
==== Front
Heliyon
Heliyon
Heliyon
2405-8440
Elsevier

S2405-8440(24)12735-1
10.1016/j.heliyon.2024.e36704
e36704
Research Article
On the use of thermomechanical processing to enhance the strength-ductility-toughness balance of plain low-carbon steel
Khorasani Fatemeh
Jamaati Roohollah jamaati@nit.ac.ir
⁎
Jamshidi Aval Hamed
Department of Materials Engineering, Babol Noshirvani University of Technology, Shariati Ave., Babol, 47148–71167, Iran
⁎ Corresponding author. jamaati@nit.ac.ir
24 8 2024
15 9 2024
24 8 2024
10 17 e3670410 3 2024
9 8 2024
21 8 2024
© 2024 The Authors
2024
https://creativecommons.org/licenses/by-nc-nd/4.0/ This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
This study aimed to fabricate a dual-phase (DP) steel with a combination of high strength-ductility-toughness by thermomechanical processing for industrial applications. Accordingly, the effects of 40 % cold deformation and intercritical annealing temperature on the microstructural evolution and tensile properties of low-carbon steel were studied. The microstructure of dual-phase steels consisted of ferrite (α) and martensite (α'), however, the morphology of martensite was different. With increasing the intercritical annealing temperature, the martensite fraction increased gradually from 0.41 at 770 °C to 0.51 at 860 °C. Also, the increase of martensite fraction from 0.41 to 0.51 caused a decrease in the martensite carbon concentration from 0.174 to 0.142 wt%. All dual-phase steels had a larger hardness than the as-received steel. When the temperature of annealing increased from 770 °C to 860 °C, the yield strength enhanced from 370.4 to 496.3 MPa, the ultimate tensile strength improved from 642.0 to 809.2 MPa, the total elongation decreased slightly from 31.8 % to 29.8 %, and the energy absorption increased from 190.0 to 217.6 J/cm3. The work hardening rate of all DP samples was considerably higher than other samples. Unlike initial, quenched, and rolled samples, the DP steels exhibited three-stage strain hardening behavior with increasing the true strain. For all dual-phase steels, a perfect ductile fracture was observed, with numerous uniform deep and coarse dimples. The dominant mechanism in the fabricated dual-phase steels was interface decohesion.

Keywords

Dual-phase steel
Thermomechanical processing
Intercritical annealing temperature
Microstructure evolution
Mechanical behavior
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pmc1 Introduction

Low-carbon steels are widely used in the automotive industry owing to their large formability, low cost, and excellent machinability [1,2]. However, they exhibit low strength and yield point phenomenon. It is a great challenge to evade the trade-off between the ductility and strength of low-carbon steels when considering large-scale production and cost. To attain a good combination of strength, ductility, and toughness, plain and low-cost techniques can be useful in the automotive industry. Using thermomechanical processing (rolling + heat treatment) to achieve ferritic-martensitic microstructure in low-carbon steels is a promising technique. Ferritic-martensitic dual-phase (DP) steels have a good balance of strength and ductility, as well as a high strain hardening rate, and reveal a continuous yield behavior [[3], [4], [5], [6]]. Martensite fraction, the morphology of martensite, martensite's carbon content, and the size of α and α' are effective on the mechanical behavior of DP steels [[7], [8], [9], [10], [11], [12], [13], [14], [15]].

The morphology of α' in dual phase steels depends on the partially reverted austenite. It is well-accepted that both thin acicular and coarse globular austenite can be formed during the reversion of austenite, which is transformed to the needle and blocky α', respectively, during quenching. The globular and acicular austenite can be created simultaneously and the amount of each morphology depends on several factors such as initial microstructures, transformation temperatures, and heating rate. Besides the dual phase steels with needle and blocky α', microstructures with elongated (fibrous) martensite were processed by changing the parameters of heat treatment and thermomechanical processing. It was reported that if the α' is chained (interconnected) then this remarkably enhances the yield and tensile strength without sacrificing the ductility. The morphology and distribution of the martensite affect the load transfer between the ferrite and martensite. The formation of martensite islands at the grain boundaries hinders the plastic flow in the ferrite and improves the yield strength. In addition, the formation of chained martensite enhances the portion of the load carried by the α' [[7], [8], [9], [10], [11], [12], [13], [14], [15], [16]].

