Assessment of the Structural Integrity of a Laser Weld Joint of Inconel 718 and ASS 304L

Abstract

For high-temperature industries operating at nearly 750 °C (advanced ultra-super critical boilers), dissimilar welding between Inconel alloys and austenitic stainless steel (ASS) are commonly adopted. The high-temperature resistive properties of Inconel and ASS alloys are highly qualified for high-temperature applications. In this experimental study, dissimilar autogenous laser beam welding (LBW) between Inconel 718 and ASS 304L is investigated. This paper explains the detailed study on the microstructural and mechanical behavior of the LBW dissimilar joint. The microstructural study indicates the presence of laves phases in the weld zone. Additionally, the weld zone shows heterogeneous microstructural formation, owing to the non-uniform welding heat in the different areas of the weld zone. The optical images show the presence of mixed dendrites, i.e., equiaxed, cellular, and columnar morphology, in the weld zone and in the fusion zones of either side. The energy-dispersive spectroscopy (EDS) results show the presence of segregated elements (Nb, Mo, Cr, and Ti) at the weld center. These segregated elements are the reason for the occurrence of the laves phases in the weld zone. The presence of Nb and Mo may form the laves phase (Fe, Ni, Cr)₂ (Nb, Mo, Ti) along with Fe, Ni and Cr. The presence of an unmixed zone is observed in the HAZ of the Inconel 718, whereas the HAZ of the ASS 304L shows the presence of an unmixed zone (UZ) and a partially mixed zone (PMZ), as observed on the optical and SEM images. To obtain the mechanical properties of the laser weld, the tensile test, microhardness test, and impact test were measured at room temperature. The tensile specimens show a brittle failure at the ASS 304L side, which was initiated from the weld top, with average tensile stress of 658.225 MPa. The reason for the ASS 304L fracture is because of the presence of UZ and PMZ, and the lower hardness value of the ASS side. The UZ and PMZ lead to the fracture of the tensile specimen along the ASS 304L side’s HAZ. The measurement of microhardness carried out along the transverse length indicates an average microhardness of 214.4 HV, and the value is 202.9 HV along the weld depth. The mixed morphology of the microstructure promotes the variation in hardness in both directions. The hardness along the length shows a high hardness value in the weld zone and uniformly decreases along the base materials. The Charpy impact test of the weld zone shows the brittle fracture of the impact specimens. From the microstructural and mechanical results, the LBW dissimilar weld between Inconel 718 and ASS 304L is qualified for safe use in high-temperature end applications, such as AUSC power plants.

1. Introduction

Inconel 718 (IN 718) is a Ni-based precipitation-hardened superalloy that exhibits high performance properties, such as excellent weldability, high thermal strength, superior mechanical properties, high creep resistance (up to 860 °C), stress rupture resistance, and high-temperature corrosion resistance. The excellent high-temperature properties of this superalloy is because of the presence of gamma prime (γ′ [Ni₃(Al, Ti)]) and gamma double prime (γ″ (Ni₃Nb)) phases in the austenitic matrix of IN 718 [1]. Due to its excellent weldability and high-temperature strength, this superalloy is widely adapted in applications where high-temperature mechanical properties are required. IN 718 is widely used in jet engines, gas turbines, rotor shafts, the aerospace industries, the oil and gas industries, the automotive industries, submarines, and nuclear reactors [2,3]. The excellent properties and characteristics of IN 718 are because of high alloying elements and precipitation hardening technique, which increases the production cost and limits for wide and general applications. From an economic aspect, it is highly expensive to use Ni-based superalloys to produce single-structure or industrial parts. Hence, considering the economic cost and quality of superalloy structures, less expensive materials with almost similar properties and characteristics can be utilized for dissimilar weld [4,5]. Austenitic stainless steel is one of the materials that has similar properties and is less expensive, which is widely used in high-temperature industries along with superalloys.

