grain boundary (LAGB) fraction increases abruptly after 30 % cold rolling 5

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and then drops down and remains practically constant after 50 % and 80 % of cold rolling. The reverse is the case with the high angle grain boundary (HAGB) fraction as depicted in Fig. 10b. After 30 % cold rolling the HAGB number fraction suddenly drops, followed by an increase after 50 % cold rolling which practically remains unchanged after 80 % cold rolling as well. As far as the coincidence site lattice (CSL) boundaries are concerned, the initial hot rolled material shows the highest number fraction of such boundaries. It decreases significantly after 30 % cold rolling, and increases again after 50 % cold rolling and then remains unaltered after 80 % cold rolling. It is interesting to note here that the number fractions of LAGB, HAGB and CSL

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Fig. 8. u2 = 458 section of the ODFs of (a) 30 %, (b) 50 % and (c) 80 % initially cold rolled (d) 30 %, (e) 50 % and (f) 80 % initially cold rolled and annealed at 500 8C and (g) 30 %, (h) 50 % and (i) 80 % initially cold rolled and annealed at 700 8C steels.

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(c) Fig. 9. Variation in c-fibre for the different cold rolled and annealed steels.

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Fig. 10. Grain Boundary Character Distribution (GBCD) of the steel: (a) LAGB, (b) HAGB; and (c) CSL distribution in the initial hot rolled condition, after 30 %, 50 % and 80 % cold rolling.



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boundaries individually are practically the same for both 50 % and 80 % cold rolled materials. Figure 11a – c represent the number fraction distribution of the different types of boundaries in the 30 %, 50 % and 80 % cold rolled materials which were subsequently annealed at 500 8C and 700 8C. In the annealed condition, the LAGB number fractions are always higher in the 500 8C annealed material than in the 700 8C annealed samples. Reverse is the case with the HAGB number fractions. The CSL number fractions are always on the lower side for 500 8C annealing temperature in comparison with the 700 8C annealing temperature.

Figure 12a – c represent the frequency % of different CSL boundaries ranging from R3 to R29b in the 30 %, 50 % and 80 % cold rolled as well as the subsequently annealed (at 500 8C and 700 8C) materials. In general, the largest frequency of CSL boundaries in all the cases is of R3. Invariably, the R3 fraction is much higher after annealing at 700 8C than after annealing at 500 8C. After annealing at 700 8C the R3 boundary frequency is highest in the 80 % cold rolled steel followed by the 50 % and 30 % cold rolled materials. The fraction of CSL boundaries other than the R3 is also much higher in the 700 8C annealed materials in comparison to the 500 8C annealed samples. The boundaries con-

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Fig. 11. The number fraction distribution of different types of boundaries in the (a) 30 %; (b) 50 % and (c) 80 % cold rolled and annealed at 500 8C and 700 8C steel.

Fig. 12. The frequency percentages of CSL boundaries in the (a) 30 %; (b) 50 % and (c) 80 % cold rolled and annealed at 500 8C and 700 8C steel.


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cerned are R5, R7, R9, R11, R17b, R21b and R25b. Some new CSL boundaries also get created after annealing at 700 8C temperature, predominantly in the 80 % cold rolled material. 3.4. Mechanical properties Figure 13a – c represents the engineering stress–strain diagrams of the 30 %, 50 % and 80 % cold rolled materials which were subsequently annealed at 200 8C, 300 8C,

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(c) Fig. 13. Engineering stress-strain diagram of (a) 30 %; (b) 50 % and (c) 80 % cold rolled and annealed at 200 8C, 300 8C, 400 8C, 500 8C, 600 8C and 700 8C steel.

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(d) Fig. 14. Mechanical properties of the annealed steels (a) UTS, (b) % elongation, (c) failure-ductility value and (d) hardness.



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400 8C, 500 8C, 600 8C and 700 8C. All these plots clearly indicate the expected trend of decreasing UTS and increasing % elongation values with increase in the annealing temperature. All the plots show a distinct yield point behaviour. Figures 14a and b show the UTS and percentage elongation values of the cold rolled steels after annealing at different temperatures. The UTS value at any particular annealing temperature is the highest for the 80 % cold rolled steel and lowest for the 30 % cold rolled material. The UTS of the 50 % cold rolled steel lies within these two extremes. Understandably, the UTS values show a progressive decline with annealing temperature for any particular cold rolling level. Again, the % elongation achieved in the 80 % cold rolled steel is also the highest whereas the 30 % cold rolled steel shows the lowest % elongation values. The 50 % cold rolled steel exhibits % elongation values which lie within the two extremes. Obviously, the % elongation values, for any of the cold rolled level, increase with the increase in annealing temperature. Figure 14c shows a plot of failure-ductility values (UTS · % Elongation) for the steel, at the three different cold rolling levels, as a function of annealing temperature. It is clear from this figure that the highest values for this parameter at any annealing temperatures are achieved for the 80 % cold rolled material. The bulk hardness values of the steel rolled at different levels, as a function of annealing temperature, are shown in Fig. 14d. The hardness plot shows a trend similar to the UTS plot, shown in Fig. 14a.