There are many published works about the influence of rolling followed by intercritical annealing on the microstructural evolution and tensile behavior of DP steels. Mirzadeh et al. [17] studied the microstructure and tensile properties of DP steel fabricated by cold rolling and intercritical annealing. They found that the martensite fraction is very important to achieve a good balance between strength and elongation. Also, martensite microstructure is a better initial microstructure compared to tempered martensite and normalized microstructures. With performing the cold rolling, the grain size of ferrite and martensite reduced to 5 μm and 3 μm, respectively. Jia et al. [18] investigated the impact of cold rolling on the degree of homogeneity of deformation and ferrite recrystallization of dual-phase steel. They stated that increasing the homogeneity of the initial strain has a significant impact on the kinetics of ferrite recrystallization and grain refinement by decreasing space for growth. Soleimani [19], studied the effect of unidirectional- and cross-rolling followed by annealing on the microstructural evolution and tensile properties of low-carbon steel. They found that this process can improve mechanical properties and generate a good balance between strength and ductility. Kim et al. [20] studied the changes in tensile properties and microstructural evolution of dual-phase steel produced by 0 %, 20 %, 30 %, 40 %, 50 %, and 60 % cold rolling followed by intercritical annealing treatment. They found that with increasing the thickness reduction, yield strength (YS), strain hardening rate, and ultimate tensile strength (UTS) increase. Pan et al. [21] achieved the ultrafine-grained dual-phase steel with average sizes for ferrite and martensite of 2.7 μm and 2.9 μm, respectively, by using rolling and annealing. Also, the tensile strength and uniform elongation enhanced to 1267 MPa and 8.2 %, respectively. Paparao [22] fabricated ultrafine-grained dual-phase steel with a great balance between strength and ductility, 1295 MPa and 13 %, respectively, by intercritical annealing and cold rolling.

In most previous works on the dual-phase (ferritic-martensitic) steels, symmetric rolling was used. Recently, Yaghoobi et al. [23,24] used asymmetric rolling followed by intercritical annealing to manufacture dual-phase steel. They stated that asymmetric rolling is more effective on mechanical properties than symmetric rolling. Asymmetric rolling provides more shear strain compared to symmetric rolling and affects the recrystallization and transformation during intercritical annealing treatment. Therefore, the current study's objective was to manufacture ferritic-martensitic steel using asymmetric rolling followed by intercritical annealing. Also, the influence of annealing temperature on the microstructural evolution and mechanical behavior is investigated.

2 Materials and methods

The steel used in this investigation was an AISI 1008 sheet (Fe-0.08C-0.76Mn-0.17Si-0.02Cr-0.02Ni-0.025P-0.027Al-0.002Mo). The thickness of the initial sheet was 5 mm. Firstly, the initial sheet was cut into 50 mm (length) × 25 mm (width) × 5 mm (thickness). Then, the samples were heat-treated at 1050 °C for 10 min, followed by quenching in water. Afterward, a 40 % reduction in thickness was applied to the martensitic sample using the asymmetric rolling process (single-roll drive mode) with a laboratory rolling mill (with a roll diameter of 150 mm and a width of 200 mm) without lubricant at ambient temperature. Ac1 and Ac3 were estimated by Ref. [25]:(1) Ac1 = 723–10.7 Mn - 16.9 Ni - 29.1 Si + 16.9 Cr + 290 As + 6.38 W

(2) Ac3 = 910 - 203C0.5 - 15.2 Ni + 44.7 Si + 31.5 Mo + 104 V + 13.1 W

Based on the temperature of Ac1 (720 °C) and Ac3 (902 °C), four different annealing temperatures consisting of 770, 800, 830, and 860 °C were selected. The samples at four temperatures were intercritical annealed for 10 min followed by fast cooling. In Fig. 1, the processing route is shown schematically.Fig. 1 The processing route used in this work.

Fig. 1

The sample was cut into two different sections, RD–ND and RD–TD. The microstructures were observed with a light microscope (Dewinter, DG victory) and a scanning electron microscope at different sections. The surfaces were firstly ground by silicon carbide abrasive paper (80, 360, 600, 1200, and 2500 grit), then alumina suspension was used to polish them, and finally, 1.5 % Nital etchant was used. Using ImageJ software, the microstructures were quantitatively examined. The SANTAM STM–250 machine was used to conduct tensile tests on samples with gauge dimensions of 12 mm × 3 mm (length × width) at a testing speed of 1 mm/minute at room temperature. The total elongation of samples was determined by measuring the difference in the gauge length before and after the tensile test. SEM was used to determine the fracture mechanism. A Koopa UV1 Vickers hardness device was used to conduct hardness testing on the RD–TD plane under a 10 kg load. For each sample, hardness measurements were taken at 11 different points and the average value was calculated.