Austenitic stainless steel is widely employed in the manufacturing industries, the aviation industries, nuclear reactors, heat exchangers, chemical reactors, pressure vessels and many more applications. Austenitic stainless steel 304L (ASS 304L) is a nickel- and chromium-based alloy. The addition of Cr in ASS 304L assists in achieving excellent corrosion resistance, and Cr is also considered a good ferrite stabilizer (α and δ ferrite). A high percentage of Ni content provides an austenitic matrix to this alloy to stabilize the austenitic structure at a low temperature. Additionally, the low carbon content restricts the formation of carbides at a high temperature [6]. Therefore, the dissimilar combination of Inconel and austenitic stainless steel is highly qualified for high-temperature applications where dissimilar welding is essential. To achieve a sound dissimilar weld, it is required to have suitable welding methods, optimum parameters, and welding environments. Many published articles have provided brief information related to different welding methods used in the past. Each welding method reveals different responses on Inconel and austenitic stainless steel combination. Although IN 718 is widely employed, its welding problem is the segregation of Nb element and formation of brittle intermetallic compounds (γ′-precipitates and Nb-rich compounds (laves phases)) during solidification. Laves phases are brittle intermetallic compounds (Ni, Cr, Fe)2(Nb, Mo, Ti) that occur during the solidification of Ni alloys with Nb elements [3]. The segregation of Nb and formation of intermetallic compounds cause the depletion of hardening alloys from the IN matrix, promoting crack initiation and propagation [7]. Numerous research studies have published different ways to bring down the formation of laves phases and their effect on weld mechanical properties. Some of the possible ways are employing (1) high cooling rates, (2) less heat inputs, (3) pulsing techniques, (4) steep thermal gradient, (5) Ni-based fillers with low Nb content, etc. [8,9].

An investigation was made on GTAW dissimilar weld of IN 604 and ASS 304 (3 mm each), comparing an autogenous weld and a weld using austenitic stainless steel (ER304) filler metal. The obtained grain morphology of the welds is dependent on their solidification rate. The autogenous weld and the pulsed current-gas tungsten arc weld (PC-GTAW) produced a fine grain structure because of a higher solidification rate when compared to the weld with the filler metal and the use of continuous current. The faster cooling rate promotes less segregation in the weld area [10]. Another study examined welding of IN 718 and SS 316l using the fillers ER2553 (SS) and ERNiCu-7 with continuous and pulsed current techniques. The weld produced using the Ni-based filler ERNiCu-7 (Nb free) produces fully austenitic weld matrix with no laves phase formation. The weld obtained by pulsed current GTAW using the ERNiCu-7 filler produces finer weld grains without laves phases, which enhances the mechanical properties. However, the use of the filler ER2553 is not recommended as it does not show effective result in surpassing the laves phases [11]. Furthermore, an examination on dissimilar weld between IN 625 and SS 904l using the GTAW process with the Ni- and Mo-based fillers, ERNiCrMo-4 and ERNiCrCoMo-1, was carried out. The fillers used in this study were Nb free, and this was performed to provide an Ni matrix along with suppressing the formation of laves phases in the weld zones and HAZ. Both the fillers are equally effective in suppressing the formation of deleterious laves phases at the weld center [12]. One study examined the GTAW dissimilar weld of IN 718 and SS 310 with the fillers IN 625, IN 82, and SS 310. The dissimilar weld with IN 625 experiences segregation of Nb and Mo elements, which is related to the presence of high amount of Fe in the IN 625 filler, whereas the presence of a lower amount of Fe element in IN 82 reduces the segregation of Nb element. The dissimilar weld with the filler SS 310 experiences high heat input, thereby increasing its solidification rate, generating the solidification sub-grain boundaries (SSGB) and solidification grain boundaries (SGB), and promoting solidification cracking in the weld zones [13]. Identical results have been reported in dissimilar welded study of IN 718 and SS 316l using the fillers ER2594, ERNiCrMo-4, and ERNiCrCoMo-1 by GTAW process. The formation of laves phases are successfully restricted in the HAZ of the IN 718 alloys using Ni-based fillers. Additionally, NbC phases were noticed in the HAZ of IN 718, and this is because of the high percentage of carbon in the fillers. Carbon shows a high penetration rate in the Ni matrix [14]. Hence, it is concluded that the use of Ni-based (Nb-free) fillers is beneficial to suppress the segregation effect of Nb and Mo elements in promoting the formation of laves phases in the HAZ and in the weld area. Additionally, the application of high heat input intensifies the solidification rate of the weld area, providing enough time for the elements to have high segregation rates. Generally, the GTAW process is used to weld this combination because of its easy setup and cost effectiveness, but the generation of high heat input is a major disadvantage of this process. To overcome the high heat input issue of GTAW, a modified PC-GTAW with less heat input and high solidification rate has been successfully performed by many researchers.

