Microstructure and properties of wire arc additive manufactured 304/316L stainless steel welded joint

ABSTRACT

This study investigates the microstructural evolution and mechanical properties of a dissimilar tungsten inert gas (TIG) welded joint fabricated from wire arc additively manufactured (WAAM) 304 and 316L stainless steels (SS). The TIG welding was performed at 12.8 V, 70 A, with a wire feed rate of 2.5 mm/s and a 15 L/min Ar shielding gas flow. The WAAM deposition used a current of 180 A and a travel speed of 1 mm/s. Electron backscatter diffraction (EBSD) analysis revealed significant microstructural asymmetry across the joint. Due to the higher austenitic stability and lower melting point of 316L SS, a wider heat-affected zone (HAZ) and coarser grains were observed on the 316L SS side. Epitaxial growth at both fusion lines led to columnar grain formation, but competitive solidification resulted in an asymmetric fusion zone. The welding thermal cycle induced a pronounced “thermo-mechanical” effect, The coarse columnar grains formed by epitaxial growth contain a large number of low-angle grain boundaries (LAGB). Quantitatively, the proportion of LAGBs increased from 18% in the 304 SS base metal to 34% in the adjacent HAZ, while on the 316L SS side, it increased from 15% to 22%, particularly in the HAZ of the 304 SS side. This transformation is attributed to thermal strain-induced dislocation rearrangement and subgrain formation. Mechanical testing showed an ultimate tensile strength of 462 MPa and elongation of 24.5%, with fracture consistently initiating in the coarse-grained HAZ of the 316L SS side. Post-fracture analysis revealed that the 316L SS side underwent significantly greater plastic deformation, as evidenced by more pronounced slip bands and a final length of 15.07 mm compared to 9.47 mm on the 304 SS side. The 316L SS side, defined by both its lower base metal strength and welding-induced microstructural degradation, ultimately controls the integrity of WAAM dissimilar stainless steel joints.

Keywords:
WAAM; Stainless steel; TIG welding; Microstructure.

1. INTRODUCTION

While wire arc additive manufacturing (WAAM) research has extensively studied single-material microstructures and properties, systematic work on joining dissimilar WAAM-fabricated austenitic stainless steels (SSs)—and their multi-scale behavior under complex thermal histories—remains lacking [¹,²,³,⁴,⁵,⁶, ⁶]. This knowledge gap limits the potential of WAAM technology in the manufacturing of multi-material and functionally graded components [⁷,⁸,⁹,¹⁰]. as recently demonstrated in solid-state additive manufacturing of dissimilar material systems where interfacial microstructure control is critical for achieving sound bonding, e.g., Al/steel bimetallic components via wire-based friction stir additive manufacturing [¹¹], and in joining of high-temperature intermetallics such as TiAl alloys via novel brazing filler metals [¹²].

While both 304 and 316L SSs belong to the austenitic system, the addition of Mo in 316L SS significantly enhances its resistance to pitting corrosion, particularly in chloride-containing environments [¹]. However, the presence of Mo also alters the phase stability and high-temperature behavior of the material. Studies have shown that 304 SS is more prone to strain-induced martensitic transformation under low temperatures or strain conditions, while the austenitic phase of 316L SS is more stable [⁷, ¹³,¹⁴,¹⁵,¹⁶]. The thermal stability of metastable austenite plays a critical role in determining the transformation characteristics and microstructure evolution during complex thermal processes, as systematically reviewed by Fydryc et al. [¹⁷] from both experimental and theoretical modeling perspectives. Factors affecting thermal stability—including chemical composition, grain size, and prior deformation—exhibit different sensitivities, and their corresponding mechanisms have been extensively discussed [¹⁸]. In addition, physical metallurgy principles have been increasingly integrated with machine learning approaches to guide the design of advanced steels, where intermediate physical parameters (e.g., driving force for precipitation, equilibrium volume fraction) are incorporated to enrich data information and improve prediction accuracy, particularly for small-sample problems [¹⁹]. These fundamental insights into austenite stability and the PM-guided design methodology provide essential guidance for understanding the dissimilar joining behavior of 304 and 316L SSs, particularly regarding the asymmetric microstructural evolution across the fusion zone during welding thermal cycles. Thus, the welding of these two different austenitic stainless steels are often required. Arc welding processes, namely metal inert gas (MIG) and tungsten inert gas (TIG) welding processes, are widely sued in industrial applications to join stainless steels including dissimilar grades [¹⁷, ²⁰,²¹,²²]. When these WAAM-fabricated materials—with their unique textures and substructures—undergo TIG welding, solidification, heat-affected zone (HAZ) phase transformations, and potential carbide precipitation become exceptionally complex [²³, ²⁴]. Welding heat input not only alters the inherent columnar grain structure of WAAM but may also induce complex elemental diffusion and redistribution near the interface of the dissimilar materials, directly affecting the mechanical properties and corrosion resistance of the joint [⁷, ⁸, ²⁵].

