Laser Welding of UNS S33207 Hyper-Duplex Stainless Steel to 6061 Aluminum Alloy Using High Entropy Alloy as a Filler Material

: The high entropy alloy (HEA) ﬁller used during the fabrication method determines the reliability of HEAs for steel-aluminum dissimilar alloy conﬁguration. HEAs have a direct impact on the formation of intermetallic compounds (IMC) formed by the interaction of iron (Fe) and aluminum (Al), and inﬂuence the size of the joint’s interaction zone. A novel welding process for Fe-Al alloy joints was developed to prevent the development of a brittle iron-aluminum interface. This research involved investigation of the possibility of using HEA powdered ﬁller. Fe 5 Co 20 Ni 20 Mn 35 Cu 20 HEAs was used as a ﬁller for the laser joining lap conﬁguration joining hyper-duplex stainless steel UNS S33207 to aluminum alloy 6061. This HEA has unique properties, such as high strength, good ductility, and high resistance to corrosion and wear. A tiny portion of the stainless-steel area was melted by varying the welding parameters. The high-entropy alloy (HEA) with slow kinetic diffusion and large entropy was employed to aid in producing solid solution structures, impeding the blending of iron and aluminum particles and hindering the development of Fe-Al IMCs. The weld seam was created without the use of Fe-Al IMCs,. The specimen broke at the HEAs/Al alloy interface with a tensile-shear strength of 237 MPa. The tensile-shear strength achieved was 12.86% higher than for the base metal AA 6061 and 75.57% lower than for the UNS S33207 hyper-duplex stainless steel.