Additionally, the joining of thick plates by the GTAW process creates wide bead width and requires more welding passes, which generates high thermal cycle and alters the mechanical and metallurgical properties. Application of high thermal cycles promotes the formation of deleterious phases, grain coarsening, liquation cracking, and high degree of residual stresses [15]. To avoid the wide bead width and high thermal cycle, effective welding processes with concentrated heat source, higher depth of penetration, and large multipass are required. Laser beam welding and electron beam welding (high energy beam) are joining processes that provide the benefits of low heat inputs and high solidification rates [5]. Electron beam welding process provides the benefits of low heat inputs, high solidification rates (localized cooling), narrow weld bead, and HAZ, restricting the induction of residual stresses along the weld zone. A defect-free weld joint is obtained between IN 625 and ASS 304L using the electron beam welding process, providing the advantages of low heat input and faster cooling rate [16]. It is concluded that the high heat input, concentrated heat, and high solidification rate in laser beam welding provides welding advantages over conventional welding processes. The high solidification rate protects the weld zone from metallurgical defects and solidification cracks that occur during the solidification processes. The narrow weld bead achieved in laser beam welding protects the weld zone from weld defects.

In this work, the authors investigated the dissimilar welding between IN 718 and ASS 304L of thick plates (8 mm) using the laser beam welding (LBW) process. From previously published articles on dissimilar laser-welded, Ni-based super alloys and austenitic stainless steels, laser beam welding differentiates from conventional welding methods in terms of narrow weld bead width, deep penetration, high cooling rates, high weld speed, low heat input, and narrow HAZ [17]. These essential features of laser welding process make it an important method to weld dissimilar Ni-based superalloys and austenitic stainless steels. An experimental investigation on IN 718 and ASS 304 rods using the laser beam welding process concludes that the precipitation of laves phases are effectively lowered due to the fast cooling rate when compare to conventional welded joints. No grain coarsening is observed in HAZ, whereas due to the lack of strengthening elements, FZ suffers low hardness. Heat treatment of the joint helps in increasing the hardness due to retention of γ″ [15]. In another dissimilar laser welded joint using IN 718 and SS 316 plates of 3 mm, the authors confirm the segregation of elements (laves phases) through EDS in HAZs due to the concentrated heat source in the small area. Additionally, the relation between welding speed and cooling is found to be linear, i.e., weld with higher speed solidifies faster. Faster welding speed gives finer grains, enhancing the microhardness of the weld [18]. A laser beam welding of 10 mm P22 steel similarly reports that higher heat input results in an increase in the weld metal and HAZ hardness value [19]. In a laser beam welded dissimilar study using P92 martensitic and ASS 304L, due to the high heat input, a high level of mixed microstructure was observed, altering the mechanical property, including tensile strength, hardness, and impact toughness. Post weld heat treatment was carried out in this study to improve the mechanical properties [20].

Although there are very limited published articles available on laser welded dissimilar Ni-based superalloys and austenitic stainless steels, the authors made the conclusion that it is advantageous to obtain dissimilar joint for this bimetallic combination using laser welding in reference to its excellent mechanical and metallurgical properties. There is no study published on the bimetallic combination of IN 718 and ASS 304L plates welded by using laser welding. Hence, this study aims to extract some experimental conclusions related to the mechanical and metallurgical responses of laser-welded dissimilar joints of IN 718 and ASS 304L. We aim to extract some useful conclusion about the changes in microstructural morphology in laser-welded dissimilar joints of IN 718 and ASS 304L. The impact of microstructural variation on the mechanical properties of the weld will be further discussed. The mechanical characterization includes a tensile test and microhardness and impact test correlation. Additionally, this study includes a fractography study of the tensile strength and impact of two fractured samples. The results from the microstructural, mechanical, and fracture behavioral study provide notable insights to end users and also help researchers understand the laser welded behavior of this bimetallic combination.