Although existing studies have focused on welding dissimilar austenitic stainless steels in their conventionally rolled condition, the distinct initial microstructures of as-fabricated WAAM materials prevent the direct extrapolation of welded joint behavior [²⁵,²⁶,²⁷,²⁸]. In particular, the dissimilar combination of WAAM 304 and 316L SS—each with different austenite stability and thermal response—has not been systematically explored, and it remains unclear how these WAAM-specific microstructures govern asymmetric weld zone evolution. Therefore, the core objective of this study is to reveal the microstructural characteristics and performance of WAAM dissimilar SS joints. It involves a detailed analysis of the microstructural gradient from the WAAM base metal (BM), through the heat-affected zone (HAZ), to the fusion zone (FZ), with particular emphasis on elemental diffusion behavior, phase evolution, and interfacial bonding characteristics. In addition, microhardness and room-temperature tensile tests are performed to correlate the observed microstructural characteristics with mechanical performance, including strength, ductility, and fracture behavior.

2. MATERIALS AND EXPERIMENTAL METHODS

The schematic of the WAAM and TIG welding procedures is shown in Figure 1. A Yaskawa MA1440 six-axis robot with an RD350 inverter pulsed welding unit was used to fabricate 304 SS and 316L SS components under 99.99% Ar. The deposition current and travel speed were 180 A and 1 mm/s, respectively. The WAAM components were then machined into 1.5 mm thick plates. The faying surfaces were ground, degreased with ethanol, and dried. Autogenous oscillating TIG welding was used with optimized parameters: 12.8 V, 70 A, a wire feed rate of 2.5 mm/s, and a 15 L/min flow of 99.99% Ar, producing a sound 304/316L SS joint. Cross-sectional specimens were prepared for analysis. Microstructure was examined using an optical microscope. EBSD analysis was performed on an Oxford Symmetry S2 system (step size: 50 nm) to obtain grain morphology and crystallographic data. Microhardness across the weld was measured using an Vicks hardness tester (300 gf load, 15 s dwell) followed ASTM E384-17. Tensile tests were conducted on a universal testing machine at a crosshead speed of 0.5 mm/min according to ASTM E8/E8M-21. An extensometer was attached to the specimen to record strain during the elastic and early plastic deformation stages.

Figure 1
Schematic diagram of the WAAM and welding process.

3. RESULTS

Figure 2 displays the microstructure of the 304/316L SS welded joint. Figure 2 (a) and (b) show the 304 SS BM and the fusion line region adjacent to the 304 SS side, respectively. Figure 2(c) and (d) present the 316L SS BM and the fusion line region adjacent to the 316L SS side, respectively. The BMs microstructure consists of equiaxed austenitic grains, which is characteristic of a typical WAAM structure [²⁹,³⁰,³¹].

Figure 2
Microstructure of the 304/316L SS joint: (a–d) Corresponding details.

Due to the different BMs on either side of the weld joint, their physical properties differ [²⁸]. The melting point of 316L SS is lower than that of 304 SS, making it more susceptible to thermal effects during welding. The fusion metal is predominantly derived from the 316L SS. Consequently, the HAZ on the 316L SS side is wider (approximately 320 µm), and its grain structure undergoes coarsening due to the welding heat (average grain size of ~45 µm) [³², ³³]. In comparison, the HAZ on the 304 SS side exhibits a width of ~180 µm and an average grain size of ~28 µm. Figure 2 (d) reveals that columnar dendrites initiate growth, with their growth direction approximately perpendicular to the fusion line, extending toward the FZ center. This growth direction follows the temperature gradient, as columnar dendrites advance by progressively consuming ferrite. Due to the high temperature gradient from the fusion line toward the weld center, the columnar dendrites grow in that direction.

Combined with the EBSD results shown in Figure Figure 3(a) and Figure 3(c) represent the BMs of the 304 and 316L SSs, respectively. Both exhibit microstructure characteristics typical of WAAM [³¹]. However, due to the presence of Mo and typically higher Ni content in 316L SS, its austenitic structure exhibits greater stability compared to 304 SS, leading to differences in initial grain size and substructure density, subsequently affecting their thermal response. Key changes occur near the fusion line, as shown in Figure 3(b) and Figure 3(d). Under the influence of the welding heat source, the BM near the fusion line undergoes partial melting and intense thermal cycling. Epitaxial growth occurs at the fusion line for both materials. This means that the metal in the weld pool uses the unmelted grains on the surface of the BM as a substrate and continues to grow along the same crystallographic direction, forming a coarse columnar grain zone [³²]. However, due to differences in thermal conductivity between 304 SS and 316L SS, the interfacial temperature gradient and solidification front morphology on the 304 SS of the weld pool differ from those on the 316L SS [³³]. This results in different growth rates, sizes, and competitive selection processes for the columnar grains on each side, ultimately leading to an asymmetric microstructure in the FZ.