3 Results and discussion

3.1 Microstructure

Fig. 2(a–d) demonstrates the microstructure of the initial sheet at the RD–ND and RD–TD sections. Microstructure includes two phases, α and P (pearlite), with volume fractions of about 88 % and 12 %, respectively. The average α grain size is ∼38 μm. Fig. 3(a–d) and 4(a-d) show quenched and deformed microstructures from two different sections. The as-quenched microstructure consists of mainly lath martensite. As can be seen in Fig. 4, the α' stretched along the RD after the 40 % deformation at room temperature. Certainly, by applying 40 % rolling deformation, the density of dislocations in the sample has increased. The increase in dislocation density and the refining of the martensite phase induced by 40 % cold rolling can be quite effective on the morphology of the martensite phase of the final dual-phase steels.Fig. 2 Two different magnifications of the initial microstructures: (a,b) RD–TD and (c,d) RD–ND.

Fig. 2

Fig. 3 Two different magnifications of the martensite microstructures: (a,b) RD–TD and (c,d) RD–ND.

Fig. 3

Fig. 4 Two different magnifications of the deformed microstructures: (a,b) RD–TD and (c,d) RD–ND.

Fig. 4

The microstructures of DP samples produced by different annealing temperatures of 770, 800, 830, and 860 °C are shown in Fig. 5, Fig. 6, Fig. 7, Fig. 8, respectively. All DP steel's microstructure consists of ferrite and martensite (strictly speaking martensite + retained austenite). However, the distribution and morphology of α' are different. Fig. 5 shows martensite in the form of a very narrow band near the grain boundaries of α. In the microstructure of the 800-10, 830-10, and 860-10 samples, however, there are martensite islands that are uniformly distributed within fine ferrite grains. After cold rolling, the density of dislocation in martensite increases dramatically. The deformed lath martensite with a high density of dislocation is a favorable site for the nucleation of both austenite and ferrite grains during annealing treatment. Before Ac1 temperature, lath martensite transforms into ferrite and carbide. Between Ac1 and Ac3, fine γ grains are nucleated at the ferrite grain boundaries and also in the vicinity of carbides. The austenite grain grows as the intercritical annealing temperature rises, and eventually connects to form a chained network structure [[26], [27], [28]]. During the water quenching, austenite transforms into martensite. The morphology of lath martensite changes to chain martensite. It should be emphasized that the asymmetric rolling process causes martensite islands to be distributed uniformly in the ferrite matrix in all DP samples except for the 770-10 steel. At such a low intercritical annealing temperature, the holding time of 10 min is not high enough to form chain martensite. It was reported that, compared to symmetric rolling, asymmetric rolling generates larger shear plastic deformation and produces more uniform plastic strain through the sheet thickness [[29], [30], [31], [32], [33], [34], [35], [36], [37], [38], [39]]. This can be responsible for the uniform distribution of martensite in the produced dual-phase steels.Fig. 5 Two different magnifications of the 770-10 sample microstructures: (a,b) RD–TD and (c,d) RD–ND.

Fig. 5

Fig. 6 Two different magnifications of the 800-10 sample microstructures: (a,b) RD–TD and (c,d) RD–ND.

Fig. 6

Fig. 7 Two different magnifications of the 830-10 sample microstructures: (a,b) RD–TD and (c,d) RD–ND.

Fig. 7

Fig. 8 Two different magnifications of the 860-10 sample microstructures: (a,b) RD–TD and (c,d) RD–ND.

Fig. 8

Fig. 9 demonstrates the variation of α' fraction and the carbon concentration of martensite as a function of annealing temperature (see Fig. 8). As the intercritical temperature increases, the martensite fraction becomes higher. This is due to increasing the fraction of the γ phase during the annealing treatment. After the lowest tested annealing temperature (770 °C), the fraction of martensite is around 0.41. With increasing the intercritical annealing temperature, it increases gradually to reach 0.51 at 860 °C. The amount of C in the α' phase is affected by the martensite fraction. This has a considerable impact on the mechanical behavior of dual-phase steels and can be computed using the following equation:(3) CM=(C‐CF(1‐VM))/VM

where C is the as-received AISI 1008 sheet's carbon weight percent (0.08 %), CF (0.015 %) is the solubility limit of carbon in the ferrite phase at ambient temperature, and VM is the α' fraction. According to Fig. 9, the fraction of α' has a reverse relationship with the carbon concentration of martensite. The increase of martensite fraction from 0.41 to 0.51 caused a decrease in the martensite carbon content from 0.174 to 0.142 wt%. Consequently, increasing the annealing temperature decreases the hardness of the martensite phase in the produced dual-phase steel.Fig. 9 Variations in the martensite percentage and carbon content at various intercritical annealing temperatures.