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(b) Fig. 15. Variation of n-values in 30 %, 50 % and 80 % cold rolled steels and annealed at (a) 500 8C and (b) 700 8C temperature.


The plots of the strain hardening exponent (n) values, calculated from the tensile stress–strain data, of the three cold rolled steels which were annealed at 500 8C and 700 8C are shown in Fig. 15a and b, respectively. These clearly shows that the 80 % cold rolled and annealed steel shows the highest n values amongst the three different cold rolled steels, after annealing at 500 8C as well as at 700 8C.

4. Discussion The experimental steel shows a ferrite–bainite dual phase microstructure in the hot rolled condition. On annealing the cold deformed material there is a gradual breakdown of the bainitic plates, as shown in the microstructures of Figs. 3, 4 and 5. The net result is that the ferritic areas increase at the expense of bainite as the annealing temperature increases. Annealing at the higher temperatures of 600 8C and 700 8C leads to the formation of mainly ferritic grains, with the remnants of the bainite and also some isolated bainitic carbides decorating the grain boundaries. Some fine precipitates (of carbides) are also sometimes observed in the grain interiors. Thus the annealed materials also show an essentially ferrite–bainite dual phase microstructure along with some carbides. Saha Podder et al. [5] have reported that the presence of small amounts (2 % and above) of martensite in the microstructure of hot rolled ferrite–bainite dual phase steels is sufficient to produce the continuous yielding behaviour in tensile tests, which is characteristic of the conventional ferrite–martensite dual phase steels. The La Pera etched microstructures of the steel used in the present investigation did not reveal the presence of any trace of martensite in the matrix either in the hot rolled or in the cold rolled and annealed condition (Fig. 2). This could be the reason why none of the stress–strain diagrams of the present steel shows any continuous yielding behaviour (Fig. 13). It is therefore quite clear that the ferrite–bainite dual phase steels behave differently from the ferrite–martensite dual phase steels. In fact, similarly treated ferrite–martensite dual phase steels have shown continuous yielding behaviour [17]. The initial hot rolled steel possesses yield strength of 420 MPa, UTS of 570 MPa and has 24 % elongation. Subjecting the steel to cold rolling and annealing yields substantial improvement in strength with some decrease in ductility (Fig. 14a and b). However, depending on the requirement, a suitable combination of high strength coupled with reasonable ductility can be achieved. For example, when the 80 % cold rolled material is annealed at 400 8C, a strength-elongation combination of 950 MPa-11 % can be obtained. Again, a strength-elongation combination of 875 MPa-15 % can be achieved when the 80 % cold rolled steel is annealed at 500 8C. Thus, such steels can be used in those applications which require high strength but not a very high level of ductility. The strain hardening exponent (n) values (Fig. 15) show that for the materials annealed at both 500 8C and 700 8C, there is a general increase in the n-value with an increase in the level of cold deformation. A reasonably high value of n = 0.23 has been obtained for the 80 % cold rolled steel when annealed at 500 8C. A slightly higher n = 0.25 has been obtained in the same material when annealed at 700 8C. This shows that such steels will exhibit a reasonable amount of necking resistance during tensile loading. 9

