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Microstructure and high-temperature oxidation behaviour of AISI 304L stainless steel welds produced by gas tungsten arc welding using the Ar–N2–H2 shielding gas

  • Sompong Chueaprakha , Thammaporn Thublaor , Thamrongsin Siripongsakul , Panya Wiman , Walairat Chandra-ambhorn and Somrerk Chandra-ambhorn EMAIL logo
Published/Copyright: September 9, 2025
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High Temperature Materials and Processes
From the journal High Temperature Materials and Processes Volume 44 Issue 1

Abstract

In the present work, we added 1.5 and 5% v/v hydrogen gas to an Ar–7% N2 shielding gas for the gas tungsten arc welding of AISI 304L stainless steel to investigate its effect on the weld microstructure and oxidation rate. By using Ar–7% N2 as a shielding gas, the weld metal contained 2.1% of delta ferrite in an austenite matrix. The addition of 1.5 and 5% hydrogen gas in the shielding gas provided the welds with a higher ratio of delta ferrite to austenite matrix, ranging from 3.8 to 6.9%, thus helping reduce the risk of hot cracks. The weld metals were further subjected to an oxidation test in synthetic air at 700°C, and the parabolic oxidation kinetics were observed. The parabolic rate constant of the weld metal produced using the Ar–7% N2 shielding gas was 5.44 × 10–13 g2 cm–4 s–1. When 1.5 and 5% hydrogen gas was added to the Ar–7% N2 shielding gas, the rate constants were reduced to 64% and 24% of that of the weld produced using only Ar–7% N2 shielding gas, indicating the promising role of the presence of hydrogen in the Ar–7% N2 shielding gas on improving the weld metal oxidation resistance for high temperature services.

Keywords: welding; hydrogen; oxidation; stainless steel; Brouwer diagram

1 Introduction

Austenitic stainless steel accounts for more than 65% of global stainless steel production due to its excellent corrosion and oxidation resistance, superior mechanical properties at high temperatures, and strong weldability [1]. Austenitic stainless steel is widely used across various industries, including defence technology, as well as a component in missile structures, casings, and engines [2,3,4]. In manufacturing processes, these parts are often welded together, and welding quality is critically important for ensuring the reliability and performance of high-performance equipment [5,6,7]. AISI 304L stainless steel, for instance, is used extensively in various appliances, tools, manufacturing machines, and industrial facilities due to its excellent mechanical properties and high corrosion resistance at room and high temperatures. However, its material performance changes when exposed to heat treatment processes during fabrication, particularly welding [6,7].

When welding austenitic stainless steel, the delta ferrite phase may be formed in the fusion zone of the weld metal and affect the weld properties [8]. The presence of delta ferrite in austenitic stainless steel welds prevents hot cracking during solidification. An appropriate amount of delta ferrite is typically 3–10%, which is known to reduce micro-fissuring in welds [9,10,11,12]. However, excessive amounts can sensitize welds to embrittlement [13,14]. With respect to the corrosion resistance, delta ferrite content exceeding 10% may increase the susceptibility of the weld to galvanic corrosion between the austenitic and ferritic phases due to the electrochemical potential difference between these microstructural constituents [15]. There can also be a negative impact on mechanical properties when excessive delta ferrite is present, as well as deterioration in the resistance to stress corrosion cracking, further compromising the integrity and longevity of the weld.

The use of combined shielding gases can also alter the composition and phase formation within welds in much the same way that filler metals or shielding flux cores are applied in the welding process. For instance, N2 in Ar shielding gas leads to dissolved nitrogen. This influences phase stability and promotes austenitic phase formation [16,17,18,19]. Previous studies have shown that the addition of N2 to Ar shielding gas could increase the austenite content. This suppresses delta ferrite content in welds as a result [6,20,21,22].

Furthermore, one of the shielding gas designs involves adding a small proportion of hydrogen to the shielding gas [23,24,25,26,27]. Adding hydrogen gas in the shielding gas is expected to enhance welding productivity by facilitating more efficient heat transfer due to its relatively high thermal conductivity compared to commonly used shielding gases such as argon or helium [24,25,26,27]. This increased thermal conductivity can lead to higher energy input, faster melting rates, and improved penetration depth, ultimately contributing to greater welding efficiency and reduced processing time [23,24,25,26]. The dependence of H2 in Ar shielding gas on delta ferrite content in stainless steel welds is still unclear. Brauser and Kannengieser reported on the increasing trend of delta ferrite content in duplex stainless steel welds [27]. This suggests that the combination of N2–H2 in Ar shielding gas may be able to balance the delta ferrite content in welds. However, the incorporation of hydrogen into the shielding gas must be carefully controlled to optimise welding productivity while ensuring the desired microstructure and mechanical properties of the weld. As previously mentioned, excessive hydrogen content can lead to adverse effects, such as increased porosity, hydrogen-induced cracking, and alterations in phase balance.

