Texture and mechanical properties of quenching and partitioning steel

Advanced high-strength steels, such as quenching and partitioning (Q&P) steels, are of considerable interest in the automotive industry owing to their desirable mechanical properties. However, further research is required to elucidate the relationship between the microstructure and mechanical properties of Q&P steels. Therefore, this study investigated the microstructure, texture, and mechanical properties of Fe0.23C1.55Si1.92Mn0.04Al Q&P steel prepared under different conditions. After tensile deformation, the intensity of the H {001}〈110〉 orientation in the samples increased from 2.47 to 3.51. The sample partitioned at 375 °C for 60 s had the highest ultimate tensile strength of 1446 MPa. The sample partitioned at 450 °C for 180 s achieved the highest elongation at fracture of 17.7%. As the partitioning temperature (PT) increased, the width of the lath martensite increased, the volume fraction of primary martensite decreased, and the secondary martensite of the III sample increased. As the partitioning time (Pt) increased, the size and content of secondary martensite decreased.


Introduction
In recent years, advanced high-strength steels (AHSSs) have attracted considerable attention from the automotive industry owing to their potential to simultaneously improve safety and reduce weight [1]. In particular quenching and partitioning (Q&P) steels, first proposed by Speer et al in 2003 [2], have high strength and ductility, which makes them an excellent choice for automotive parts. The Q&P process consists of several steps. First, heat treatment is used to austenitise the steel. Then, the steel is quenched to produce martensite. Next, during the partitioning step, the steel is held at a certain temperature and carbon diffuses from the supersaturated martensite into the residual austenite (RA) [3]. Finally, when the sample returns to 25°C, the microstructure of the Q&P steel consists of high-carbon RA and low-carbon martensite. However, when the steel is deformed austenite transforms into martensite, which results in stress relaxation and increases the ductility of the material. Stress can also accelerate the transformation of pearlite, which increases the hardness of the material [4]. Thus, the strength and plasticity of the steel are improved.
The mechanical properties of automotive steels depend on the textures produced by different manufacturing processes. For example, alignment with the {332}〈113〉 component is a favourable texture orientation for deep drawing processes. This is followed by alignment with the {554}〈225〉 component and then the {111}〈112〉 and {111}〈110〉 components, which belong to the γ-fibre component [5]. Thus, in order to improves the deep drawability of low-carbon steel, it is very important to increase the γ-fibre [6]. It has been shown that, as the intensity of the γ-fibre increases and the orientation difference decreases, the formability of the steel under bidirectional tension decreases significantly [7]. The Young's modulus of the crystal is highly dependent on the orientation and it is greatest in the 〈111〉 direction and lowest in the 〈001〉 direction [8].
Recently, the relationship between phase and mechanical properties of Q&P steel is of great interest. Q&P process is one of the most promising methods to produce the third-generation advanced high-strength steel with improved strength and elongation. However, the optimal time and temperature of partitioning treatment are not clear enough, only few studies have acknowledged that Q&P will affect the texture of the steel and change the mechanical properties of the steel. Therefore, this research discusses the effects of partitioning temperature and time on the microstructure, texture evolution and mechanical properties under Q&P process, and further investigates the volume fraction, phase type, morphology and distribution of martensite and austenite, and the variations of tensile mechanical properties under Q&P process, and found that at different partitioning temperatures, the elongation of the sample will increase dramatically, and the sample strength will decrease, these results directly reflect the impact of temperature on the test samples, and provide the manufacturing design for the third-generation advanced high-strength steel.