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The textural data reveal that a strong c-fibre forms in the 80 % cold rolled material; the c-fibres obtained for the 50 % and the 30 % cold rolled samples have only random intensity (Fig. 9). Annealing of the 80 % cold rolled sample leads to a decrement in the overall intensity of the c-fibre. In general, for any of the cold rolling level, the 500 8C annealed material shows a stronger c-fibre than the 700 8C annealed sample. Attainment of a moderately strong c-fibre is considered a pre-requisite for obtaining a high r value. Although reasonably strong c-fibres have been obtained for the 80 % cold rolled as well as for the 80 % cold rolled and 500 8C annealed materials, the calculated r values for these as well as for the other heat treated samples have always been found to be less than 1.0. This clearly shows that these ferrite–bainite steels have poor deep drawing characteristics after cold rolling and annealing, similar to those of ferrite–martensite dual phase steels [18]. Therefore, such materials may not be considered in applications where a high level of deep drawing is required. In recent years extensive research on obtaining ultrafine grained steels with minimum alloy costs has been actively pursued. One such important method consists of advanced thermomechanical processing involving deformation-induced grain subdivision. Several studies have been reported on the grain refinement and strengthening of cold rolled and annealed Fe–C based martensitic steels at room [17, 19] and elevated temperatures using this technique [20]. The possibility of formation of submicron sized grains by the combined actions of cold rolling and annealing has also been studied in the present ferrite–bainite dual phase steel. In fact, TEM studies have revealed the formation of grains in the size range 150 – 350 nm in the material cold rolled 80 % and then annealed at 400 8C (Fig. 7d). Saha and Ray [21] claim to have obtained submicron to nano sized (< 100 nm) grains by subjecting a hot rolled IF steel to 98 % and 99.5 % cold rolling. Severe cold rolling of this steel from 90 % to 98 % led to an increment of the HAGB fraction from 0.39 to 0.50 in their case. Similarly, cold rolling of a martensitic steel by 80 % led to the formation of grains as small as *50 nm size and corresponding HAGB fraction has been found to be *0.6 [17]. In both the above cases, formation of nano sized grains takes place after the process of cold deformation itself. Extensive TEM work has failed to reveal the presence of nano sized grains (< 100 nm) in the present steel after heavy cold working. The HAGB fraction of the present experimental 80 % cold rolled steel has been found to have a maximum value of *0.4, which is less than the values of HAGB fractions obtained for the two previous cases. It is therefore quite likely that the HAGB fraction of the present cold worked steel is not of a sufficiently high value to trigger the formation of nano sized grains (< 100 nm). The GBCD plots (Fig. 10) reveal an interesting phenomenon. Although microstructural observations (Fig. 2) have shown that the grain structure becomes finer and more aligned with the increase in cold rolling deformation, the number fractions of LAGB, HAGB and CSL boundaries remain remarkably similar (almost constant) for the 50 % and the 80 % deformation levels. It is quite expected that as deformation increases from 50 % to 80 %, more and more HAGBs will form due to dislocation accumulation. However, at the same time, a remarkable sharpening of the crystallographic texture has also been observed as the amount 10

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(c) Fig. 16. Variation in (a) % Elongation; (b) UTS and (c) n-values of 30 %, 50 % and 80 % cold rolled steels annealed at 500 8C and 700 8C temperature.



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of deformation increases from 50 % to 80 % (Figs. 8 and 9). It is known that a larger number fraction of LAGBs is associated with a stronger texture, where the misorientations between the neighbouring grains are expected to be much less. It therefore appears that as the cold deformation level increases from 50 % to 80 %, both HAGBs and LAGBs increase in such a manner that their relative number fractions do not change much. It must be kept in mind, however, that the misorientation evolution of existing boundaries is essentially controlled by grain rotations due to orientation and compatibility requirements, and that new LAGBs are continuously formed filling up the statistics at the lower edge. A second interesting observation emanating from the GBCD plots (Fig. 11) is that the number fractions of the LAGBs for the annealed samples always show a larger value after annealing at 500 8C than at the higher temperature of 700 8C. This can also be explained by the fact that the overall texture and especially the c-fibre is much sharper in the former, as compared to in the latter case. Finally, a word must be said about the distribution of the CSL boundaries in the deformed and in the annealed materials. The relevant plots (Figs. 10 and 12) clearly indicate that the CSL boundary number fractions are nearly the same for both the 50 % and the 80 % cold rolled materials, and this value is more than twice the value for the 30 % cold rolled steel. The most prominent CSL boundary in all the cases is R3. The frequency % of R3 boundaries is nearly the same in both the 50 % and the 80 % cold rolled materials, and this value is nearly twice the value for the 30 % cold rolled steel. Thus, the variation in the number fraction of CSL boundaries in the different cold rolled materials is almost directly related to the variation in the frequency % of the R3 boundaries. Randle and Davies [22] have suggested that a higher proportion of R3 boundaries in alpha-brass leads to a general increase in the strain-to-failure (ductility) values, with only a minor drop in strength values. In Fig. 16a – c the % elongation, UTS and strain hardening exponent (n) values for the heat treated samples have been plotted against the fraction of R3 boundaries, the largest fraction among all the CSL boundaries. Both the % elongation and strain hardening exponent show higher values for higher R3 fractions, whereas UTS decreases with higher R3 fractions. The utility of these steels can be vastly improved if by some means ductility can be enhanced without sacrificing the strength to a large extent. Obviously, annealing at higher and higher temperatures cannot be a feasible solution. In this connection, application of grain boundary engineering (GBE), as suggested by Randle and Davies [22] using iterative strainrecrystallization cycles is worth investigating. It is envisaged that subjecting the present steel to iterative processing may lead to a beneficial engineering of the grain boundaries, thereby increasing the overall ductility of the steel, without sacrificing the strength to a large extent.

