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BY 4.0 license Open Access Published by De Gruyter Open Access December 5, 2022

Deformation behavior and formability of friction stir processed DP600 steel

  • Imren Ozturk Yilmaz , Mumin Yilmaz and Onur Saray EMAIL logo
From the journal Open Chemistry

Abstract

The effect of friction stir processing (FSP) on the formability of DP600 steel was experimentally investigated and the basic relationships between biaxial deformation behavior and FSP-induced evolutions in microstructural and mechanical properties were established. FSP formed a microstructure that consists of lath martensite with increased volume fraction compared to as-received (AR) microstructure that mainly consisted of well-distributed fine martensite particles in a ferrite matrix. Consequently, AR yield strength (301 MPa) and ultimate tensile strength (621 MPa) increased to about 811 and 1054 MPa, respectively. This strength enhancement achieved accompanied by adequate uniform elongation and elongation to failure values of 6.3 and 13.0%, respectively. Under biaxial loading conditions, good strain hardenability of the AR DP600 steel brought about a large membrane stretching regime leading to high punch force for biaxial flow. After FSP, both punch displacements within the membrane stretching regime decreased due to the increased volume fraction of lath martensite leading to higher cracking tendency. In result, cup depth and peak punch force of FSPed DP600 decreased from 8.7 mm and 33.2 kN to 7.1 mm and 28.1 kN, respectively. The obtained results simply indicate that FSP can be employed to enhance the strength of dual-phase steels with a reasonable level of formability.

1 Introduction

Increasing expectations in crash safety and environmental harmony have placed lightweighting of vehicle bodies to the highest level of consideration among the automotive design criteria [1]. Development and implementation of the advanced high strength steel (AHSS) sheets can be regarded as a milestone in the field of lightweight body-in-white (BIW) design and manufacturing [2]. Dual-phase (DP) steel is one of the widely used members in the first generation of AHSS, which provides various combinations of strength and formability through its unique microstructure of fine martensite particles dispersed in a ferrite matrix [3]. This microstructure consists of hard martensite embedded in soft ferrite, hence the name given as DP steels. Ductile ferritic matrix satisfies the needs of the ductility and formability, while martensite phase particles provide strength of the DP steel [4]. Therefore, increase in martensite volume fraction increases the strength of DP-steel series at the expense of formability [5]. It has been well established that, friction stir processing (FSP) is one of the effective tools that provides a simultaneous increase in uniaxial strength and/or ductility of metals via grain refinement down to micron levels [6,7]. Consequently, FSP can be regarded as an alternative approach to improve the performance of DP steel in lightweight BIW applications. It is well known that, forming is an essential manufacturing tool for BIW parts. On this point of view, compatibility of the FSPed regions to sheet metal forming applications needs to be clarified under multiaxial loading conditions [711]. However, DP steels mostly have been the subject of studies on joining of similar/dissimilar materials by friction stir welding (FSW) [1217] and friction stir spot welding [1822]. The formability of friction stir welded (FSWed) plates has been studied in several studies. Miles et al. [10] indicated that plain strain formability FSWed DP590 sheets with dissimilar thickness were 20% higher than that of the laser welded counterpart. Aktarer et al. [11] reported declined Erichsen index and increased punch force in miniaturized Erichsen testing of the FSPed DP600 steel. They concluded that the roughening with the orange peel effect on free surfaces of stretched metal can be eliminated by FSP [11]. Lee et al. [13] investigated the feasibility of FSW in forming for tailor welded blanks, which are made from individual sheets of steel of different thickness, strength, and coating, which are joined together by laser welding. They showed that, formability of the DP590 steel was strongly reduced by the FSW due to the inferior strength at the heat-affected zone [13]. Mahmoudiniya et al. [14] studied the FSW of DP700 steel. They focused on the effects of process parameters on the microstructure. Aktarer et al. studied the FSP of DP600. They also concluded that optimization of process parameters is necessary to further improve the performance of DP600 steel [14]. Different results obtained from these studies, simply indicate the need of further investigations on relationships between formability and evolution of microstructural and mechanical properties of steels, especially AHSS. Besides, it is well known that geometrical welding problems like root and lap defects may arise during joining of two separate surfaces via FSW [22]. One may expect that such geometrical problems may cause misleading results on multiaxial testing of relatively large-scale specimens. To eliminate such joining problems, FSP may be considered as a more suitable method to identify the effect of friction stir technologies on uniaxial biaxial flow behavior and formability.