Experimental procedures
First, 0.23C-1.55Si-1.92Mn-0.04Al (wt%) was cold-rolled to a thickness of 1.5 mm. Then, samples were prepared, as shown in figure 1. The samples were subjected to Q&P heat-treatment cycles in a thermomechanical simulator (GLEEBLE ® 3800). The heat treatment involved austenitisation at 900°C for 90 s, quenching at 260°C for 90 s, isothermal treatment (partitioning) at 375 or 450°C (PT) for 60, 120, or 180 s (Pt), and quenching to 25°C. This process is summarised in figure 2. Hereinafter, the samples treated at x°C for y s are labelled x-y.
Tensile tests were conducted using a thermo-mechanical simulator (GLEEBLE ® 3800) at a strain rate of 0.001 s −1 . Then the microstructures of the samples were characterised using scanning electron microscopy (SEM; S-3400N). The samples were prepared for analysis by grinding, mechanical polishing, and etching with 4 vol% HNO 3 in ethanol.
X-ray diffraction (XRD) analysis was conducted and the orientation distribution functions (ODFs) of the samples were obtained using an x-ray diffractometer (D8 Advance, Brucker) with 1.5418 Å of the Cu-    I  I  I   I  I  I  I  I   200  220  311  3  200  220  2   200 220 where V γ is the volume fraction of austenite, I γ is the integrated intensity per angular diffraction peak in austenite, and I α is the integrated intensity per angular diffraction peak in ferrite or martensite. The concentration of carbon in the RA was calculated using the equation where C γ is the carbon concentration of the RA and α γ is the lattice parameter of the austenite phase [9].
The samples were prepared for electron backscatter diffraction (EBSD) analysis using standard metallographic techniques. First, the samples were mechanically polished. Then, to eliminate the internal stress, the remaining samples were immersed in a solution of 10-15 vol% perchloric acid in alcohol and ion-polished using an argon-ion-beam polishing system (GATAN Ilion II 697). Field-emission SEM (FE-SEM; FEI Nova NanoSEM 450) was used to characterise the samples with a voltage of 20 kV, beam current of 6.0 nA, tilt angle of 70°, and working distance of 9.8 mm. Areas of 25 × 30 μm in the centre of each sample were scanned with steps of 0.06 μm. The characterised data were post-processed using the HKL Channel 5 software.

Results and discussion
3.1. Microstructural characterisation 3.1.1. Evolution of the V γ and carbon content During the isothermal treatment, carbon diffused from the supersaturated martensite to the RA, when the partitioning temperature is higher than the initial martensite temperature (Ms). This will result in a bainitic phase transition where ferrite nucleates and grows in low-carbon RA, contributing to carbon enrichment into austenite; therefore, the austenite had greater high-temperature stability and a larger volume at 25°C. The XRD results for the samples are shown in figure 3(a). Equations (1) and (2) were used to calculate the volume fraction V γ and average carbon concentration C γ of the RA, as shown in figure 3(b). The results show that the Q&P steels consisted of martensite and RA. Some researchers have reported that bainite forms during partitioning and that the martensite mainly forms after the second quenching to 25°C [10,11]. However, ferrite and bainite were not detected in this study although their volume fractions may have been too low to detect using XRD.
In the samples treated at 375°C, the V γ decreased as the Pt increased. This indicates that carbon can transfer from martensite to austenite even when the Pt is short. The concentration of carbon in the RA increased considerably in 180 s. The trends in V γ and C γ between 60 and 200 s were identical to those reported in the literature [12]. However, in the samples treated at 450°C, V γ increased as the Pt increased. This may be caused by the nucleation and growth of ferrite in the low-carbon austenite during the formation of bainite. As the bainite grew, carbon diffused from supersaturated ferrite to the austenite; hence, the carbon content of the austenite increased, and when the sample was quenched, more austenite was retained at 25°C.