Fig. 9

The SEM microstructure and results of EDS analysis for 770-10 and 860-10 DP steels are shown in Fig. 10(a–g) and 11(a-g), respectively. Scanning of the etched 770-10 and 860-10 DP steels was performed by secondary electron mode. The microstructure of both DP samples consists of ferrite and martensite, which were exhibited as dark gray and light gray in the SEM images, respectively. As seen in Fig. 10, Fig. 11, the surface of the samples includes a low and a high surface indicating the different capacities of ferrite and martensite phases against corrosion in the etchant solution. In fact, the martensite phase has protrusions while the ferrite phase has depressions. As can be seen in Fig. 10, the martensite phases located at the ferrite grain boundary (point 1) reveal a higher carbon concentration (4.08 wt%). Also, point 2 shows a martensite island with a lower carbon content (0.89 wt%). This result shows that during the intercritical annealing treatment, the narrow austenite formed in the grain boundaries has a higher carbon content compared to the equiaxed austenite. This can be related to more diffusion of C towards the grain boundaries with a lower atomic density of iron atoms. It should be noted that there is no lath martensite in the microstructure of the 770-10 DP sample due to the presence of a large carbon content in the martensite phase. From Fig. 10, the ferrite phase (points 1 and 2) exhibits a very low carbon concentration (0.05 and 0 wt%), as expected. As can be observed in Fig. 11, lath martensite is formed in the microstructure of 860-10 DP steel owing to the presence of a small carbon content in the martensite. As seen, points 1 and point 2 have a low carbon weight percent (0.11 wt%). Points 3 and 4 reveal the ferrite phase with 0.06 and 0 wt% carbon.Fig. 10 The SEM images of the 770-10 sample and EDS spot analysis.

Fig. 10

Fig. 11 The SEM images of the 860-10 sample and EDS spot analysis.

Fig. 11

3.2 Mechanical properties

The bar chart (Fig. 12) shows the variations of hardness for all samples. As seen, martensitic and deformed martensitic structures exhibit a higher hardness than the initial steel. The hardness of the quenched and 40 % cold-rolled samples increased significantly, a little less than 1.7 and 2.5 times in contrast to the as-received sample, respectively. The reason for the remarkable increase in hardness of around 218 HV for the quenched sample is the formation of martensite phases. Also, the effect of strain hardening causes the hardness to rise sharply to around 320 HV for the deformed sample. By applying the intercritical annealing on the deformed α', the hardness is dramatically reduced due to a reduction in dislocation density and the formation of the soft ferrite phase. All dual-phase steels have a larger hardness than the as-received steel owing to the presence of α' in the microstructure of DP steels. Compared to the as-received steel, the hardness value of the DP samples increased by 145 % and 186 % in the minimum and maximum range, respectively. For the 770-10, 800-10, 830-10, and 860-10, the hardness is 185.7, 208.0, 226.0, and 239.2 HV respectively. The main reason for this change is increasing the fraction of hard α' by increasing the temperature. Another reason is decreasing the mean free path of dislocations in the α grains according to the theory of Bag and Ray [40].Fig. 12 The hardness of as-received, as-quenched, cold-rolled, 770-10, 800-10, 830-10, and 860-10 samples.

Fig. 12

The representative stress-strain curves of the as-received, as-quenched, cold-rolled, and DP steels and the corresponding tensile data (yield and ultimate tensile strength, energy absorption, and total elongation) are shown in Fig. 13, Fig. 14, respectively. From Fig. 13, the yield point elongation (YPE) is observed in the initial sample. This has been linked to the Cottrell atmospheres created by interstitial atoms (carbon) near dislocations. When a dislocation locks in, a phenomenon known as the YPE occurs in low-carbon steels [41]. The necessary stress for unlocking the dislocations is greater than that for the movement of the dislocation. It is obvious that there is an abrupt yield drop and an upper and lower yield point. However, continuous yield behavior is visible in the tensile curves of dual-phase steels. As mentioned before, the transformation of γ to α' is associated with the formation of many unpinned geometrically necessary dislocations in the α grains. The movement of these unpinned dislocations along the boundaries of α/α' during the early stage of the plastic deformation leads to the elimination of YPE.Fig. 13 The stress-strain (engineering) curves of samples.