In the view of above, the current study aims to reveal the impact of FSP on the formability of DP steels. Moreover, the mechanical behaviors of FSPed AHSS under biaxial stretching conditions have not been investigated before. This study also aims to establish the relationships between the FSP-induced microstructural evolution and biaxial flow behavior of DP600 steel as a model material to common AHSS.

2 Materials and methods

The material utilized in the current study was the commercially available 1.1 mm thick DP600 steel in the chemical composition of 0.14 C, 0.50 Mn, 0.50 Si, 0.10 P, 0.03 Al, 0.015 S, and balance Fe (wt%). As-received (AR) sheets were cut in dimensions of 200 mm × 60 mm and subjected to FSP. The process was applied using a WC tool with shoulder diameter, pin diameter, and pin length of 14, 6, and 0.8 mm, respectively. Based on the results of the several experimental trials, tool rotation speed, processing speed, and tilt angle were determined as 1,000 rpm, 1.6 mm/s, and 3°, respectively, as optimum repeatable process parameters. Processing temperature was monitored using an infrared thermal camera. Microstructures of AR and FSPed DP600 steel samples were observed using scanning electron microscope (SEM). Metallographic examinations were performed on samples sectioned perpendicular to the FSP direction. Metallography samples were prepared by following the standard procedures and then etched in 2% Nital solution (Figure 1(a)). Uniaxial tensile tests were employed to determine the mechanical properties. Dog-bone-shaped specimens with gauge length of 12 mm and gauge width of 6 mm (Figure 1(b)) were extracted using wire electric discharge machining technique as shown in Figure 1(a). Quasi-static tests with a strain rate of 1 × 10−3 s−1 ran on a universal test machine equipped with video extensometer (Figure 1(a)). Repeatability of the mechanical test results was ensured with at least three accompanied tests. Stretch formability of DP600 steel before and after FSP was determined by cupping tests according to ASTM E643-15. FSPed formability test specimens are placed into a die to coincide the central line of the process (tool center movement line) to the center of the semi-spherical punch and fixed using a holder force of 10 kN. Tests were performed without lubrication using a 20 mm diameter hemi-spherical punch moving with a velocity of 0.167 mm/s. During the tests, punch force (F) and punch displacement (X) data were collected using a computer. Punch displacement where the peak punch force (FCD) reached is determined as cup depth (CD). Cupping tested specimens were sectioned and prepared for metallographic examinations to reveal the effect of microstructural properties on the fracture behavior.

Figure 1 
               Schematic illustrations of (a) locations of the test specimens in FSPed sheet, (b) dimensions of tensile test specimen, and (c) testing position of the FSPed sample in the cupping test die (SZ: stir zone, TMAZ: thermo-mechanically affected zone, and HAZ: heat-affected zone).
Figure 1

Schematic illustrations of (a) locations of the test specimens in FSPed sheet, (b) dimensions of tensile test specimen, and (c) testing position of the FSPed sample in the cupping test die (SZ: stir zone, TMAZ: thermo-mechanically affected zone, and HAZ: heat-affected zone).