Introduction
A hybrid joint made of dissimilar metals has been a focus of research in battery case manufacturing, and in the automotive and marine sectors for a long time [1,2]. The synergistic qualities of the metals, as well as the creation of design efficient structures, are the primary driving forces behind this [3]. However, joining dissimilar metal combinations can be challenging because of inferior chemical compatibility and dissimilarities in physical properties. For example, when welding aluminum to steel, the interface between aluminum and (Al) iron (Fe) atoms throughout the welding process forms intermetallic compounds (IMCs) with high brittleness [4]. As the presence of these IMCs limits the strength of the joint, the amount of these compounds should be kept to a minimum.
Furthermore, all fusion welding methods occur in transient conditions, and the time and temperature cannot be separately controlled. Moreover, the heat cycle applied to the weldment determines the thickness of the IMC layer and the width of the weld [5,6]. Therefore, to regulate these two components, it is first necessary to understand the heat cycle and its effects.
The advantages of iron (Fe) and aluminum alloys (Al) include excellent mechanical strength and corrosion resistance. As a result, they have a wide range of uses in maritime, automotive, and space applications [7]. Consequently, Fe-Al composite structures have received increasing interest as a means to amalgamate the positive characteristics of Fe with Al metal alloys. However, due to significant variations in material parameters, such as thermal properties, melting point, material characteristics, and grain morphology, effective Fe/Al welding conditions are hard to formulate and can generate Fe-Al phases with high brittleness [8,9].
The main benefits of laser welding compared to arc welding are: (1) by adjusting the laser metal interaction, the dimension and melt pool geometry can be controlled, which would be difficult to achieve with an arc-welding method; (2) because of the high energy density in laser welding, a nominal heat-affected zone is created; (3) higher cooling and solidification levels compared to other fusion welding methods; and (4) localized and focused heating [10][11][12][13][14]. Furthermore, with gas tungsten arc welding (GTAW), managing the temperature ranges is not feasible as the temperature required to melt the aluminum will only wet the steel [15][16][17]. Traditional non-consumable techniques, such as GTAW or plasma, would not provide such control.
Lasers are used for the cladding layers in many industries, e.g., the metal industry [18][19][20] and medicine [21,22]. The process is also used in water environments [23,24]. In addition, lasers are widely used in the fusion welding process to fabricate joints between duplex steel and aluminum alloys. Most studies have focused on minimizing Fe-Al IMC formation by regulating the beam offset position and distance to control the volume of the melted base material [25]. Other investigations have found that dual-beam welding improves the uniformity of IMC thickness, and that silicon helps decrease fractures in the weld. Nevertheless, none of the strategies listed above can prevent IMC development in Fe/Al interfaces. Furthermore, in dissimilar metal welding, a portion of the non-melted transitional layer remaining in the weldments butting surface will entirely avoid blending with parent materials [26].
The approach is non-ideal for lap-welding configuration with lasers [27]. As a result, novel welding processes must be investigated in order to efficiently regulate the development and growth of Fe-Al IMCs in the weld zone of the lap interface. Based on their high entropy impact, HEAs can enhance the development of solid solutions [28]. The mechanical characteristics of high-entropy alloys (HEAs) are directly connected to their structural stability. Therefore, proper monitoring of component composition is essential in understanding structural stability and determining mechanical characteristics. Co-Cr-Fe-Ni HEAs have been built using a single face-centered cubic (FCC) solid solution. Elemental components have a noticeable influence on their structure and performance; the Cr and Fe elements, in particular, have an apparent effect on structural stability and the equilibrium lattice constant. In addition, the Ni components have a noticeable impact on rigidity. According to the Pugh ratios, Cr and Ni addition may enhance ductility, but Co and Fe addition may lower it. Therefore, the structural stability and stiffness of Co-Cr-Fe-Ni HEAs are enhanced by increasing Co and Fe concentrations or reducing Cr and Ni contents. The structural stability and mechanical characteristics of Co-Cr-Fe-Ni HEAs may be attributed to the strength of the metallic bonding and covalent bonding, which is influenced by the element content variation [29].
In previous studies, the mechanical properties of the weldments have been related only to the thickness of the IMC layer, and the weld interface area among the stainless steel and aluminum has not been considered. Mechanical strength should evolve linearly with the bonding area. However, the continuous IMC layer generated at the weld surface prevents the linear trend from being observed. A thickness of the IMC layer up to 8 µm would result in sound Fe-Al joints [30][31][32]. Although utilizing shallow energy in the welding process to restrict IMC layer formation has a favorable impact, this method also results in the formation of a narrow bonding area with lower weld strength [33,34]. In dissimilar metal welding, the sluggish kinetic diffusion property is exploited to prevent the blending of adjacent layers. For example, using powdered HEAs as a filler metal leads to the effective joining of aluminum and austenitic stainless steel in butt configuration [35]. Nevertheless, using HEAs as filler metals in the laser welding of dissimilar metals, such as aluminum and steel, has been little explored.
Since none of the standard filler metal compositions meet the stricter requirements of HEAs, further investigations into the competence of such alloys to perform as fillers are worthwhile, and focus has shifted in this direction. In the following, we use the Pickering and Jones convention [36,37] and label HEAs to reflect the atomic number of the elements. However, there are non-equiatomic configurations with more predominant elements in the composition over others; this results in the renaming of selected alloys. The elements are rearranged from how they were provided in their initial form.
For the welding of dissimilar alloys, such as hyper-duplex stainless steel and aluminum alloy, vanadium and niobium are frequently utilized as corresponding layers that have a substantial repressive influence upon the generation of Fe-Al IMCs. The welding stress can lessen the exceptional ductility of copper, and the coefficient of diffusion of Ni and Fe is much greater than that of Al and Fe. Supplementing Ni can interfere with the formation of IMCs formed by the combination of Fe and Al. According to HEA theory, the atomic radius of Cu, Ni, Mn, and Co are identical. Therefore, they can produce a solid solution with ease. This research presents a laser welding technique for hyper-duplex stainless steel (UNS S33207) and aluminum alloy (AA 6061) using Fe 5 Co 20 Ni 20 Mn 35 Cu 20 HEAs. Fe-Al IMCs are predicted to be repressed using HEAs fillers in Fe/Al lap joint configuration and to limit the melting of parent metal on the hyper duplex stainless-steel side. To date, this type of laser welding has been rarely performed.

