- Email: [email protected]
Laser beam welding of DP980 dual phase steel at high temperatures
Optics and Laser Technology xxx (xxxx) xxxx
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Optics and Laser Technology journal homepage: www.elsevier.com/locate/optlastec
12Laser beam welding of DP980 dual phase steel at high temperatures Caroline Cristine de Andrade Ferreira, Vagner Braga, Rafael Humberto Mota de Siqueira, ⁎ Sheila Medeiros de Carvalho, Milton Sergio Fernandes de Lima Photonics Division, Institute for Advanced Studies, Trevo Amarante 1, São José dos Campos, SP 12228-001, Brazil
H I GH L IG H T S
is proposed a new technique for laser welding of dual phase steels. • ItKeyhole welded DP980 steels presents a ferritic-austenitic microstructure. • The hardness, residual stresses and bending of the sheet were reduced. •
A R T I C LE I N FO
A B S T R A C T
Keywords: Laser beam welding Dual phase steels Heat treatment Hardness testing
Dual phase steels have been used in safety components for the automotive industry and represents a standard in terms of energy absorbing steel alloys. After welding, these components present a martensitic transformation in the fusion and heat-affected zones with an intrinsic brittleness. The current contribution presents a procedure where the DP980 steel sheet is kept at a given temperature during and after the laser weld in order to generate bainite instead of martensite. The results using an isothermal treatment at 527 °C shown a microstructure composed by grain boundary and bainitic ferrite and retained austenite with a constant hardness in the base material and heat-affected and fusion zones of 280 HV. The room temperature welds present hardness values between 320 and 500 HV. The high-temperature welds also shown a decrease in the maximum residual stress at the weld centerline by 1/3, with a consequential reduction in the warp of the joined component.
1. Introduction More and more automobile manufacturers are using advanced high strength steels for energy absorption in the case of a car accident [1]. One of most promising alloys for the safety components, such as pillars and roof bows, is called dual phase (DP) steels. Dual phase steels are composed of two phases, martensite and ferrite, which in a given morphology and size, could attain tensile strengths near to 1 GPa. DP980 is especially interesting steel typically surpassing a maximum stress of 980 MPa during tensile tests [2]. Although this alloy fulfills most of the materials selection criteria for the safety components, welding is still a challenge. Any type of welding procedure changes the well-balanced phase distribution and stress concentrators appear. Gould et al. [3] indicated that laser beam welding (LBW) and electric resistance weld (ERW) of DP980 results in a large amount of martensite in the fusion zone (FZ) as a result of rapid cooling. In the same way, Sreenivasan et al. [4] shown that LBW DP980 coupons lose their formability, using limiting dome height tests as standard, due to the martensitic weld bead. These authors also indicated a soft zone in
⁎
the heat-affected zone (HAZ), where necking occurred well before the maximum deformation of the cup [4]. The soft zone is located far from the fusion line (FL) in HAZ as a result of the temper of the base material (BM) martensite. Lima et al. [5] proposed a method to overcome the problems associated to the rapid quenching of advanced high strength steels by using inductive heating during LBW. The inductive heating has the advantage to decrease the cooling rate of the FZ and HAZ in order to produce a given phase or mixture of phases, in particular bainite. The same method could be used to reduce the martensite amount in the welds for the DP980 steel. The objective of the present contribution is to study the inductive heating of a DP980 steel sheet during laser welding in terms of microstructural evolution and hardness. 2. Experimental The current material is a 1-mm thick dual phase class DP980 sheet with the chemical composition giving in Table 1 (source ArcelorMittal).
Corresponding author at: EFO, IEAv, Trevo Amarante 1, 12228-001 Sao Jose dos Campos, SP, Brazil. E-mail address: [email protected] (M.S.F.d. Lima).
https://doi.org/10.1016/j.optlastec.2019.105964 Received 2 August 2019; Received in revised form 16 September 2019; Accepted 24 November 2019 0030-3992/ © 2019 Elsevier Ltd. All rights reserved.