3 Results and discussion

Microstructure of the AR DP600 steel consisted of fine martensite grains placed at the vicinity of the ferrite phase grains with a mean grain size of 15 ± 5 μm (Figure 2(a) and (b)). FSP was successfully applied to DP600 sheets without the formation of micro- and/or macro-scaled cracks (Figure 2(a)). Processing temperature reached the range of 915 ± 20°C under steady-state conditions. As a result of the applied intense shear strain at this elevated temperature range, FSP strongly affected the microstructure of DP600 steel and formed well-characterized deformation zones known as stir zone (SZ), thermo-mechanically affected zone (TMAZ), and heat-affected zone (HAZ) (Figure 2(a)) [6,12,21,23]. At SZ, FSP strongly affected the microstructural organization and the morphological features of AR steel (Figure 2(b)). As can be seen from Figure 2(c), microstructure was strongly dominated by lath martensite phase with a volume fraction of about 94% (Figure 2(c)). Also, primary ferrite phase existed in the microstructure with a grain size of about 4 µm (Figure 2(c)). Transformation of AR fine martensite particles into lath martensite may be attributed to the high processing temperatures exceeding Ac3 [19]. In this temperature range initial microstructure is expected to transform into austenite phase. This may be followed by dynamic recrystallization of the austenite phase around the rotating tool [22]. In this stage, austenite grains were refined more effectively with higher strain rates and lower processing temperatures [12]. It is also shown that, apart from thermal cycles, severe plastic deformation of the processed material may be effective on the phase transformations [18]. In this context, dislocation accumulation points may act as preferred sites for ferrite nucleation [19]. In the current case, processing conditions were selected as the highest processing speed in the lowest possible rotating speed to enhance austenite grain refinement. Consequently, in the cooling cycle of the process refined ferrite grains were formed at the austenite grain boundaries and the interior of the grains transformed into lath martensite (Figure 2(c)) [19]. In TMAZ an obvious decrease in the volume fraction of the lath martensite was evident (Figure 2(d)). In this zone, peak-processing temperature in the range of Ac1–Ac3 may result in transformation of carbon-rich martensite phase into austenite, while undissolved ferrite remained unchanged [9,12]. Consequently, ferrite volume fraction was higher in the microstructure of TMAZ compared to that of the SZ. On the same point of view, increasing distance from the processing center may increase the volume fraction of the undissolved ferrite (Figure 2(d)) [9,12]. As can be seen in Figure 2(e), microstructure of HAZ mainly consisted of martensite particles rather than lath martensite grains. This may occur due to low peak processing temperatures below Ac1 as an expected result of higher distance to the processing line. This may somehow cause tempering of the initial ferrite phase and martensite islands without any phase transformation.

Figure 2 
               SEM microstructures of DP600 before and after FSP: (a) Low magnification microstructure of processed region, (b) AR-steel, (c) SZ, (d) TMAZ and (e) transition from TMAZ to HAZ.
Figure 2

SEM microstructures of DP600 before and after FSP: (a) Low magnification microstructure of processed region, (b) AR-steel, (c) SZ, (d) TMAZ and (e) transition from TMAZ to HAZ.

