Next Article in Journal
Response and Failure Behavior of Square Tubes with Varying Outer Side Lengths Under Cyclic Bending in Different Directions
Previous Article in Journal
Modeling Mechanical Properties of Industrial C-Mn Cast Steels Using Artificial Neural Networks
Previous Article in Special Issue
Effect of Post-Weld Heat Treatment Cooling Strategies on Microstructure and Mechanical Properties of 0.3 C-Cr-Mo-V Steel Weld Joints Using GTAW Process
 
 
Search for Articles:
Title / Keyword
Author / Affiliation / Email
Journal
Article Type
 
 
Advanced Search
 
Section
Special Issue
Volume
Issue
Number
Page
 
Journals
Metals
Volume 15
Issue 7
10.3390/met15070791
metals-logo
Submit to this Journal Review for this Journal Propose a Special Issue
► Article Menu

Article Menu

share Share announcement Help format_quote Cite question_answer Discuss in SciProfiles
first_page
settings
Order Article Reprints
Font Type:
Arial Georgia Verdana
Font Size:
Aa Aa Aa
Line Spacing:
Column Width:
Background:
Article

Effect of the Heat Affected Zone Hardness Reduction on the Tensile Properties of GMAW Press Hardening Automotive Steel

by
Alfredo E. Molina-Castillo
1,
Enrique A. López-Baltazar
1,*,
Francisco Alvarado-Hernández
1,
Salvador Gómez-Jiménez
1,
J. Roberto Espinosa-Lumbreras
1,
José Jorge Ruiz Mondragón
2 and
Víctor H. Baltazar-Hernández
1,*
1
Materials Science and Engineering Program, Universidad Autónoma de Zacatecas, Zacatecas 98000, Zacatecas, Mexico
2
Innovabienestar de México S.A.P.I de CV, Saltillo 25290, Coahuila, Mexico
*
Authors to whom correspondence should be addressed.
Metals 2025, 15(7), 791; https://doi.org/10.3390/met15070791
Submission received: 10 June 2025 / Revised: 4 July 2025 / Accepted: 8 July 2025 / Published: 13 July 2025
(This article belongs to the Special Issue Welding and Joining of Advanced High-Strength Steels (2nd Edition))
Browse Figures
Versions Notes

Abstract

An ultra-high-strength press-hardening steel (PHS) and a high-strength dual-phase steel (DP) were butt-joined by the gas metal arc welding (GMAW) process, aiming to assess the effects of a high heat input welding process on the structure-property relationship and residual stress. The post-weld microstructure, the microhardness profile, the tensile behavior, and the experimentally obtained residual stresses (by x-ray diffraction) of the steels in dissimilar (PHS-DP) and similar (PHS-PHS, DP-DP) pair combinations have been analyzed. Results indicated that the ultimate tensile strength (UTS) of the dissimilar pair PHS-DP achieves a similar strength to the DP-DP joint, whereas the elongation was similar to that of the PHS-PHS weldment. The failure location of the tensile specimens was expected and systematically observed at the tempered and softer sub-critical heat-affected zone (SC-HAZ) in all welded conditions. Compressive residual stresses were consistently observed along the weldments in all specimens; the more accentuated negative RS were measured in the PHS joint attributed to the higher volume fraction of martensite; furthermore, the negative RS measured in the fusion zone (FZ) could be well correlated to weld restraint due to the sheet anchoring during the welding procedure, despite the presence of predominant ferrite and pearlite microstructures.