2. Materials and Method

IN 718 and ASS 304L are dissimilar base materials used for the present experimental investigation. The autogenous dissimilar weld between IN 718 and ASS 304L plates with a dimension 150 mm × 65 mm × 8 mm was welded using the LBW process. The laser used to obtain the dissimilar weld is a LDF6000-30 Diode Laser with a power of 6 kW and a weld speed of 1 m/min. The laser beam produced has a size of 175 mm and a focal length of 0.2 mm. Additionally, pure argon gas (99.99%) with a flow rate of 15 L/min was used to protect the fusion zone. The elemental composition of the base materials was extracted using an optical emission spectroscopy (OES), and the analysis is tabulated in Figure 1. Figure 1 shows the schematic diagram of the autogenous dissimilar weld plates and the testing coupons’ dimensions. Figure 1a represents the laser beam welding setup, and Figure 1b–d show the extracted coupons of tensile, impact, and metallography, respectively, after the completion of the weld. The testing coupons were extracted after the completion of laser welding using electrical discharge machining (EDM) wire cut. The extracted metallographic and mechanical coupons show no surface defects, and there are no internal defects as verified through the test results (

Table 1. Composition of the base materials as received.

Elements (%)	C	Cr	Co	Mo	Al	Ti	Fe	Ni	Si	Mn	Nb
ASS 304L	0.11	0.92	0.46	0.92	-	-	Bal	0.28	2	8.08	-
IN 718	-	19.3	-	2.9	0.6	0.9	Bal	53.2	0.1	-	5.7

).

All metallographic and mechanical coupons were extracted following the ASTM standards. The metallographic cross-section coupons were prepared using the standard procedure to bring out the microstructure of the weld. The coupons were initially ground, followed by being polished with SiC paper to a grit size of 2000 and with a velvet cloth while applying 1 µm alumina powder. Thereafter, the polished surface was etched to reveal the microstructure of the fusion zone, HAZ, and base materials. The microstructure of the Inconel 718 base and fusion zone was revealed through electrolytic etching for 10 s in a 10% oxalic acid solution at 8 V DC supply. For the ASS 304L side’s HAZ and base material, an aqua regia (3HCl + 1HNO₃) etchant was used. An optical microscope and a scanning electron microscope (SEM) equipped with EDS were employed to reveal the microstructure across the weld. The EDS provides the elemental mapping from different zones of the etched samples. A Vickers microhardness tester was employed to measure the microhardness of the metallographic coupons across the weld in the transverse direction following the ASTM E384 standards. The hardness test was carried out with a loading of 1000 g for 10 s per indentation. Standard tensile samples as per the ASTM E8/E8M were extracted across the cross section to obtain the weld zone’s tensile strength (Figure 1). The tensile test was managed at room temperature with a uniform speed of 1 mm/min. For the Charpy V-notch test, two coupons with a dimension of 55 mm × 10 mm × 7.5 mm were extracted as per the ASTM E370 standards.

3. Results and Discussion

3.1. Microstructural Characterization

Figure 2 represents the optical and SEM images of the IN 718 and ASS 304L base materials in their received conditions. From the microstructural images of the IN 718 base material (Figure 2a,b), it clearly reveals the coarse grains of the Ni-rich austenitic matrix with uniformly distributed Nb and Ti carbide precipitates in the grains and grain boundaries. Along with these strengthening precipitates, twin boundaries also appear in the IN 718 matrix, which collectively enhances the dislocation density and mechanical strength of the 718 alloy. Additionally, the microstructural images of the ASS 304L show austenitic grains and uniformly distributed annealed twins (Figure 2c,d), whereas Figure 3b shows the weld cross section of the laser-welded IN 718 and ASS 304L. The weld bead from the IN 718 side shows a varying width from top to bottom in the shape of a “wine glass”. No porosity and other metallurgical defects are observed in the weld bead after employing the optimum laser weld parameters.