Engineering strain versus engineering stress curves showing the effects of FSP on the uniaxial deformation behavior of DP600 steel are represented in Figure 3. Mechanical properties and hardening parameters calculated from these curves are outlined in Table 1. It is understood from Figure 3 that fracture of the AR sample occurs by plastic instability after a considerable amount of strain accumulation. Flow stress in tensile test showed an effective increase after the yield strength (σ y) of 301 MPa and reached an ultimate tensile strength (UTS) of 621 MPa accompanied with a high uniform elongation of 21.3% (Table 1). During plastic deformation, strain hardening exponent and strain hardening coefficient took values of 0.24 and 1,135 MPa, respectively (Table 1). This is mainly related to the existence of the ferrite phase grains (Figure 2(b)) as a suitable environment for dislocation interactions leading to effective strain hardening. After UTS, ductile fracture with plastic instability and neck formation occurred till an elongation to failure value of about 34.7%. Figure 3 indicates that FSP caused a contraction in ductility of the steel with an increase in strength. The deformation behavior mostly remained unchanged after FSP, namely strain hardening of the steel was still evident after the process. The σy and UTS of the processed state were determined to be almost two-fold higher than those of the AR condition as 811 and 1,054 MPa, respectively. More importantly, such high strength values are obtained along with a moderate strain hardening behavior (Figure 3). As represented in Table 1, the strain hardening exponent and the strain hardening coefficient of FSPed DP600 steel took values of 0.14 and 1714 MPa, respectively. Uniform elongation (ε u) and elongation to failure values are determined to be 6.3 and 13.0%, respectively. These values indicate that strain hardening-induced plastic deformation covers nearly half of the total elongation. This may postpone the formation of the early onset of necking and may give rise to high strength and an acceptable level of ductility [2426]. Effective strength enhancement achieved by the FSP can be attributed to the higher volume fraction of the martensite phase in the SZ as the high strength component of the micro-constituents (Figure 2(c)). Such microstructural transformation may also be responsible for the decrease in ductility of the FSPed DP600 steel. As an expected result, the decreased volume fraction of ferrite phase in the microstructure of FSPed steel is expected to strongly reduce the contribution of the dislocation interactions [27]. Consequently, FSPed steel reflects lower ductility (uniform elongation and elongation to failure) and decreased strain hardenability. Strain hardening parameters also support this idea. It is understood from Table 1 that, strain hardening coefficient of the processed steel is considerably higher than that of the AR ones. However, strain hardening exponent is decreased after applied FSP (Table 1). An increase in strain hardening coefficient indicates that the plastic flow of the FSP-induced microstructure requires higher force. This may be a result of higher resistance of martensite phase to dislocation interactions and more pronounced in cracking tendency. A decrease in strain hardening coefficient, on the other hand, may be related to the decrease in the volume fraction of the ferrite phase as the only media available for dislocation interactions in the microstructure of the DP steel.

Figure 3 
               Engineering strain–engineering stress curves of AR and FSPed DP600 steel.
Figure 3

Engineering strain–engineering stress curves of AR and FSPed DP600 steel.

Table 1

Mechanical properties of AR and FSPed DP600 steel samples

Conditions σ y (MPa) σ UTS (MPa) ε u (%) ε f (%) K (MPa) n
AR 301 ± 6 621 ± 13 21.3 ± 0.2 34.7 ± 2 1,135 0.24
FSPed 811 ± 48 1054 ± 56 6.3 ± 0.1 13.0 ± 2 1,714 0.14

* Note: σ y: yield strength, σ UTS: ultimate tensile strength, ε u: uniform elongation, ε f: elongation to failure, K: strain hardening coefficient, n: strain hardening exponent.

The evolution of the mechanical properties after FSP may be affected by several factors including, tool geometry and processing conditions. Hence, various combinations of strength and ductility can be obtained by variating the process conditions [28]. In the current study, 1.1 mm thick sheets were processed with a displacement-controlled FSW machine using a feed rate (welding speed/tool rotation speed) of 0.096 mm/rev. In a previous study, 1.4 mm thick DP600 steel was subjected to FSP using a constant downforce of 6 kN and a feed rate of 0.11 mm/rev [17]. In the current study, comparable and slightly higher strength values were achieved with nearly the same uniform elongation compared to previous studies due to the lower feed rates leading to decreased heat input. This comparison shows that displacement-controlled FSP may be more effective in FSP of thin AHSS sheets to avoid substrate effects, especially at elevated processing temperatures. It may be important to note that, above-mentioned strength and ductility values achieved with the displacement-controlled FSP of DP600 steel are comparable and slightly higher than those of the commercial DP-980 steel [24], which is a cold-forming grade steel sheet mainly developed to satisfy the mechanical performance needs of the cold stamped high strength BIW parts [9].