1. Introduction

Mitigation of air pollution through the reduction of automobile gas consumption is an imperative global issue nowadays. The reasonable introduction of advanced high-strength steel (AHSS) in autobody car assembly is helping to accomplish such environmental demand by noticeable vehicle weight reduction [1,2,3]. The press hardening steel (PHS) (it belongs to the family of AHSSs) [4,5], also called hot press forming steels (HPF), boron steel, hot-formed steel (HF), 22MnB5 or 1.5528 PHS, is commonly employed in automobile body structures, such as door beams, impact beams, bumpers, pillars, roof rails, and tunnels [6,7,8,9,10]. PH steel is predominantly comprised of a martensite microstructure, along with a reduced amount of retained austenite and/or bainite, which enables PHS with an ultra-high tensile strength and outstanding uniform elongation [7,11,12].
PHS can be effectively joined using gas metal arc welding (GMAW), especially when high load-bearing capacity is required [13,14,15,16,17,18]. However, GMAW is characterized as a high heat input welding process; therefore, essential attention has been taken on the suitable selection of the welding parameters—such as current, voltage, travel speed, shielding gas composition, wire filler metal, among others—[13,14,15,16,19,20] in order to effectively control the unavoidable post-welding phase transformations at the fusion zone (FZ) and the heat affected zone (HAZ), thereby on understanding their influence on the resultant mechanical behavior [21,22,23,24]. Overlapped PHS steel sheets of 1.1 mm in thickness were GMAW joined by employing three different filler metal wires; a considerable drop in hardness (softening) was measured at the subcritical-HAZ (SC-HAZ) due to martensite tempering. According to the tensile-shear testing results, a transition of the failure location was observed from the FZ to the SC-HAZ when increasing the throat thickness and leg length, along with a considerable increase in the tensile shear strength [13]. The weldability of 1.5 GPa martensitic steel (MS) in a lap-joint configuration was evaluated using different welding power types, as well as variation in the shielding gas composition (Ar and CO2). It has been shown that the tensile HAZ fracture predominantly occurs in the softening region over a wider range of DC pulse welding currents, as compared to the BM-HAZ failure location observed in inverted DC power type weldments. Interestingly, the failure mode switches from the weld region when using independently Ar or CO2 shielding gas to the SC-HAZ failure mode when utilizing the mixed shielding gas composition of 80% Ar and 20% CO2 [15]. Preheating at 350 °C was performed on a 22MnB5 steel GMAW welding with three different filler metal compositions and a mixture of shielding gas of 82% Ar and 18% CO2. Although preheating resulted in fusion zone hardness reduction for all types of filler metals, which clearly reduced the overall tensile strength, a non-measurable hardness effect was observed at the SC-HAZ [20]. 22MnB5 steel of 1.6 mm sheet thickness was butt-welded by employing an ER70S-6 welding filler metal. The steel sheets were further austenitized at 860 °C and rapidly cooled to room temperature. The fracture location of the tensile specimens was located at the HAZ-BM interface (softened region), whereas for the heat-treated samples, the failure switched to the FZ-HAZ interface [14].
On the other hand, the effect of softening in wire arc welding PHS has been compared between tensile and fatigue assessments [16,25,26]. For instance, GMAW-controlled short-circuit and plasma arc-welding (PAW) PHS specimens of 3 mm in sheet thickness have been butt-welded; results indicated that the failure location of the reinforcement removed tensile specimens; randomly occurred either at the FZ, SC-HAZ, and/or intercritical-HAZ (IC-HAZ). For the fatigue trials, damage always occurred in the weld toe region for specimens without weld reinforcement removal, as well as inside the FZ and BM in the ground-flush samples [16]. A comparative analysis of the fatigue behavior of butt-welded joints by GMAW on dual-phase (DP) steel (DP440 and DP590) and martensitic steel (MS1500) was carried out. The study revealed that the fracture locations during the tensile testing were expected to be encountered at the softened SC-HAZ for MS1500 and DP590 steels. Interestingly, all failures occurred at the weld toe in the fatigue specimens for all the analyzed steels. Additionally, the fatigue tests for specimens without weld beads revealed that fatigue life correlates well with the lowest hardness location [25].
Several reports on AHSS joined to other dissimilar steel grades utilizing different arc welding processes, such as GMAW, pulsed welding processes (P-GMAW, P-GTAM, and P-PAW), laser beam welding (LBW), or hybrid welding, have been recently disclosed [27,28,29]. For instance, a comprehensive review by focusing on arc welding (GMAW and GTAW) and laser welding of AHSS with strength levels above 800 MPa pointed out the dissimilar pairs: S700MC-S960QC [22], TWIP-TRIP [24], and TRIP780-DP980 [30]. The main results indicated a decrease in strength and hardness reduction at the HAZ, grain coarsening or microstructural changes at the HAZ, and slight softening at the HAZ of the TRIP steel side [27]. A dissimilar Yb: YAG disc-laser welding butt-joint of a high-strength low-alloy carbon steel and medium-Mn low-Ni austenitic stainless steel was carried out; the main finding indicated that the tensile strengths and flow behaviors of the dissimilar joints were as low as the softer steel [28]. TRIP800 and DP800 steel sheets with a thickness of 1.2 mm were paired together using GMAW and LBW; the failure location switched from the SC-HAZ in the GMAW samples to the BM failure in the LBW samples, with both failures found at the TRIP steel side [29]. In particular, PHS steel has been joined to HSLA [31] and DP steel [32,33] using various welding processes. Boron steel was paired with HSLA steel using GMAW, and three different welding configurations were employed: single lap-shear, double lap-shear, and start-stop. The maximum tensile strength of the weldment was reported and compared among the assessed configurations; however, no details were provided with regard to failure location [31]. A laser-GMA hybrid welding process was used to pair 22MnB5 and DP800 galvanized steel sheets in an overlap configuration. The tensile testing results suggested an average strength of 800 MPa for the range of gap sizes, utilized between 0.4 and 0.6 mm; undoubtedly, the overall strength of the weldment was in agreement with the DP800 steel strength [32]. A continuous fiber laser was utilized for tailor welding PHS (1500 MPa) to DP980 steel, with/without Al-Si coating; the sheet thickness in both steels was 1.5 mm. The drop in hardness at the SC-HAZ of the DP side was reported to be 47 HV (from 325 HV up to 278 HV), whereas in the case of the PHS was 147 HV (from 460 HV up to 313 HV); as indicated by the tensile testing results, the fracture location occurred at the base metal of the DP980 side in both conditions (with/without Al-Si coating) along with a maximum strength range in between 966 MPa to 975 MPa [33]. A DP980 and 2 GPa PHS of 1.8 mm in thickness steel sheets have been welded via LBW and cold metal transfer (CMT) welding; the tensile strength was slightly higher by 50 MPa in the LBW specimen, whereas the total elongation was higher in the CMT dissimilar welded specimens [34]. According to the aforementioned, it is clear that PHS steel has been barely joined to other steel grades in dissimilar steel sheet combinations. In addition, no reports have been disclosed so far on pairing PHS steel with DP steel under a high heat input welding process, such as GMAW.
The analysis of residual stress (RS) is of remarkable significance for a better comprehension of the mechanical performance of weldments; understanding RS distribution may lead to predicting possible failures when subjected to tension, fatigue, and/or formability [29]. Investigations made on experimental RS measurements across a weldment of AHSSs have indicated that the RS distribution strongly depends on: (a) the type of welding processes [29], (b) the setup of the welding parameters (due to the heat input control) [35,36], (c) the post-welding heat input (heat treatment, paint baking, etc.) [37,38], (d) post-welding mechanical cold deformation [39], and (e) the mismatch in steel sheet thickness [40]. Although many efforts have been made to measure and analyze the experimental surface residual stresses of AHSS weldments, no major discussion has been provided with regard to the resultant RS at the SC-HAZ, in particular for PHS steels.
Many attempts have been conducted to analyze the microstructure and the mechanical properties of the PHS steel subjected to different types and conditions of welding thermal cycles; however, no efforts have been made to systematically study the structure-property relationship coupled with the experimental residual stress analysis when pairing ultra-high and high-strength steel sheets (i.e., 1500 MPa and 980 MPa) and by employing a high heat input welding process. Therefore, this work aims to evaluate the effect of the GMAW process on the microstructure, hardness, tensile behavior, and surface residual stress of butt-joined welded PHS and DP steel in similar and dissimilar combinations.

2. Materials and Methods

PHS and DP steel of 1.2 mm in sheet thickness were utilized in this work. Table 1 lists both steel chemistries obtained through optical emission spectroscopy (OES) using a Bruker Q4 TASMAN (Bruker, Mexico City, Mexico); the calculated carbon equivalent (CEy), as per Yurioka’s equation, has been additionally calculated [41].
PHS and DP steel sheets of 100 mm length and 50 mm width were paired in similar (DP-DP and PHS-PHS) and dissimilar (PHS-DP) combinations through butt-joint welding, single square groove configuration, and by utilizing an Infra MM 300-E (Infra, Naucalpan, Mexico) GMAW machine operated at 10.1 kVA/7.7 kW with non-pulsed short circuit metal transfer mode. An ER70S-6 solid welding wire of 0.9 mm in diameter was used; a travel and work angle of 90° and a direct current electrode positive were configured. The selected GMAW welding process parameters are provided in Table 2.
The welding torch motion was performed by a Panasonic YA-1CA161 robotic arm and a GII-016 controller (Zacatecas, Mexico); all sheets were fixed in copper-alloy jigs, and a single-pass welding was applied to one side of a joint, as shown in Figure 1.
The welded samples were sectioned in the transverse direction for microstructure and hardness analysis and prepared according to the standard practice for the preparation of metallographic specimens [43]; the etching process was realized with a 2% Nital solution. Microhardness indentations were performed along the transverse cross-section of the weldment using an HVM Shimadzu (Shimadzu, Zacatecas, Mexico) hardness tester with an applied load of 300 g, dwell time of 10 s, and an indentation spacing of 200 µm.
A wire electric discharge machine (EDM) was employed for cutting all specimens perpendicular to the welding direction, as well as to machine the tensile specimens. A total of three tensile specimens were prepared per welded condition (PHS-PHS, DP-DP, and PHS-DP) and similarly per BM (PHS and DP). The tensile sample dimensions were 25 mm in gage length, 6 mm in width, 6 mm in the radius of the fillet, and 100 mm in the overall length of the sample, according to the standard [44]. It is important to mention that the weld line was positioned at the center of the gauge length. The uniaxial tensile tests were carried out in a Shimadzu AG1 universal testing machine (Shimadzu, Zacatecas, Mexico).
Samples of 10 mm in width by 50 mm in length (i.e., the length dimension was transverse to the welding direction) for the residual stress measurements were cut from the original welded specimens by EDM. Two measurements per location (longitudinal and transverse) were executed at the FZ, UC-HAZ, SC-HAZ, and the BM. The sin2ψ procedure was applied in this investigation to determine the surface residual stress through x-ray measurements along the welded specimens at different locations, namely FZ, HAZ, and BM. Further details regarding the sin2ψ procedure have been described elsewhere [29]. Cu radiation and a micro-diffraction pin of 1.0 mm in size, adapted in a Bruker Advanced D8 Eco™ diffractometer (Bruker, Zacatecas, Mexico), were employed for measuring the residual stresses from 134° ≤ 2θ ≤ 140°. The α-Fe (222) diffraction peak was evaluated by setting a step size of 0.005° and a time of 3 s per step. The residual stresses were calculated from the experimental data (six different ψ angles starting at 0° up to 45°) through the slope of the least-squares line method by supposing material constant values E = 247 GPa and Poisson ratio v = 0.25. The lattice spacing (do) in this experiment was measured at Ψ = 0 due to elastic strains, which do not introduce more than a 0.1% difference between the true do and d at any ψ [40]. One of the most acceptable approaches for providing the best linear regressions with high precision is the gravity peak positioning method [41] for quantitative XRD stress evaluation; therefore, the accuracy of the stress measurements in this work lies within the range of ±41~52 MPa.

3. Results and Discussion

3.1. Microstructure and Hardness

Figure 2 shows the transverse cross-section macrostructures of the GMAW samples for (a) PHS-PHS, (b) DP-DP, and (c) PHS-DP steels, indicating the various regions of the weldment, i.e., the fusion zone (FZ), the heat-affected zone (HAZ) conveniently subdivided as upper-critical (UC-HAZ), inter-critical (IC-HAZ), and sub-critical (SC-HAZ), as well as the base metal (BM). The extension of the abovementioned regions is quite consistent among the welded specimens; furthermore, a full penetration without visible defects (such as porous and/or slag inclusions) is evidently observed in the FZ for all samples.
The high magnification images of the FZ microstructure for PHS-PHS (Figure 3b), DP-DP (Figure 3d), and PHS-DP (Figure 3f) joints were taken from the indicated regions (by the blue rectangles) in Figure 3a–e, respectively. A balanced content of fine and equiaxed grains of pearlite (P) and ferrite (α) phases is revealed in the FZ of the PHS-PHS joint (Figure 3b); whereas larger grains of P and α phases were found at the FZ of the DP-DP specimen along with a reduced fraction of Widmanstätten ferrite morphology (Figure 3d). Interestingly, the larger volume fraction of α as compared to P phase is clearly observed in the FZ of the PHS-DP steel joint combination (Figure 3f).
Figure 4 shows the coarse and fine regions (within the UC-HAZ) and the IC-HAZ microstructures of both PHS and DP steels. Predominant martensite (α’) along with bainite (B) aggregates is clearly revealed at the UC-HAZ of both steels; however, a larger fraction of bainite is distinctly observed at the UC-HAZ of the DP steel side. As expected, the IC-HAZ is composed of α, P, and a reduced fraction of α’ in both steels owing to partial transformation to austenite during the heating cycle and further rapid cooling, resulting in α’ formation.
The BM microstructure at high magnification for PHS and DP steels is provided in Figure 5a,b, respectively. The base metal microstructure of PHS is predominantly composed of martensite (α’) phase along with a reduced fraction of bainite (B) with an average hardness of 507 HV. The base metal microstructure of the DP steel consists of a ferrite (α) matrix, martensite (α’) islands, and a reduced volume fraction of retained austenite phase (γ or RA) with an average hardness of 323 HV. The decomposition of the original BM microstructure, owed to the tempering of martensite [45] at the SC-HAZ, is clearly seen in Figure 5c,d for PHS and DP steels, respectively. The recovery of the α phase and the precipitation of cementite carbides [46] as a result of the tempering of martensite (TM) have been pointed out in both steels. However, due to the lower volume fraction of martensite (49 pct.) on the DP steel side, the SC-HAZ hardness was measured at 205 HV (a drop of 118 HV), whereas the SC-HAZ hardness in PHS steel was 245 HV (a considerable drop in hardness of 262 HV).
The hardness profile distribution throughout the weldment for the steel combinations is given in Figure 6. The average FZ hardness resulted in 280 HV, 278 HV, and 276 HV for the PHS-PHS, DP-DP, and PHS-DP steel combinations, respectively. The average FZ hardness is quite consistent among the welded specimens and, thereby, in agreement with the predominant presence of ferrite and pearlite microstructures previously described. The peak hardness value within the weldment was measured at the BM followed by the UC-HAZ in the PHS steel (i.e., 470 HV) due to the high volume fraction of martensite, whereas the UC-HAZ of the DP steel side averaged 419 HV due to the high martensite content along with a considerable volume fraction of bainite. A transitional reduction in hardness up to the SC-HAZ is clearly seen along the IC-HAZ of both steel sides.

3.2. Uniaxial Tensile Behavior

The engineering stress-strain curves for the welded specimens in both similar and dissimilar steel combinations are provided in Figure 7a; in addition, the BM uniaxial tensile curves (dashed lines) have been incorporated for benchmarking. Diminished tensile properties (both ductility and strength) are clearly revealed in the welded specimens as compared to the BM samples of both steels. According to the results listed in Table 3 (indicating the average value and the standard deviation), the ultimate tensile strength (UTS) of the PHS-PHS steel weldment is drastically reduced by about 46%, whereas in the case of the DP-DP steel weldment, the UTS dropped by almost 29% as compared to their respective BM tensile strength. On the contrary, the uniform elongation was drastically reduced in the DP steel, at approximately 58%, whereas for the PHS steel, the ductility dropped by only 10%. Interestingly, the UTS resulted close to the DP-DP, and the ductility was similar to that of the PHS-PHS for the dissimilar weld combination PHS-DP, thereby showing that the dissimilar pair combination resulted in limited mechanical performance, inherited by both PHS and DP steels, reduced properties.
On the other hand, the failure location occurred at the softened SC-HAZ in all welded specimens, as pointed out by the yellow arrow in Figure 7b. Despite the hardness reduction was being severe in the PHS steel, the hardness was lower at the SC-HAZ of the DP steel (205 HV) as compared to the PHS steel (245 HV); this is also in agreement with the failure location consistently encountered at the SC-HAZ of the DP steel in the dissimilar joint PHS-DP.
In order to further explain the failure mechanism of the tensile specimens, Figure 8 illustrates a representative image of (a) a transverse cross-section of the failed specimen, (b) a representative SEM micrograph of the crack nucleation, and (c) a fractography image. A plastically deformed and reduced cross-section area (necking) confirms a ductile failure location just at the SC-HAZ in the specimen, as depicted in the macrostructure in Figure 8a. The representative microstructure in Figure 8b reveals that crack nucleation clearly originated at the interface between ferrite (α) and the tempered martensite (TM), as pointed out by the arrow. The fracture path seems to be aligned crosswise to the elongated α and the TM banded structure. Figure 8c shows the representative fracture mode at the SC-HAZ, exhibiting deep conical equiaxed and elongated dimples of various size and sparsely distributed probably due to the high volume fraction of soft microstructures (ferrite), along with smaller regions of sheared surface near the center of the image, which indicates a different failure mechanism in the region of fracture.
A distinct correlation is confirmed between the softening at the SC-HAZ (due to the tempering of martensite) and the resultant tensile resistance of the weldments; for instance, the hardness loss in the SC-HAZ with respect to the BM was calculated at about 52% for the PHS steel, and, correspondingly, the UTS loss of the PHS weldment resulted in about 46%. By following a similar trend, the hardness loss in the SC-HAZ of the DP steel was about 34%, and the corresponding UTS loss was 29%. On the other hand, the lowest hardness value was measured at the SC-HAZ of the DP steel, which is the location of failure in the dissimilar PHS-DP weldment, and, thereby, the resultant UTS value was similar to that obtained in the DP-DP weldment.
While the abovementioned trend between the hardness loss in the SC-HAZ and the UTS of the weldment seems to be congruous for similar steel type joints (PHS-PHS and DP-DP), the resultant poor properties in the dissimilar PHS-DP joint can be explained in terms of the inherited resultant microstructure at the SC-HAZ in both steels. For instance, the overall strength in the dissimilar PHS-DP steel was comparable to that of the welded DP steel, indicating that the load-carrying capacity of the dissimilar weldment is closely correlated to the weaker region, i.e., the SC-HAZ of the DP side, in which a larger volume fraction of ferrite and lower fraction of the precipitated carbides within the tempered martensite (inherited by the lower volume fraction of the prior martensite phase) was revealed in Figure 5. It is important to recall here that HAZ-softening in AHSS strongly depends on the martensite volume fraction [47]. On the other hand, the uniform elongation in the dissimilar joint (PHS-DP) was comparable to that of the welded PHS steel; thus, suggesting that the volume fraction of the hard martensite phase at the unaffected BM in the PHS steel side, the hard martensite at the HAZ in both steels, and the corresponding volume fraction (49%) of martensitic islands in the DP side, withstands substantial amount of deformation in the dissimilar weldment.

3.3. Residual Stress Analysis

Figure 9 shows the experimental x-ray diffraction measured residual stress (RS) profiles along the weldments. Whereas the ordinate displays the residual stress value, the abscissa corresponds to the RS measured location points, the zero value corresponds to the FZ location and moving toward the right/left of the graph, and the values correspond to other locations, such as the HAZ and the BM of either side of the welded specimen. Negative longitudinal (LRS) and transverse (TRS) residual stresses (all below −200 MPa) were measured along the various zones of the PHS-PHS steel joint (Figure 9a). A particular trend was also observed in the DP-DP steel joint, as the RS values range between −244 and +109 MPa (Figure 9b). Interestingly, the dissimilar PHS-DP steel joint resulted in a high RS on the DP steel side (in a range between −73 MPa and 157 MPa), followed by negative TRS and LRS at the FZ (−102 MPa and −120 MPa, respectively), and even lower TRS and LRS on the PHS steel side (in a range between −219 MPa and −468 MPa), as shown in Figure 9c. Hence, it is clearly demonstrated that a more negative RS is achieved in the PHS steel side in both similar and dissimilar welding joints. It is important to note that in similar steel joints, the LRS are consistently lower if compared to TRS in both sheets of steel along the weldment, except for the BM. The previously mentioned behavior is in agreement with other contributions in which LRS are consistently reported as lower than TRS in arc-welding butt-joints [29,40].
Although the residual stress distribution along a sheet weldment strongly depends on various key factors, such as the weld joint restraint [48], the sheet geometry (thickness, width, and length) [40], the type of joint (butt, overlap, spot, etc.) [49], the resultant weld geometry (reinforcement size, width, penetration, etc.) [50], and the net heat input (welding parameters) [51]; all previously mentioned factors have been deliberately fixed in all weldments in this investigation. Therefore, the effect of the dissimilar steel grade (between PHS and DP steels) seems to have an important influence on the final RS distribution profile due to the mismatch properties of the studied steels; in particular, the prior microstructure (BM) and the post-weld localized phase transformations (i.e., FZ and HAZ) in both steels. According to this, there is a well-known trend to form compressive stresses in steel due to the martensite phase transformation in steel. The non-isotropic volume change makes displacive transformations in steel inherently associated with a dilatation that occurs normal to the habit plane, hence, promoting micro residual stresses [52]. The effects of martensite phase transformation on the development of residual stresses in steel have been extensively studied. For instance, during the transformation stage of the parent austenite (in which stresses are comparable to the yield strength of austenite) into martensite, the stresses are reduced to almost zero at the beginning of the transformation to martensite. When all the austenite phase is depleted, the thermal contraction continues, thereby causing the accumulation of stress. On the ongoing phenomenon, the martensite start is about ~250 °C, thus favoring variants that comply with the stress that lead to the possibility of turning them into compressive stresses [53,54,55]. As previously mentioned, PHS steel is comprised of a predominant martensite phase, whereas DP steel is composed of 49% martensite; this is quite consistent with the resultant RS values measured at the BM of both steels (it is worth noting that more negative RS are measured in the PHS steel). Although the post-welding phase transformations developed along the weldment mostly ended in martensite (i.e., HAZ), which further justifies the presence of compressive RS, two other important zones resulted in tempered martensite (SC-HAZ) and ferrite-pearlite as preponderant microstructures (FZ) in all welded conditions; despite that, compressive LRS were measured in both zones (FZ and SC-HAZ). In one sense, the lowest hardness value at the SC-HAZ of the DP steel is a quite narrow region (~800 µm), whereas in the PHS steel side is solely a local point (see Figure 6), which could further explain the greater negative RS value measured in the PHS steel side; in fact, the is because of the SC-HAZ surroundings are composed of martensite phase. On the other hand, the measured compressive stresses in the FZ, regardless of the presence of ferrite and pearlite microstructures, could be explained in terms of weld restraint (sample anchoring) during the solidification stage [49,52,56]. Similar contributions have reported the influence of the weld restraint on the final residual stress distribution [40,57,58], and negative RS at FZ due to clamping was similarly obtained [57]. The grain growth of the primary solidification structure is restricted by the hard surroundings of the base metal sheet, which is simultaneously strongly fixed to the copper-alloy jigs. It is important to mention two key issues: (a) the compressive residual stresses were consistently measured in the FZ of all three welded conditions (PHS-PHS, DP-DP, and PHS-DP), and (b) the negative residual stresses in the FZ of GMAW of thin AHSS sheets have been consistently reported in other contributions [29,40].

4. Conclusions

The dissimilar steels PHS and DP were butt-joined by employing the GMAW process. The post-weld microstructures of the different regions, along with the uniaxial tensile behavior and the residual stress distribution across the weldment, have been analyzed. Therefore, the following conclusions are obtained:
  • A drastic hardness reduction of 262 HV was measured at the SC-HAZ of the PHS steel, which contrasts with the drop in hardness of 118 HV at the SC-HAZ of the DP steel. Correspondingly, the prior martensite content in the PHS steel had a marked influence on post-weld softening, as the higher the amount of martensite, the higher the volume of martensite tempering at the SC-HAZ of the PHS steel.
  • The tensile behavior of the weldments resulted in a clear reduction of strength and ductility as compared to their corresponding BM. While the UTS was drastically reduced with respect to the PHS steel (i.e., 46%), the associated elongation was severely diminished in the DP steel (i.e., 58%). The strength of the dissimilar welding pair, PHS-DP steel (699 MPa), was similar to the DP-DP weldment, whereas the elongation was comparable to that of the PHS-PHS weldment (7.9%).
  • All failure locations systematically occurred at the SC-HAZ in similar steel joints, whereas for dissimilar pair PHS-DP, the failure was consistently found at the SC-HAZ of the DP steel side.
  • Compressive residual stresses resulted along the weldments of all specimens. The more negative stresses observed at the PHS steel have been attributed to the presence of preponderant martensite. It is believed that the negative stresses measured at the FZ were promoted by the weld restraint (sheet anchoring during the welding procedure) in spite of the presence of soft ferrite and pearlite phases.

Author Contributions

Conceptualization, A.E.M.-C., E.A.L.-B. and V.H.B.-H.; methodology, F.A.-H., S.G.-J. and V.H.B.-H.; validation, E.A.L.-B., J.R.E.-L. and J.J.R.M.; formal analysis, S.G.-J. and V.H.B.-H.; investigation, E.A.L.-B., F.A.-H. and V.H.B.-H.; data curation, A.E.M.-C., E.A.L.-B., J.R.E.-L., J.J.R.M. and V.H.B.-H.; writing—original draft preparation, A.E.M.-C., E.A.L.-B., F.A.-H., S.G.-J., J.R.E.-L., J.J.R.M. and V.H.B.-H.; writing—review and editing, A.E.M.-C., E.A.L.-B., F.A.-H., S.G.-J., J.R.E.-L., J.J.R.M. and V.H.B.-H. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the Universidad Autónoma de Zacatecas, and Innovabienestar de Mexico S.A.P.I de CV.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding authors.

Acknowledgments

Alfredo E. Molina-Castillo wants to thank the Maestría en Ciencia e Ingeniería de los Materiales, Universidad Autónoma de Zacatecas.

Conflicts of Interest

The authors declare no conflicts of interest.

References

  1. Schmitt, J.-H.; Iung, T. New developments of advanced high-strength steels for automotive applications. Comptes Rendus Phys. 2018, 19, 641–656. [Google Scholar] [CrossRef]
  2. Matlock, D.; Speer, J.G.; De Moor, E.; Gibbs, P.J. Recent developments in advanced high strength sheet steels for automotive applications: An overview. JESTECH 2012, 15, 1–12. [Google Scholar]
  3. Bonnen, J.; Mandapati, R.; Kang, H.; Khosrovaneh, A.K.; Amaya, M.; Citrin, K.; Shih, H.-C. Durability of Advanced High Strength Steel Gas Metal Arc Welds. SAE Int. J. Mater. Manuf. 2009, 2, 155–171. [Google Scholar] [CrossRef]
  4. Kuziak, R.; Kawalla, R.; Waengler, S. Advanced high strength steels for automotive industry. Arch. Civ. Mech. Eng. 2008, 8, 103–117. [Google Scholar] [CrossRef]
  5. Mendiguren, J.; Herrero-Dorca, N.; Saenz de Argandoña, E.; Galdos, L. Press hardening of alternative materials: Conventional high- strength steels. Int. J. Mater. Form. 2018, 11, 663–670. [Google Scholar] [CrossRef]
  6. Wei, X.; Chai, Z.; Lu, Q.; Hu, J.; Liu, Z.; Lai, Q.; Wang, J.; Xu, W. Cr-alloyed novel press-hardening steel with superior combination of strength and ductility. Mater. Sci. Eng. A 2021, 819, 141461. [Google Scholar] [CrossRef]
  7. Fan, D.; Kim, H.; De Cooman, B. A Review of the Physical Metallurgy related to the Hot Press Forming of Advanced High Strength Steel. Steel Res. Int. 2010, 80, 241–248. [Google Scholar] [CrossRef]
  8. WorldAutoSteel AHSS Application Guidelines. Press Hardened Steels. Available online: https://ahssinsights.org/metallurgy/steel-grades/phs-grades/ (accessed on 16 October 2024).
  9. Cheng, Z.; Gao, M.; Liu, J.; Wang, S.; Wu, G.; Gao, J.; Wu, H.; Mao, X. Multi-Scale Microstructural Tailoring and Associated Properties of Press-Hardened Steels: A Review. Materials 2023, 16, 3799. [Google Scholar] [CrossRef]
  10. Hannula, J.; Porter, D.A.; Kaijalainen, A.; Kömi, J. Evaluation of Mechanical Properties and Microstructures of Molybdenum and Niobium Microalloyed Thermomechanically Rolled High-Strength Press Hardening Steel. JOM 2019, 71, 2405–2412. [Google Scholar] [CrossRef]
  11. Votava, F.; Jirková, H.; Kučerová, L.; Jeníček, Š. Study of Transition Areas in Press-Hardened Steels in a Combined Tool for Hot and Cold Forming. Materials 2023, 16, 442. [Google Scholar] [CrossRef]
  12. Seo, E.J.; Cho, L.; De Cooman, B.C. Application of Quenching and Partitioning (Q&P) Processing to Press Hardening Steel. Metall. Mater. Trans. A 2014, 45, 4022–4037. [Google Scholar] [CrossRef]
  13. Kim, D.-Y.; Lee, T.H.; Kim, C.; Kang, M.; Park, J. Gas metal arc welding with undermatched filler wire for hot-press-formed steel of 2.0 GPa strength: Influence of filler wire strength and bead geometry. Mater. Today Commun. 2023, 34, 105244. [Google Scholar] [CrossRef]
  14. Yaakob, K.I.; Ishak, M.; Idris, S.R.A.; Aiman, M.H.; Quazi, M.M. Characterization of heat-treated gas metal arc-welded boron steel sheets. Int. J. Adv. Manuf. Technol. 2018, 94, 827–834. [Google Scholar] [CrossRef]
  15. Hwang, I.; Yun, H.; Kim, D.; Kang, M.; Kim, Y.-M. Gas metal arc weldability of 1.5 GPa grade martensitic steels. Met. Mater. Int. 2018, 24, 149–156. [Google Scholar] [CrossRef]
  16. de Lima, E.H.; de Almeida, D.T.; Souza, D.; Bianchi, K.E.; Mohrbacher, H. Evaluation of Fatigue Performance of Press Hardening Steel Joints Welded by GMAW-CSC and PAW Processes. Metals 2022, 12, 2131. [Google Scholar] [CrossRef]
  17. Shome, M.; Tumuluru, M. (Eds.) Introduction to welding and joining of advanced high-strength steels (AHSS). In Welding and Joining of Advanced High Strength Steels (AHSS); Woodhead Publishing: Cambridge, UK, 2015; pp. 1–8. [Google Scholar]
  18. Duchet, M.; Haouas, J.; Gibeau, E.; Pechenot, F.; Honecker, C.; Munier, R.; Weber, B. Improvement of the fatigue strength of welds for lightweight chassis application made of Advanced High Strength Steels. Procedia Struct. Integr. 2019, 19, 585–594. [Google Scholar] [CrossRef]
  19. Führ, T.; Almeida, D.; Bertol, D.; de Souza, J.; Lopes, A. Mechanical performance of heavy gauge press hardened steel modified with Nb and Mo additions on GMAW welded joints. Tecnol. Metal. Mater. Min. 2021, 18, e2123. [Google Scholar] [CrossRef]
  20. Pfeifer, T.; Węglowski, M.S. Characteristic of MAG welded joints of 22MnB5 steel grade for automotive industry. Adv. Manuf. Sci. Technol. 2014, 38, 71–80. [Google Scholar]
  21. Varol, F.; Ferik, E.; Ozsarac, U.; Aslanlar, S. Influence of current intensity and heat input in Metal Inert Gas-brazed joints of TRIP 800 thin zinc coated steel plates. Mater. Des. 2013, 52, 1099–1105. [Google Scholar] [CrossRef]
  22. Njock Bayock, F.; Kah, P.; Mvola, B.; Layus, P. Effect of Heat Input and Undermatched Filler Wire on the Microstructure and Mechanical Properties of Dissimilar S700MC/S960QC High-Strength Steels. Metals 2019, 9, 883. [Google Scholar] [CrossRef]
  23. Netto, A.; Njock Bayock, F.M.; Kah, P. Optimization of GMAW Process Parameters in Ultra-High-Strength Steel Based on Prediction. Metals 2023, 13, 1447. [Google Scholar] [CrossRef]
  24. Májlinger, K.; Kalácska, E.; Russo Spena, P. Gas metal arc welding of dissimilar AHSS sheets. Mater. Des. 2016, 109, 615–621. [Google Scholar] [CrossRef]
  25. Ahiale, G.K.; Oh, Y.-J. Microstructure and fatigue performance of butt-welded joints in advanced high-strength steels. Mater. Sci. Eng. A 2014, 597, 342–348. [Google Scholar] [CrossRef]
  26. Svoboda, H.G.; Nadale, H.C. Fatigue Life of GMAW and PAW Welding Joints of Boron Microalloyed Steels. Procedia Mater. Sci. 2015, 9, 419–427. [Google Scholar] [CrossRef]
  27. Perka, A.K.; John, M.; Kuruveri, U.B.; Menezes, P.L. Advanced High-Strength Steels for Automotive Applications: Arc and Laser Welding Process, Properties, and Challenges. Metals 2022, 12, 1051. [Google Scholar] [CrossRef]
  28. Hamada, A.; Ali, M.; Ghosh, S.; Jaskari, M.; Keskitalo, M.; Järvenpää, A. Mechanical performance and formability of laser-welded dissimilar butt joints between medium-Mn stainless steel and high-strength carbon steel. Mater. Sci. Eng. A 2022, 831, 142200. [Google Scholar] [CrossRef]
  29. Baltazar-Hernández, V.H.; López-Baltazar, E.A.; Alvarado-Hernández, F.; Gómez-Jiménez, S.; Ruiz-Mondragón, J.J.; Biro, E.; Zhou, N. Surface Residual Stress Analysis in GMAW and LBW of the Dissimilar TRIP-DP Steels Joint: An Experimental Approach. Metals 2022, 12, 880. [Google Scholar] [CrossRef]
  30. Sharma, R.S.; Molian, P. Yb:YAG laser welding of TRIP780 steel with dual phase and mild steels for use in tailor welded blanks. Mater. Des. 2009, 30, 4146–4155. [Google Scholar] [CrossRef]
  31. Kang, H.; Khosrovaneh, A.; Link, T.; Bonnen, J.; Amaya, M.; Shih, H.-C. The Effect of Welding Dimensional Variability on the Fatigue Life of Gas Metal Arc Welded Joints. SAE Int. J. Mater. Manuf. 2011, 4, 298–313. [Google Scholar] [CrossRef]
  32. Möller, F.; Kügler, H.; Goecke, S.-F. Laser-GMA-Hybrid Welding of high strength multi-material joints. In Proceedings of the Lasers in Manufacturing Conference, Center München, Germany, 22–25 June 2015. [Google Scholar]
  33. Di, H.S.; Sun, Q.; Nie, X.K.; Wang, X.N.; Chen, X.M. Mirostructure and properties of laser welded joints of dual phase and press-hardened steels. Procedia Eng. 2017, 207, 1665–1670. [Google Scholar] [CrossRef]
  34. Chen, X.; Jin, Z.; Liang, J.; Wang, Y.; Yang, D.; Yi, H.; Xie, G. A comparative study of the galvanized DP980 steel and 2 GPa PHS welded joint via laser beam welding and cold metal transfer welding. Mater. Lett. 2024, 369, 136726. [Google Scholar] [CrossRef]
  35. Kong, F.; Kovacevic, R. Measurement of surface residual stresses and testing mechanical properties of high-strength steel butt joints obtained by hybrid laser/gas metal arc welding. J. Strain Anal. Eng. Des. 2013, 48, 437–445. [Google Scholar] [CrossRef]
  36. Colombo, T.C.A.; Rego, R.R.; Otubo, J.; de Faria, A.R. Mechanical reliability of TWIP steel spot weldings. J. Mater. Process. Technol. 2019, 266, 662–674. [Google Scholar] [CrossRef]
  37. Kang, H.T.; Lee, Y.-L.; Sun, X.J. Effects of residual stress and heat treatment on fatigue strength of weldments. Mater. Sci. Eng. A 2008, 497, 37–43. [Google Scholar] [CrossRef]
  38. Shamsujjoha, M.; Enloe, C.M.; Chuang, A.C.; Coryell, J.J.; Ghassemi-Armaki, H. Mechanisms of paint bake response in resistance spot-welded first and third generation AHSS. Materialia 2021, 15, 100975. [Google Scholar] [CrossRef]
  39. Ren, S.; Ma, N.; Tsutsumi, S.; Watanabe, G.; He, H.; Cao, C.; Huang, L. Post-weld cold working for fatigue strength improvement of resistance spot welded joint of advanced high-strength steel. J. Mater. Process. Technol. 2022, 299, 117364. [Google Scholar] [CrossRef]
  40. Mujica, L.; Weber, S.; Pinto, H.; Thomy, C.; Vollertsen, F. Microstructure and mechanical properties of laser-welded joints of TWIP and TRIP steels. Mater. Sci. Eng. A 2010, 527, 2071–2078. [Google Scholar] [CrossRef]
  41. Kasuya, T.; Yurioka, N. Determination of Necessary Preheat Temperature to Avoid Cold Cracking under Varying Ambient Temperature. ISIJ Int. 1995, 35, 1183–1189. [Google Scholar] [CrossRef]
  42. AWS A5.18/A5.18M:2023; Specification for Carbon Steel Electrodes and Rods for Gas Shielded Arc Welding. American Welding Society: Doral, FL, USA, 2023.
  43. ASTM E3-95; Standard Practice for Preparation of Metallographic Specimens. ASTM International: West Conshohocken, PA, USA, 1995.
  44. ASTM E8/E8M-09; Standard Test Methods for Tension Testing of Metallic Materials. ASTM International: West Conshohocken, PA, USA, 2013.
  45. Nayak, S.S.; Baltazar Hernandez, V.H.; Zhou, Y. Effect of Chemistry on Nonisothermal Tempering and Softening of Dual-Phase Steels. Metall. Mater. Trans. A 2011, 42, 3242–3248. [Google Scholar] [CrossRef]
  46. Baltazar Hernandez, V.H.; Nayak, S.S.; Zhou, Y. Tempering of Martensite in Dual-Phase Steels and Its Effects on Softening Behavior. Metall. Mater. Trans. A 2011, 42, 3115–3129. [Google Scholar] [CrossRef]
  47. Khan, M.S.; Soleimani, M.; Midawi, A.R.H.; Aderibigbe, I.; Zhou, Y.N.; Biro, E. A review on heat affected zone softening of dual-phase steels during laser welding. J. Manuf. Process. 2023, 102, 663–684. [Google Scholar] [CrossRef]
  48. Elmesalamy, A.; Francis, J.A.; Li, L. A comparison of residual stresses in multi pass narrow gap laser welds and gas-tungsten arc welds in AISI 316L stainless steel. Int. J. Press. Vessel. Pip. 2014, 113, 49–59. [Google Scholar] [CrossRef]
  49. Bate, S.; Green, D.; Buttle, D. A Review of Residual Stress Distributions in Welded Joints for the Defect Assessment of Offshore Structures; HSE Books: London, UK, 1998. [Google Scholar]
  50. Hu, X.; Jiang, H.-Y.; Luo, Y.; Jin, Q.; Peng, W.; Yi, C.-M. A Study on Microstructure, Residual Stresses and Stress Corrosion Cracking of Repair Welding on 304 Stainless Steel: Part II-Effects of Reinforcement Height. Materials 2020, 13, 2434. [Google Scholar] [CrossRef]
  51. Balakrishnan, J.; Vasileiou, A.N.; Francis, J.A.; Smith, M.C.; Roy, M.J.; Callaghan, M.D.; Irvine, N.M. Residual stress distributions in arc, laser and electron-beam welds in 30 mm thick SA508 steel: A cross-process comparison. Int. J. Press. Vessel. Pip. 2018, 162, 59–70. [Google Scholar] [CrossRef]
  52. Francis, J.A.; Bhadeshia, H.K.D.H.; Withers, P.J. Welding residual stresses in ferritic power plant steels. Mater. Sci. Technol. 2007, 23, 1009–1020. [Google Scholar] [CrossRef]
  53. Bhadeshia, H.K.D.H. TRIP-Assisted Steels? ISIJ Int. 2002, 42, 1059–1060. [Google Scholar] [CrossRef]
  54. Jones, W.K.C.; Alberry, P.J. Paper 78 The role of phase transformations in the development of residual stresses during the welding of some fast reactor steels. Ferritic Steels Fast React. Steam Gener. 1978, 2, 471–475. [Google Scholar] [CrossRef]
  55. Nichols, R.W. Residual Stresses in Welded Construction and Their Effects, Proceedings of the An International Conference, London, UK, 15–17 November 1977; Welding Institute: Cambridge, UK, 1978. [Google Scholar]
  56. Bartlett, J.L.; Li, X. An overview of residual stresses in metal powder bed fusion. Addit. Manuf. 2019, 27, 131–149. [Google Scholar] [CrossRef]
  57. Wu, X.I.N.; Wang, Z.; Yu, Z.; Liu, S.; Bunn, J.; Kolbus, L.; Feng, Z. Control of Weld Residual Stress in a Thin Steel Plate through Low Transformation Temperature Welding Consumables. Weld. J. 2020, 99, 124–134. [Google Scholar] [CrossRef]
  58. Liu, C.; Yan, Y.; Cheng, X.; Wang, C.; Zhao, Y. Residual stress in a restrained specimen processed by post-weld ultrasonic impact treatment. Sci. Technol. Weld. Join. 2019, 24, 193–199. [Google Scholar] [CrossRef]
Metals 15 00791 g001
Figure 1. Experimental fixture for GMAW.
Figure 1. Experimental fixture for GMAW.
Metals 15 00791 g001
Metals 15 00791 g002
Figure 2. Transverse cross-section macrostructure of the GMAW samples showing the various zones along the weldment for: (a) PHS-PHS, (b) DP-DP, and (c) PHS-DP joints.
Figure 2. Transverse cross-section macrostructure of the GMAW samples showing the various zones along the weldment for: (a) PHS-PHS, (b) DP-DP, and (c) PHS-DP joints.
Metals 15 00791 g002
Metals 15 00791 g003
Figure 3. Microstructure of the fusion zone (FZ) indicating the presence of pearlite (P) and ferrite (α) for: (a,b) PHS-PHS, (c,d) DP-DP, and (e,f) PHS-DP steels weldments. The blue boxes within the micrographs (a,c,e) correspond the location of the obtained microstructures shown in (b,d,f), respectively.
Figure 3. Microstructure of the fusion zone (FZ) indicating the presence of pearlite (P) and ferrite (α) for: (a,b) PHS-PHS, (c,d) DP-DP, and (e,f) PHS-DP steels weldments. The blue boxes within the micrographs (a,c,e) correspond the location of the obtained microstructures shown in (b,d,f), respectively.
Metals 15 00791 g003
Metals 15 00791 g004
Figure 4. The microstructure development along the HAZ of the PHS and DP steels side is comprised of martensite (α′), bainite (B), pearlite (P), and ferrite (α), located at: (a,b) UC-HAZ coarse grain, (c,d) UC-HAZ fine grain, and (e,f) IC-HAZ.
Figure 4. The microstructure development along the HAZ of the PHS and DP steels side is comprised of martensite (α′), bainite (B), pearlite (P), and ferrite (α), located at: (a,b) UC-HAZ coarse grain, (c,d) UC-HAZ fine grain, and (e,f) IC-HAZ.
Metals 15 00791 g004
Metals 15 00791 g005
Figure 5. High magnification base metal microstructure consisted of martensite (α′), bainite (B), retained austenite (RA), and ferrite (α) phases for (a) PHS Steel and (b) DP Steel. Microstructure development at the sub-critical heat-affected zone (SC-HAZ) for (c) PHS steel and (d) DP Steel, showing tempered martensite (TM) regions composed of precipitated cementite carbides (Fe3C) and ferrite (α).
Figure 5. High magnification base metal microstructure consisted of martensite (α′), bainite (B), retained austenite (RA), and ferrite (α) phases for (a) PHS Steel and (b) DP Steel. Microstructure development at the sub-critical heat-affected zone (SC-HAZ) for (c) PHS steel and (d) DP Steel, showing tempered martensite (TM) regions composed of precipitated cementite carbides (Fe3C) and ferrite (α).
Metals 15 00791 g005
Metals 15 00791 g006
Figure 6. Microhardness Vickers profile along the weldment for (a) PHS-PHS, (b) DP-DP, and (c) PHS-DP steel combinations.
Figure 6. Microhardness Vickers profile along the weldment for (a) PHS-PHS, (b) DP-DP, and (c) PHS-DP steel combinations.
Metals 15 00791 g006
Metals 15 00791 g007
Figure 7. (a) Engineering stress-strain curves for base metal and welded specimens and (b) the failure location of the tensile specimens (as indicated by the arrow).
Figure 7. (a) Engineering stress-strain curves for base metal and welded specimens and (b) the failure location of the tensile specimens (as indicated by the arrow).
Metals 15 00791 g007
Metals 15 00791 g008
Figure 8. (a) Transverse cross-section macrostructure showing the failure location at the SC-HAZ of the welded specimen, (b) microstructure revealing the crack nucleation located at the interface of ferrite and the tempered martensite (as indicated by the arrow), and (c) fractography image showing the presence of equiaxed and elongated dimples of varied size unfolding the fracture mode.
Figure 8. (a) Transverse cross-section macrostructure showing the failure location at the SC-HAZ of the welded specimen, (b) microstructure revealing the crack nucleation located at the interface of ferrite and the tempered martensite (as indicated by the arrow), and (c) fractography image showing the presence of equiaxed and elongated dimples of varied size unfolding the fracture mode.
Metals 15 00791 g008
Metals 15 00791 g009
Figure 9. Experimentally measured (by x-ray diffraction) residual stress profile along the weldment for (a) PHS-PHS, (b) DP-DP, and (c) PHS-DP steel combinations.
Figure 9. Experimentally measured (by x-ray diffraction) residual stress profile along the weldment for (a) PHS-PHS, (b) DP-DP, and (c) PHS-DP steel combinations.
Metals 15 00791 g009
Table 1. Chemical composition (OES) of PHS and DP steel (wt.%).
Table 1. Chemical composition (OES) of PHS and DP steel (wt.%).
MaterialCSiMnPCrMoNiCuAlBCEy
DP0.150.3141.410.0140.030.0140.010.030.05-0.41
PHS0.2570.2060.9390.0360.0220.0370.1600.0440.0420.00100.44
ER70S-6 *0.06–0.150.8–1.151.40–1.850.0250.150.150.150.5---
* Chemical composition as per the American Welding Society (AWS) [42].
Table 2. Selected welding parameters for GMAW.
Table 2. Selected welding parameters for GMAW.
Current
(A)		Voltage
(V)		Shield GasGas Flow
(L/min)		Travel Speed
(m/min)
7516.575% Ar + 25% CO2180.6
Table 3. Uniaxial tensile properties for BM and welded specimens.
Table 3. Uniaxial tensile properties for BM and welded specimens.
SampleYS
(MPa)		UTS
(MPa)		Uniform Elongation
(%)
BM
PHS902 ± 281520 ± 209.1 ± 0.7
DP445 ± 39986 ± 1426 ± 2.0
Steel Weld Combination
PHS-PHS467 ± 37824 ± 197.9 ± 1.1
DP-DP413 ± 45703 ± 1310.8 ± 0.3
PHS-DP453 ± 20699 ± 167.9 ± 1.5
Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

Share and Cite

MDPI and ACS Style

Molina-Castillo, A.E.; López-Baltazar, E.A.; Alvarado-Hernández, F.; Gómez-Jiménez, S.; Espinosa-Lumbreras, J.R.; Ruiz Mondragón, J.J.; Baltazar-Hernández, V.H. Effect of the Heat Affected Zone Hardness Reduction on the Tensile Properties of GMAW Press Hardening Automotive Steel. Metals 2025, 15, 791. https://doi.org/10.3390/met15070791

AMA Style

Molina-Castillo AE, López-Baltazar EA, Alvarado-Hernández F, Gómez-Jiménez S, Espinosa-Lumbreras JR, Ruiz Mondragón JJ, Baltazar-Hernández VH. Effect of the Heat Affected Zone Hardness Reduction on the Tensile Properties of GMAW Press Hardening Automotive Steel. Metals. 2025; 15(7):791. https://doi.org/10.3390/met15070791

Chicago/Turabian Style

Molina-Castillo, Alfredo E., Enrique A. López-Baltazar, Francisco Alvarado-Hernández, Salvador Gómez-Jiménez, J. Roberto Espinosa-Lumbreras, José Jorge Ruiz Mondragón, and Víctor H. Baltazar-Hernández. 2025. "Effect of the Heat Affected Zone Hardness Reduction on the Tensile Properties of GMAW Press Hardening Automotive Steel" Metals 15, no. 7: 791. https://doi.org/10.3390/met15070791

APA Style

Molina-Castillo, A. E., López-Baltazar, E. A., Alvarado-Hernández, F., Gómez-Jiménez, S., Espinosa-Lumbreras, J. R., Ruiz Mondragón, J. J., & Baltazar-Hernández, V. H. (2025). Effect of the Heat Affected Zone Hardness Reduction on the Tensile Properties of GMAW Press Hardening Automotive Steel. Metals, 15(7), 791. https://doi.org/10.3390/met15070791

Note that from the first issue of 2016, this journal uses article numbers instead of page numbers. See further details here.

Article Metrics

Back to TopTop