Materials
Plates of UNS S33207 hyper-duplex stainless steel and AA 6061 aluminum alloy with 2 mm thickness and 100 mm length and 70 mm width were used; the length and width were chosen according to the ASTM D8172-18 standard for shear testing resistance seam welds [38]. The materials' mechanical properties and chemical composition were as per the materials certificate provided by the supplier, as detailed in Tables 1 and 2. First, the oxide layer and other surface impurities on the plates were removed by grounding the hyper duplex stainless steel and linishing the aluminum alloy. Afterwards, all plates were degreased with acetone solution. Materials with a purity of 99.9% were used. The raw materials used were Fe, Co, Ni, Cu, and Mn. These all occupy the fourth period of the periodic table of elements. They share atomic radius, mechanical characteristics, and chemical properties. These elements were anticipated to create a simple phase structure in the HEAs. It has been suggested that HEA systems which include Fe, Co, and Ni often have a high strength-to-ductility ratio. As a result, Fe, Co, Ni, Cu, and Mn were chosen for this experiment. The measured elements were blended using a ball mill. The powder-mixing tank was cleaned with acetone solution before conducting the experiment. The powder was stirred for more than 10 h. For the proper filling of the HEAs powder, a 30 • bevel angle groove was fabricated on the parent metal.

Laser Welding
In the welding process, a fiber continuous laser welding machine was used (Dongguan Sanhe, China, model: 3HE-MF1500W, date of manufacture: 2018), and a continuous welding approach was applied. With a purity of 99.9%, argon was used as a shielding gas during welding. The weld parameters were optimized before the experiment with the Fe 5 Co 20 Ni 20 Mn 35 Cu 20 HEAs powders. An ideal set of process parameters was determined and used for welding dissimilar metals of AA 6061 and UNS S33207 alloys, with a 3 mm laser spot diameter, 640 W laser power, and 8 mm/s scanning speed [39,40]. Different parameter combinations were tested in trial runs before finalizing the best combinations [41][42][43]. Figure 1a shows the laser welding schematic with UNS S33207 hyper-duplex stainless steel (DSS) as the skin and AA 6061 aluminum alloy as the stinger with HEAs filler metal. Figure 1b shows the welded specimen. The laser beam was focused on the butting surface centerline of the lap-joint arrangement. This welding method prevents vaporization, which results in poor weld quality [44]. As a result, the heat generated when the laser beam fell on the HEAs area was effectively carried across the surface of the hyper-duplex steel.

Mechanical Testing
Tensile-shear testing was performed to determine the strength of the joints. Straightsided specimens of 25 mm width and 100 mm length were evaluated according to the ASTM D1002-10 standard. The tests were performed at room temperature with a 1 mm/min crosshead speed. A Zwick Roell tensile testing machine was used to conduct the experiment (Model: ProLine Z005 TN: 2018). Two specimens were prepared using the same parameter combinations, and the average value was taken. Microhardness was also determined at room temperature using a Mitutoyo Micro Vickers hardness testing system (Model: HM-102/103, 2018) as per ASTM E384 standard [45]. Microhardness testing was performed with a load of 25 g for 15 s.