Microstructure
SEM images of the samples are presented in figure 4; they show a tempered martensite matrix, bainite, and RA. The martensite block was observed near the prior-austenite grain boundaries (figures 4(a), (b), (d), and (e)) after 60 and 120 s of isothermal treatment, and inside the martensite packet boundary (figures 4(b) and (e); dark dotted line) after 120 s of isothermal treatment. Irregular polygonal grains were also observed and they became larger and more numerous as the Pt increased (figures 4(c) and (f)). Moreover, as the Pt increased, the microstructures of the samples gradually changed from the original quenched martensite to tempered martensite.
In the samples treated at both 375 and 450°C, carbon diffused from martensite to austenite, and the austenite transformed into bainitic and ferrite. When the austenite was treated at a higher temperature, carbon partitioning resulted in the formation of regions with high and low concentrations of carbon. Ferrite nucleation occurred preferentially in the low-carbon regions. After a certain time, carbides precipitated between the ferrite and formed a mechanical mixture (bainite) of ferrite and carbides. Akbary et al suggested that, although samples with a higher austenite fraction quench to higher temperatures, this is not the main reason for the increased bainite content [13]. This supports the assumption that the higher fractions of bainite were related to higher fractions of unstable austenite. As the PT increased from 375 to 450°C with a fixed Pt, the fraction of bainite increased. Similarly, as the Pt increased at a fixed temperature, the fraction of bainite increased. The EBSD analysis results (figure 5) revealed that film-like inter-lath RA and blocky RA existed at the martensite packet boundaries and prior-austenite grain boundaries. The film-like RA was usually located between the martensite laths. From figure 5(a), the V γ was estimated to be approximately 4.8%, which is slightly lower than the value obtained from the results of XRD.
Regions of the same colour had similar crystallographic orientations. Austenite and adjacent martensite in the same region were the same colour. Within a martensite package, there was a residual film-like morphology.  The morphology of the martensite was approximately the same as that of the surrounding lath martensite. The austenite grains within a region originated from the same single prior-austenite grain and shared a single crystallographic orientation. The main orientation plane was {110}, as shown in figure 5. The presence of {110} austenite, contributed to the martensitic displacive shear transformation and increased the ductility and toughness [14]. Figure 6 shows the combined IQ map and IPFs of various samples after tensile deformation. The RA fractions in the deformed samples (figures 6(a)-(c)) were lower than those in the corresponding undeformed samples (figures 5(a)-(c)) owing to the martensitic transformation of RA during plastic deformation, which is called the transformation-induced plasticity effect [15]. Most of the blocky RA was transformed but some of the film-like RA was not transformed. Therefore, the film-like RA was more resistant to large plastic deformation. Li et al suggested that the residual stress surrounding lamellar RA parallel to lath martensite increases the hydrostatic pressure, which increases V γ , and it does not meet the conversion requirements. In the early stages of deformation, the RA could only adapt to the surrounding phase and applied stress by changing the grain orientation. Hence, the film-like RA parallel to the lath martensite is more favourable for improving the ductility of the material than the block-shaped RA.
In figures 6(d)-(f), the preferred orientations of the martensite were distinct and the martensite grain decreased in size. The main orientations of sample distribution were associated with the (110) crystal face, which is consistent with the tensile results reported by Peng [16]. Figure 7(a) shows the martensite matrix and grain boundaries obtained via EBSD. In figure 7(a), the red and black lines represent misorientation angles of less than 15°. (low-angle boundaries) and more than 15°(highangle boundaries), respectively. Local misorientations of 0-5°can be used to evaluate local strain gradients and stress concentration areas, and to illustrate the potential for stress concentration to cause cracks. The  misorientation angle distributions of the samples are shown in figure 7(b). Samples 375-60 and 450-60, which contained martensite and RA, had more 1-5°areas than sample 450-180, which contained bainite, martensite, and RA. This indicates that the former samples had a higher probability of stress concentration than the latter sample.
During deformation, cracks are likely to initiate from areas of high stress concentration. Therefore, the relatively low stress concentration in sample 450-180 inhibited microcrack generation and effectively improved the toughness of the material. Furthermore, the high-angle grain boundaries in sample 450-180 were stronger than those in the other samples. As the Pt increased to 180 s, the initial lath martensite (low-angle boundaries) was gradually tempered to form ferrite (high-angle boundaries). Moreover, as shown in figure 7, there were more high-angle boundaries in the undeformed samples than in the deformed samples. Thus, tensile deformation decreases the proportion of high-angle grain boundaries and increases the number of low-angle grain boundaries. Low-angle boundaries had little effect on the cleavage microcracks, whereas medium-angle boundaries changed their propagation direction and high-angle boundaries prevented them from propagating. This occurred because line defects built up at the grain boundaries when the angle between the optimal glide planes was large.   orientations of the τ-fibre are have 〈110〉 parallel to the transverse direction (〈110〉 || TD) and the texture components include H {001}〈110〉, Cu {112}〈111〉, F {332}〈113〉, and G textures. The γ-fibre is characterised by 〈111〉 parallel to the sheet-normal direction (〈111〉 || ND) and the related texture components are E and F textures.

