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J Weld Join > Volume 38(4); 2020 > Article
Yang, Kil, Moon, Nam, and Kang: Tensile and Microstructural Behaviors of Austenitic Stainless Steel GTA Welds for Cryogenic Application


This study investigates the weldability of gas tungsten arc (GTA) on stainless steel (STS) 304 using STS 308L filler and elucidates the cryogenic properties in STS 304 weld. GTA welding were performed with welding current of 90 A, welding velocity of 10.5 cm/min, Ar shielding gas, and in a single-pass. The GTA welding conditions were configured to ensure a full penetrating weld. For full penetration weld metal (WM), a small amount of δ-ferrite formed. Additionally, the WM of STS 304 exhibited larger dendrite packets compared to the grain size of base metal (BM). Therefore, the drop of hardness in WM was more significant compared to BM. The GTA welds had tensile properties that were inferior to BM at 298 K, and the tensile fracture occurred near the centerline of the transverse welds. The cryogenic tensile strength of GTA welds was superior to the room-temperature property due to the significant phase transformation of martensite at 77 K. However, the elongation decreased.

1. Introduction

With the recent increase in global LNG usage owing to various environmental regulations, the demand for cryogenic facilities such as LNG ships, pipelines, and LNG storage tanks has been growing1). In particular, the requirement for cryogenic strength and toughness has expanded as the development of natural gas in the deep sea or polar regions has been actively underway2,3). Austenitic stainless steel, which is currently mostly applied to cryogenic equipment, especially such as valves and pipelines, has been known to have excellent cryogenic properties such as strength and toughness even at low temperatures because of its properties of transformation induced plasticity (TRIP)4,5). When pipelines of cryogenic structures are welded, gas tungsten arc welding (GTAW) is mainly applied, which can minimize the deformation of welds with low heat input, does not require re-welding inside pipelines thanks to back bead formation, and has a beautiful weld appearance. The type of welding consumables according to the base metal (BM) has also an influence on weldability.
When austenitic stainless steel is welded, high temperature and solidification cracks may occur because of impure elements (P, S) in the steel that influence cracking susceptibility, and the integrity of welds are affected by whether or not such cracks occur6,7). It is known that a small amount of δ-ferrite formed during welding of stainless steel has high solubility of P and S compared to austenite, thus reducing the sensitivity of solidification cracking8,9), and by applying this effect, it is possible to reduce solidification cracks that may occur during welding. It is also known that stainless steel undergoes phase transformation from austenite to martensite in cryogenic regions, thus leading to nonlinear hardening10-12). That is, after the first yield, a phenomenon is seen where strength and elongation significantly increase while going through threshold strain and then additional second hardening11,12). Additionally, since austenite stabilizing elements are reduced for stainless steel because of the main Fe component, deformation- induced martensite is formed in the γ-austenite structure even at room temperature when external stress is exerted. It is known that the increase in deformation induces reduction in the stacking fault energy (SFE), thus the deformation mechanism changes from slip → twin → ε-martensite transformation behavior13-15). When deformation increases rapidly due to the decrease in test temperature, the fraction of the deformation mechanism is dominantly increased and ε-martensite is transformed into α-martensite. Therefore, the variation in the SFE, which is a function of the chemical composition of a material and temperature, plays a crucial role in the cryogenic mechanism of stainless steel. Although studies on the cryogenic strength of this stainless steel have been conducted mainly related to the BM, there has been no research on the mechanism for the cryogenic strength of stainless steel welds in relation to changes in microstructure. Therefore, in this study, we aim to evaluate the weldability of austenitic stainless steel, and identify the mechanism of cryogenic properties by comparing the tensile strength properties at room and cryogenic temperatures.