Figure 4(a) shows load (F)–displacement (X) curves obtained from the cupping tests. As a common feature of the cupping tests, deformation of a sample occurs in an inhomogeneous manner including various stages due to the continuously variating contact conditions between semi-spherical punch surface and sample surface. It has been shown in the previous studies that these deformation stages directly affect the first-order derivative of the punch force with respect to displacement (dF/dX) [8,25]. To indicate this, variation of the dF/dX through the punch travel are plotted as shown in Figure 3(b). As can be seen in Figure 4(b), deformation of a sample forms four deformation stages known as “elastic–plastic bending” (A), “plastic bending” (B), “membrane stretching” (C), and “deformation localization” (D). In stage (A), increase in contact area of the semi-spherical punch and cupping sample causes local yielding and subsequently yield surface propagation [19]. In this stage dF/dX values continuously decrease as shown in Figure 4(b). This is followed by transition of the deformation mode into stage (B) leading to increase in dF/dX values. By the end of stage (B), increase in dF/dX is considerably saturated and further punch displacements caused minor variations in the value through stage (C). In stage (C), thickness of the sample decreases in a uniform manner. Also, an expected decrease of punch force due to the thinning is compensated with strain hardening. Hence, dF/dX values remain nearly stable through stage (C) [25]. Consequently, this stage represents a deformation regime in which strain hardening is one of the most determinant mechanical behaviors that directly affect the formability. As a result of the strain accumulation within stage (C) [8], necking takes place leading to a sharp and continuous decrease in dF/dX values at stage (D). At the end of stage (D), the sample reach CD and the dF/dX value is equal to zero due to the crack initiation (Figure 4(b)).

Figure 4 
               (a) Punch displacement (X)–punch force (F) curves and (b) dF/dX–X curves and deformation regions of AR and FSPed DP600 steel.
Figure 4

(a) Punch displacement (X)–punch force (F) curves and (b) dF/dXX curves and deformation regions of AR and FSPed DP600 steel.

Macrographs and SEM micrographs showing cross-section of the fracture site of AR and FSPed DP600 steel are represented in Figures 5 and 6, respectively. It is obvious from the macrographs of the cupping tested AR samples that, fracture occurred with cracks that propagated through a crescent-shaped path (Figure 5(a)). On the other hand, FSPed DP600 steel fractured with cracks initiated at SZ and propagated perpendicular to the processing direction (Figure 6(a)). This mainly indicates that mechanical behaviors of the SZ dominated formability of the FSPed steel. Hence, effect of deformation behaviors of the TMAZ and HAZ was less pronounced compared to that of SZ (Figure 6(a)).

Figure 5 
               (a) Top macro-view, (b) cross-sectional macro-view, and (c) and (d) fracture zone microstructure of the AR cupping test specimen.
Figure 5

(a) Top macro-view, (b) cross-sectional macro-view, and (c) and (d) fracture zone microstructure of the AR cupping test specimen.

Figure 6 
               (a) Top macro-view, (b) cross-sectional macro-view, and (c) and (d) fracture zone microstructure of the FSPed cupping test specimen.
Figure 6

(a) Top macro-view, (b) cross-sectional macro-view, and (c) and (d) fracture zone microstructure of the FSPed cupping test specimen.

Fracture of the AR and FSPed cupping test samples occurred after neck formation (Figures 5(b) and 6(b)). Thickness reduction in the neck region was more pronounced in the AR sample (Figure 5(b)) compared to that of the FSPed counterpart (Figure 6(b)). Fracture site microstructure of AR cupping sample shows that, cracks propagated in ferritic matrix through the interface of ferrite–martensite phases (Figure 5(c) and (d)). Moreover, several cracks (indicated with white arrows) that partly propagated with similar characteristics were evident at the fracture site of the AR cupping sample (Figure 5(d)). Cross-sectional micrographs of FSPed cupping test samples show that, crack propagation path underwent radical direction changes. Also, cracks tend to propagate at the ferrite–martensite interface through the longitudinal direction of the needle-shaped martensite grains (Figure 6(d)).