Please cite this article as: Caroline Cristine de Andrade Ferreira, et al., Optics and Laser Technology, https://doi.org/10.1016/j.optlastec.2019.105964
Optics and Laser Technology xxx (xxxx) xxxx
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Table 1 Chemical composition of the DP980 steel (wt.%). Fe as the balance [2]. C
Si
Mn
Mo
Cr
Al
0.13
0.03
1.90
0.30
0.20
0.06
The last hot rolling occurred at 650 °C and then the coil was quenched to the room temperature in order to produce both ferrite and martensite. According to the material data sheet, the current DP980 contains 37 ± 6 vol% ferrite and 63 ± 6 vol% martensite [2]. The sheets were sectioned by spark erosion in squares of 40 × 40 mm, ground and washed to remove contaminants and then butt-welded using a fiber laser. The fiber laser (IPG, model YLR-2000) is composed of an ytterbium-doped glass host (wavelength 1080 nm) and a 100 µm diameter delivery fiber. The laser beam was focused using OptoSkand collimator and a focal lens (f 160 mm) to a plate inserted in an induction furnace. The induction furnace was conceived to heat the steel coupons to a given temperature, as reported previously [5]. The temperature in the middle of the fusion zone was recorded using a pyrometer Metis MP25 with emissivity calibrated to the range of temperatures between 100 and 700 °C. After a number of trials, the following parameters were chosen for a full penetration weld of the DP980 steel: laser power of 800 W, weld speed of 50 mm/s and focus on the top of the sheet. The spot diameter was 0.1 mm, as an image of the optical fiber with the same outer diameter. The temperatures of the welds were considered after the MUCG83 software from Peet and Bhadeshia works [6]. The software calculates Widmanstatten ferrite, bainite and martensite start temperatures. The thermodynamic data were obtained using ThermoCalc [7], version M, using FEDAT database. For the estimation of temperatures above 700 °C finite element software (SysWeld, ESI-Group) was used. The same dimensions of the real component had been used in simulations and the mesh resolution around the laser line was about 50 µm node-to-node. The simulation results of the simulation included the temperature evolution in the middle of the fusion zone, the final deformation after the weld and the Von Mises residual stresses. The obtained welds were analyzed using optical microscopy (Zeiss, model Imager2M) and scanning electron microscopy (Hitachi, model TM-3000). For the etching, a solution of 2% nitric acid in ethanol was used. Vickers hardness values were obtained using FutureTech FM-800 equipment. The indentation load was 100 gf for 10 s.
Fig. 1. TTT curves for the current alloy, indicating Widmanstatten ferrite (Ws), Bainite (Bs) and Martensite (Ms) start temperatures. The calculated temperature profiles RT and HTs are plotted together with measured temperatures HTm.
Ms-line. Under these conditions, it is expected almost fully martensitic FZ-grains. Fig. 1b presents the same TTT plot, but now with the high temperature (HT) weld procedure. The furnace temperature was kept at 526 °C, i.e. 100 °C above Ms, and the sheet was kept at this temperature for 10 min before cooling. An overheating of 100 °C is necessary to assure the sheet is kept above Ms during the time the furnace cover is opened for the laser interaction. In Fig. 1b, two curves are plotted, HTs and HTm, representing the simulated and measured temperatures. As could be seen, simulated and measured temperatures are in good agreement in Fig. 1a. This good fit is a result of a proper choice of the input parameter and in particular the laser absorptivity in the keyhole cavity as 85%. According to the estimations, the fusion zone of HT welds precipitated ferritic bainite after 13 s of cooling with an average cooling rate of 3.5 °C/s, which is much lower than the cooling rate when RT crosses Ms, Fig. 1a, i.e. 850 °C/s. The first part of the HTm curve, representing the period where the heat was on, took about 513 s (8.55 min), within bainite is supposed to growth. After this period, a sharp decrease of the HTm curve indicates the sheet was cooling down to 230 °C with a cooling rate of 0.24 °C/s.
3. Results 3.1. Temperature profiles Fig. 1 presents the time–temperature-transformation (TTT) of the current alloy calculated after the MUCG83 software [6]. Widmanstatten ferrite is expected to occur in the range 590–760 °C and the bainite range is 420–590 °C. The software calculated the martensite start temperature as 426 °C, which is in good agreement with Andrews predictions [8]. The Ms-curve presented two paths, one at the initial composition of the alloy (426 °C) and other at 278.6 °C. This last refers to the ThermoCalc [7] calculations of the residual austenite composition when the balance is 1/3 ferrite and 2/3 austenite and might not represent the right position in the TTT plot. Also, according to ThermoCalc, Ms temperature drops subzero when the volume phases fractions are above 50/50%. For the current study, one considers the TTT diagram is good enough to provide an estimation of the phases loci in continuous cooling at different cooling rates. Fig. 1a includes a plot of the room temperature (RT) temperature profile estimated at the middle of the fusion zone (FZ). The time origin is considered at the peak temperature and hence after 0.24 s the primary austenite crosses the
3.2. Microstructural features Fig. 2 presents micrographs of the room temperature (RT) samples. 2
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Fig. 2. Room temperature welded sample: (a) overview of the weld bead showing the fusion zone (FZ), recrystallized grains (RG) region, partial transformation (PT) region and tempered region (TR). (b) Detail of the martensite micro-constituent in FZ. (c) Detail of the PT/TR boundary.