2. Experimental method

For the BM used in this study, among the various types of austenite stainless steel (STS), we performed solution treatment of STS 304 cold-rolled steel sheet at 1040°C for 1 hour, and then water-cooled it to prepare the sheet with a thickness of a 1.5 mm. The size of the STS 304 sheet is 55 mm (W) × 100 mm (L) × 1.5 mm (T). We also prepared a groove angle of 30°, a root face length of 0.5 mm, and a root gap of 0.5 mm for feeding filler metal. In addition, in order to fill the space of V-groove, we applied STS 308L filler metal. The chemical composition of the BM and filler metal used in this study is shown in Table 1. In this study, we also performed single-pass GTAW by applying a welding current of 90 A, welding speed of 10.5 cm/min, and filler metal of 2.0 Φ, which were optimal conditions to achieve fully penetrated welds.
Table 1
Chemical composition of the base metal and filler metal
C Si Mn Ni Cr P S Mo Fe
Base metal 0.06 0.44 1.04 8.03 18.1 0.03 0.01 - Bal.
Filler metal 0.01 0.41 1.60 10.1 20.0 - - 0.1 Bal.
We observed the macro and microstructure of the weld through the backscattered electron (BSE) mode of scanning electron microscopy (SEM), and analyzed the crystal structure of the weld using STS 308L filler metal and STS 304 BM through X-ray diffraction (XRD). We observed the microstructural behavior, grain size and phase fraction of the weld through electron backscattered diffraction (EBSD), and performed the quantitative analysis and mapping of electron-probe microanalysis (EPMA) to confirm the component behavior of the weld.
In order to identify the mechanical properties of the BM and the weld, we observed the hardness of the weld for 10 seconds with a load of 2.942 N (300 gf) using a micro vickers hardness tester, and measured the position 0.5 mm below the surface of the weld at an interval of 0.5 mm. After fabricating a sub-size specimen of ASTM E816) in the vertical direction of the welding line by taking into account the thin sheet BM, we performed the tensile test with the strain rate of 8.3 × 10-4 s-1 at room temperature (298 K) and cryogenic temperature (77 K). In addition, by observing the microstructure around the fracture surface after the tensile test, we analyzed the mechanism of cryogenic properties.