The combination of two shielding gases, Ar–H2 or Ar–N2, and the mechanical and chemical properties of the weld metal are of research interest. In the case of Ar–N2, the quantity of delta ferrite in the weld decreases because the addition of nitrogen stabilises the austenite [16]. The combination of three shielding gases, Ar–N2–H2, was investigated by Phakpeetinan et al. on 316L welds [28]. They found that adding hydrogen to Ar–N2 shielding gas increased the depth-per-width ratio of the weld, indicating the potential to apply the Ar–N2–H2 shielding gas to increase the welding speed, i.e. the productivity without sacrificing the acceptable weld bead shape. Recently, an investigation was also reported concerning the microstructure and oxidation resistance of the welds produced by the gas tungsten arc welding (GTAW) technique with Ar–N2–H2 shielding gas for Fe–15.7 wt% Cr–8.5 wt% Mn stainless steel [29].

Although the addition of hydrogen to the shielding gas exhibits a promising effect on welding efficiency, it also increases the dissolution of hydrogen in the steel weld [15,30]. The dissolution of hydrogen could further impact the properties of the steel weld [31,32] as well as the characteristics of the oxide scale formed at high temperatures. It has been widely reported that hydrogen plays a significant role in modifying the defect structure of the oxide [33,34,35,36,37] and affects the oxide formation rate [38,39,40]. The high temperature oxidation behaviour of the AISI 304L weld produced using Ar–N2–H2 shielding gas has not been investigated extensively. Thus, the objective of this work is to investigate the influence of adding hydrogen to Ar–N2 shielding gas on the microstructure of the welds, oxidation kinetics, and oxide scale formation at high temperature. Understanding these factors is essential for optimising welding parameters, improving material performance, and enhancing the long-term oxidation resistance of AISI 304L welds in high-temperature applications.

2 Materials and methods

In this study, AISI 304L austenitic stainless steel was used as the base material. The chemical composition of this alloy is presented in Table 1. Sample preparation was done by cutting a 2 mm-thick AISI 304L stainless steel plate into a size of 200 mm width and 100 mm length, configured as square butt joints. Welding operations were performed using an automatic GTAW process without the addition of a filler metal. A schematic diagram of the welding process is shown in Figure 1. The specific welding parameters utilised in this investigation are summarised in Table 2. Further details regarding the experimental procedures and methodologies adopted in the present study are also provided. The non-consumable electrode used for the joints investigated in this study was a thoriated tungsten electrode (EWTh-2) with a diameter of 2.4 mm. The nozzle size was 8 mm, and the tip angle was ground to 60°. Welding was performed with DC electrode positive polarity. The mixed shielding gas and backing gas were flowed at rates of 10 and 9 dm3·min–1, respectively. Ar, N2, and H2 gases with purity of 99.995% were mixed as shielding gases with the combinations in percentage by volume of Ar–7% N2, Ar–7% N2–1.5% H2, and Ar–7% N2–5% H2. After welding, the specimens were allowed to cool to room temperature. All welds were examined under a stereo microscope. The widths of the weld metal on both the face and root sides were determined using metallographic methods. For microstructural observation, the weld metal was sectioned perpendicular to the welding line. The cross-sectional surface labelled as face A was prepared for microstructure analysis. Figure 2 shows macroscopic views of the samples welded using the shielding gas with different compositions. The weld bead shapes were assessed and qualified according to ISO 5817 quality class B [41].

Table 1

Chemical composition of the AISI 304L stainless steel (wt%)

C Mn Si P S Cr Ni Cu Mo Si Fe
0.02 1.81 0.38 0.03 0.38 18.00 7.60 0.38 0.26 0.38 Bal.
Figure 1 Schematic drawing of the workpiece illustrating weld beads formation.
Figure 1

Schematic drawing of the workpiece illustrating weld beads formation.

