Microstructure and Tensile Behaviors of Laser-Arc Hybrid Welds Comparing to Submerged Arc Welds of High-Manganese Steel for Cryogenic Applications
Article information
Abstract
The weldability and the relationship between microstructure and tensile properties at 298 and 110 K produced by laser-arc hybrid welding (LAHW) in high-Mn steel welds were thoroughly investigated. In the laser zone of LAHWs using filler wire, more Mn vaporization in the laser zone was observed than at the arc zone of LAHWs. The arc zone showed a decrease in Mn content of ~0.6 wt%, while the laser zone showed a decrease of ~0.9 wt%. The arc and laser zones of the LAHWs showed stacking fault energies (SFEs) of 17.8 and 17.3 mJ/m², respectively. The tensile deformation of LAHWs at 298 K was conducted with a deformation twins mode, while it was shifted to deformation twins + ε-martensite transformation at 110 K. The yield strength was slightly higher in the laser zone, which had a finer grain size compared to the arc zone. The formation of ε-martensite with deformation twins preceded necking during tensile testing, therefore increasing the yield strength at 110 K. In terms of performance, the LAHW process demonstrated a 25% increase in productivity compared to the conventional submerged arc welding (SAW) process, with a yield strength exceeding 400 MPa, comparable to that of SAW. These findings indicate that LAHW is a highly effective welding method for high-Mn steels, particularly in cryogenic applications.
1. INTRODUCTION
The excavation of shale gas in polar regions and the deep sea, as well as the demand for hydrogen energy, is increasing. Cryogenic materials are essential for related industries, and both interest and demand for these materials are rising[1-3]. Cryogenic materials are commonly used with 9% Ni and austenitic stainless steels[4-7]. However, the base and welding materials containing high amounts of Ni, while offering excellent properties, are costly. Recently, high-Mn steels have received significant attention due to their outstanding mechanical properties at cryogenic temperatures and high price competitiveness[8-11]. High-Mn steels exhibit an excellent balance of strength and toughness due to the twinning-induced plasticity (TWIP) effect during tensile deformation at room temperature[12-15]. The deformation mechanism of high-Mn steel varies depending on the stacking fault energy (SFE), which is influenced by alloying composition, temperature, and grain size[16-20].
Welding accounts for 30-40% of fabrication time and 30-50% of shipbuilding costs[21]. In particular, improving welding productivity is crucial for welding steel plates having a thickness above 20 mm. Multi-pass arc welding is normally used in the production field. However, due to significant welding deformation and low welding efficiency, narrow gap arc welding, electron beam welding, and laser welding have been proposed for welding thick plates[22-26]. Narrow-gap arc welding provides slightly increased efficiency by reducing the number of welding passes. However, welding efficiency and quality remain unsatisfactory due to shallow penetration and arc instability[25]. Electron beam welding offers high energy density and low heat input, resulting in deep penetration and minimal deformation. However, its applicability is limited due to the requirement of a high vacuum environment[23]. Laser beam welding is attracting attention as a technology for welding thick plates due to its low heat input, deep penetration, and high energy efficiency in air[24]. During laser beam welding, keyholes are formed, causing the vaporization of volatile alloying elements such as Mn and Mg[27-29]. This vaporization may lead to deterioration of the weld zone properties[30-32]. Laser-arc hybrid welding (LAHW) technology combines laser and arc welding, with the two heat sources interacting to create a single high-intensity energy source. LAHW allows for increased gap tolerances compared to laser welding while maintaining the deep penetration and high welding speeds necessary for efficient welding[33-36].
Weldability is essential in the industrial application of high-Mn steels. Numerous studies focused on evaluating weldability have been reported including friction stir welding, arc welding, and laser beam welding on the high-Mn steel[37-41]. The authors previously reported the weldability of 24% Mn steel using submerged arc welding (SAW) for cryogenic applications[39]. Weld metals with a low SFE exhibited excellent tensile properties at both 298 and 110 K, which is attributed to deformation twinning and occasional ε-martensite transformation. Xie et al.[41] reported that strength and plasticity were enhanced by controlling the nugget zone microstructure through preheating in the friction stir welding of high manganese steel.
This study investigates the microstructural and tensile properties combined with Mn vaporization of high-Mn steel subjected to LAHW. Specifically, we focus on the tensile properties of the welds at 110 and 298 K, and their relationship with microstructural behaviors as compared to the authors’ previous studies on the conventional SAW process[39].
2. EXPERIMENTAL PROCEDURE
Table 1 shows the composition of high-Mn steel and filler wires. The high-Mn steels were produced using a blast furnace. The cast ingot underwent solution treatment, followed by hot-rolling to a thickness of 30 mm and controlled cooling. To meet the cryogenic toughness requirements of the weld metal in high-Mn steel, undermatched filler wires with low Mn content compared to the base plate was used.
Chemical composition (wt%) of the high-Mn steel and filler wires of LAHW and SAW
Clean-up with acetone was conducted on the base plate before welding. Figure 1a shows a schematic diagram of the LAHW. The test pieces of the LAHW were prepared with dimensions of 200 × 600 mm and assembled with a zero-gap square groove using a hydraulic jig (Fig 1a).
Schematic diagram of (a) LAHW and (b) groove morphology of SAW
A total of two passes were applied for welding: the first pass on the top and the second pass on the bottom. The LAHW system consists of a single-mode 20 kW fiber laser (IPG YLS-20,000-S2T) and a 400 A arc welding power source (Fronius TransPuls Synergic-4000). The laser beam had a diameter of 20 μm, wavelength of 1080 nm, and a beam parameter product of 8 mm·mrad. The laser focusing lens had a focal length of 300 mm. The laser defocusing position was set to -8 mm above the surface, and the arc-to-laser distance was 6 mm. Figure 1b shows a double-V groove morphology of the SAW, which was conducted with one pass inside and two passes outside, for a total of three passes. The LAHW and SAW conditions of the study are summarized in Table 2.
Conditions of LAHW and SAW
The specimens were polished and etched in a solution of ethanol (100 mL), hydrochloric acid (5 mL), and picric acid (2 g) and the resulting microstructure was observed by light optical microscopy (LOM) and scanning electron microscopy (SEM). The crystal structures were identified by X-ray diffraction (XRD) using a Cu-Kα target and electron back-scattered diffraction (EBSD). XRD was performed at a scan speed of 1°/min from 20 to 100°, 40 kV, and 40 mA. The diffraction data for the EBSD experiment were obtained as orientation maps with step sizes of 0.1-0.2 µm at 20 kV. Mapping and a quantitative analysis using electron probe microanalysis (EPMA) were applied to measure the chemical composition of the weld. The EPMA was conducted with a probe current of 100 nA and an accelerating voltage of 20 kV. The tensile specimens were prepared in the transverse direction of the welded specimens and tensile testing was performed with standard-size specimens (ASTM E8). The extensometer was set on an Instron testing machine in a low temperature chamber and a loading speed of 1 mm/min was applied.
3. RESULTS AND DISCUSSION
3.1 Microstructure of the LAHW and SAW specimens
Figure 2a shows an inverse pole figure (IPF) map of the base metal measured by EBSD.
Microstructure and phase identification: (a) IPF image of the base metal, (b) XRD patterns of the base and weld metals
The grain size of the base metal is primarily distributed in the range of 5-60 µm, with an average size of 15±17 µm. Figure 2b shows the XRD patterns of the base and weld metals. Both the base and weld metals consist of a single austenite phase, with no other phases present.
Figure 3 shows the cross-sectional microstructure of LAHW specimens produced without defects such as microcracks or macropores.
Macrostructure of (a) LAHW specimens, microstructure of (b) arc zone weld metal at the centerline, (c) CGHAZ near arc weld, (d) laser zone weld metal, (e) CGHAZ near laser weld, and EBSD-IPF maps of (f) arc zone weld metal and (g) laser zone weld metal
The weld metals in the arc and laser zones observed in the top pass exhibit columnar grains formed along the weld centerline (Figs 3b and 3d, respectively). Figures 3c and 3e show the fusion boundary (black-dotted lines) and coarse grain heat-affected zone (CGHAZ) of the arc and laser zones. Grain growth occurred in the CGHAZ due to the welding heat input, with a wider CGHAZ area in the arc zone (Fig 3c), where the heat input is higher than the laser zone. Figure 3f shows the IPF map of the arc zone weld metal, where the average columnar grain width is 82±72 µm. Figure 3g shows the IPF map of the laser zone weld metal, with an average columnar grain width of 27±26 µm. The columnar grain width in the arc zone is larger than that in the laser zone due to the large heat input of the arc zone.
Figure 4 shows the microstructure of SAW specimen produced without defects.
Macrostructure of (a) SAW specimens, microstructure of (b) CGHAZ near the 1st pass weld, (c) CGHAZ near the 1st and 2nd pass boundaries, (d) EBSD-IPF maps of 2nd pass weld metal
Figures 4b shows the 1st pass CGHAZ. Figure 4c shows the CGHAZ near the crossing boundary between the 1st and 2nd passes, where significant grain growth occurred due to the overlapping heat input. Figure 4d shows an IPF map of the 2nd-pass weld metal. The columnar grain width was primarily distributed in the range of 40-140 µm, with an average of ~ 86 µm, which was mostly same as that of the arc zone weld metal (Fig 3f).
3.2 Mechanical properties of the LAHW and SAW specimens
Figure 5 shows the hardness distribution of the LAHW and SAW specimens.
Distribution of Vickers hardness for (a) arc and laser zones of LAHWs and (b) each pass of SAWs
The hardness in the arc zones of top and bottom passes was measured at a depth of 1 mm from each weld bead surface, while the laser zone was measured at half the thickness (1/2 t). The base metal exhibited an average hardness of ~250 Hv. The hardness decreased in the CGHAZ, with the weld metal showing ~193 Hv in both the arc and laser zones (Fig 5a). The hardness of weld metals is lower than that of the base metal due to the use of undermatched filler wire and the coarse grain size caused by the welding heat input. The hardness of the 2nd pass of SAWs was approximately 23 Hv higher than that of the 1st and 3rd passes due to the significant dilution of the base metal in the 2nd pass root face (Fig 5b). The hardness distribution for LAHW and SAW showed the same tendency of hardness in the base and weld metals.
Figure 6 shows the stress-strain curves of the arc and laser zones of LAHWs and SAWs at 298 and 110 K.
Engineering tensile stress–strain curves of the LAHWs and SAWs specimens tested at room (298 K) and cryogenic (110 K) temperatures
The arc zone was sampled from the top pass, and the laser zone at 1/2 t. Yield strength (YS), tensile strength (TS), and elongation (EL) were measured from the curves, and the results of the LAHW specimens and conventional SAWs are summarized in Table 3. The arc and laser zones of the LAHWs exhibited a YS of 431 and 449 MPa, a TS of 798 and 793 MPa, and an EL of 39 and 24%, respectively, at 298 K. Compared to the arc zone, the laser zone indicated a slightly increased YS and decreased EL. The arc and laser zones exhibited a YS of 542 and 603 MPa, a TS of 917 and 1050 MPa, and an EL of 20 and 16%, respectively, at 110 K. The laser zone exhibited higher YS and TS than the arc zone at 110 K. At cryogenic temperature, the YS and TS of all specimens increased and the EL decreased as compared to room temperature.
Tensile properties summarized for arc zone, laser zone of LAHWs and conventional SAWs at various testing temperatures (298 and 110 K)
Figure 7 shows SEM fractographs near the weld centreline after tensile fracture.
Fractography of (a, b) arc zone and (c, d) laser zone at 298 and 110 K, respectively
All specimens fractured near the weld centerline, regardless of the test temperature and arc/laser zone. Figure 7a and 7c show fractographs of the arc and laser zones, respectively, tested at 298 K. At room temperature, all specimens exhibit ductile dimple fracture. At cryogenic temperatures (110 K), all specimens primarily exhibit dimple fractures with some quasi-cleavage (QC) fractures (Figs 7b and 7d). The SAWs also indicate the same fracture location in the transverse weld joint and fractographs at 298 and 110 K, as reported in the authors’ previous study[39].
3.3 Microstructural evolution and tensile properties of LAHW specimens as a function of temperature
Figure 8 shows the EPMA mapping and compositional line profile of the top-pass arc zone for LAHWs. The EPMA analysis was focused on the main alloying elements such as Mn and Cr. Insignificant variation of color contrast between the weld and base metals was observed in the EPMA mapping, although some micro-segregation due to hot rolling was present in the base metal (Fig 8a).
EPMA analysis of arc zone for LAHWs: (a) mapping for Fe, Mn, and Cr and (b) quantitative analysis measured along the white-dotted line
A quantitative analysis of the composition was conducted along the white-dotted line in the BSE image (Fig 8b), and a slight decrease (~0.6 wt%) in Mn content was observed in the weld zone due to the undermatched filler wire applied and base-metal dilution in the study.
Figure 9 shows the EPMA mapping and compositional line profile of the laser zone for LAHWs.
EPMA analysis of laser zone for LAHWs: (a) mapping for Fe, Mn, and Cr and (b) quantitative analysis measured along the white-dotted line
The laser-zone welds exhibited a slightly more yellow color in Mn composition as compared to the base metal (Fig 9a). A slight decrease in the Mn composition was quantitatively observed in the weld zone (Fig 9b) and the laser zone indicated a ~0.9 wt% lower Mn content than that of the base metal (Table 4). The weld zone was more significantly affected by Mn vaporization due to laser keyhole formation than the arc zone (Fig 8b). The laser zone exhibited slightly more Mn loss than the arc zone and the laser zone was mainly affected by the laser beam with ineffective influence of the undermatched filler wire. It was also confirmed from the authors’ previous study that the full keyhole mode provides insignificant Mn vaporization as compared to the partial-keyhole mode[42]. Therefore, the laser zone of the LAHW specimen showed the most significant Mn loss due to the partial-keyhole mode in this study. Furthermore, the weld zone of the SAW had minimal variation of Mn content (23.7 wt%) as compared to the base metal due to base metal dilution, even though undermatched filler wire with the lowest Mn content (20 wt%) was used[39]. Therefore, the laser zone of the LAHWs exhibited slightly more significant loss of Mn content than the arc zone of LAHWs and SAWs.
The average composition (wt%) measured by EPMA quantitative analysis in the base metal, SAWs, arc zone and laser zone of LAHWs
Figure 10 shows the cross-sectional EBSD analysis results of the arc zone for the LAHW specimen observed near the center of tensile fracture at 298 and 110 K. EBSD mapping includes IPF, coincidence site lattice (CSL) boundaries, and phase maps. The CSL map shows a Σ3 boundary (deformation twin boundary) indicated by a red line. After tensile fracture at 298 K, only deformation twins occurred in the arc zone, without any phase transformation (Fig 10a).
EBSD-IPF, CSL, and phase maps observed in the centerline of the cross-sectional arc zone welds for LAHWs post to the tensile testing at various temperatures: (a) 298 and (b) 110 K
As the temperature was decreased to 110 K, the deformation mode shifted to deformation twins with some γ →ε martensite transformation (7%), as indicated in Fig 10b.
Figure 11 presents the results of the cross-sectional EBSD analysis of the laser zone for the LAHW specimen observed near the 1/2 t after tensile fracture at 298 and 110 K.
EBSD-IPF, CSL, and phase maps observed in the centerline of the cross-sectional laser zone welds for LAHWs post to the tensile testing at various temperatures: (a) 298 and (b) 110 K
After tensile fracture at 298 K, only deformation twins occurred in the laser zone, without any phase transformation (Fig 11a). As the temperature was decreased to 110 K, the deformation modes included deformation twins and some γ→ε martensite transformation (6%), as shown in Fig 11b. In both the arc and laser zones of the LAHW specimens, the deformation mode showed deformation twins at 298 K and deformation twins + ε-martensite at 110 K. Some εmartensite was formed in the arc and laser zones during tensile deformation at 110 K, leading to some quasi-cleavage fracture (Figs 7b and 7d)[43].
The SAW specimen showed deformation twins after fracture at 298 K, while deformation twins and γ→ε martensite transformation occurred at 110 K[39]. Both SAWs and LAHWs exhibited the same deformation modes at 298 and 110 K.
The deformation mode of high-Mn steel during tensile deformation varies depending on the SFE, leading to deformation twins or ε- and α'-martensite transformation. In Fe-Mn-C-based steels, the SFE is greatly influenced by Mn content and temperature[16,17]. To evaluate these effects, the SFE was calculated using the Olson-Cohen thermodynamic model, expressed by the following equation[44]:
SFE = 2ρ△Gγ→ε 2σγ/ε
where △Gγ→ε is the free energy of the phase transformation (γ → ε-martensite), ρ is the molar concentration of atoms in the {111} planes, and σγ/ε is the interfacial energy per unit area of the interphase interface. △Gγ→ε was calculated using the model for Fe-Mn-C based steel presented by Dumay et al.[45].
The SFEs were calculated using the composition values observed in the study (Table 4). At 298 K, the SFEs of the SAWs, arc zone, laser zone, and base metal were 17.2, 17.8, 17.3, and 18.7 mJ/m², respectively. The SFE in the laser zone, where Mn vaporization occurred, was the lowest in the LAHWs despite their minimal variation. In the SAWs, Mn vaporization occurred insignificantly, but the SFE was nearly the same as that of the laser zone due to the use of undermatched filler wire. The lower limit where deformation twins occur is known to be ~20 mJ/m²[46]. However, De Cooman et al.[47] reported a limit higher than 13 mJ/m² in Fe-18Mn-0.6C steel. According to Remy et al.[48], this range was 10–15 mJ/m² in a Fe-Mn-Cr-C system.
The deformation mode for the SAWs, arc, and laser zones of LAHW specimens at 298 K is deformation twinning, and the SFE range calculated in this study is reasonably consistent with the range proposed by Remy et al[48]. Since the SFE decreases by 20–60% as the testing temperature decreases to cryogenic levels (−196 °C)[49], the deformation mode of the arc and laser zones at 110 K shifts to deformation twins + ε-martensite. The arc and laser zones of the LAHW specimens have almost the same SFE and deformation mode, but the YS was slightly higher in the laser zone. This was due to the strengthening effect from its smaller columnar grain width (27±26 µm) than that in the arc zone (82±72 µm), including insignificant SFE variation caused by Mn vaporization of the LAHWs (17.3−17.8 mJ/m²) and SAWs (17.2 mJ/m²). Ultimately, the LAHWs exhibited the slightly higher Mn vaporization and lower level of strength as compared to the conventional SAWs. However, the LAHWs achieved a YS exceeding 400 MPa and 25 % higher productivity than the SAW.
4. CONCLUSIONS
This study investigated the tensile properties and microstructural evolution of LAHW specimens for high-Mn steels at room- and cryogenic-temperatures. The key conclusions are summarized as follows:
1) LAHWs showed no defects, such as microcracks or voids, in the as-welded state and exhibited an austenite single phase.
2) The arc and laser zones of the LAHW specimens showed decreased hardness due to the undermatched filler wire and coarse grain structure as compared to the base metal. Therefore, all specimens fractured near the weld centerline during the tensile test regardless of the testing temperature.
3) The SFE calculated at 298 K indicated that both the arc and laser zones had very similar SFE values of 17.8 and 17.3 mJ/m², respectively, and the tensile deformation exhibited deformation twinning at 298 K. As the temperature was decreased to 110 K, the deformation mode shifted to deformation twins + ε-martensite.
4) The LAHWs exhibits nearly the same level of Mn vaporization and strengths as compared to the conventional SAW. A YS exceeding 400 MPa and 25 % higher productivity for the LAHWs were achieved as compared to the SAWs.
Acknowledgements
This work was supported by the Technology Innovation Program (Grant No. 20022454) funded by the Ministry of Trade, Industry, and Energy (MOTIE, Korea) and the korea Institute of Machinery and Materials (NK250C).