Comparison of deformation stages of the AR and FSPed samples indicates that, punch displacements required for transition of the deformation mode from stage (A) to stage (B) and stage (B) to stage (C) slightly affected from FSP (Figure 4). Normally, considerable increase in yield strength of FSPed DP600 sample may be expected to shift punch displacement to higher values due to the enhanced elastic deformation range compared to that of AR counterpart (Figure 4). The matter can be explained by means of changes in geometrical properties of the processed region. According to the detailed information given in experimental procedure section, FSP was applied to DP600 steel with a constant shoulder penetration depth of 0.1 mm (Figure 1). Consequently, the FSPed samples are somehow thinner than that of the AR ones. As an expected result of lower thickness, displacements in elastic deformation range may be lower than the AR case with initial thickness of 1.1 mm. As an expected result of such a thickness reduction, elastic displacements required to shift deformation mode from stage (A) to stage (B) and from stage (B) to stage (C) slightly change.

It is obvious from Figure 4(b) that, punch displacements within stage (C) considerably decreased after FSP. This may be related to decrease in allowable uniform thickness strain due to the decrease in uniform elongation. Transformation of the well-distributed fine martensite phase particles of the AR microstructure into lath martensite and refined ferrite grains have been found responsible for strength enhancement after FSP (Table 1). Considering the results of the uniaxial tension tests, these microstructural evolutions also decreased the uniform elongation and/or strain hardenability of the DP steel due to the decreased continuity of the ductile and formable ferritic matrix. Moreover, lath martensite grains of the FSPed DP600 steel may form large stress gradient at several stress concentration points due to the distinct plastic deformation behaviors of the microstructural constituents. As a result of formation of such stress concentrations, the magnitude of biaxial plastic strains that can be imposed to FSPed steel may expected to be limited compared to that of the AR counterpart (Figure 4). This may be responsible for decrease in punch displacement within the membrane stretching regime. These phenomena can also be effective on decrease in punch displacement within stage (D) regime (Figure 4(b)). It is well known that the main mechanism of deformation localization of the ductile metals is formation and coalescence of micro-voids [11]. Micro-void formation requires an increase in dislocation density at the local sites to eliminate continuity of the grains that undergo plastic deformation. It has been shown that, micro-voids are nucleated at the ferrite/martensite interfaces due to interfacial decohesion in DP steels [26]. This behavior become more pronounced under biaxial deformation conditions as in the case of the current study [26]. Magnitude of the stress concentration may be expected to be lower in the AR microstructure due to the fine particle size and comparatively equiaxed martensite morphology. So, to propagate initiated crack, any possible paths pass through the ferritic matrix. Hence, crack propagation and void coalescence require ferritic matrix to undergo further plastic deformation. In the case of cupping test of the AR steel, plastic strain that needed to coalescence voids also enlarges punch displacement in stage (D). Microstructure of fracture site of the AR cupping samples also supports this idea (Figure 5(d)). Namely, cracks indicated with white arrows in the microstructure of AR cupping sample (Figure 5(d)) may initiate after uniform thinning of the sample by flow of ferritic matrix [26,27]. On the other hand, in the microstructure of the FSPed steel, cracks mostly propagated through the length of the high aspect-ratio lath martensite grains (Figure 6(d)). This may indicate that, lath martensite–ferrite interface acts as a more preferred crack propagation sites compared to martensite particles of the AR microstructure (Figure 6(d)). Moreover, higher volume fraction of the lath martensite, mostly eliminates contribution of the ferritic matrix to the punch displacement within the “deformation localization” (D) stage in the cupping test of the FSPed sample. Consequently, punch displacement within the deformation localization state is lower and decrease in dF/dX values is sharper in FSPed steel (Figure 4(a) and (b)).