Table 2

Welding parameters and heat input for various shielding gas compositions

Specimens Shielding gas (vol%) Weld speed (mm·s−1) Current (A) Voltage (V) Heat input (kJ·mm−1)
Ar–7% N2 93% Ar + 7% N2 3.97 85 11 ± 1 14.13
Ar–7% N2–1.5% H2 91.5% Ar + 7% N2 + 1.5% H2 3.97 85 12 ± 1 15.42
Ar–7% N2–% 5H2 88% Ar + 7% N2 + 5% H2 3.97 85 13 ± 1 16.70
Figure 2 The weld beads produced using (a) Ar–7% N2, (b) Ar–7% N2–1.5% H2, and (c) Ar–7% N2–5% H2 shielding gases.
Figure 2

The weld beads produced using (a) Ar–7% N2, (b) Ar–7% N2–1.5% H2, and (c) Ar–7% N2–5% H2 shielding gases.

For the oxidation test, each welded sample was prepared by cutting along the welding line to obtain only the weld metal with a length of 10 mm and a width of 3 mm, as shown in Figure 1. All samples were ground on SiC paper at up to 1,200 grit, cleaned in ethanol with an ultrasonic wash for 5 min, and then dried in forced air. The samples were placed in a tube furnace where synthetic air was flowed at a rate of 1 dm3·min–1. The atmosphere in the furnace was heated from room temperature to 700°C for 30 min. The oxidation test was conducted for 8 h, after which the samples were cooled in the furnace. After ending each cycle, the mass change of each sample was measured using a five-digit balance. This cyclic oxidation was performed for ten cycles. The samples were characterised using a scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS) for elemental analysis, and an X-ray diffraction (XRD) technique using a glazing mode with a diffraction angle of 2° and a regular mode was used to identify the formed oxides and phases of the weld metal.

For microstructural analysis, each welded sample was cut across the welding line. Samples were ground on SiC paper at up to 1,200 grit. The polishing process was carried out using alumina powder, followed by electrochemical polishing in an oxalic acid solution at a DC voltage of 6 V. At the weld zones, the delta ferrite content was analysed using a ferrite content analyser (Feritscope FMP30), with a 2 mm probe. Measurement was conducted using a magnetic induction method to measure the amount of the magnetisable phase, like the ferrite contained in the austenitic stainless steel weld. Measurement was done after calibration with the provided standard samples to construct a calibration curve. The results are directly expressed as a percentage of ferrite content. Additionally, an oxygen–nitrogen analyser was employed to determine the dissolved nitrogen content in the weld, and measurements were conducted on the weld bead after calibration.

3 Results

Figure 3 shows the microstructures of the samples produced using Ar–7% N2, Ar–7% N2–1.5% H2, and Ar–7% N2–5% H2 shielding gases at different zones, i.e. the weld metal area and the fusion boundary zone between the heat-affected zone (HAZ) and the base metal (BM). As shown on the left side of Figure 3, the width of the HAZ increased with increasing hydrogen content in the Ar–7% N2 shielding gas. This can be attributed to the increased heat input associated with hydrogen addition, which enhances the thermal energy delivered to the weld pool. Similar observations have been reported in other studies investigating the effect of hydrogen addition to the shielding gas [42]. Figure 4 shows the microstructure of the welds taken from the centre part of the microstructure at high magnification. For the weld metal area, the formation of dendritic delta ferrite and austenite matrix can be observed. The ferrite analyser was used to measure the ferrite content of the weld metals, and values of 2.1, 3.8 and 6.9% were obtained with standard deviations of 0.2, 0.3, and 0.4% in the weld metals produced using Ar–7% N2, Ar–7% N2–1.5% H2, and Ar–7% N2–5% H2 shielding gases, respectively, as shown in Figure 5. The dendritic structure in the matrix was still observed in the HAZ, while the BM showed a twin structure as a typical microstructure of the austenitic stainless steel. Figure 5 also shows that the nitrogen content in the weld metal is in the range of about 0.06–0.14 wt%. The nitrogen dissolution in the weld was less when more hydrogen gas was added to the shielding gas.

Figure 3 Microstructure observed from optical microscope of the weldments produced using (a) Ar–7% N2, (b) Ar–7% N2–1.5% H2, and (c) Ar–7% N2–5% H2 shielding gases.
Figure 3

Microstructure observed from optical microscope of the weldments produced using (a) Ar–7% N2, (b) Ar–7% N2–1.5% H2, and (c) Ar–7% N2–5% H2 shielding gases.

Figure 4 SEM micrographs at higher magnification of the welds produced using (a) Ar–7% N2, (b) Ar–7% N2–1.5% H2, and (c) Ar–7% N2–5% H2 shielding gases.
Figure 4

SEM micrographs at higher magnification of the welds produced using (a) Ar–7% N2, (b) Ar–7% N2–1.5% H2, and (c) Ar–7% N2–5% H2 shielding gases.

Figure 5 The delta ferrite content and dissolved nitrogen of the welds produced using shielding gas with different compositions.
Figure 5

The delta ferrite content and dissolved nitrogen of the welds produced using shielding gas with different compositions.

Figure 6 shows the XRD results of the welds obtained after welding. As expected, the patterns of all samples exhibited the presence of austenite (ICDD 31-0619) and ferrite (ICDD 87-0722). The major XRD peak (110) at 2θ of 44.8° represents the ferrite phase and is obvious. In addition, its subsequent peaks (200) and (211) are slightly visible at 65.2° and 82.5°, respectively. For austenite, the representative peaks (111), (200), and (220) were detected at 2θ of 43.5°, 50.7°, and 74.7°. The addition of hydrogen to the shielding gas resulted in a change in the diffraction patterns compared to the one without hydrogen. The XRD peak of the (200) austenite was the major one instead of the (111) one.

Figure 6 XRD patterns of the welds produced using shielding gas with different compositions.
Figure 6

XRD patterns of the welds produced using shielding gas with different compositions.

The oxidation test in synthetic air at 700°C was performed to investigate the oxidation behaviour of the welds. The mass gains of the welds produced using the shielding gases with different compositions were plotted as a function of the oxidation cycle, as shown in Figure 7(a). The experiments were repeated twice for each condition. It was found that the mass gains of the welds produced using the Ar–N2–H2 shielding gas tended to be lower than those of the welds produced without the addition of hydrogen, while increasing the hydrogen content in the shielding gas resulted in the reduction of mass gain. The mass gain values in Figure 7(a) were further squared and plotted as a function of the oxidation cycle, as shown in Figure 7(b). The linear relations between the square of mass gain and the oxidation cycle of the studied samples were observed with R 2 in the range of 0.969–0.998. By estimating that one cycle corresponds to 8 h of oxidation, the parabolic rate constant ( k p ) can be calculated using equation (1), where ( Δ m / A ) is the mass gain and t is the oxidation period. The average parabolic rate constants of the welds produced using Ar–7% N2, Ar–7% N2–1.5% H2, and Ar–7% N2–5% H2 shielding gases were 5.44 ± 1.29 × 10–13, 3.47 ± 0.54 × 10–13, and 1.32 ± 0.17 × 10–13·g2 cm–4·s–1, respectively. Each dashed line in Figure 7 was plotted using the rate constant obtained from each tested sample:

(1) Δ m A 2 = k p × t .

Figure 7 (a) Mass gain and (b) square of mass gain as a function of the oxidation cycle for the cyclic oxidation test at 700°C in synthetic air.
Figure 7

(a) Mass gain and (b) square of mass gain as a function of the oxidation cycle for the cyclic oxidation test at 700°C in synthetic air.

Figure 8 illustrates the SEM cross-section with the EDS line-scan results of the sample, which had the largest mass gain in this study, i.e. the weld metal produced using Ar–7% N2 after the cyclic oxidation for ten cycles. The average scale thickness measured from five different locations in this figure was 2.07 μm with a standard deviation of 0.54. When hydrogen was added to Ar–7% N2 shielding gases, the average scale thickness significantly decreased to 1.53 μm with a standard deviation of 0.17 and 0.76 μm with a standard deviation of 0.15 for the weld metal produced using Ar–7% N₂–1.5% H₂ and Ar–7% N₂–5% H₂ shielding gases, respectively. It was noted that the scale thickness was significantly thinner when 1.5 and 5% hydrogen were added to the Ar–7% N₂ shielding gas.

Figure 8 SEM cross section and EDS line scanning results of the welds produced using Ar–7% N2, Ar–7% N2–1.5% H2 and Ar–7% N2–5% H2 shielding gases after cyclic oxidation at 700°C in synthetic air for 10 cycles.
Figure 8

SEM cross section and EDS line scanning results of the welds produced using Ar–7% N2, Ar–7% N2–1.5% H2 and Ar–7% N2–5% H2 shielding gases after cyclic oxidation at 700°C in synthetic air for 10 cycles.

The EDS results in Figure 8 show the increased signals of Cr, Mn, and O in the oxide scale. The XRD technique in a glazing mode was applied, and the obtained results in Figure 9(a) confirm the presence of Cr2O3 (ICDD 82-1484) and MnCr2O4 (ICDD 75-1614) for all studied samples. The XRD analysis using a regular mode was also performed. The obtained results in Figure 9(b) show that Cr2O3 (ICDD 82-1484) and MnCr2O4 (ICDD 75-1614) were still detected, though with weak signals for the latter one. The presence of the austenite and ferrite phases was confirmed. It was found that, despite being exposed at 700°C for 10 cycles, the more intense (200) austenite peaks were still detected for the welds produced using hydrogen-added shielding gases, Ar–7% N2–1.5% H2 and Ar–7% N2–5% H2, compared with the welds produced using the shielding gas without hydrogen.

Figure 9 XRD patterns of the welds after the oxidation for 10 cycles characterised using the (a) grazing incidence and (b) regular modes.
Figure 9

XRD patterns of the welds after the oxidation for 10 cycles characterised using the (a) grazing incidence and (b) regular modes.

The addition of hydrogen also significantly increased the volume of molten material in the weld pool due to the higher thermal conductivity of argon–hydrogen mixtures at temperatures at which hydrogen molecules dissociate.

4 Discussion

4.1 Weld microstructure

From the macroscopic observation of the weldment in Figure 2, the addition of hydrogen in the Ar–7% N2 shielding gas tends to yield a weld metal with larger volume, evidently for the weld metal produced using Ar–7% N2–5% H2. This observation is in agreement with the study by Tusek [43], which applied the GTAW on AISI 304 stainless steel and found that the addition of hydrogen content in Ar shielding gas resulted in a large volume of weld metal. The mixing of hydrogen in Ar shielding gas for the GTAW of AISI 316L stainless steel plate was also found to result in a greater volume of the weld metal [44,45,46,47]. At a room temperature of 27°C, the thermal conductivity of hydrogen was found to be higher than that of Ar and N2 by about ten and seven times, respectively [48]. Under the welding conditions, it was also reported that the thermal conductivity of the Ar–H2 mixture was higher than that of Ar [23], consequently resulting in more heat input to the weld pool [49], and thus providing a larger volume of the melt metal, which ultimately presented as weld metal.

As for the microstructure of the weld metal, though the AISI 304L stainless steel is an austenitic stainless steel when it is welded and solidifies, the primary delta ferrite and secondary austenite could be formed during cooling [50,51], as shown in Figure 3. For the weld produced using the Ar–7% N2 shielding gas, the delta ferrite content was 2.1%. It was reported that the presence of less than 3% delta ferrite made the weld susceptible to a solidification crack due to the formation of an intergranular liquid film by sulphur present in the steel during solidification [52]. A higher delta ferrite content, in the range of 3–10%, was recommended to avoid a solidification crack because sulphur was more dissolved in ferrite than in austenite [52]. In the present work, it was found that the mixing of 1.5 and 5% hydrogen gas in the Ar–7% N2 shielding gas could increase the delta ferrite content to about 4–7%, making it safe from a solidification crack. Phakpeetinan et al. [28] reported that the addition of hydrogen in Ar–N2 shielding gas could help reduce the dissolution of nitrogen in the welds and consequently reduce the formed delta ferrite, as also shown in Figure 5. To relate the weld structure to its corrosion property, Singh Raman et al. [52] and Mittal and Sidhu [53] suggested that the formation of delta ferrite could degrade the high-temperature oxidation resistance of the weldment since Cr could leave the matrix for the ferrite phase, resulting in the poor oxidation resistance of the alloy matrix. However, the present results showed that the weld that contained the higher delta ferrite content did not exhibit poorer oxidation resistance but rather improved it. To explain this result, we first checked the Cr content in the austenite matrix by EDS and found that this value was slightly decreased when Ar–N2–H2 was used as a shielding gas instead of Ar–N2, i.e. the Cr content was 20.3% with the standard deviation of 0.6 for the weld produced using the Ar–7% N2 shielding gas while it was 19.9% with a standard deviation of 0.2 for the welds produced using the Ar–7% N2–5% H2 shielding gas. As a result of the almost unchanged Cr content in the matrix of the studied welds, the adverse effect on the weld oxidation rate was not observed. The reason for the improved oxidation resistance should be related to the mixing of hydrogen gas in the shielding gas, which will be explored in the following section.

4.2 Hydrogen effect on the weld oxidation rate

When the shielding gas containing hydrogen was applied during welding, the hydrogen could be dissolved into the steel, as observed by the change in the XRD patterns of the matrix. As shown in Figures 6 and 9(b), the XRD peak of (200) austenite becomes the majority with respect to the (111) peak for the welds produced in hydrogen-added shielding gas. Similar XRD patterns, which showed the shift of the (111) peak from a majority to a minority, were also found during cathodic hydrogen charging in AISI 304 stainless steel [54]. Chew and Willgoss [55] reported that the concentration of hydrogen dissolved in the steel welded by GTA was proportional to the square root of the partial pressure of hydrogen gas, i.e. obeying Sievert’s law. The hydrogen dissolution reaction may then be described by reaction (2) with Sievert’s law (equation (3)):

(2) 1 2 H 2 = H ̲ ,

(3) K 2 = a H ̲ p H 2 1 / 2 ,

where H ̲ is the hydrogen dissolved in the steel, a H ̲ is the activity of H2, and p H 2 is the partial pressure of H2 [15]. During the oxidation test, the dissolved hydrogen in the steel might be released from the substrate and diffused into the oxide scale [56]. Kofstad [33,34,35] and other researchers [36,57] suggested that hydrogen gas in the atmosphere could dissolve into the oxide as hydrogen interstitial with a unit effective charge ( H i ⋅ ) . Thus, it was suggested that the dissolved hydrogen in the steel dissolved into the oxide by reaction (4) with an equilibrium constant in equation (5) [56]:

(4) H ̲ = H i ⋅ + e ′ ,

(5) K 4 = [ H i ⋅ ] [ e ′ ] a H ̲ .

The presence of the hydrogen interstitial defect could affect the defect structure of the oxide. As for the structure of defects of chromia thermally grown on Cr, Kurokawa et al. [58] expressed the parabolic rate constant as a function of the chemical potential of oxygen gas and temperature. We have recently proposed a solution to explicitly express the parabolic rate constant, k p , according to equation (1), as a function of the oxygen partial pressure and temperature [40,59]. The expression shows that the total k p is composed of two terms: the first contributed from chromium interstitial diffusion ( k pi ) equalling 8 α 2 A ( p O 2 ,int . − 3 / 16 − p O 2 − 3 / 16 ) , and the second contributed from the chromium vacancy diffusion ( k pv ) equalling 8 α 2 B ( p O 2 3 / 16 − p O 2 ,int . 3 / 16 ) [40,59], where α is the ratio of the mass gain to the oxide thickness [40,59], A and B are constants in the diffusion coefficient in the form of A p O 2 − 3 / 16 + B p O 2 3 / 16 [60], and the subscript int. indicates that the interested value is the one at the interface between Cr and Cr2O3. At a fixed oxygen partial pressure of 0.21 bar, the log k p vs ( 1 / T ) diagram was plotted to evaluate the contribution of each term, k pi and k pv , to the total k p , as shown in Figure 10(a). The diagram was constructed using thermodynamic data from Kubaschewski et al. [61] for the calculation of the equilibrium dissociation pressure of Cr2O3, and using the diffusivity of chromia thermally grown on a single crystal of Cr given by Yurek et al. [58,60]. It can be seen that k pv is dominant at relatively high temperatures, i.e. higher than 1,048°C, and the total k p value tends to coincide with k pv . Thus, the major defect in this temperature range should be the chromium vacancy with the effective charge of −3. On the contrary, k pi is dominant at temperatures lower than 1,048°C, e.g. the studied temperature of 700°C in this work, and the log k p line approaches the log k pi line, indicating the dominance of the chromium interstitial defect. At a fixed temperature of 700°C, the log k p vs log p O 2 diagram was also plotted, as shown in Figure 10(b). Similar to the situation at a slightly higher temperature of 800°C reported in the literature [40,58,62], k pi is found to be dominant over a wide range of the oxygen partial pressure and significantly higher than k pv . For example, k pi is about 105 times larger than k pv at an oxygen partial pressure of 10–10 bar. As a result, the total k p is approximately equal to k pi . This analysis indicates that the dominant defects in a single crystal of Cr grown thermally on Cr at 700°C in oxygen at an oxygen pressure of 0.21 bar at Cr2O3 dissociation are chromium interstitials with an effective charge of 3. Notably, the calculated k p is significantly lower than the one obtained from the experiment, which should be from the grain boundary diffusion of defects at relatively low temperatures, e.g. 800°C [58,59,62] or lower.

Figure 10 Calculated parabolic rate constant for a single crystal of chromia thermally grown on Cr (a) as a function of temperature at an oxygen pressure of 0.21 bar, and (b) as a function of oxygen partial pressure at 700°C.
Figure 10

Calculated parabolic rate constant for a single crystal of chromia thermally grown on Cr (a) as a function of temperature at an oxygen pressure of 0.21 bar, and (b) as a function of oxygen partial pressure at 700°C.

In the case of stainless steel, Roy et al. [63] proposed that the defects responsible for the growth of the oxide scale on AISI 441 stainless steel at 900°C were also chromium interstitial, but additionally with manganese interstitial for the Mn–Cr spinel formation and oxygen vacancy. The relevant defect reactions with their equilibrium constants were reported [40,63,64,65,66], as summarised in equations (6)–(11):

(6) 2Cr i ⋅ ⋅ ⋅ + 3 2 O 2 + 6 e ′ = 2 Cr Cr + 3 O O ,

(7) K 6 = 1 [ Cr i ⋅ ⋅ ⋅ ] 2 [ e ′ ] 6 p O 2 3 / 2 ,

(8) Mn i ⋅ ⋅ + Cr 2 O 3 + 1 2 O 2 + 2 e ′ = MnCr 2 O 4 ,

(9) K 8 = 1 [ Mn i ⋅ ⋅ ] [ e ′ ] 2 p O 2 1 / 2 ,

(10) V O ⋅ ⋅ + 1 2 O 2 + 2 e ′ = O O ,

(11) K 10 = 1 [ V O ⋅ ⋅ ] [ e ′ ] 2 p O 2 1 / 2 .

If the hydrogen dissolved in the steel is injected into the oxide and turns reaction (4) dominant, the electroneutrality condition will be as follows [29,37]:

(12) [ H i ⋅ ] = [ e ′ ] .

By inserting equation (12) into equations (7), (9), and (11), the concentrations of defects can be described as a function of the hydrogen interstitial concentration and oxygen partial pressure as follows:

(13) [ Cr i ⋅ ⋅ ⋅ ] = 1 K 6 1 / 2 [ H i ⋅ ] 3 p O 2 3 / 4 ,

(14) [ Mn i ⋅ ⋅ ] = 1 K 8 [ H i ⋅ ] 2 p O 2 1 / 2 ,

(15) [ V O ⋅ ⋅ ] = 1 K 10 [ H i ⋅ ] 2 p O 2 1 / 2 .

Based on equations (13)–(15), the authors constructed the Brouwer diagrams, as shown in Figure 11, for the situation when the hydrogen dissolved in the steel diffuses to the oxide, thus producing the hydrogen interstitial and turning the electroneutrality condition to [ H i ⋅ ] = [ e ′ ] . As seen in the figure, at a given oxygen partial pressure, the concentrations of Cr i ⋅ ⋅ ⋅ , Mn i ⋅ ⋅ , and V O ⋅ ⋅ decrease when more hydrogen defect is injected to the oxide. The increased concentration of hydrogen defect in the oxide is due to the higher activity of hydrogen dissolved in the steel, i.e. [ H i ⋅ ] ∝ a H ̲ 1 / 2 obtained from equations (12) and (5). The hydrogen activity in the steel is increased because of the mixing of hydrogen gas in the shielding gas, i.e. a H ̲ ∝ p H 2 1 / 2 , from equation (3). This means that increasing the content of hydrogen gas in the shielding gas can result in an increase in the hydrogen defect in the oxide, i.e. [ H i ⋅ ] ∝ p H 2 1 / 4 , which consequently reduces the concentrations of Cr i ⋅ ⋅ ⋅ , Mn i ⋅ ⋅ , and V O ⋅ ⋅ defects. Because these defects are the reactants for the oxide formation according to reactions (6), (8), and (10), the decreases in their concentrations consequently help reduce the oxide formation rate, as observed in the present work.

Figure 11 Brouwer diagrams for Cr i ⋅ ⋅ ⋅ {\text{Cr}}_{\text{i}}^{\cdot \cdot \cdot } , Mn i ⋅ ⋅ {\text{Mn}}_{\text{i}}^{\cdot \cdot } and V O ⋅ ⋅ {\text{V}}_{\text{O}}^{\cdot \cdot } when [ H i ⋅ ] = [ e ′ ] {[}{\text{H}}_{\text{i}}^{\cdot }]={[}e^{\prime} ] .
Figure 11

Brouwer diagrams for Cr i ⋅ ⋅ ⋅ , Mn i ⋅ ⋅ and V O ⋅ ⋅ when [ H i ⋅ ] = [ e ′ ] .