In a general view of the weld bead, Fig. 2a, the microstructure could be divided into four regions: a fusion zone (FZ), a recrystallized grains region (RG), a partially transformed region (PT) and a tempered region (TR). The heat-affected zone is then comprised in three different subzones RG, PT and TR. Fig. 2a shows only one pore in FZ, which could eventually be a result of a base material porosity as also seen in PT. The width of the weld zone is about 0.3 mm and a typical columnar growth is perceived from the fusion line (FL). The microstructure inside FZ is completely martensitic as presented in Fig. 2b. Recrystallized grains are noted within a 0.1 mm region from FL and PT measures 0.1 mm wide, 0.3 mm far from FL. According to FEM results the peak temperatures in the middle of these sub-regions of HAZ are: RG 1200 °C, PT 720 °C and TR 400 °C.
Fig. 3. High temperature welded sample: (a) overview of the weld bead showing the fusion zone (FZ) and the heat-affected zone (HAZ). (b) Detail of the interface between FZ and HAZ. (c) FZ- columnar grains region. (d) FZ – equiaxed grains region. Legend: BF: ferritic bainite; A: austenite; GBF: grain boundary ferrite.
3
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because of the local softening. The boundary between PT and TR is presented in Fig. 2c and it is marked by a darkening of the martensite micro-constituent as long as the ferrite phase unchanged. Fig. 3a shows the weld bead FZ and HAZ for the high temperature condition, i.e. following to the temperature profile in Fig. 1b. Comparing Figs. 3a–2a, it could be noted that the share of the HAZ in subzones is lost in the first case. Also, the annealing of FZ and HAZ resulted in a more homogeneous contrast using the same etching. As can be seen in Fig. 3b the transition FZ to HAZ is smoother and dominated by ferrite (light gray) and austenite (dark gray) phases. These phases are best observed using scanning electron microscopy of the columnar (Fig. 3c) and equiaxed grains (Fig. 3d) in FZ. In both cases, the microstructure is composed of bainitic ferrite and retained austenite as noted in Fig. 3d. Grain boundary ferrite (GBF) is a thin 0.3 µm layer and gave rise to some Widmanstatten growth, although at lower distances (Fig. 3d). Bainitic ferrite (BF) is also noticed in the middle of the grains and corresponds to about 44% of the transformed volume. The rest is austenite (A) stabilized at ambient temperature due to ferrite segregation of elements, in particular C and Mn, according to ThermoCalc.
Fig. 4. Hardness profiles of the RT and HT conditions as a function of the distance of the centerline (CL). The tempered region (TR) is marked in the RT plot.
3.3. Hardness and strain-stress behavior These temperatures are in good agreement with the reported data [9,10], although these data refer to DP600 steel. Antunes et al. [10] cited TR as a possible stress concentrator in tensile strength tests
Fig. 4 presents the Vickers hardness profile for the RT and HT conditions measured from the weld centerline (CL). FZ corresponds approximately to the three points around CL as marked by a dark gray
Fig. 5. FEM of Von Mises residual stresses for (a) RT and (b) HT conditions. The numbers next to the laser path (middle of the sheet) represent residual stresses in FZ and HAZ. The scale at left represents the range of stresses in the entire simulated volume. 4
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welded coupon was also reduced from 0.6° to 0.2° for HT. These results indicate the high-temperature procedure could be interesting for the DP980 steel for thermomechanical processing and use in safety components.