Figure 3
IPF images and distribution diagrams of HAGBs and LAGBs of the welded joints at different positions: (a) 316L SS BM, (b) fusion line near the 316L SS, (c) 304 SS BM, (d) fusion line near the 316L SS.

The grain boundary network in the BM is predominantly composed of high-angle grain boundaries (HAGBs), which is a typical characteristic of the rapid solidification process in WAAM. In contrast, the proportion of grain boundaries in the regions near the fusion line undergoes noticeable changes. The epitaxially grown columnar grains contain numerous low-angle grain boundaries (LAGBs), which arise from slight orientation variations among dendrite arms. More importantly, the welding thermal cycle exerts a ‘thermo-mechanical’ coupling effect on the BM just outside the fusion line. For 304 SS, it is more prone to strain-induced martensitic transformation or carbide precipitation during thermal cycling, potentially introducing more phase boundaries or hindering grain boundary migration. For 316L SS, its higher high-temperature stability may lead to a microstructure dominated by recovery and subgrain formation. Consequently, the evolution paths of grain boundaries in the HAZ for the two materials distinctly different, resulting in steep gradients of mechanical properties on either side of the fusion line.

Vickers hardness testing was conducted at a total of 40 sampling points spaced 0.5 mm apart across the entire weld cross-section. The hardness distribution is shown in Figure 4(c). The Vickers hardness of the 304 SS BM is approximately 220 HV, while that of the 316L SS BM is approximately 180 HV. The hardness of the FZ is slightly higher than that of the HAZ on the 316L SS.

Figure 4
Mechanical properties of the 304/316L SS weld joint: (a) tensile sample size, (b) fractured sample size,(c) microhardness, (d)stress-strain curve and fracture morphology.

The post-fracture morphological characteristics of the gauge section of the tensile specimen for the 304/316L SS welded joint are shown in Figure 4(b), with Figure 4(a) depicting the original dimensions of the gauge section. Ridges and valleys were observed on the sample surface, which clearly correspond to slip bands. These are typical features of plastic deformation during tensile testing, where dislocation motion is obstructed. The fracture occurred in the HAZ of the 316L SS, close to the weld joint. Following tensile deformation, the lengths of the 304 SS region, 316L SS region, and fusion zone (FZ) were measured as 9.47 mm, 15.07 mm, and 3.76 mm, respectively. Compared with the original FZ length of 3.18 mm, the FZ exhibited an elongation of 18.2%. After fracture, the elongated length of the 316L SS base metal (BM) was significantly greater than that of the 304 SS BM. Moreover, more pronounced and wider slip bands were observed in the 316L SS, indicating that this region experienced greater and more heterogeneous plastic deformation during tensile loading. This suggests that the strength of the 316L SS was lower than that of the 304 SS during the test. The 316L SS yielded first, undergoing plastic deformation. Subsequent work hardening increased its strength, after which plastic deformation propagated throughout the entire sample. The HAZ of the 316L SS exhibited coarse grains, which significantly reduce material toughness and became the initiation site for fracture. Figure 4(d) shows the stress-strain curve. The 304/316L SS welded joint exhibited an ultimate tensile strength of 462 MPa and an elongation of 24.5%. The fracture surface consisted of dimples of varying sizes and tear ridges, indicating a mixed-mode fracture combining ductile and toughness characteristics.

4. DISCUSSION

The microstructural asymmetry observed across the WAAM 304/316L SS welded joint is a critical factor governing its mechanical performance. As reviewed by LIAN et al. [³⁴], thermal gradient control is central to regulating grain morphology and defect suppression in additive manufacturing, and similar principles apply to the TIG welding of WAAM-fabricated materials where the moving heat source imposes complex thermal cycles on the deposited material. This asymmetry originates from the distinct chemical compositions and thermophysical properties of the two base materials. The 316L SS, with its higher nickel and molybdenum content, exhibits greater austenitic stability and a lower melting point compared to 304 SS. These characteristics lead to a wider HAZ and more pronounced grain coarsening on the 316L SS side during TIG welding [³⁵,³⁶,³⁷]. The thermal conductivity differences further influence the temperature gradient across the FZ, resulting in asymmetric epitaxial growth and competitive grain selection. Such findings align with previous studies on dissimilar austenitic SS welds, where compositional gradients directly affect solidification behavior and final microstructure [³⁸, ³⁹]. The presence of a coarsened grain structure in the 316L SS HAZ is particularly significant, as it reduces the local resistance to plastic deformation and serves as a preferential site for fracture initiation. This microstructural gradient, coupled with the variation in grain boundary characteristics, underscores the need for careful control of welding parameters to minimize asymmetry and optimize joint performance.

The evolution of grain boundary networks under the influence of welding thermal cycles further contributes to the mechanical heterogeneity of the joint. EBSD analysis revealed a marked increase in LAGBs within the HAZ, particularly on the 304 SS side (from 18% in the BM to 34% in the HAZ). This transformation is attributed to the ‘thermo-mechanical’ coupling effect during welding, where thermal strains induce dislocation rearrangement and subgrain formation. In 304 SS, which is more susceptible to strain-induced martensitic transformation, the thermal cycle may also promote the formation of phase boundaries, further complicating the grain boundary network. In contrast, the 316L SS side, with its higher stacking fault energy and austenitic stability, tends to undergo recovery processes that stabilize subgrain structures without significant phase transformation [⁴⁰, ⁴¹]. These divergent evolution paths result in a steep gradient in mechanical properties across the fusion line, influencing both hardness distribution and deformation behavior. The net effect of LAGB accumulation, however, is twofold: while LAGBs contribute to localized strengthening by impeding dislocation motion, the coarse grains that accompany them provide fewer barriers to crack propagation, thereby reducing the strain hardening capacity of the HAZ. A similar grain boundary engineering concept has been reported by WANG et al. [⁴²] in laser-directed energy deposited NiTi shape memory alloys, where herringbone grain architectures with jagged grain boundaries were shown to buffer stress concentrations, homogenize deformation, and suppress intergranular crack propagation, offering a valuable comparative perspective on how non-planar grain boundary morphology can enhance mechanical performance in additively manufactured alloys.

The mechanical testing results confirm that the microstructural asymmetry directly dictates the failure behavior of the welded joint. The ultimate tensile strength of 462 MPa and elongation of 24.5% indicate a ductile- dominant fracture mode, as evidenced by the dimpled fracture surface. However, the localization of fracture in the coarse-grained HAZ of the 316L SS side highlights the role of microstructural softening in that region. During tensile loading, the 316L SS section underwent greater plastic deformation, as reflected by the more pronounced slip bands and greater elongation compared to the 304 SS side. This suggests that the 316L SS HAZ yielded first, followed by work hardening and subsequent strain transfer to the fusion zone and 304 SS side. The coarsened grains in the 316L SS HAZ, which reduce the Hall–Petch strengthening effect, combined with a less pronounced increase in LAGB density compared to the 304 SS HAZ, created a zone of reduced toughness that became the fracture initiation site [³⁷, ⁴³]. These observations emphasize that in dissimilar WAAM joints, The 316L SS side, defined by both its lower base metal strength and welding-induced microstructural degradation, ultimately controls the integrity of WAAM dissimilar stainless steel joints. Therefore, future work should focus on tailoring the welding heat input and post-weld heat treatments to refine the HAZ microstructure and balance the mechanical properties across the joint. Beyond experimental parameter optimization, the multi-physical process simulation framework developed by LV et al. [⁴⁴] for resistance spot welding—which enables simultaneous prediction of thermal fields, contact behavior, and microstructural evolution—offers a promising methodology for future parametric studies of WAAM dissimilar joining, where the complex interplay between thermal gradients and microstructural response could be systematically investigated.

5. CONCLUSIONS

In summary, the WAAM 304/316L SS dissimilar joint fails from the 316L SS side because its lower melting point and higher austenitic stability promote greater HAZ softening—a wider HAZ (~320 µm, ~180 µm) and coarser grains (~45 µm, ~28 µm)—which lowers local deformation resistance and localizes fracture (15.07 mm, 9.47 mm plastic strain on the 316L side). The compositional asymmetry thus dictates both the microstructural evolution and the failure pathway.

6. ACKNOWLEDGMENTS

This work was supported by the Support Program of Heilongjiang Province for Basic Research of Outstanding Young Teachers(YQJH2025044), and the Selective Support Program of Heilongjiang Province for Returned Overseas Scholars.

7. DATA AVAILABILITY

Data will be made available on request.

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Publication Dates

• Publication in this collection
  10 Aug 2026
• Date of issue
  2026

History

• Received
  20 Mar 2026
• Accepted
  01 July 2026