Notably, the Brouwer diagrams in our recent study that exhibited the role of water vapour in the O2–H2O atmosphere on concentrations of defects responsible for stainless steel oxidation were constructed [40]. The defects considered in that case are similar to the ones treated here, but the injection of a hydrogen defect into the oxide by water vapour is based on reaction (16) [40]. By combining that reaction with the equilibrium reaction among hydrogen gas, oxygen gas, and water vapour, it is possible to obtain the reaction showing the relationship between the hydrogen interstitial, electron, and hydrogen gas, as in reaction (17). This reaction is the same as the one obtained from the summation of reactions (2) and (4) in the present work. This indicates that while the present study concerned the oxidation of weld metals produced using the Ar–N2–H2 shielding gas and the previous one focussing on the oxidation in O2–H2O [40] are physically different, the dependence of concentrations of defects potentially responsible for the oxide growth on the hydrogen interstitial concentration in both cases is mathematically identical. The only difference is that the hydrogen interstitial in the previous study [40] was from water vapour in an oxidising atmosphere, while the hydrogen defect in the present work was from the hydrogen dissolved in the steel, which was due to the use of hydrogen gas as a component in the shielding gas. Thus, the present study confirmed the possible role of hydrogen dissolved in the oxide on reducing the concentrations of defects potentially responsible for stainless steel oxidation and, therefore, slowing the oxidation rate, as observed in the present work and also in previous work for oxidation in humidified oxygen [40]:

(16) 1 2 H 2 O = H i ⋅ + 1 4 O 2 + e ′ ,

(17) 1 2 H 2 = H i ⋅ + e ′ .

5 Conclusion

AISI 304L stainless steel was welded using the GTAW technique using Ar–7% N2 without and with 1.5 and 5% H2 as the shielding gases. In terms of the microstructural aspect, the austenite matrix of all studied weld metals was embedded with delta ferrite. The addition of 1.5 and 5% H2 helped increase the ferrite content from 2.1% for the weld metal produced using the Ar–7% N2 shielding gas to 3.8–6.9%, thus reducing the risk of a hot crack. The increased percentage of ferrite was accompanied by a lower content of an austenite stabiliser, i.e. nitrogen dissolved in the weld metals produced using the shielding gas mixed with H2. For the high temperature corrosion behaviour, the oxidation of all studied weld metals in synthetic air at 700°C followed a parabolic rate law with the reduced rate constant for the weld metals produced using the Ar–7% N2 shielding gas mixed with H2. When the H2 content was increased from 1.5 to 5% in the Ar–7% N2 shielding gas, the rate constant was further reduced. These findings suggest the beneficial role of mixing H2 in a shielding gas in improving the corrosion resistance of weld metals for high-temperature service.

Acknowledgments

This research was funded by the National Science, Research, and Innovation Fund (NSRF) and King Mongkut’s University of Technology North Bangkok with Contract No. KMUTNB-FF-65-31.

  1. Funding information: The National Science, Research, and Innovation Fund (NSRF) and King Mongkut’s University of Technology North Bangkok with Contract No. KMUTNB-FF-65-31.

  2. Author contributions: Sompong Chueaprakha: writing-original draft, writing-review & editing, investigation, data curation, and visualisation; Thammaporn Thublaor: writing-original draft, writing-review and editing, formal analysis, validation, and visualisation; Thamrongsin Siripongsakul: writing-original draft, formal analysis, visualisation, resources, and funding acquisition; Panya Wiman: formal analysis; Walairat Chandra-ambhorn: writing – review and editing; Somrerk Chandra-ambhorn: conceptualisation, methodology, writing – original draft, writing – review and editing, formal analysis, project administration, and supervision.

  3. Conflict of interest: Authors state no conflict of interest.

  4. Data availability statement: The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.

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Received: 2025-04-29
Revised: 2025-07-28
Accepted: 2025-08-13
Published Online: 2025-09-09

© 2025 the author(s), published by De Gruyter

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https://doi.org/10.1515/htmp-2025-0095
Keywords for this article
welding; hydrogen; oxidation; stainless steel; Brouwer diagram
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