bar. For RT, the hardness value at CL presented a peak compared to the neighboring data as a result of the martensite transformation (Fig. 4). In the RT case, the hardness in RG and PT sub-zones (light gray bars in Fig. 4) are relatively lower, around 300 HV, as a result of the ferrite amount. However, the tempered region (borders if the light gray bars) did not present a drop at the left and right of HAZ as previewed in the literature [11]. This result may be due to the very short time in laser welding, leading to a limited tempering of the martensite in the HAZ. The HT profile is flatter than RT and represents a ferrite-austenite balance in these regions with an average hardness of 280 HV. In Fig. 4, HT side, a little depletion could be noticed near to CL. This depletion could be associated to the centerline segregation during the final stages of solidification as reported in the literature [12]. In these regions, more ferrite is reported and then the hardness drops. Fig. 5 presents the calculated residual stresses according to the Von Mises criterion for the RT and HT conditions. The maximum internal residual stresses attained 780 and 271 MPa for the RT and HT conditions respectively. According to the calculations, RT residual stress attained a level too near to the maximum tensile stress of the alloy, and thus could be considered a threat. The numbers next to the laser line of Fig. 5 represents the residual stresses on the surface of the welded coupons. For the RT case, the surface stresses were 360 and 470 MPa for FZ and HAZ, respectively. In the HT case, these values were 120 and 140 MPa, respectively. These results indicated a reduction to 1/3 of residual stresses when the inductive heating is used instead of conventional welding. The estimated bending of the sheet after the weld is also related to the temperature. The simulated sheet edge rose 0.40 mm in RT condition compared to 0.13 mm for the HT case. For a 40 mm × 80 mm coupon, the final bending angle was 0.6° and 0.2° for RT and HT, respectively. These values are low and difficult to validate using metrology because the actual contortions of the steel sheet.
Declaration of Competing Interest The authors declared that there is no conflict of interest. Acknowledgements CCAF thanks Brazilian National Council for Scientific and Technological Development (CNPq) for an under graduation scholarship (PIBIC-IC). VB thanks São Paulo Research Foundation FAPESP for the doctoral scholarship (grant#2018_23884-5). The authors acknowledge the financial support of São Paulo Research Foundation FAPESP grants#2016/16683-8 and 2016/11309-0. References [1] Auto/Steel Partnership Joining Project (2007) Advanced high strength steel (AHSS) weld performance study for autobody structural components. A/SP Joining Technologies Committee Report. https://www.a-sp.org/-/media/files/asp/ enabling-programs/ahss-joining—ahss-weld-performance-study-for-autobodystructural-components.ashx (accessed 01 August 2019). [2] Arcelor Mittal (2019) Dual Phase steels. https://automotive.arcelormittal.com/DP (accessed 26 July 2019). [3] J.E. Gould, S.P. Khurana, T. Li, Predictions of microstructures when welding automotive advanced high-strength steels, Welding J (2006) 111s–116s. [4] N.N. Sreenivasan, M.M. Xia, S.S. Lawson, Y.Y. Zhou, Effect of laser welding on formability of DP980 steel, ASME J. Eng. Mater. Technol. (4) (2008) 1300410041–041004-9. [5] M.S.F. Lima, D. Gonzalez, S. Liu, Microstructure and mechanical behavior of induction assisted laser welded AHS steels, Welding J. 96 (2017) 376-s–388-s. [6] M. Peet, Bhadeshia HKDH Phase Transformations and Complex Properties Group, Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, U.K., 2019 (accessed 26 July /2019). [7] J.O. Andersson, T. Helander, L. Höglund, P.F. Shi, B. Sundman, Thermo-calc and DICTRA, computational tools for materials science, Calphad 26 (2002) 273–312. [8] K.W. Andrews, Empirical formulae for the calculation of some transformation temperatures, J. Iron Steel Inst. 203 (1965) 721–727. [9] G.C.C. Correard, G.P. Miranda, M.S.F. Lima, Development of laser beam welding of advanced high-strength steels, Int. J. Adv. Manuf. Technol. 83 (2016) 1967–1977. [10] W.D. Antunes, M.S.F. Lima, Experimental development of dual phase steel laser-arc hybrid welding and its comparison to laser and gas metal arc welding, Soldagem Inspeção 21 (3) (2016) 379–386. [11] J. Wang, L. Yang, M. Sun, T. Liu, H. Li, A study of the softening mechanisms of laserwelded DP1000 steel butt joints, Mat. Design 97 (2016) 118–125. [12] I.H. Brown, The role of microsegregation in centreline cold cracking of high strength low alloy steel weldments, Scr. Mater. 54 (3) (2006) 489–492.