5. Conclusions The following conclusions can be drawn from the present study: 1. Dual phase microstructures consisting of ferrite and bainite, with different distributions and morphologies, can be produced by annealing of the cold rolled steel. Int. J. Mat. Res. (formerly Z. Metallkd.) 101 (2010) 10

2. With the increase in annealing temperature decrease in strength and simultaneous increase in ductility occur, as expected. 3. The 80 % cold rolled material and its 500 8C annealed counterpart show a reasonably sharp c-fibre, but the r values are quite low (< 1.0) in all the cases. 4. The ferrite–bainite dual phase steels do not show continuous yielding behaviour, in contrast with the behaviour of ferrite–martensite dual phase steels. 5. Favourable strength-elongation combinations can be obtained in these steels by suitable adjustment of the cold rolling level and the annealing parameters. 6. These steels can be used in those applications which require high strength but not a very high level of ductility. In addition, such steels may not be considered in applications where a high level of deep drawability is required. 7. Grain boundary engineering may be a good option to produce better strength-ductility combination in such steels. The authors wish to thank the management of Tata Steel, Jamshedpur, India for allowing this work to be published. References [1] A. Murugaiyan, A. Saha Podder, A. Pandit, S. Chandra, D. Bhattacharjee, R.K. Ray: ISIJ Int. 46 (2006) 1489. DOI:10.2355/isijinternational.46.1489 [2] A.P. Coldren, G. Tither: JOM 30 (1978) 6. [3] M. Rashid: SAE Paper 770164 presented at Int. Automotive Engineering Cong. and Exposition, Detroit, Michigan, (1977). [4] S. Hayami, T. Furukawa: Proc. Microalloying Conf., Vanitec, London, (1975) 78. [5] A. Saha Podder, D. Bhattacharjee, R.K. Ray: ISIJ Int. 47 (2007) 1058. DOI:10.2355/isijinternational.47.1058 [6] M. Sudo, S. Hashimoto, T. Hosoda, Z. Shibata, K. Hirata: 1983 R and D: Research and Development Kobe Steel Engineering Reports 33 (4) 49. [7] T. Waterschoot, B.C. De Cooman, D. Vanderschueren: Ironmaking and Steelmaking. 28 (2001) 185. DOI:10.1179/030192301677948 [8] M. Cai, H. Ding, J. Zhang, L. Li, Z. Tang, C.Y. Xuebao: Chinese Journal of Materials Research. 23 (2009) 83. [9] M.M. Karimi, S. Kheirandish: Steel Research Int. 80 (2009) 160. DOI:10.2374/SRI08SP082 [10] Hongtao, Zhang, Chunfu, Huang, Ganyun, Pang: Int. Symp. on Low-Carbon Steels for the 90’s, 1993, 367. [11] M.H. Cai, H. Ding, J.S. Zhang, Z.Y. Tang, Wang, W.W. Kang: T’ieh/Iron and Steel (Peking) 43 (2008) 77. [12] G.P. Yun: Zhuzao Jishu/Foundry Tech. 27 (2006) 347. [13] O. Yakubovsky, N. Fonstein, C. Cheng, D. Bhattacharya: Galvatech ’04: 6th International Conference on Zinc and Zinc Alloy Coated Steel Sheet – Conference Proceedings, 2004, 547. [14] H.J. Bunge: Z. Metallkd. 56 (1965) 872. [15] D.G. Brandon: Acta Mat. 14 (1966) 1479. DOI:10.1016/0001-6160(66)90168-4 [16] P. Van Houtte: The MTM-FHM Software System, Version 2, Department of Metallurgy and Materials Engineering, Katholieke Universiteit Leuven, Belgium. [17] C. Ghosh, A. Haldar, P. Ghosh, R.K. Ray: ISIJ Int. 48 (2008) 1631. [18] R. Padmanabhan, A.J. Baptista, M.C. Oliveiva, L.F. Menezes: J. Mater. Process Tech. 184 (2007) 288. DOI:10.1016/j.jmatprotec.2006.11.051 [19] R. Ueji, N. Tsuji, Y. Minamino, Y. Koizumi: Acta Mater. 50 (2002) 4177. DOI:10.1016/S1359-6454(02)00260-4 [20] R. Song, D. Ponge, D. Raabe, R. Kaspar: Acta Mater. 53 (2005) 845. DOI:10.1016/j.actamat.2004.10.051 [21] R. Saha, R.K. Ray: Scr. Mater. 57 (2007) 841. DOI:10.1016/j.scriptamat.2007.06.064 [22] V. Randle, H. Davies: Metal. Mater. Trans. A 33 (2002) 1853. DOI:10.1007/s11661-002-0193-3

(Received May 13, 2009; accepted July 30, 2010) 11

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Bibliography DOI 10.3139/146.110406 Int. J. Mat. Res. (formerly Z. Metallkd.) 101 (2010) 10; page 1 – 12 # Carl Hanser Verlag GmbH & Co. KG ISSN 1862-5282

Correspondence address Mr. Chiradeep Ghosh Research and Development Division Tata Steel, Jamshedpur, India, Pin – 831001 Tel.: +91 657 664 8945 Fax: +91 657 227 1510 E-mail: [email protected]

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