Fig. 13

Fig. 14 Tensile properties obtained from tensile tests: (a) yield stress, (b) ultimate tensile stress, (c) elongation, and (d) energy absorption.

Fig. 14

As can be observed in Fig. 13, Fig. 14, the as-received steel has the smallest YS and UTS, but its elongation is the largest. The 40 % cold rolled sample has the highest YS and UTS, however, its TE and EA are the lowest compared to other samples. The presence of α' and the large dislocation density are primarily responsible for the tensile behavior of the 40 % deformed steel. In DP steels, the impact of strain hardening is eliminated and also ferrite is formed, causing a drop in strength and a dramatic rise in ductility and toughness. However, owing to the creation of α-α′ microstructure, the resulting DP steels exhibit much greater strength and toughness than the as-received steel. As can be observed, the trend of YS, UTS, and EA for DP steels is increasing, and that of TE is decreasing. When the temperature is increased from 770 to 860 °C, the yield strength enhanced from 370.4 to 496.3 MPa, the ultimate tensile strength improved from 642.0 to 809.2 MPa, the total elongation decreased slightly from 31.8 % to 29.8 %, and the energy absorption increased from 190.0 to 217.6 J/cm3. The enhancement in the martensite fraction is responsible for these outcomes. The localization of strain in an α grain cannot propagate to the neighbor α grains because all α grains are surrounded by the chain martensite. Consequently, this microstructure can contribute to high ductility even at large levels of strength. On the other hand, the difference in the hardness of α and α' can be decreased by increasing the martensite fraction. This leads to achieving better ferrite/martensite interface cohesion. Therefore, the plastic instability is postponed, and finally, a larger ductility is achieved.

Reducing the hardness difference between α and α' can also affect the work hardening rate. For better understanding, the dσ/dε vs. ε plots of all samples are depicted in Fig. 15. Generally, there is a remarkable difference in strain hardening rate between the dual-phase steels and other samples. The work hardening rate of all DP samples is considerably higher than other samples, and the result is attributed to the distribution of the hard α' particles in the soft α matrix. Unlike initial, quenched, and rolled samples, dual-phase steels show three-stage work hardening behavior with increasing the ε as indicated by dashed line arrows in Fig. 15. This means the different values of n can describe the work hardening behavior of produced DP steels. In the first stage, the work hardening rate decreases remarkably, which is attributed to the plastic deformation of the ferrite. In the second stage, the slope of the strain hardening rate plot reduces that is due to the α deformation restrained by the α'. The 3rd stage with the lowest slope shows the poor capability of work hardening, which is owing to the co-deformation of α and α'. The difference in work hardening capability between ferrite and martensite phases leads to the dominance of the interface decohesion mechanism during the tensile test. It should be noted that the hardness difference between α and α' is decreased owing to increasing the annealing temperature (α' fraction), therefore, the plasticity of α' is greatly improved. For this reason, the plastic deformation of martensite in the 860-10 sample starts earlier and moves to the 2nd stage. Consequently, the work hardening rate of the second stage for the 860-10 sample is larger than that for the other dual-phase steels as shown in Fig. 15. From Fig. 15, it is clear that when the intercritical annealing temperature rises, the overall work hardening rate enhances as a result of an enhance in the martensite fraction and density of geometrically necessary dislocations, and also enhancement of strain hardening ability of martensite caused by a reduction in the C concentration. According to Pickering [42] and Balliger and Gladman [43], the plastic deformation of 2nd stage of work hardening is mainly governed by α', and the rate of strain hardening depends on the volume fraction of α' (VM) and the size of α' (DM): dσ/dε∝VM/DM. Therefore, an enhancement in the work-hardening rate is expected with increasing the temperature of intercritical annealing treatment (from 770 to 860 °C) due to the increase in the volume fraction of α' (from 0.41 to 0.51). Accordingly, larger strain hardening results in maintaining ductility as indicated in Fig. 13, Fig. 14.Fig. 15 The dσ/dε vs. ε plots of all samples.