4. Conclusions A dual phase DP980 steel sheet was welded in room temperature and high temperature (526 °C) in order to observe microstructural and hardness changes. The room temperature fusion and heat-affected zones are composed of martensite with a hardness range between 320 and 500 HV. The high temperature weld fusion zone is composed by grain boundary and bainitic ferrite and austenite and the hardness profile is flatter than the previous case at approximately 280 HV. The residual stresses, after Von Mises criterion, decrease to 1/3 when using HT procedure compared to RT. In addition, the final bending of the
5
Contents lists available at ScienceDirect
Optics and Laser Technology journal homepage: www.elsevier.com/locate/optlastec
12Laser beam welding of DP980 dual phase steel at high temperatures Caroline Cristine de Andrade Ferreira, Vagner Braga, Rafael Humberto Mota de Siqueira, ⁎ Sheila Medeiros de Carvalho, Milton Sergio Fernandes de Lima Photonics Division, Institute for Advanced Studies, Trevo Amarante 1, São José dos Campos, SP 12228-001, Brazil
H I GH L IG H T S
is proposed a new technique for laser welding of dual phase steels. • ItKeyhole welded DP980 steels presents a ferritic-austenitic microstructure. • The hardness, residual stresses and bending of the sheet were reduced. •
A R T I C LE I N FO
A B S T R A C T
Keywords: Laser beam welding Dual phase steels Heat treatment Hardness testing
Dual phase steels have been used in safety components for the automotive industry and represents a standard in terms of energy absorbing steel alloys. After welding, these components present a martensitic transformation in the fusion and heat-affected zones with an intrinsic brittleness. The current contribution presents a procedure where the DP980 steel sheet is kept at a given temperature during and after the laser weld in order to generate bainite instead of martensite. The results using an isothermal treatment at 527 °C shown a microstructure composed by grain boundary and bainitic ferrite and retained austenite with a constant hardness in the base material and heat-affected and fusion zones of 280 HV. The room temperature welds present hardness values between 320 and 500 HV. The high-temperature welds also shown a decrease in the maximum residual stress at the weld centerline by 1/3, with a consequential reduction in the warp of the joined component.
1. Introduction More and more automobile manufacturers are using advanced high strength steels for energy absorption in the case of a car accident [1]. One of most promising alloys for the safety components, such as pillars and roof bows, is called dual phase (DP) steels. Dual phase steels are composed of two phases, martensite and ferrite, which in a given morphology and size, could attain tensile strengths near to 1 GPa. DP980 is especially interesting steel typically surpassing a maximum stress of 980 MPa during tensile tests [2]. Although this alloy fulfills most of the materials selection criteria for the safety components, welding is still a challenge. Any type of welding procedure changes the well-balanced phase distribution and stress concentrators appear. Gould et al. [3] indicated that laser beam welding (LBW) and electric resistance weld (ERW) of DP980 results in a large amount of martensite in the fusion zone (FZ) as a result of rapid cooling. In the same way, Sreenivasan et al. [4] shown that LBW DP980 coupons lose their formability, using limiting dome height tests as standard, due to the martensitic weld bead. These authors also indicated a soft zone in
⁎
the heat-affected zone (HAZ), where necking occurred well before the maximum deformation of the cup [4]. The soft zone is located far from the fusion line (FL) in HAZ as a result of the temper of the base material (BM) martensite. Lima et al. [5] proposed a method to overcome the problems associated to the rapid quenching of advanced high strength steels by using inductive heating during LBW. The inductive heating has the advantage to decrease the cooling rate of the FZ and HAZ in order to produce a given phase or mixture of phases, in particular bainite. The same method could be used to reduce the martensite amount in the welds for the DP980 steel. The objective of the present contribution is to study the inductive heating of a DP980 steel sheet during laser welding in terms of microstructural evolution and hardness. 2. Experimental The current material is a 1-mm thick dual phase class DP980 sheet with the chemical composition giving in Table 1 (source ArcelorMittal).
Corresponding author at: EFO, IEAv, Trevo Amarante 1, 12228-001 Sao Jose dos Campos, SP, Brazil. E-mail address: [email protected] (M.S.F.d. Lima).
https://doi.org/10.1016/j.optlastec.2019.105964 Received 2 August 2019; Received in revised form 16 September 2019; Accepted 24 November 2019 0030-3992/ © 2019 Elsevier Ltd. All rights reserved.
Please cite this article as: Caroline Cristine de Andrade Ferreira, et al., Optics and Laser Technology, https://doi.org/10.1016/j.optlastec.2019.105964
Optics and Laser Technology xxx (xxxx) xxxx
C.C.d.A. Ferreira, et al.