3. Experimental results

3.1 Component behavior and microstructure of GTA weld

Fig. 1a shows the shape of the weld cross-section of the GTA weld using STS 308L filler. No macro-defects such as pores and cracks were observed in the GTA weld. The upper bead was flat, but the lower bead achieved fully penetrated welds with a somewhat convex shape. Fig. 1b shows the components of the GTA weld through the EPMA line analysis. As seen in Fig. 1a, Fig. 1 the EPMA line analysis was performed along the area indicated by the red dotted line from the STS 304 BM to the centerline of the weld metal (WM). As clearly seen in Table 1, the chemical composition of the main alloy components (C, Mn, Ni and Cr) of the STS 308L filler metal is higher than that of the STS 304 BM. Therefore, it confirmed that the C, Mn, Ni, and Cr component behavior of the WM was slightly higher than that of the BM. This has an influence on the macro and micro component segregation of the weld. In addition, the reason for the irregularity of the main component behavior near the fusion line is considered to be the dilution of the main components in the heat-affected zone (HAZ) and the WM.
Fig. 1
GTA weld using STS 308L filler: (a) Weld pool shape and (b) Component behavior along the red-dotted line noted in Fig. 1a
Figs. 2a and 2b show the macro and micro component segregation behavior of the GTA weld using the STS 308L filler, respectively. Through macro and micro mapping of EPMA, we analyzed the segregation behavior of Ni, Cr, and Fe, which were the main components of the STS 304 BM and WM. Since there was little color difference between the BM and the WM, there was little macro segregation behavior of each component (Fig. 2a). The microstructure of the HAZ and the WM shows equiaxed and columnar shapes, respectively. It confirmed that the Cr component, which is a ferrite- stabilizing element, is mainly segregated at the grain boundary and interdendritic region, that the Ni component, which is an austenite-stabilizing element, is mainly segregated at the dendrite core, and that the Fe component is consistently present in all regions. Therefore, it can be predicted that the dendrite core contains austenite while the interdendritic regions and grain boundaries contain δ-ferrites17,18). In addition, in order to confirm the component behavior of each region, the results of the component analysis of the main components of dendrite core and the interdendritic region are shown in Table 2.
Fig. 2
Component segregation of the GTA weld using STS 308L filler: (a) macro- and (b) micro-segregation behavior
Table 2
Chemical composition of elements at the marked area (a, b) in Fig. 2b [wt. %]
C Mn Ni Cr Si Fe
a-dendrite core 0.08 2.4 9.7 18.3 0.32 69.2
b-interdendritic region 0.1 1.5 3.5 26.0 0.4 68.5
For the GTA weld of stainless steel, by applying the quantitative components (Creq, Nieq) of the WM to the Schaeffler diagram, it is possible predict the phase formed in the weld. Creq and Nieq near the centerline of the WM in this study were calculated as follows19,20).
Fig. 3 shows the result of the Schaeffler diagram of the STS 304 WM using the STS 308L filler. Austenite and δ-ferrite could be formed in the WM9), and δ-ferrite was predicted to be 7%.
Fig. 3
Formation of δ-ferrite calculated using Schaeffler diagram
To prove the predicted results through the Schaeffler diagram, we performed phase analysis through XRD. Fig. 4 shows the XRD patterns of the STS 304 BM and weld using the STS 308L filler. The diffraction peaks (2θ = 43.4 °, 51.6 °, 74.7 °) of the FCC single phase were observed in the STS 304 BM, but the diffraction peak of the FCC and BCC composite phase coexisted in the GTA weld. Normally, when the STS 304 BM is welded using the STS 308L filler, a certain amount of δ- ferrite can be formed in the weld9,10). Therefore, as predicted by the Schaeffler diagram, we were able to confirm that austenite and δ-ferrite coexisted in the GTA weld.
Fig. 4
XRD patterns of STS 304 BM and GTA weld using STS 308L filler
Fig. 5 shows the microstructural behavior from the STS 304 BM to the centerline of the weld using the STS 308L filler. The STS 304 BM has a microstructure containing approximately 20 ± 2 μm equiaxed grains and a small amount of twin because of recrystallization and grain growth during solution heat treatment. As seen in Fig.4, we confirmed that it consisted mainly of FCC single phases (Fig. 5a). The microstructure of the HAZ near the fusion line was similar to that of the BM, but the grain size became slightly coarser than that of the BM (approximately 22 ± 2 μm). In particular, near the fusion line, it confirmed the formation of δ-ferrite along the grain boundary caused by welding heat. The WM near the fusion line has the same crystal orientation based on the fusion line, and columnar grains grown in one direction are observed. The division between these columnar grains is distinguished by the color difference according to the crystal orientation in the inverse pole figure (IPF) map of Fig. 5b, Fig. 5 and the interface between the grains is represented by the grain boundary. In addition, δ-ferrite was observed along the interdendritic region classified as low-angle grain boundary within the grains. In general, through the austenite single-phase solidification (A-mode), and two-phase solidification of austenite-δ-ferrite (AF- mode) and δ-ferrite- austenite (FA-mode) caused by the welding heat generated during welding of stainless steel, δ-ferrite is crystallized8). In addition, ε-martensite was observed locally in the WM and the HAZ near the fusion line, which was caused by the occurrence of tensile/compressive stress due to the fusion and solidification process generated during welding (Fig. 5(b)). The columnar grains were grown in one direction toward the centerline of the weld, and the formation of equiaxed grains was observed at the centerline of the weld. The microstructure at the centerline of the weld had coarse grain size, and consisted of austenite and approximately 6.8% vermicular δ-ferrite (Fig. 5(c)). This result is almost the same as the ferrite formation level predicted to be 7% in Fig. 3.
Fig. 5
Microstructural behavior of the GTA weld using STS 308L filler: (a) STS 304 BM, (b) fusion line, (c) STS 308L WM centerline

3.2 Mechanical properties of the GTA weld

3.2.1 Hardness properties of the GTA weld

Figs. 6a and 6b show the hardness distribution and grain size of each region in the GTA weld using STS 308L filler, respectively. As seen in Fig. 5a, Fig. 5 the STS 304 BM was composed of the austenite single-phase, and the average hardness of the BM was measured to be approximately 188 ± 1 Hv0.3 since it has the grain size of approximately 20 ± 2 μm. As it gets closer to the fusion line, the grain size of the HAZ becomes slightly coarse, but the hardness distribution tends to increase due to the formation of δ-ferrite by welding heat. In addition, the hardness distribution that tends to increase rapidly in the HAZ near the fusion line is observed, which is considered to be attributed to the presence of Cr-rich carbides21) formed based on the fusion line and ε-martensite formed locally (Fig. 6c). The hardness distribution decreases rapidly from the fusion line toward the centerline of the WM, and the average hardness of the WM is 180 ± 1 Hv0.3, which is approximately 8 Hv lower than the BM. The main reason for this result is considered to be because the grain size of the BM is approximately 6 times finer than the columnar grains (approximately 120 ± 2 μm) of the WM (Fig. 6b). Since the columnar dendrites grew in one direction from the grains of the HAZ while going through the fusion and quenching steps by the welding heat source, the grain size of the HAZ and the dendrite arm spacing of the WM were almost similar. However, the grain size for the longitudinal direction of the columnar dendrite showed a difference.
Fig. 6
Characteristics of GTA weld using STS 308L filler: (a) Hardness distribution, (b) grain size in BM/weld centerline, and (c) presence of carbide and ε-martensite in HAZ

3.2.2 Tensile properties of the GTA weld at room and cryogenic temperature

Figs. 7a and 7b show the results of stress-strain curves and tensile fracture positions of the STS 304 BM and the GTA weld using the STS 308L filler. We performed tensile tests at room (298 K) and cryogenic (77 K) temperatures. With the decrease in test temperature from 298 K to 77 K, the yield strength (YS) and tensile strength (TS) of the STS 304 BM increased by 43% and 55%, respectively, but the elongation decreased by approximately 30%. Regardless of tensile test temperature, the tensile properties of the weld showed more degradation than the BM. As described in Fig. 6b, Fig. 6 this is because tensile fracture occurred at the centerline of the relatively weak WM by the coarse grain size of the WM compared to the BM. However, the cryogenic tensile properties of the BM and the weld showed better results than the room temperature tensile properties. Specially, the secondary hardening region after the first yield at the cryogenic temperature was clearly seen for both the BM and the weld test specimen, thus the weld was found to show the same cryogenic temperature tensile properties as the BM11,22).
Fig. 7
Tensile properties of STS 304 BM and GTA welds using STS 308L filler tested at 298 and 77 K: (a) tensile stress-strain curves and (b) fracture positions in the welds
Figs. 8a and 8b show microstructures in the vicinity of the tensile fracture region at room (298 K) and cryogenic (77 K) temperatures of the GTA weld using the STS 308L filler. We analyzed the microstructure of each region by applying the inverse pole figure (IPF), coincidence site lattice (CSL) boundary, and phase map of EBSD. Regardless of tensile test temperature, all weld specimens were fractured at the centerline of the WM. In order to investigate the mechanism of excellent cryogenic tensile properties of the weld, it observed whether deformation twin and martensite phase transformation by transformation induced plasticity (TRIP) occurred. Whether deformation twin was formed or not was represented by the CSL boundary of Σ3 indicated by a red line in the IQ image. In addition, as for martensite phase transformation, by observing the formation of martensite and the fraction through the phase map, we confirmed main strengthening mechanism. In general, deformation twin and strain-induced martensite transformation occur in austenitic stainless steels when stress is exerted, and these structures depend on the stacking fault energy (SFE). The SFE generated by the destruction of the crystal structure of the atomic arrangement present in the metal varies depending on the alloy element, temperature, grain size, etc. In particular, it is known that the SFE is greatly reduced depending on the change in temperature (289K - 77 K), and the fraction of deformation twin and martensite phase transformation increases sharply at cryogenic temperatures23,24).
Fig. 8
EBSD microstructures near the tensile fracture position of the weld using STS 308L filler at various temperatures: (a) 298 K and (b) 77 K
A small amount of deformation twin and martensite phase transformation were formed in the microstructure of the weld at room temperature. In the case of the WM using the STS 308L filler, deformation twin and martensite phase transformation may be formed when stress is exerted even at room temperature depending on the alloy component and grain size. On the other hand, we confirmed that the fraction of deformation twin and martensite phase transformation for the microstructure in the weld at cryogenic temperature was higher than for that of the room temperature weld. In the room temperature tensile test, with the increase in the strain rate, the deformation mechanism changes from slip → twin → ε-martensite by inducing the reduction in the SFE. However, the strain rate increases sharply with the decrease in test temperature, thus the fraction of deformation mechanism increases and the ε-martensite transforms into α-martensite. Therefore, as shown in 8(b), deformation twin and α-martensite phase transformation are dominantly formed at the same time. In addition, δ- ferrite, which was present regardless of the change in temperature, was indicated by a yellow dotted line in the IPF and phase map of each test temperature, and it confirmed that the δ-ferrite of the vermicular shape composed in the weld metal before the tensile test was present in a fragmented form. However, since the elongation is reduced compared to the room temperature test specimen, it is considered that the mechanism of excellent cryogenic tensile properties of the GTA weld using STS 308L filler is mainly a strengthening mechanism by martensite phase transformation.
Figs. 9a and 9b show the fracture surface of the GTA weld that was tensile fractured at room and cryogenic temperature, respectively. Coarse dimples were formed in the weld at room temperature, and the brittle fracture surface was observed locally due to the formation of martensite phase transformation (Fig. 9a). However, the dimple size of the fracture surface of the cryogenic weld was finer than that of the room temperature test specimen, and the fraction of the brittle fracture surface also tended to be higher in the cryogenic weld (Fig. 9b). This is considered to be because the formation of deformation twin and martensite phase transformation predominantly occurred in the cryogenic tensile test. This is also related to the reduction of the SFE due to the decrease in test temperature.
Fig. 9
Fracture surface morphology of the weld using STS 308L filler tested at various temperatures: (a) 298 K and (b) 77 K

4. Conclusion

In this study, we investigated the GTA weldability of STS 304 and its applicability in a cryogenic environment of weld using STS 308L filler, and drew the following conclusion.
  • 1) Austenitic and δ-ferrite coexisted in the weld of STS 304 using STS 308L filler. 7% of δ-ferrite was formed in the interdendritic region of columnar dendrites grown in one direction from the grains of the HAZ near the fusion line by welding heat.

  • 2) The reason that the average hardness of the WM is lower than that of the STS 304 BM is that the grain size of the BM is finer than that of the columnar grain of the WM. The HAZ near the fusion line has higher hardness than the BM because of the presence of local ε-martensite and the formation of carbide.

  • 3) Regardless of the change in tensile test temperature (298 K - 77 K), since all the GTA welds showed lower tensile properties than the STS 304 BM, tensile fracture of all the welds occurred at the centerline of the WM. However, the STS 304 weld using STS 308L filler showed excellent cryogenic tensile properties compared to those at room temperature as with the STS 304 BM.

  • 4) The improvement of cryogenic tensile properties was attributed to the fact that deformation twin and martensite phase transformation dominantly occurred. However, since the elongation is reduced in the cryogenic GTA weld, the main strengthening mechanism is considered to be the phase transformation of deformation-induced martensite.


This work was supported by the Ministry of Trade, Industry and Energy (MOTIE, Korea) under the Industrial Technology Innovation Program Grant No. 10080728, and the World Class 300 Project R&D (S2482209) of the MOTIE and MSS (Korea).


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