The SEM images in Figure 3a,c, which present the weld interface of the IN 718 and ASS 304L, display a microstructure with no metallurgical defect or any cracking. The optical images of the weld interfaces of the IN 718 and ASS 304L are shown in Figure 4a,b. From Figure 4, it is concluded that the weld interface is free from any defects and displays a cellular and columnar solidified microstructure. The grains in the fusion zone of the IN 718 show fine grain structures in comparison to the ASS 304L. The fusion zone of the IN 718 shows considerable unmixed zone (UZ), whereas the fusion zone of the ASS 304L shows an UZ and a partially melted zone (PMZ). The occurrence of UZ, PMZ, and islands are the outcomes of the macro-segregations happening at the fusion line of the weld and base materials (Figure 4). Although the weld interface of the IN 718 (Figure 5) exhibits a small amount of micro-fissures or liquation cracking effect, which is not of much concern. The occurrence of a micro-fissuring effect is due to coarser grains observed in the grain boundaries near the HAZ of the IN 718. Identical observations regarding micro-fissures were noted in a study of Co₂ laser-welded IN 718. Additionally, this study concluded that welding parameters have a significant effect on the formation of HAZ micro-fissures [21]. In contrast, in another study, it was mentioned that the formation of HAZ micro-fissures cannot be restricted during the welding of Ni superalloys with a high welding speed and low heat inputs [22]. Some of the common factors that are responsible for development of HAZ micro-fissures or liquation cracking in the grain boundaries are mechanical restraints, tensile stresses, microstructure being prone to cracking, segregation of elements, and constitutional liquation. Among the listed factors, constitutional liquation is the most crucial factor for the occurrence of micro-fissures or the liquation cracking at the HAZ grain boundaries. The large difference in the composition of the base material and the weld zone results in a decrease in the ductility of the alloys, leading to liquation of the grain boundaries. This reduced ductility promotes the grain boundaries’ embrittlement. Mechanical restraints and welding parameters cause some shrinkage stress and solidification cracking, promoting liquation cracking. Additionally, elemental segregation of Nb and Mo decreases the melting temperature through the grain boundaries, causing local melting and promoting liquation cracking at the HAZ grain boundaries [23]. From Figure 3b, we can clearly observe that micro-fissures occur rarely in a few HAZ areas, and, hence, this is not of much concern; however, the HAZ of the ASS 304L side shows an austenitic morphology with annealed twins, coarse austenitic grains, and ferritic stringers (Figure 3c). The presence of course austenitic grains shows the effect of welding heat adjacent to the weld zone.

From the optical image (Figure 4b), the ASS 304L interface zone shows different zones, namely a weld zone, an unmixed zone, and an HAZ. Epitaxial grain growth is clearly observed near the interface region. In the HAZ of the ASS 304L, lathy δ ferrite and austenitic grains are clearly visible in the SEM image (Figure 6a,b). The formation of the ferrites in the HAZ is possibly due to the welding thermal cycle, which promotes the transformation of austenite to ferrite [24]. Additionally, the high weld heat input provides the unmixing of the precipitates and leads to the formation of coarse grains in the HAZ. The weld zone microstructure formation is due to the solid phase transformations at a high welding heat input and solidification behavior, which is mainly driven by the cooling rate and the composition of the materials to be welded.

The microstructure morphology of the laser-welded IN 718 and ASS 304L (Figure 7) confirms that the weld zone is completely free from porosity and any other metallurgical defects. Fine-grained mixed dendritic microstructures are observed all over the weld zone. The SEM/EDS color mapping of the weld zone center shown in Figure 8 reveals the presence of elements in the weld zone. This result reveals that the amounts of niobium (Nb) and moybdenum (Mo) elements in the weld zone are 5% and 4%, respectively. At the time of re-solidification of the weld zone, Nb and Mo are elements that show strong micro-segregation, which can be related to the equilibrium coefficient value (k), k = C_s/C_L. where C_s and C_L are the compositional values of the solid and the liquid at the solid–liquid interface, respectively [25,26]. Elements having an equilibrium coefficient less than one show a high tendency toward micro-segregation during solidification [25,27]. From the EDS result, it is possible micro-segregation of Nb and Mo occurs in the austenitic-rich weld matrix of Ni, Fe, and Cr, indicating the presence of laves phases. Laves phases are brittle, topologically closed packed (TCP) phases having the chemical formula (Fe, Ni, Cr)₂ (Nb, Mo, Ti), which appear as the terminal solidification product under non-equilibrium conditions. It is also stated that the higher atomic radius of Nb and Mo restrict their complete diffusion into the Ni-rich austenitic matrix. The segregation occurs in the inter-dendritic region of the weld zone. Additionally, micro-segregation has been recorded at the weld interface due to the presence of a solute-rich liquid at the solid/liquid interface dendrites [8]. The elemental segregation in the weld zone is a time-dependent phenomenon, which is highly affected by the welding heat input and solidification rate. It is also mentioned that strong elemental segregation occurs if the solidification rates are slower, providing additional time for the solute rearrangement of the alloying elements in the weld zone [8]. The EDS point analysis of the dendritic and interdendritic region shown in Figure 9 confirms the presence of Nb elements (spot 4) in the interdendritic region, which might confirm the presence of laves phase in the interdendrites.

Additionally, the transition of the dendritic cores along the length and depth is recorded at different locations, i.e., at the weld top, weld center, and weld bottom, and in the fusion line of the two interfaces. Figure 10 shows the different morphologies of the microstructures in the weld zone from the top to the bottom of the weld, including equiaxed, cellular, and columnar dendrites. In Figure 10, the weld zone shows the heterogeneity of the weld microstructures from the top to the bottom of the weld. The formation of a heterogeneous microstructure is due to variations in the cooling rate, which changes the mode of solidification from the fusion line toward the weld center and from the top surface of the weld to the bottom of the weld zone. In Figure 4a and Figure 10, the microstructure at the IN 718 interface shows coarse and fine cellular and columnar dendrites. In contrast, in Figure 4b and Figure 10, the weld interface of the ASS 304L shows a similar microstructure after solidification. The microstructure morphology at the weld center solidifies at a different rate and shows different kinds of microstructure morphology, including equiaxed, cellular, and columnar dendrites (Figure 7 and Figure 11). The formation of a different kind of microstructure can be related to the difference in the thermo-mechanical properties of the IN 718 and ASS 304L.

The changed morphology of the dendrites can be correlated with the G/R ratio, where G represents the temperature gradient and R represents the solidification rate. It has been recorded that the value of R (solidification rate) is maximum at the center of the weld zone and minimum near the fusion line, thereby (G/R)_CL < (G/R)_FL. This change in the G/R ratio helps in the transition of columnar dendrites to equiaxed dendrites. Additionally, the high welding heat input promotes a reduction in the G/R ratio at the center of the weld zone and increases the supercooling at the solid–liquid interface [28]. This increase in the supercooling leads to the production of equiaxed dendrites at the weld center, as shown in Figure 7b. The presence of columnar dendrites is fairly observed near the fusion line of the IN 718 and ASS 304L (Figure 5), whereas equiaxed dendrites are clearly observed at the center of the weld. Additionally, in Figure 10, the interface region of the IN 718 shows mostly cellular equiaxed dendrites and, less columnar dendrites are observed. This might be due to the small value of the G/R ratio at the weld interface, which intensifies the supercooling at the solid–liquid interface. The microstructure at the weld center from the top, middle, and bottom also shows different grain morphology, as shown in Figure 11. The weld center microstructure shows the mixed behavior of the microstructure’s dendrites. The uniform and symmetrical solidification at the top of the weld center promotes cellular and equiaxed dendrites, and very few areas show the presence of columnar dendrites. The majority of the area in the middle of the weld center is covered with cellular and equiaxed dendrites, rather than columnar dendrites. In Figure 7, segregated elements can also be observed in the interdendritic regions, which has been discussed in previous section. At the bottom of the weld zone, cellular dendrites are in the majority compared to the interface region, which might be due to the uniform and symmetric solidification. Hence, a mixed microstructural morphology is also observed at the bottom of the weld zone.

3.2. Tensile Testing

Tensile tests were performed on two different transverse standard tensile specimens at room temperature, and the stress–strain graph is plotted in Figure 12. Two fractured tensile samples under examination, as shown in Figure 12a, initiated the fracture next to the weld zone in the HAZ of the ASS 304L side. The two fractured samples show an avg. value of tensile stress of 658.225 MPa, which falls in between the tensile strength of the welded base materials IN 718 (801 MPa) and ASS 304L (632 MPa). The obtained tensile strength of the laser-welded joints follows the order of ASS 304L < laser-welded joint < IN 718. Both the fractured samples initiated the fracture at the weld zone and ended up fracturing in the ASS 304L base material. Additionally, the nature of the fracture is completely brittle. The reason for the initiation of fracture from the weld zone might be the presence of the laves phase in the weld zone, whereas the presence of considerable PMZ and UZ in the HAZ of the ASS 304L promotes the fracture propagation from the weld zone to the ASS 304L. Similar tensile fracture trend was concluded in a study of a GTAW joint of IN 617 and SS 304H using the filler IN 617 [29]. The fracture of the tensile samples shows strong agreement with the result of the hardness value, which is explained in the next section. The dissimilar laser-welded joint between IN 718 and ASS 304L is confirmed for use in high-temperature applications due to its considerable high tensile strength.