Table 1 Chemical composition of the DP980 steel (wt.%). Fe as the balance [2]. C
Si
Mn
Mo
Cr
Al
0.13
0.03
1.90
0.30
0.20
0.06
The last hot rolling occurred at 650 °C and then the coil was quenched to the room temperature in order to produce both ferrite and martensite. According to the material data sheet, the current DP980 contains 37 ± 6 vol% ferrite and 63 ± 6 vol% martensite [2]. The sheets were sectioned by spark erosion in squares of 40 × 40 mm, ground and washed to remove contaminants and then butt-welded using a fiber laser. The fiber laser (IPG, model YLR-2000) is composed of an ytterbium-doped glass host (wavelength 1080 nm) and a 100 µm diameter delivery fiber. The laser beam was focused using OptoSkand collimator and a focal lens (f 160 mm) to a plate inserted in an induction furnace. The induction furnace was conceived to heat the steel coupons to a given temperature, as reported previously [5]. The temperature in the middle of the fusion zone was recorded using a pyrometer Metis MP25 with emissivity calibrated to the range of temperatures between 100 and 700 °C. After a number of trials, the following parameters were chosen for a full penetration weld of the DP980 steel: laser power of 800 W, weld speed of 50 mm/s and focus on the top of the sheet. The spot diameter was 0.1 mm, as an image of the optical fiber with the same outer diameter. The temperatures of the welds were considered after the MUCG83 software from Peet and Bhadeshia works [6]. The software calculates Widmanstatten ferrite, bainite and martensite start temperatures. The thermodynamic data were obtained using ThermoCalc [7], version M, using FEDAT database. For the estimation of temperatures above 700 °C finite element software (SysWeld, ESI-Group) was used. The same dimensions of the real component had been used in simulations and the mesh resolution around the laser line was about 50 µm node-to-node. The simulation results of the simulation included the temperature evolution in the middle of the fusion zone, the final deformation after the weld and the Von Mises residual stresses. The obtained welds were analyzed using optical microscopy (Zeiss, model Imager2M) and scanning electron microscopy (Hitachi, model TM-3000). For the etching, a solution of 2% nitric acid in ethanol was used. Vickers hardness values were obtained using FutureTech FM-800 equipment. The indentation load was 100 gf for 10 s.
Fig. 1. TTT curves for the current alloy, indicating Widmanstatten ferrite (Ws), Bainite (Bs) and Martensite (Ms) start temperatures. The calculated temperature profiles RT and HTs are plotted together with measured temperatures HTm.
Ms-line. Under these conditions, it is expected almost fully martensitic FZ-grains. Fig. 1b presents the same TTT plot, but now with the high temperature (HT) weld procedure. The furnace temperature was kept at 526 °C, i.e. 100 °C above Ms, and the sheet was kept at this temperature for 10 min before cooling. An overheating of 100 °C is necessary to assure the sheet is kept above Ms during the time the furnace cover is opened for the laser interaction. In Fig. 1b, two curves are plotted, HTs and HTm, representing the simulated and measured temperatures. As could be seen, simulated and measured temperatures are in good agreement in Fig. 1a. This good fit is a result of a proper choice of the input parameter and in particular the laser absorptivity in the keyhole cavity as 85%. According to the estimations, the fusion zone of HT welds precipitated ferritic bainite after 13 s of cooling with an average cooling rate of 3.5 °C/s, which is much lower than the cooling rate when RT crosses Ms, Fig. 1a, i.e. 850 °C/s. The first part of the HTm curve, representing the period where the heat was on, took about 513 s (8.55 min), within bainite is supposed to growth. After this period, a sharp decrease of the HTm curve indicates the sheet was cooling down to 230 °C with a cooling rate of 0.24 °C/s.
3. Results 3.1. Temperature profiles Fig. 1 presents the time–temperature-transformation (TTT) of the current alloy calculated after the MUCG83 software [6]. Widmanstatten ferrite is expected to occur in the range 590–760 °C and the bainite range is 420–590 °C. The software calculated the martensite start temperature as 426 °C, which is in good agreement with Andrews predictions [8]. The Ms-curve presented two paths, one at the initial composition of the alloy (426 °C) and other at 278.6 °C. This last refers to the ThermoCalc [7] calculations of the residual austenite composition when the balance is 1/3 ferrite and 2/3 austenite and might not represent the right position in the TTT plot. Also, according to ThermoCalc, Ms temperature drops subzero when the volume phases fractions are above 50/50%. For the current study, one considers the TTT diagram is good enough to provide an estimation of the phases loci in continuous cooling at different cooling rates. Fig. 1a includes a plot of the room temperature (RT) temperature profile estimated at the middle of the fusion zone (FZ). The time origin is considered at the peak temperature and hence after 0.24 s the primary austenite crosses the
3.2. Microstructural features Fig. 2 presents micrographs of the room temperature (RT) samples. 2
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Fig. 2. Room temperature welded sample: (a) overview of the weld bead showing the fusion zone (FZ), recrystallized grains (RG) region, partial transformation (PT) region and tempered region (TR). (b) Detail of the martensite micro-constituent in FZ. (c) Detail of the PT/TR boundary.