Fig. 15

Fig. 16 illustrates the α' fraction - mechanical properties relationship. Through linear fitting, the following equations are proposed to describe the correlation between the mechanical properties and martensite fraction:(4) YS(MPa)=1301.6fM–162.96

(5) UTS(MPa)=1704.5fM–56.787

(6) Hardness(HV)=523.28fM–25.982

Fig. 16 The correlation between the mechanical properties and martensite fraction.

Fig. 16

As depicted in Fig. 16, the corresponding R2 values for yield strength, ultimate tensile strength, and hardness are 0.991, 0.997, and 0.984, respectively, which indicates a good correlation between martensite fraction and mechanical properties. Yaghoobi et al. [24] and Khorasani et al. [44] observed a similar relationship for the dual-phase steels.

3.3 Fracture mechanism

The fracture surface images of 40 % cold-rolled and DP steels are demonstrated in Fig. 17(a–j). As can be seen in Fig. 17(a) and (b) the fracture mode of deformed steel is ductile, however, the diameter of dimples is low. For all dual-phase steels, a perfect ductile fracture is observed, with numerous uniform deep and coarse dimples. The ductile fracture in the dual-phase steels consists of three stages: microvoid nucleation, microvoid growth, and voids coalescence by cross-linkage [7,[45], [46], [47], [48], [49]]. The incompatibility of plastic deformation in α and α' in dual-phase steels induces the nucleation-growth-coalescence of microvoids at the α/α' interfaces resulting in interface decohesion as depicted by blue arrows in Fig. 17(d). The main mechanisms that cause the formation of microvoids in dual-phase steels are the decohesion of the α/α' interface and martensite cracking. It should be noted the dominant mechanism in small-to-moderate fractions of α' is interface decohesion while at the large fraction of martensite, martensite cracking is dominant [13,44,[48], [49], [50], [51]]. According to Fig. 17, the dominant mechanism in the fabricated DP steels is interface decohesion.Fig. 17 The fractographs of (a,b) 40 % cold-rolled, (c,d) 770-10, (e,f) 800-10, (g,h) 830-10, and (i,j) 860-10 samples at two different magnifications.

Fig. 17

4 Conclusions

1) All DP steel consisted of ferrite and martensite. However, the distribution and morphology of α' were different. The narrow martensite formed in the grain boundaries had a higher carbon content compared to the equiaxed martensite.

2) By increasing the annealing temperature, the α' fraction enhanced gradually from 0.41 at 770 °C to 0.51 at 860 °C. Also, the increase of martensite fraction from 0.41 to 0.51 caused a decrease in the martensite carbon content from 0.174 to 0.142 wt%.

3) All DP steels had a larger hardness than the as-received sample owing to the presence of α' in the microstructure. Compared to the as-received steel, the hardness of the DP samples increased by 145 % and 186 % in the minimum and maximum range, respectively.

4) When the temperature increased from 770 to 860 °C, the yield strength enhanced from 370.4 to 496.3 MPa, the ultimate tensile strength improved from 642.0 to 809.2 MPa, the total elongation decreased slightly from 31.8 % to 29.8 %, and the energy absorption increased from 190.0 to 217.6 J/cm3. The enhancement in the martensite fraction was responsible for these outcomes.

5) The work hardening rate of all DP steels was considerably higher than other samples, which was due to the distribution of the hard α' particles in the soft α matrix. Unlike initial, quenched, and rolled samples, DP steels showed three-stage work hardening behavior with increasing the true strain.

6) For all dual-phase steels, a perfect ductile fracture was observed, with numerous uniform deep and coarse dimples. The dominant mechanism in the produced DP steels was interface decohesion at α/α'.

From the obtained results, it can be concluded that the asymmetric cold rolling followed by intercritical annealing at a high temperature is capable of remarkable strength enhancement and maintaining ductility, and therefore significant toughness improvement in dual phase steel, which is required by the automobile industry in the manufacturing of the automotive body. Additional research that can be performed includes evaluating the influence of the carbon content and annealing time on the microstructure and mechanical properties.

Data availability

All data included in this study are available upon request by contact with the corresponding author.

CRediT authorship contribution statement

Fatemeh Khorasani: Writing – original draft, Software, Resources, Investigation. Roohollah Jamaati: Writing – review & editing, Supervision, Methodology, Conceptualization. Hamed Jamshidi Aval: Writing – review & editing, Supervision, Conceptualization.

Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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