Materials Characterization
The weld cross-sections were examined to evaluate the weld profile. The cross-sections were placed in epoxy resin and polished using P2000, P1500, P1000, P800, P400, P240 and P100-grit sandpapers with a drum polisher [46]. Keller's and 2% Nital etch solutions were used to etch the aluminum and hyper-duplex stainless-steel surfaces, respectively [47]. To determine the phases present in the base material and weld regions, X-ray diffraction (XRD) was performed (Model No: DW-XRD-27, Make: 2017, Company: Chongqing Drawell Instrument Co., Ltd., Shanghai, China), at 45 kV and 35 mA loaded with a Cu source, which was used as line 0.05 • of step size 2θ and 2 s/step. Figure 2 shows a cross-section of a laser-welded joint. The composition, fluidity, and surface tension of molten liquid aluminum and high-entropy alloy varied. As a result, it was not necessarily fully mixed, resulting in macro segregation [48,49]. In addition, the shape of the fusion zone on the steel plate indicated conduction welds. The bead geometry also revealed that the depth of diffusion on the steel plate was less than half its thickness, indicating that no molten aluminum was mixed with molten steel during the joining process. The breadth of the aluminum fusion zone, as shown in Figure 2, determined the bonding area between steel and aluminum. The parent material melted acceptably on the hyper-duplex stainless-steel side [50]. As a result, no welding imperfections, such as fractures or pores were prolonged in the joint interface of the weldment. However, after solidification, an HEA zone was formed at the weld interface due to the melting of the HEAs during welding.   [51,52]. As a result, the large combining entropy of HEAs was able to improve elemental solubility. In addition, the high blending entropy effect of the HEAs powder produced random solid solutions over IMCs. As a result, the HEAs solidification phase comprised DSS and V solid solutions after welding.

Results and Discussion
As illustrated in Figure 4, the interaction zone of HEAs/Al is represented in dark grey color; in Figure 5, the HEAs/Fe zone was relatively planar.
The EDS curves at the joint interfaces of the HEAs/Al and HEAs/Fe are shown in Figures 6 and 7. Consistent with the Al-V phase diagram analysis, the EDS curve revealed the substrate at the HEAs/Al joint region, indicating a more significant presence of Al and V. The Al-V IMCs may be dispersed there [53]. The atom intensity of the HEAs/Fe interaction region remained constant. There was no base, which indicated that the atoms in the HEAs/Fe diffused together to generate a layer of diffusion throughout the welding process. In each zone, using EDS analysis, the configurations of the resultant material were measured, and the findings are given in Table 3. At area A in Figure 2, Al and V were 72.34 and 18.27 (wt. %), respectively, and the microstructure was Al 3 V. As a result, Al 3 V made up most of the multiple layers (nearly all were 0.7 µm) at the interaction zone [54].  The Al-V IMCs performed effectively at room temperature. No Al element was identified at area B in Figure 2, suggesting that Al element diffusion had been inhibited. Because the HEAs contained many elements, this resulted in substantial lattice deformation in the grain configuration. The kinetic slow diffusion effect, caused by distortion in the structural lattice of HEAs, will obstruct atom mobility [55,56]. Fe atoms visible in area C (Figure 2), but not in area E (Figure 2), indicated a tiny quantity of Fe atoms diffusing to the HEAs with a limited diffusion distance. The HEAs larger mixing entropy can improve element compatibility. A small amount of hyper-duplex stainless steel dispersed into the HEAs without creating any IMCs.
Moreover, all materials in the HEAs powder on the Al and Fe sides (indicated as v-vi in Figures 6 and 7) were similar to the HEAs powder used as the filler. The elements Al and Fe were not identified in these areas, showing that using HEAs as a filler prevents the interaction of Al and Fe. The small dark dots in Figures 6 and 7 are vanadium particles (indicated by vii in Figure 7). A small amount of non-melted material particles were identified in the weld interface, which did not influence the joint strength [57]. Figure 8 shows the weld's tensile shear fracture, along with the HEAs/Al interface; this area was the weakest in this joining [58]. The influence of heat and the weak IMC formed between HEAs and Al in this region led to this failure. In Figure 9, all the tensile fractures occurred around the weld interface between the HEAS/Al alloy, and dimples could be seen on the fracture surface, indicating a ductile fracture mechanism for the specimen. When comparing the tensile characteristics of nonfiller laser lap welding to the HEAs powder filler induced laser-beam welding, filler-less laser lap welding had a lower tensile shear strength [59,60]. In addition, the on-filler laser lap weld joint's ductility was higher than the HEAs filler-induced laser welding. The latter finding, and the investigation results, suggest that the production of coarse columnar grains in the weld joint reduces ductility. The fracture locations of the tensile-tested parent metal and welds showed that the failure modes were ductile with adequate plastic deformation. Figure 9 shows a significant difference in the size and placement of the dimples on the fracture surface of the joint. The existence of voids and the size variation of the dimples indicated that the ductile fracture occurred with a significant drop in the tensile strength of the laser weld joint [61][62][63]. Several micro-pores were also identified, which might have been the reason for the lower tensile strength compared to the parent metals. Figure 9 shows a coarse and extended dimple fracture that could break under ductile shear.
The elongated cavities and coarse dimples indicated localized slip, resulting in increased ductility, without lowering the tensile strength of the HEA-induced laser joint. Figure 10 reveals a minor brittle cleavage fracture in the tensile-tested fractography. However, there was a noticeable variation in fracture patterns. The predominant properties of the HEAs filler and fine secondary dimples enhanced the tensile strength of the HEAs-filler-induced laser joint. The fractography (Figures 8-10) illustrates the specimen's ductile and brittle fracture properties [64]. The presence of the Al 3 V phase was verified using XRD data from the fracture cross-section, as shown in Figure 11. Figure 11. XRD data from fracture cross-section. Figure 12 shows the tensile stress-strain curve at room temperature for the HEAspowder-filler-induced laser welding. The tensile fracture was prolonged at the weldment's HEAs/Al metal interface. The weld had a tensile strength of 237 MPa. This highlighted that the weld joint had lower tensile strength than one of the base metals UNS S33207 hyper-duplex stainless steel and higher tensile strength than the aluminum alloy Al 6061. The noise observed in the stress-strain graph was caused by the HEAs' constant emergence of different phases. These different regions were: a linear elastic area, a tiny plastic region, and a further linear elastic region followed by plastic deformation [65]. As previously observed, coarse columnar grains in the joint region resulted in poorer ductility than fine equiaxed refined microstructure [66,67]. The Vicker's microhardness profile for the joint cross-section is given in Figure 13. The results showed that the average microhardness achieved in the weld zone where HEAs filler was deposited was 674 HV. However, the highest microhardness achieved in the weld zone where the HEAs/Fe interaction occurred was 693 HV, and in the HEAs region was 703 HV. The HEAs/Al weld interface evidenced a microhardness of 207 HV. Hence, the shear-transformed DSS phase with Al 3 V and V solid solutions transformed the post-weld microstructure of the weld zone [68,69]. Higher microhardness values in the weld zone might have been due to this phase transformation. The welded joints' hardness profiles (shown in Figure 13) are consistent with the joint's resultant tensile shear strength characteristics.