Values of CD and punch force at CD (FCD) deduced from FX curves are given in Table 2. As can be understood from Table 2, AR DP600 steel showed a continuously increasing FX curve characteristic with a high CD of about 8.7 ± 1.3 mm and FCD of 33.2 ± 1.9 kN, FSP of DP600 steel did not considerably affect the curve characteristics but caused a slight decrease in CD and FCD values to about 7.1 ± 0.5 mm and 28.1 ± 1.2 kN, respectively (Table 2). Simultaneous decrease in CD and FCD may be attributed to the above-mentioned biaxial deformation behavior variations after FSP of DP steel (Figure 4). Decrease in CD can be attributed to decrease in strain hardening capability of the FSPed steel and formation preferred crack propagation sites at the lath martensite interfaces (Figure 6(d)). These variations also decrease punch displacements within stages (C) and (D) of the cupping test, respectively (Figure 4). In result, CD of the FSPed steel takes limited values compared to that of the AR counterpart. Although AR steel reflects lower strength, peak force required for plastic deformation under biaxial stretching takes higher values (Tables 1 and 2). Lower FCD of the FSPed steel may be attributed to higher cracking tendency of the FSPed microstructure. These stress concentrations are mainly sourced from the formation of the lath martensite dominated microstructure (Figure 3(b)). In this microstructure, decrease in ferrite phase volume fraction, also decreases the contribution of this phase against crack propagation. Consequently, increased stress concentrations initiate crack with lower force requirement and cracks can propagate through the interfaces of the lath martensite as a preferred path. In result, both crack initiation and crack propagation require lower punch forces leading to limited FCD in the cupping test of FSPed steel.

Table 2

CD and FCD of AR and FSPed samples

Conditions CD (mm) FCD (kN)
AR 8.7 ± 1.3 33.2 ± 1.9
FSPed 7.1 ± 0.5 28.1 ± 1.2

4 Conclusion

AR microstructure of the DP600 steel that consists of homogeneously distributed network of martensite particles within ferrite matrix is eliminated by FSP. At SZ, FSP formed a microstructure that consists of lath martensite with a volume fraction of 94% and refined mean ferrite grain size down to about 4 µm.

Microstructural evolutions achieved with the process gradually enhance the strength of the DP600 steel along with ductility loss. After FSP, more than two-fold increase from 301 to 811 MPa is obtained in yield strength of the steel. UTS of steel reached 1,054 MPa as a comparable value to AHSS grades mostly preferred in lightweight automobile component design. Improvement in strength of the FSPed DP600 steel occurs primarily due to the increase in volume fraction of the martensite phase via transformation from well-distributed particles into lath martensite at SZ. This transformation also decreases uniform elongation (21.3%) and elongation to failure (34.7%) values to 6.3% and 13.0%, respectively, after FSP.

AR DP600 steel shows a high CD of about 8.7 ± 1.3 mm and FCD of 33.2 ± 1.9 kN. FSP of DP600 steel decreased both CD and FCD values to about 7.1 ± 0.5 mm and 28.1 ± 1.2 kN, respectively. Such a variation in CD and FCD is mainly attributed to decrease in strain hardenability of FSPed steel leading to lower displacement in membrane stretching stage and higher cracking tendency due to decrease in continuity of ferrite phase, which is the micro-constituent that provides formability of the DP steels.

FSP is an effective tool to improve the strength of DP steel with acceptable formability. These improvements can contribute lightweight body component design and manufacturing for automotive industry.


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Acknowledgments

The authors thank Borçelik Steel Company for supplying DP600 sheets used in the study.

  1. Funding information: This study was funded by Scientific and Technical Research Council of Turkey TUBITAK under grant number 115M649.

  2. Author contributions: All authors have equal contribution in this article.

  3. Conflict of interest: The authors declare that they have no conflict of interest.

  4. Ethical approval: The conducted research is not related to either human or animal use.

  5. Data availability statement: The datasets generated during and/or analyzed during the current study are available from the corresponding author on reasonable request.

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Received: 2022-11-07
Revised: 2022-11-13
Accepted: 2022-11-14
Published Online: 2022-12-05

© 2022 the author(s), published by De Gruyter

This work is licensed under the Creative Commons Attribution 4.0 International License.

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