3.3. Microhardness Discussion

The variation in the Vickers hardness of the LBW IN 718 and ASS 304L is shown in Figure 13a,b. The displayed microhardness is obtained along the transverse weld length and weld depth, respectively, by applying 1000 g load for 10 s to obtain each indentation. The obtained average hardness value for the transverse weld zone along the transverse length is 214.4 HV and along the weld depth is 202.9 HV. The hardness of the weld zone is attributed to the mixed microstructure obtained in the weld due to non-uniform solidification. Additionally, the presence of laves phase can be one of the reasons for the weld zone hardness. The microstructural change can be related to the weld zone’s non-uniform heating and cooling rates. The non-uniform weld cycle promotes the mixed-dendrite morphology containing equiaxed, cellular, and columnar dendrites. These mixed microstructural behaviors promote the variation in hardness along the depth of the weld zone, as shown in Figure 13b.The highest value of hardness along the weld length is 220.2 HV and the lowest is 181.5, showing a difference of 39 HV. The hardness value in the IN 718 due to the weld metal shows a uniformly increasing trend. The hardness value at the fusion line of the IN 718 and the weld zone shows the highest value of hardness due to the elemental segregation and formation of the laves phase at the grain boundaries. At the same time, the huge variation in hardness value along the weld from the top to the bottom can be related to the microstructural inhomogeneity. The ASS 304L shows a little variation in the HAZ, and this due to the recrystallization of the microstructure in between the HAZ and the fusion zone.

3.4. Impact Toughness

Figure 14 shows the impact of the two fractured specimens. The fracture obtained is totally brittle in nature. The impacted sample’s result shows a very low average value of 81.5 J at the weld center. The obtained low impact value is related to the Nb- and Mo-rich lave phase present in the weld zone. The brittle laves phase is present in the interdendritic core. This result shows an agreement with earlier published results, which conclude that the presence of laves phase in the weld zone reduces the impact strength of the weld. The weld between SS 316L and AH36 steel shows similar results [30]. The results of the impact fracture are in line with the tensile and hardness results, showing a uniform brittle fracture. The hardness variation is also affected in the weld zone by the non-uniform solidification. The presence of the brittle laves phase in the weld promotes the brittle fracture and the high variation in the mechanical properties.

4. Conclusions

The dissimilar LBW between IN 718 and ASS 304L is successfully achieved through applied laser beam parameters. The weld produced is free from weld defects, and its interfaces are also well bonded.

The SEM and optical images of the microstructure show heterogeneity in the weld zone due to the non-uniform solidification of the weld metal. The non-uniform solidification rate promotes the formation of different morphological dendrites, i.e., equiaxed, cellular, and columnar dendrites. In the weld zone, equiaxed and cellular dendrites cover the majority of the area, whereas columnar dendrites appear near the interfaces of the ASS 304L and at the bottom of the weld.

The elemental segregation is recorded using an EDS analysis of the center of the weld zone. The results show the presence of Nb and Mo elements. These elements form a brittle laves phase in the weld zone, altering the mechanical properties of the LBW IN 718 and ASS 304L.

The interface of the IN 718 shows grain coarsening and the presence of an UZ, whereas the interface of the ASS 304L shows the presence of an UZ, a PMZ, and very little changes in the grains.

The mechanical properties of the weld show the brittle nature of the weld zone. From the tensile results, the fracture was initiated at the weld zone and propagated to the ASS 304L, which matches with the hardness result. The average tensile stress is 658.225 MPa, and the average microhardness value along the transverse length is 214.4 HV and along the weld depth is 202.9 HV. The obtained hardness results indicate that the ASS 304L base material has the lowest hardness value. The hardness value along the weld depth shows high variation due to the presence of mixed dendrite structures at the center of the weld zone from the top to the bottom. The average value of impact toughness along the weld zone is 81.5 J, and this low value of impact toughness is due to the presence of Nb and Mo segregation and the formation of laves phases.

Hence, further parametric study may be required to bring down the formation of lave phases in the weld zone, which may shift the tensile fracture location from the weld zone, improving the mechanical properties of the LBW combination.