In a general view of the weld bead, Fig. 2a, the microstructure could be divided into four regions: a fusion zone (FZ), a recrystallized grains region (RG), a partially transformed region (PT) and a tempered region (TR). The heat-affected zone is then comprised in three different subzones RG, PT and TR. Fig. 2a shows only one pore in FZ, which could eventually be a result of a base material porosity as also seen in PT. The width of the weld zone is about 0.3 mm and a typical columnar growth is perceived from the fusion line (FL). The microstructure inside FZ is completely martensitic as presented in Fig. 2b. Recrystallized grains are noted within a 0.1 mm region from FL and PT measures 0.1 mm wide, 0.3 mm far from FL. According to FEM results the peak temperatures in the middle of these sub-regions of HAZ are: RG 1200 °C, PT 720 °C and TR 400 °C.
Fig. 3. High temperature welded sample: (a) overview of the weld bead showing the fusion zone (FZ) and the heat-affected zone (HAZ). (b) Detail of the interface between FZ and HAZ. (c) FZ- columnar grains region. (d) FZ – equiaxed grains region. Legend: BF: ferritic bainite; A: austenite; GBF: grain boundary ferrite.
3
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because of the local softening. The boundary between PT and TR is presented in Fig. 2c and it is marked by a darkening of the martensite micro-constituent as long as the ferrite phase unchanged. Fig. 3a shows the weld bead FZ and HAZ for the high temperature condition, i.e. following to the temperature profile in Fig. 1b. Comparing Figs. 3a–2a, it could be noted that the share of the HAZ in subzones is lost in the first case. Also, the annealing of FZ and HAZ resulted in a more homogeneous contrast using the same etching. As can be seen in Fig. 3b the transition FZ to HAZ is smoother and dominated by ferrite (light gray) and austenite (dark gray) phases. These phases are best observed using scanning electron microscopy of the columnar (Fig. 3c) and equiaxed grains (Fig. 3d) in FZ. In both cases, the microstructure is composed of bainitic ferrite and retained austenite as noted in Fig. 3d. Grain boundary ferrite (GBF) is a thin 0.3 µm layer and gave rise to some Widmanstatten growth, although at lower distances (Fig. 3d). Bainitic ferrite (BF) is also noticed in the middle of the grains and corresponds to about 44% of the transformed volume. The rest is austenite (A) stabilized at ambient temperature due to ferrite segregation of elements, in particular C and Mn, according to ThermoCalc.
Fig. 4. Hardness profiles of the RT and HT conditions as a function of the distance of the centerline (CL). The tempered region (TR) is marked in the RT plot.
3.3. Hardness and strain-stress behavior These temperatures are in good agreement with the reported data [9,10], although these data refer to DP600 steel. Antunes et al. [10] cited TR as a possible stress concentrator in tensile strength tests
Fig. 4 presents the Vickers hardness profile for the RT and HT conditions measured from the weld centerline (CL). FZ corresponds approximately to the three points around CL as marked by a dark gray
Fig. 5. FEM of Von Mises residual stresses for (a) RT and (b) HT conditions. The numbers next to the laser path (middle of the sheet) represent residual stresses in FZ and HAZ. The scale at left represents the range of stresses in the entire simulated volume. 4
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welded coupon was also reduced from 0.6° to 0.2° for HT. These results indicate the high-temperature procedure could be interesting for the DP980 steel for thermomechanical processing and use in safety components.