Texture evolution
In figure 9, the distribution of the crystal orientations is divided into three sections: j 2 = 0°, 45°, and 65°in the Euler space. All the important orientations in the material were within these sections. High-intensity orientations were located near the H and G textures in the j 2 = 0°section. By contrast, low-intensity orientations were located near the B and Cu textures. The partitioning time texture evolution did not depend on the PT regardless of the and the ODF intensity did not change. However, as the Pt increased, the maximum intensity of the ODF increased. Combined with the results in figures 6(a) and (c), this shows that bainite grew along the shear plane of the specific crystal plane {225} of the prior-austenite grain, which improved the strength of the texture. Hence, the texture components can be modified by grain growth during cooling. Figure 10 shows the intensity of the γand τ-fibres in various samples. There were small fluctuations in the intensity over the entire range. The E and F components were evenly distributed. The intensity of the γ-fibre fluctuated between 1.42 and 1.74. The intensities of samples 375-60 and 450-60 fluctuated at the same amplitude, whereas the intensity of sample 450-180 was approximately the opposite. As the Pt increased, the grain orientations of the E and F components switched. Moreover, sample 450-60 exhibited the lowest {332} 〈113〉 orientation intensity and the highest H {001} 〈110〉 orientation intensity among the τ-fibres in the samples. Figure 11 shows the ODFs representing the macro-texture components of the samples after tensile deformation. In comparison to those of the undeformed samples, the texture components of the deformed samples were relatively unchanged; however, there were dramatic differences in the intensities of some textures. After deformation, the max ODF values of samples 375-60 and 450-60 increased from 2.0 and 2.08 to 3.33 and 3.47, respectively. The intensities of the RtG and H textures in samples 375-60 and 450-60 in the j 2 = 0 images increased, whereas those in sample 450-180 changed slightly. Figure 12 shows the intensities of the texture components after tensile deformation in terms of the γ and τ-fibres. The texture of the deformed sample 450-60 was characterised by the γ-fibre texture and the maximum ODF value of the E {111}〈110〉 texture component was 2.1. There were no changes in the other samples. Furthermore, as shown in figure 12(a), the intensity of the τ-fibres in the samples increased, except for the Cu {112}〈111〉 texture component of the 450-180 sample, which decreased. Across the τ-fibres in all of the samples, the {332}〈113〉 orientation had the lowest intensity and the H orientation had the highest intensity.
Before tensile deformation, the intensity of the H orientation was between 1.2 and 1.8 for all of the samples. However, after tensile deformation, the intensity of the H {001}〈110〉 orientation increased to 2.47-3.51. The orientation of H{001}〈110〉 is perpendicular to the surface of the plate, therefore, the deformation of the steel will be caused in the orientation of thickness. This will be reflected as a decline of elongation during the tensile process. When the intensity of the sample in the orientation of H{001}〈110〉 increases to 2.47-3.51, the steel sample 450-60 increases most, which means lowest elongation. This reflects that the H{001}〈110〉 texture does not improve the impact toughness between textures [17].
The Cu {112}〈110〉, G{110}〈001〉, and {332}〈113〉 components decreased by different amounts, and the first two components reached a stable state. The {332}〈113〉 texture was derived from the deformed austenitic texture. The {332}〈113〉 texture provided the best combination of strength and toughness in the transformation texture component [18]. During deformation, the austenite volume fraction decreased owing to the deformation-induced martensite transformation; hence, the strength of the {332}〈113〉 orientation increased. Sample 450-60 had the highest V γ and the deformed sample showed the smallest decline in the {332}〈113〉 texture.Low values of the modelled mechanical driving force correspond to low transformation potentials. The martensitic transformation is more difficult for low-potential orientations than high-potential orientations [19]. The orientations with the highest potentials were observed near the H {001}〈110〉 component, where the driving forces were 0.121. The high transformation potential indicated that the texture of the RA grains was related to low transformation stability [20]. Therefore, the RA grains with this crystal orientation were more stable. The H {001}〈110〉 component of the three samples increased during tensile deformation; hence, higher strain resulted in more austenite grain transformations, which resulted in strain hardening.

Mechanical properties
The mechanical properties of the samples are listed in table 1. To estimate the strain hardening exponent n, the true stress-true strain curves were fitted and calculated using the equation    where σ t is the true stress (MPa), K is the strength coefficient (MPa), and ε t is the true strain. As shown in table 1 and figure 13, at a given PT the σ UTS of the samples decreased as the Pt increased, whereas strain hardening exponent and ductility show opposite trends, these results show that the ductility increases with the increase of Pt, but at the expense of the strength of QP steel [21,22]. The material matrix softened at higher PT, which improved its stability and facilitated greater total elongation.

Conclusion
This study investigated the effects of Q&P treatment on the texture evolution and mechanical properties of steel before and after tensile deformation. The results were as follows: