High Temperature Aging-Induced Degradation of Super 304H Stainless Steel Welds: A Microstructural and Mechanical Comparison Between TIG and SMAW

The impact toughness of austenitic stainless steel weld metal is strongly affected by the kinetics of intergranular precipitation in high-temperature use. Since this material is widely used in critical applications such as heat exchangers and pressure vessels in ultra-supercritical boiler systems, it is important to study the effects of these precipitates on its mechanical properties. Intergranular precipitation of carbides and other secondary phases at grain boundaries can significantly degrade steel toughness by facilitating localized brittle fracture. The precipitation processes, including the growth and distribution of particle-like M23C6 carbides, are central to altering the microstructure and, consequently, the resistance to impact loading. Purzyńska et al. [1] analyzed the long-term behavior of austenitic Super 304H steel after 31000 hours of operation at 570 ℃. Their microstructural examination shows severe precipitation of M23C6 carbides and sigma-phase particles at grain boundaries, along with dispersed ε-Cu and MX particles in grains. Wu et al. [2] explored strengthening and embrittling mechanisms of Super 304H steel under extended aging at 650 ℃ by considering the precipitates in the material microstructure. Their results show that the increase in hardness attributed to the M23C6 carbides at grain boundaries and intragranular nano-Cu-rich particles is critical to the increases in hardness of about 10 HV and 30 HV, respectively, after 5000 hours. Zieliński et al. [3] examined the long-term aging degradation of Super 304H steel exposed to aging for up to 50,000 hours at 650 ℃ and 750 ℃, and study the microstructure. The precipitation of various phases, such as M23C6, MX carbides, σ phase, Z phase, and ε-Cu phase, which significantly influence the creep strength of the materials, is reported in their findings. The study shows that elevated temperature is a critical factor in altering the steel's microstructure and, therefore, its mechanical behavior. Xiao et al. [4] investigated the influence of Nb addition on microstructural changes and microhardness in Super 304H austenitic stainless steel weld metals during long-term aging at 650 ℃. Their work indicates that, while Nb-free weld metal displays nanoscale M23C6 particles in the interdendritic regions, Nb-bearing weld metal has coarsened M23C6 particles at the phase interfaces between the primary Nb(C, N). Both weld metals are known to exhibit coarsening of Cu-rich particles with time. Li and Wang [5] examined the influence of second-phase precipitation on the mechanical properties of 22Cr25Ni3W3CuCoNbN steel, a new heat-resistant steel for superheaters in A-USC boilers. Their analysis reveals that exposure to aging at 700 ℃ for up to 1000 hours results in significant precipitation of the M23C6 carbides, initially preferentially at grain boundaries, thereby decreasing impact toughness. Zhao et al. [6] studied the long-term performance of Super 304H and HR3C steels after 40,000 hours of service at 630 ℃ and 3 MPa with respect to their oxidation behavior. The study shows the development of duplex oxides on the exposed steel surface, while surface deformation in Super 304H increases its corrosion resistance. M23C6 carbides, NbCrN, and ɛ-Cu precipitates are found at the grain boundaries and inside the grains for Super 304H, and Cr aggregation with nanoscale NbCrN particles in the grains is found for HR3C. Tikhonova et al. [7] examined the effect of increasing age up to 3000 h at 923 K on the structure, additional phases, strength, and intergranular detrimental corrosion of the Super 304H austenitic stainless steel. Their work shows that after 100 hours of aging, M23C6 carbides form at grain boundaries, and these boundaries broaden over time, increasing the likelihood that the material will become brittle and susceptible to intergranular corrosion. Xiao et al. [8] studied the effect of intergranular precipitation behavior on stress relaxation cracking susceptibility of Super 304H austenitic stainless steel weld metal during long-term aging. Their study shows that M23C6 carbides selectively precipitate at 45° misorientation grain boundaries, where they exhibit lower nucleation energy and higher strain concentration, making them prone to stress-relaxation cracking. Jin et al. [9] studied microstructural transformation and its role in the development of mechanical characteristics during aging of Super 304H steel at 650 ℃. The precipitation of Cu-rich, MX, M23C6, and σ particles is identified, with nano Cu-rich and MX particles distributed in the matrix and micron-sized M23C6 and σ particles formed at grain boundaries. The results show that with aging, tensile strength and hardness increase, whereas impact energy decreases, especially at two stages. Dash et al. [10] investigated the microstructural evolution and mechanical properties of Super 304H Cu-grade austenitic stainless steel (SS Super 304HCu) during long-term thermal exposure. Their work detected grain-boundary primary Nb(C, N) and M23C6 precipitates, secondary Nb(C, N) carbides, and Cu-rich phases that remained uniform over 10,000 hours of aging at 630 ℃. The results show that the creep rate increases significantly at higher temperatures, but as the Cu-rich phase precipitates, creep resistance also improves. Pulford et al. [11] examined elevated-temperature (650 ℃ and 700 ℃) uniaxial creep rupture specimens of the Super 304H austenitic stainless steel (Super 304H) and the microstructural evolution in the specimens. Their phase analysis focused on the precipitation of other second-phase products and a distinctly increased number of nanoscale Nb-rich and Cu-rich particles in the 700 ℃ specimens. Most importantly, the researchers demonstrated a direct spatial relationship between the formation of the brittle sigma phase and the formation of creep cavities. Kumar et al. [12] studied the impact of modified post-weld heat treatment (PWHT) processes on the mechanical properties and microstructure of dissimilar welded joints of T91 and Super 304H. Using Tungsten Inert Gas (TIG) welding, they performed two variants of PWHT processes: post-weld normalizing and tempering (PWNT) and post-weld direct tempering (PWDT). Results reveal enhanced microstructural uniformity and reduced Cr carbides in PWNT, whereas smaller carbides and sensitization occur on the Super 304H side in PWDT.

Hwang et al. [13] investigated the tensile and fatigue properties of Super 304H welding joints at high temperatures (500–700 ℃), a temperature range important for heat exchangers in thermal power plants. The research indicates that the tensile strength of welding joints equals that of the BM at all tested temperatures. Still, substantial shrinkage in deformation is observed due to non-uniform deformation. The high-temperature aging behavior of Super 304H austenitic steel, with particular focus on the coarsening and hardening of copper-rich precipitates at aging temperatures between 650 and 750 ℃, was investigated by Ou et al. [14]. They focused on the microstructures and revealed that these particles remain coherent and maintain a cubic-to-cubic crystallographic relationship with the matrix at longer exposure times. Still, their growth rate is primarily controlled by the diffusion volume of copper atoms. Jiang and Zhu [15] examined the strengthening mechanism of precipitates in S30432 heat-resistant steel under short-term aging at 923K, specifically in Cu-rich particles. Their work identifies Cu-rich particles as the major contributors to steel yield strength, accounting for at least 60% of the strengthening effect, ahead of Nb(C, N) and M23C6 particles. Although the degradation of the microstructures of copper-alloyed and niobium-alloyed Super 304H base materials after long periods of service (above 10,000 hours) has been well documented in the literature, there is a critical lack of understanding of the intermediate aging behavior of Super 304H weldments. The microstructural evolution of weld metal is fundamentally different from that of its BM due to the unique dendritic solidification pattern, elemental segregation, and welding thermal effects during the welding cycle. In addition, no studies have specifically identified these degradation pathways and separated the effects that the arc welding methodology has on them. The different levels of heat input and weld cooling rates associated with TIG and Shielded Metal Arc Welding (SMAW) processes result in different initial grain sizes and phase distributions, for example, localized retention of delta-ferrite. Hence, the objective of this work is to directly compare the microstructural and mechanical degradation of Super 304H stainless steel joints welded by TIG and SMAW with Super 304H and E308-16 consumables, respectively. This study subjects these joints to a controlled 500-hour thermal treatment at 650 ℃, thereby revealing the effects of the specific thermal cycles of each welding process on the precipitation of intergranular M23C6 carbide, and on the control of that carbide and its relation to the subsequent drop in impact toughness and local hardenability. In industrial applications, it is important to note that the welding methodology and the choice of filler material are closely interlinked. Accordingly, each of the welding configurations (TIG with the Super 304H filler, SMAW with the E308-16 filler) is treated as a separate 'Welding System,' and the synergistic effects of the thermal cycle and chemistry are assessed for long-term aging resistance.

3.1 Welding zone microstructure analyses

TIG welding and SMAW are welding processes that involve high heat, allowing materials to be melted together. All operational parameters, i.e., heat input, cooling rate, and shielding mode, directly influence the microstructure and mechanical behavior of the weld zone. A comprehensive understanding of these variations is necessary to optimize processes. When each technique is applied to Super 304H stainless steel, Figures 4(A) and (B) show the resulting weld zones. The welding parameters, shown in Tables 2 and 3, are also required to interpret their role in the microstructure. The following analysis explains the microstructural characteristics and formation processes of each welding region, focusing on texture, morphology, and phase formation. As illustrated in Figure 4(A), the major phase in the TIG weld zone of Super 304H stainless steel is Austenite, due to the low heat input and the high cooling rate associated with arc welding. Austenite is stable at higher temperatures, and the rapid weld-through rate of TIG welding helps maintain the phase throughout the weld region. Theoretical cooling curves for austenitic stainless steels predict that, if microsegregation localizes, even very small martensitic transformations can occur, depending on the cooling rate. The textured smoothness observed in the weld zone, however, indicates that if a transformation did occur due to the thermal cycle, the extent was insufficient to cause a significant martensitic transformation.

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Figure 4. Welding zone microstructure: (A) Tungsten Inert Gas (TIG) welding zone, (B) Shielded Metal Arc Welding (SMAW) welding zone

The dendritic shape evident in the micrograph is columnar, a form typically attributed to austenitic grains that crystallize under controlled temperature conditions. The dendritic arms are more strikingly aligned, and this observation is not misplaced, as evidenced by examples in which austenite is formed in the weld pool. The TIG welding zone exhibits a fine-grained microstructure. Such a result indicates both a low level of heat input and rapid solidification. The fracture resistance of the welded material increases as grain size decreases because smaller grains help restrict crack propagation [16]. In contrast, in the HAZ, grains tend to line out into columns and follow the direction of the heat-flow texture, improving the columnar dendritic shape that typifies the TIG-welded joint. The resulting orientation leads to improved mechanical properties, including higher tensile strength and greater fatigue resistance. TIG welding produces a dendritic structure that results from controlled, steady cooling rates. The primary dendrite arms branch perpendicular to the heat-flow direction, resulting in a more uniform texture than SMAW, which tends to exhibit irregular dendritic characteristics. The austenitic structure, which is the normal structure of the base material, Super 304H SS, may contain localized martensitic phases in the HAZ due to steep thermal gradients and the rapid solidification conditions of TIG welding, as indicated by predictive phase diagrams. Still, close control of process parameters, especially the application of argon shielding, has limited oxidation and thermomechanical phase transformations, reduced the potential to produce a large amount of martensite, and maintained a generally austenitic structure within the weld region.

Figure 4(B) shows a microstructural evaluation of SMAW welds, which demonstrates greater textural variation than other welding types. The HAZ of SMAW is highly irregular and exhibits a strong, chaotic dendritic pattern. These are the properties of SMAW because the higher heat input and slower cooling rates inherent in SMAW allow larger grains to grow. In the SMAW zone, austenite is the majority phase because Super 304H stainless steel is primarily austenitic. Austenite is likely to spread more evenly in the weld owing to the slow cooling rate associated with SMAW. However, the high heat input increases the likelihood of large-scale phase change, and in situations where cooling is inadequate or slow, ferrite may be found in isolated parts of the HAZ. Flux Coverage and the high heat provided by SMAW also ease this nucleation. Depending on the cooling rate, ferrite often forms, particularly in the HAZ, when the cooling time is sufficient for ferrite to stabilize. Although martensite formation is still possible in principle in rapid-cooling zones within the SMAW weld pool, the low cooling rate, which reflects the high heat input, reduces the likelihood of its occurrence. Instead, the retention of delta-ferrite is theoretically preferred when it is a second phase under the specific thermal dynamics of the SMAW process, thereby allowing extended exposure to elevated temperatures for compositional partitioning [18]. Microstructural observations under regulated cooling conditions indicate that, when SMAW welding is used, the microstructure of the weld region exhibits grains significantly larger than those in the TIG-welded specimen. The inverse relationship between grain size and mechanical resistance is well established: smaller grains confer greater strength and ductility. Because of its non-homogeneous nature, the SMAW process produces a grid of coarse and fine grains, which explains the preferred grain pattern in the TIG-welded microstructure. Because heat flow varies significantly across the weld section, large differences in grain shape and size result. The variable nature of heat supply and the use of flux-coated electrodes in SMAW often generate dendritic grain growth, shapes that do not exhibit as much regularity and order as in TIG welding, and inherently less consistent mechanical properties, especially tensile strength. Furthermore, slag inclusions, as byproducts of flux dispersion, can be generated in the SMAW-welded area; they can serve as crack nuclei and sites of carbide precipitation. Extended solidification, which accompanies relatively low cooling rates during SMAW welding, also facilitates large-scale phase transformations, promoting the formation of austenite as the primary solid phase in Super 304H stainless steel. This is followed by ferrite precipitation in the HAZ, which can reduce ductility and increase susceptibility to brittle intergranular failure. Finally, the elevated temperatures used during SMAW welding promote carbide precipitation at grain boundaries and further sensitize the material intergranularly. The microstructure of welded zones exhibits considerable variability in thermal history and material composition due to the welding process [19]. Although the common outcome of TIG and SMAW is an austenitic matrix, variations in heat input and cooling rate affect the formation of other potential phases, such as ferrite and martensite. These variations have been synthesized, and their influence on texture is explained in Table 4 below.

Table 4. Phase and texture comparison between Tungsten Inert Gas (TIG) and Shielded Metal Arc Welding (SMAW)

3.2 Welding zone microstructure analyses after the aging process

Aging of a welded metallic alloy, especially at high temperatures, significantly affects microstructural evolution. These effects are shown in Figures 5(A) and 5(B) for stainless steel type Super 304H welded by the TIG and SMAW processes. The material stabilizes through the precipitation of carbides and other microstructural refinements after 500 h of aging at 650 ℃. The subsequent carbide dispersion along grain boundaries and in the matrix controls the weld strength, ductility, and susceptibility to sensitization. Aging of a TIG-welded specimen at 650 ℃ in Figure 5(A) shows no structural differences, i.e., the same characteristics as before heat treatment; the microstructure remains a face-centered cubic (FCC) austenite matrix. However, an exact analysis of the interface region in Figure 6(A) reveals carbide formation that favors austenite grain boundaries. This effect implies that carbon segregation and nucleation on the dissimilar surfaces around austenite-austenite interfaces occur more rapidly during aging, limiting the distribution of carbides within the grains. As such, deleterious chromium carbides (e.g., Cr₂₃C₆ or Cr₇C₃) precipitate near the grain periphery. The aged specimen has a coarse boundary morphology, distinct evidence of disrupting the homogeneity of the initial texture relative to the as-welded microstructure, and a reorganized dendrite structure. The boundaries produced by dendrites are now much sharper and in a position to promote the precipitation of carbides. The high aging temperature of 650 ℃ greatly promotes sensitization, i.e., the formation of chromium carbides at these boundaries, thereby increasing the microstructural susceptibility to intergranular sensitization. Ductility degrades and is, in fact, abetted by accumulated carbides, which can cause brittle fracture. SMAW in Figure 5(B) shows grain boundary precipitation of carbides, similar to that found in the TIG specimen, but the texture is much more heterogeneous. SMAW microstructure is naturally coarser and more irregular, and it becomes even coarser during aging, indicating stronger carbide banding around grain boundaries [20]. However, austenite FCC remains the major phase, and aging once more leads to the precipitation of carbides at grain boundaries. Unlike the TIG microstructure, there may also be ferrite (body-centered cubic (BCC)), especially in parts with a slower cooling rate, and carbides more readily attack these ferritic areas since they have less solubility for carbon. In both SMAW and TIG, the precipitates are chromium carbides on the grain boundaries, but in SMAW, the grain size is larger. Therefore, the distribution of these carbides is less homogeneous. The formation of carbides in the SMAW joint is more prevalent than in the TIG joint, due to the lower cooling rates in the SMAW process.

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Figure 5. Welding zone microstructure after 500 h aging at 650 ℃: (A) Tungsten Inert Gas (TIG) welding zone, (B) Shielded Metal Arc Welding (SMAW) welding zone

Ferrite could also form in the SMAW weld, leading to larger carbide precipitates and making some areas prone to intergranular cracking. Irreversible grain structure and slag inclusions in the SMAW specimen also cause greater non-uniformity of carbides, thus creating uneven mechanical properties in the most heavily processed areas. The comparison of the evolution of the metallurgy of TIG and SMAW weldments after annealing at 650 ℃ at 500-hour intervals reveals significant deviations in carbide precipitation characteristics. Such observations can be explained by differences in cooling-rate contours and microstructural changes between the two welding methods. Table 5 below summarizes these key differences.

Table 5. Key differences between Tungsten Inert Gas (TIG) and Shielded Metal Arc Welding (SMAW) after aging

The authors explained the differences in carbide precipitation processes in TIG- and SMAW-welded joints after aging, which were due to the differing heat inputs. Due to the relatively low heat input and high cooling rate, TIG welding produces a finer microstructure and enables more rational carbide precipitation. In SMAW welding, the cooling rate is slower, allowing high-rate carbide precipitation, especially in ferritic zones, which are naturally prone to carbide formation [21]. In addition, TIG welding typically has a fine, homogeneous microstructure, promoting a uniform distribution of carbides. Still, SMAW welding has a coarse structure that allows ferrite to form and has an inhomogeneous distribution of carbides. The flux used in SMAW introduces impurities and further decreases the cooling rate, resulting in a more heterogeneous microstructure. TIG welding is more controlled; therefore, there is less variation in cooling rates, resulting in a more homogeneous carbide distribution.

3.3 Welding zone microstructure scanning electron microscopy, energy-dispersive X-ray spectroscopy, and transmission electron microscopy analyses after the aging process

The two SEM images in Figures 6(A) and (B) below, given in the present case, exemplify the morphology of the welding zone subjected to a 650 ℃ holding temperature over 500 hours, where the welding area has two welding processes: (a) TIG and (b) SMAW. The TIG weld zone (shown in Figure 6(A)) is fine-grained, austenitic, and homogeneous. Figure 6(B) shows the SMAW weld zone, which exhibits a rough, irregular texture due to lower cooling rates and higher heat input. Both weld zones exhibit some carbide precipitation along grain boundaries, particularly chromium carbides (M23C6), due to the high-temperature aging treatment. These morphological differences have a crucial effect on the material's mechanical properties and susceptibility to sensitization, which is why the cooling rate and heat input during welding are important. The SEM image of the TIG weld zone in Figure 6(A) shows a fine-grained, austenitic microstructure. The relatively slow heating in the controlled heat stage of the TIG process is attributed to the dendritic pattern. Dendritic solidification involves the formation of primary dendrites (trunks) and secondary dendrites (branches), in which the liquid metal forms tree-like structures. This tendency is common in austenitic stainless steel because it crystallizes in the FCC phase and has high-temperature stability. Columnar dendrite refers to a steady flow of cooling velocity and grain growth in the direction of heat flow during the process.

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Figure 6. Welding zone scanning electron microscopy (SEM) and transmission electron microscopy (TEM) after 500 h aging at 650 ℃: (A) Tungsten Inert Gas (TIG) welding zone SEM, (B) Shielded Metal Arc Welding (SMAW) welding zone SEM, (C) TIG welding zone grain boundary TEM, (D) SMAW welding zone grain boundary TEM

This is typically coupled with better mechanical properties, including higher strength and fatigue resistance. The tiny grain size impedes crack development, thereby boosting toughness. M23C6 carbide precipitation at the grain boundary requires a heat treatment of 650 ℃. Chromium from the austenitic matrix migrates to the interfaces. Consequently, the carbon remaining near it is depleted during this process, making the material more susceptible to intergranular sensitization [22]. The controlled heat input and fast cooling rates during TIG welding result in an equiaxed distribution of the fine carbides in the exposure region, which is a more desirable microstructure for retaining resistance to intergranular sensitization. However, grain boundary carbide precipitation may reduce ductility. Low-temperature exposure may also cause carbide growth, reducing tensile strength and impact toughness. Despite its ability to impart carbide precipitation effects on long-term, high-temperature strength, the low-grained microstructure retains respectable fracture toughness.

Figure 6(B) depicts the SMAW weld zone with a coarser, more irregular dendritic morphology due to the higher heat input and slower cooling rate of the SMAW process. Due to nonuniform heat during solidification, the resulting dendritic structure is less organized than that obtained in the TIG test sample. The discontinuous microstructure adversely affects the secondary phases, making the material more vulnerable to crack propagation under stress. The dominant dwell time at high temperatures also facilitates the development of large grains, reducing grain boundary strengthening, decreasing ductility, and restricting fracture toughness. An increase in the population of carbides that tend to form along the grain boundaries, including a relatively high proportion of M23C6 precipitates, also helps strengthen the grain boundaries against embrittlement. Large carbide shops, especially those found in ferritic (BCC) areas, are known to cause localized hardening and high intergranular susceptibility to cracking. The slow-cooling regime in SMAW welds also causes austenite to transform into ferrite, particularly in regions with slower cooling rates. In ferrite, carbide is expected to precipitate strongly at high temperatures, and heterogeneous carbide distribution is a major microstructural feature that contributes to susceptibility to intergranular sensitization. Moreover, ferrite precipitation reduces the weld's impact toughness, as shown in the toughness test data [23].

A substantial amount of carbide precipitation in structural steels preferentially occurs at grain boundaries. In cases where fine grains and high quenching rates can be favored by experimental arrangements of metallurgical processing conditions, as is frequently the case with high-strength steels, a beneficial result can be achieved on two counts simultaneously: high tensile strength and high impact toughness may both be retained. In comparison, alloys that undergo rapid quenching rates and large grain sizes, as found in thick-section steels, exhibit an increased concentration of carbide nucleation in the ferritic areas, leading to uneven precipitation, variability in mechanical properties, reduced ductility, and increased exposure to intergranular areas prone to brittle fracture. A major factor affecting carbide precipitation is the welding process. With high-speed TIG welding, high-temperature phase transformations are limited due to the rapid cooling rate. This fine-grained, highly homogeneous austenitic matrix gives superior mechanical properties to the TIG-welded joints and prevents the formation of massive secondary phases during initial solidification.

On the other hand, the welding processes in the SMAW family have a slightly slower cooling rate, which allows local ferrite nucleation; the austenite-ferrite transformation occurs faster than the secondary carbide precipitation reaction, causing sensitization of metals. In SMAW welds, a continuous network of large carbides is observed across grain boundaries and in ferritic areas, which are known to cause localized depletion of chromium and increase susceptibility to localized intergranular sensitization. However, TIG and SMAW welds remain carbide-precipitate-free at grain boundaries, producing a fine, homogeneous microstructure typical of high-speed TIG welding [24]. Table 6 compares the post-aging effects of both investigated welding processes.

Table 6. Tungsten Inert Gas (TIG) vs. Shielded Metal Arc Welding (SMAW) post-aging effects comparison

The presence of niobium in the EDS results shown in Table 7 suggests that the Nb-rich precipitates remained dominant in the TIG-welded joint, consistent with typical alloy behavior, Ni Carbide. The precipitates were continuous Nb-carbide precipitates at the grain boundaries, as marked by white dotted rectangles 1 and 2, though they were not yet a fully developed precipitate network. More specifically, at these boundaries, EDS identified two distinct second phases, demarcated by the white dotted rectangles 3 and 4 in Figure 6(A). Figure 6(A) illustrates this phenomenon at high magnification. The bright and the dark second phases were analyzed using the EDS points (points 3 and 4). Data indicated a large concentration of Nb in the radiant second phase and a high concentration of Cr in the dark second phase. As a result, both major Nb-carbide precipitates and Cr-enriched carbide precipitates remained on grain boundaries even after 500 hours of aging. In contrast, the Cr-carbide dissolved rapidly along grain boundaries, and the two precipitate types were evenly distributed throughout the structure's grain boundaries. Figure 6(B) for the SMAW joint indicates a substantial increase in the density of Cr-rich precipitates on grain boundaries, with a corresponding change in morphology to chain-like arrays extending across them. At the same time, the grain boundary width is much broader, which means that the precipitates have filled nearly all of the interfacial volume. EDS analysis showed that these particles were highly Nb-rich, suggesting they are still Nb-containing precipitates. There is clearly a higher incidence of intergranular precipitate coarsening during this time, with a conspicuous increase in the number and size of the precipitates, as marked at points 5, 6, and 7. These precipitates now completely cover the grain boundaries, forming a continuous network across them. The EDS results also indicated a high concentration of chromium in these intergranular precipitates, consistent with the formation of Cr-rich carbides (typically M23C6 in Super 304H aged at a suitable temperature) [25].

Table 7. Energy-dispersive X-ray spectroscopy (EDS) results for Tungsten Inert Gas (TIG) and Shielded Metal Arc Welding (SMAW) joints after aging treatment

Figures 6(C) and (D) show TEM images of the aged samples from the SMAW and TIG weld microstructures at the grain boundary. These correspond to retrieved samples of the weld metal. In the TIG sample (Figure 6(C)), large granular precipitates are observed. These grains were characterized as the main Nb carbides in the above analysis. The grain boundaries are quite narrow, and a small distance from the precipitates (also Nb carbides) has been shown to exist near the boundaries. However, the Cr carbide precipitated on the grain boundary is smaller than that precipitated on the SMAW grain boundary. Besides, the smaller second form of carbide began to appear at grain boundaries. A previous EDS examination revealed that the large carbides in the intragranular domain were secondary Nb carbides, and their identity was confirmed as NbC by SAED. Figure 6(D) illustrates large intergranular carbides precipitated on the grain boundary, which contained high Cr and minimal or no Nb in much darker zones on the grain boundary; the identity was M23C6 carbide, as determined by SAED. The Cr-bearing carbides were relatively scattered along grain boundaries due to the large number of Cr-rich carbides, the high heat input in SMAW, and the long aging time. The extent of precipitate formation along grain boundaries was very great, especially on the grain boundaries. There was greater coarsening of the precipitates in SMAW than in TIG, nearly covering the grain boundaries [26].

The SAED patterns of TIG- and SMAW-welded joints show precipitates at grain boundaries. Based on the results in Figure 7(A), a discontinuous distribution was observed at point 1 in the TEM picture in Figure 6(C), as a small quantity of carbides was distributed along grain boundaries. Based on the above analysis and EDS results, it can be concluded that M23C6 carbides were discontinuously precipitated at the grain boundaries after 500 h of aging. The precipitate at point 2 was massive and irregularly sized, and it was present both within grains as small secondary-phase islands and on grain boundaries. EDS analysis revealed that it contained significant Nb and was also identified as NbC by SAED in Figure 7(B). The SAED pattern of the bright secondary phase was indexed to the [011] zone axis, and the measured d-spacings were 2.58 Å, 2.23 Å, and 1.58 Å, consistent with the FCC structure of NbC. Likewise, from the dark outgrowths at the grain boundaries, the d-spacings were measured at 2.37 Å and 2.16 Å along the same zone axis [111], thus unambiguously assigning them to the M23C6 phase.

Figure 7. (A) and (B) Tungsten Inert Gas (TIG) selected area electron diffraction (SAED), (C) Shielded Metal Arc Welding (SMAW) SAED

SMAW TEM reveals precipitates with a continuous distribution along grain boundaries, and the degree of coarsening was also increased. EDS results indicated that they consisted of Cr, Fe, and Mo, which were identified as M23C6 based on the SAED pattern in Figure 7(C) (point 3 in Figure 6(D)). It is well known that sigma phase is a very brittle intermetallic phase that typically occurs in austenitic stainless steel weld metal [27, 28], and high temperatures can also promote its formation [29]. In this work, results at 650 ℃ show no σ-phase formation; thus, σ-phase precipitation is expected at higher aging temperatures.

3.4 Toughness test results analysis

Figure 8(A) of the TIG welding joint impact test shows a ductile fracture with a distinctive fibrous appearance. The fracture surface exhibits a small-scale fibrous structure characteristic of ductile fractures, in which the indent material undergoes high plasticity before failure. The fibrous microstructure is most likely explained by microvoid coalescence, i.e., the formation of small pores within the material and their coalescence into larger ones that eventually lead to fracture, as observed in materials with higher elongation, as seen here in metals. Fine texture is also associated with a noble microstructure within the TIG-welded joint. The heat applied by TIG welding is better controlled, resulting in a faster cooling rate and smaller grain sizes in the weld zone and HAZ. This grain-structure fineness is associated with increased resistance to crack propagation and, consequently, higher material toughness in its initial aging stages. As the carbides age, however, the microstructure is altered by their precipitation, as described below. The morphological difference shown in Figure 8(B) is due to unequal plastic deformation during solidification in metallurgical mechanics. The microvoids are becoming more prominent, and the grain texture is coarser in the SMAW specimen than in the TIG specimen. The occurrence of these holes is associated with ductile fracture conditions, which involve the breaking of grain boundaries or precipitate interfaces. Localized failures result from the coalescence of microvoids, meaning that the steel has deflected sufficiently for the voids to nucleate and grow. The deeper holes in the SMAW weld support the conclusion that there was significantly greater plastic deformation in the steel than in the TIG weld [30].

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Figure 8. Impact test fracture scanning electron microscopy (SEM): (A) Tungsten Inert Gas (TIG) welding joint, (B) Shielded Metal Arc Welding (SMAW) joint

The extrapolation of this discrepancy is the significantly higher heat input and resulting slower cooling rate inherent in the SMAW process, which allows extended time at elevated temperatures, favoring the development of larger, more brittle microstructural constituents in the HAZ. SMAW microstructure exhibits an unequal grain-size distribution and dimensions that can be said to be coarser, thus creating an irregular strain field that facilitates the formation of microvoids, which are highly pronounced but poorly distributed. The presence of MX carbides, which occur as primary carbides in long-aged Super 304H steel, is a prevailing cause of toughness loss. At 650 ℃, such carbides nucleate and grow in the microstructure after 500 h. These carbides lock in dislocations during the early stages of aging, thereby strengthening the material. In older ages, though, they grow, become hard and brittle, and act as stress concentrators. Such microstructural changes predispose the material to crack initiation at the carbide-matrix boundaries, thereby reducing overall toughness. It is clearly obvious that the formation of the MX carbides embrittles the TIG and SMAW weld zones, as evidenced by the fracture surfaces, and the noted reductions in the toughness of the steel from 213 ± 4 J/cm² before welding to 183 ± 4 J/cm² following an initial TIG weld, and from 198 ± 4 J/cm² before welding to 158 ± 4 J/cm² after an initial SMAW weld. In addition, the coarsening carbides disrupt the matrix's continuity, reducing the material's capacity to absorb impact energy. As a result, impact toughness decreases with aging. In addition, the large carbides also limit plastic deformation under impact loading. They would encourage brittle fracture behavior, as observed on the fracture surfaces of both TIG and SMAW welds [31].

Aging Super 304H stainless steel at 650 ℃ for 500 hours induces significant microstructural changes, particularly in the BM and the welded region, which eventually deteriorate fracture toughness. In such treatment, MX carbide nucleation occurs along grain boundaries and in the matrix; when they attain a coarse size, they lead to embrittlement. This leads to microstructural discontinuities that reduce a material's ductility under applied stress. The heavy precipitates weaken the weld zone under aged conditions, affecting both TIG and SMAW joints. Fracture surfaces, as revealed by SEM, show brittle carbide phases that increase in abundance with aging time. At the same time, the plasticity of steel, which measures its ability to absorb energy from impact loads, decreases because large precipitates constrain deformation. The fracture character observed throughout the alloy, with larger pores in SMAW and more stringy textures in TIG, indicates reduced toughness, as evidenced by increased brittleness [32]. The TIG welding technique also employs a controlled heat-input regime and generally produces a finer-grained microstructure in both the weld metal and the HAZ. This microstructural fineness prevents the coarse carbide precipitation usually resulting from post-weld aging, thereby reducing the concomitant loss of joint toughness. The comparatively high cooling rates of TIG welding improve the microstructure and, in turn, resistance to crack initiation and propagation. The TIG joint has an initial toughness of 213 ± 4 J/cm², which decreases only slightly to 183 ± 4 J/cm² after 500 h of thermal aging at 650 ℃, indicating minimal degradation due to carbide precipitation. (SMAW) On the other hand, it has high heat intensity and a fast cooling rate, resulting in a coarser microstructure in both the weld and the HAZ. This violent structure makes it prone to brittle failures due to the activation of larger carbide precipitates. In addition to the vulnerability to carbide-induced embrittlement, rapid cooling introduces increased residual stresses. Following 500 hours of thermal aging, the SMAW weld joint had a significant decrease in toughness from the 198 ± 4 J/cm² base value to 158 ± 4 J/cm², a 20.2% decrease. This dramatic reduction is thought to be due to intensive carbide nucleation and additional evolution of residual stresses, while the TIG joint was also reduced by 14.1% under the same conditions.

Through a comparative analysis of Super 304H stainless steel TIG- and SMAW-welded specimens, pronounced differences in fracture behavior, microstructural alterations, and reduced impact toughness are observed. When the welded joints experience ductile failure, the TIG weld exhibits a finer, more homogeneous microstructure, indicating lower toughness attenuation than the SMAW joint, which has a coarser grain size and a strong tendency toward void formation. The precipitation of carbides is a phenomenon resulting from aging treatment, leading to general embrittlement of the respective weld zones. In TIG welding, the ensuing decline is less pronounced due to faster cooling rates and relatively lower residual tensile stresses [33]. Table 8 below systematically summarizes the effect of each welding technique on material performance over time.

Table 8. Tungsten Inert Gas (TIG) and Shielded Metal Arc Welding (SMAW) welded joints after aging, impact fracture, and scanning electron microscopy (SEM) comparison

3.5 Vickers microhardness results analysis

The current study tested the microhardness of Super 304H SS BM, weld joints, aged, and HAZ at 650 °C after 500 h. The joints were produced using SMAW and TIG welding techniques; the microhardness results for the two joints are shown in Figure 9, with the corresponding BM hardness (about 190 HV) for both samples. In the HAZ, TIG-deposited material exhibited the highest hardness (205 HV) at the fusion line, which decreased toward the FZ. This hardness distribution is characteristic of a finely grained, uniformly dispersed carbide precipitate. Conversely, the hardness was less variable in the TIG-welded HAZ than in the SMAW (200 HV), which can be attributed to the relatively high heat input and slower cooling rate. In the FZ, TIG exhibited the maximum hardness of 202 HV, with a gradual reduction in hardness attributable to a compositionally uniform austenitic structure that solidified freely. SMAW joint fusion line was observed (~205 HV) due to the elevated input, which induced larger and non-homogeneous grains and an intermediate hardness compared to that of TIG welds.

Figure 9. Aged welding joints' microhardness test results

The major difference between the welding processes SMAW and TIG lies in their heat inputs and opposing cooling curves. The parameters have a vast impact on the microstructure, hardness, and mechanical performance of welded joints. The flux-covered electrode in SMAW also provides more heat than the pure tungsten electrode in TIG. Much slower cooling rates accompany the high heat input of SMAW during weld-pool solidification and post-solidification.

This extended cooling allows grain coarsening and the precipitation of secondary phases, especially carbonitrides, and favors the formation of a dendritic structure. This enhanced heat input, coupled with prolonged cooling, results in significantly coarse grains in the HAZ and even in the FZ. These larger grains compromise the material's ductility, increasing the tendency toward brittle failure. The coarsening effect is very pronounced in the HAZ and FZ, resulting in greater variation in the microhardness curves across the weld cross-section. Moreover, slower cooling facilitates the abundant precipitation of chromium carbides (M23C6) along grain boundaries. These carbide precipitates hinder the counter-movement of dislocations, consequently increasing the hardness of both the HAZ and the FZ. However, a high volume fraction of intergranular precipitates also renders it brittle and significantly enhances microstructural susceptibility to sensitization at higher exposure temperatures, such as 650 ℃. This problem is further exacerbated by the co-location of large carbides and ferrite in the SMAW microstructure, resulting in an uneven distribution of mechanical properties throughout the joint [34].

Unlike SMAW, TIG welding produces much less heat input. This difference is achieved using a stationary tungsten electrode with argon as the shielding gas, under precise control. The reduced heat input accelerates the cooling rate during solidification, causing the temperature variation to occur much faster than in SMAW—the microstructures in the HAZ and the FZ exhibit significantly smaller grain-size distributions. The close grain spacing, in turn, favors enhanced mechanical strength and elongation. The microhardness values are more consistent across all weld zones due to the uniform distribution of phases, especially austenite, and the limited extent of carbide precipitation. Although dendritic carbides can be detected in a TIG-welded joint, the low heat input and high cooling rate hinder their coarsening and dispersal. Instead, they are smaller and more evenly distributed, especially along austenite grain boundaries. The corresponding tendency is explained by high cooling rates, under which large quantities of carbides and/or ferrites are not formed, and, consequently, the overall toughness and ductility of the material are not compromised. The TIG process, therefore, produces less variation in hardness values across the BM, HAZ, and FZ than SMAW.

The microhardness results indicate that heat treatment at 650 ℃ leads to the precipitation of carbides, especially chromium carbide (M23C6), that occurs predominantly at grain boundaries. Precipitation occurs in boundary sites through carbon diffusion and chromium segregation. The tertiary forms, which develop during the process, alter mechanical properties such as hardness, ductility, and tensile strength. The microstructure of the TIG weld is fine-grained; hence, the precipitating carbides are small and located along austenite grain boundaries. The condition results in a moderate increase in hardness over the HAZ and the FZ. Regardless of the hardening, the fine-grained microstructure and the uniform distribution of carbides in the TIG-welded joint are associated with reduced embrittlement, so the resulting loss of toughness is minimal after 500 hours of aging. The moderately low effect of the precipitated carbide on the mechanical performance of the TIG joints is largely due to the small, uniformly distributed carbides that maintain ductility [35]. In contrast, SMAW welds have a coarse-grained microstructure, with larger carbide particles, leading to increased embrittlement. Bigger carbides, especially those in ferritic regions, increase stress concentrations and promote intergranular cracking. Accordingly, empirical research shows that TIG-welded joints exhibit a slight increase in microhardness, particularly in the HAZ and FZ. Enhancements in such hardness are attributed to carbide precipitation, which occurs during consecutive aging. However, TIG welding has a finer grain size and a more uniform carbide distribution, resulting in a more stable hardness profile and better overall hardness for the joint than the SMAW joint. The HAZ of TIG welds has a noticeable hardness peak. Still, the magnitude of the hardness difference is relatively small due to the high-quality microstructure and carefully controlled carbide nucleation. On the other hand, the SMAW joints with coarser grains and a high degree of carbide development exhibit significant changes in hardness following aging. The FZ and HAZ hardness values increase significantly, driven by the larger size of the carbides. However, the weld is also embrittled. The unbalanced carbides in SMAW also create large hardness gradients in the weld areas, thus providing less consistent hardness values in the joint [36].

The microhardness values for TIG- and SMAW-welded samples are given in Table 9 for comparison. This table specifically examines microstructure, cooling rate, carbide behavior, hardness distribution, and the effect of aging treatment. The main conclusion is that TIG welding yields a uniform microstructure, i.e., the hardness variation throughout the joint is relatively constant. In contrast, SMAW welding exhibits a wide variation in hardness due to coarser grains and a non-uniform distribution of carbides.

Table 9. Microhardness and microstructural characteristics comparison between Tungsten Inert Gas (TIG) and Shielded Metal Arc Welding (SMAW) welded joints after aging at 650 ℃