bar. For RT, the hardness value at CL presented a peak compared to the neighboring data as a result of the martensite transformation (Fig. 4). In the RT case, the hardness in RG and PT sub-zones (light gray bars in Fig. 4) are relatively lower, around 300 HV, as a result of the ferrite amount. However, the tempered region (borders if the light gray bars) did not present a drop at the left and right of HAZ as previewed in the literature [11]. This result may be due to the very short time in laser welding, leading to a limited tempering of the martensite in the HAZ. The HT profile is flatter than RT and represents a ferrite-austenite balance in these regions with an average hardness of 280 HV. In Fig. 4, HT side, a little depletion could be noticed near to CL. This depletion could be associated to the centerline segregation during the final stages of solidification as reported in the literature [12]. In these regions, more ferrite is reported and then the hardness drops. Fig. 5 presents the calculated residual stresses according to the Von Mises criterion for the RT and HT conditions. The maximum internal residual stresses attained 780 and 271 MPa for the RT and HT conditions respectively. According to the calculations, RT residual stress attained a level too near to the maximum tensile stress of the alloy, and thus could be considered a threat. The numbers next to the laser line of Fig. 5 represents the residual stresses on the surface of the welded coupons. For the RT case, the surface stresses were 360 and 470 MPa for FZ and HAZ, respectively. In the HT case, these values were 120 and 140 MPa, respectively. These results indicated a reduction to 1/3 of residual stresses when the inductive heating is used instead of conventional welding. The estimated bending of the sheet after the weld is also related to the temperature. The simulated sheet edge rose 0.40 mm in RT condition compared to 0.13 mm for the HT case. For a 40 mm × 80 mm coupon, the final bending angle was 0.6° and 0.2° for RT and HT, respectively. These values are low and difficult to validate using metrology because the actual contortions of the steel sheet.
Declaration of Competing Interest The authors declared that there is no conflict of interest. Acknowledgements CCAF thanks Brazilian National Council for Scientific and Technological Development (CNPq) for an under graduation scholarship (PIBIC-IC). VB thanks São Paulo Research Foundation FAPESP for the doctoral scholarship (grant#2018_23884-5). The authors acknowledge the financial support of São Paulo Research Foundation FAPESP grants#2016/16683-8 and 2016/11309-0. References [1] Auto/Steel Partnership Joining Project (2007) Advanced high strength steel (AHSS) weld performance study for autobody structural components. A/SP Joining Technologies Committee Report. https://www.a-sp.org/-/media/files/asp/ enabling-programs/ahss-joining—ahss-weld-performance-study-for-autobodystructural-components.ashx (accessed 01 August 2019). [2] Arcelor Mittal (2019) Dual Phase steels. https://automotive.arcelormittal.com/DP (accessed 26 July 2019). [3] J.E. Gould, S.P. Khurana, T. Li, Predictions of microstructures when welding automotive advanced high-strength steels, Welding J (2006) 111s–116s. [4] N.N. Sreenivasan, M.M. Xia, S.S. Lawson, Y.Y. Zhou, Effect of laser welding on formability of DP980 steel, ASME J. Eng. Mater. Technol. (4) (2008) 1300410041–041004-9. [5] M.S.F. Lima, D. Gonzalez, S. Liu, Microstructure and mechanical behavior of induction assisted laser welded AHS steels, Welding J. 96 (2017) 376-s–388-s. [6] M. Peet, Bhadeshia HKDH Phase Transformations and Complex Properties Group, Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, U.K., 2019 (accessed 26 July /2019). [7] J.O. Andersson, T. Helander, L. Höglund, P.F. Shi, B. Sundman, Thermo-calc and DICTRA, computational tools for materials science, Calphad 26 (2002) 273–312. [8] K.W. Andrews, Empirical formulae for the calculation of some transformation temperatures, J. Iron Steel Inst. 203 (1965) 721–727. [9] G.C.C. Correard, G.P. Miranda, M.S.F. Lima, Development of laser beam welding of advanced high-strength steels, Int. J. Adv. Manuf. Technol. 83 (2016) 1967–1977. [10] W.D. Antunes, M.S.F. Lima, Experimental development of dual phase steel laser-arc hybrid welding and its comparison to laser and gas metal arc welding, Soldagem Inspeção 21 (3) (2016) 379–386. [11] J. Wang, L. Yang, M. Sun, T. Liu, H. Li, A study of the softening mechanisms of laserwelded DP1000 steel butt joints, Mat. Design 97 (2016) 118–125. [12] I.H. Brown, The role of microsegregation in centreline cold cracking of high strength low alloy steel weldments, Scr. Mater. 54 (3) (2006) 489–492.
4. Conclusions A dual phase DP980 steel sheet was welded in room temperature and high temperature (526 °C) in order to observe microstructural and hardness changes. The room temperature fusion and heat-affected zones are composed of martensite with a hardness range between 320 and 500 HV. The high temperature weld fusion zone is composed by grain boundary and bainitic ferrite and austenite and the hardness profile is flatter than the previous case at approximately 280 HV. The residual stresses, after Von Mises criterion, decrease to 1/3 when using HT procedure compared to RT. In addition, the final bending of the
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