Conclusions
It is well established that intermetallic compound layer thickness, phase transformations and area of bonding have an overriding influence on the mechanical and metallurgical properties of welding using stainless steel and aluminum alloys. In this investigation, a high entropy alloys powder was chosen as a filler for laser welding. The conclusions derived from this research are as given below:

1.
In this laser welding experiment for joining UNS S33207 to aluminum alloy 6061, Fe 5 Co 20 Ni 20 Mn 35 Cu 20 HEAs powder was chosen as the lap welding filler.

2.
The Fe and Al alloy's mutual mixing and reaction were prevented by monitoring the melting quantity of the parent metal in the hyper-duplex stainless-steel side. Al/Fe intermetallic compounds were not identified in the joints weld zone. 3.
The metallographic phases identified for the HEAs in the weld region were DSS and V solid solutions.

4.
The large entropy property of HEAs favored the production of arbitrary solid solutions over IMCs. As a result, the HEAs' after-melting phase was comprised of DSS and V solid solutions.
The weld had a tensile strength of 237 MPa, and the tensile fracture was prolonged at the joint's HEAs/Al metal interface. 7.
The highest microhardness achieved in the weld zone was at the HEAs/Fe interaction area, which was 703 HV, with the HEAs/ Al weld interface having a microhardness of 207 HV.