# Effect of Nitrogen Addition to Shielding Gas on Cooling Rates and in the Microstructure of Thin Sheets of Duplex Stainless Steel Welded by Pulsed Gas Tungsten Arc Welding Process

Evandro Giuseppe Betini, Maurilio Pereira Gomes, Cristiano Stefano Mucsi, Marcos Tadeu D'azeredo Orlando, Temístocles de Sousa Luz, Marie-Noëlle Avettand-Fènoël, Jesualdo Luiz Rossi

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Evandro Giuseppe Betini, Maurilio Pereira Gomes, Cristiano Stefano Mucsi, Marcos Tadeu D'azeredo Orlando, Temístocles de Sousa Luz, et al. Effect of Nitrogen Addition to Shielding Gas on Cooling Rates and in the Microstructure of Thin Sheets of Duplex Stainless Steel Welded by Pulsed Gas Tungsten Arc Welding Process. Materials Research, 2019, Materials Research, 22 (suppl 1), 10.1590/1980-5373-mr-2019-0247. [hal-03199271](https://hal-univ-lille.fr/hal-03199271)

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# Effect of Nitrogen Addition to Shielding Gas on Cooling Rates and in the Microstructure of Thin Sheets of Duplex Stainless Steel Welded by Pulsed Gas Tungsten Arc Welding Process

Evandro Giuseppe Betiniª* , Maurílio Pereira Gomesª, Cristiano Stefano Mucsiª, Marcos Tadeu

D'Azeredo Orlandoᵇ, Temistocles de Sousa Luzᵇ, Marie-Noëlle Avettand-Fénoëlᶜ, Jesualdo Luiz Rossiª,ᶜ

Received: March 16, 2019; Revised: June 03, 2019; Accepted: June 24, 2019

The effect of the nitrogen content in the shielding gas and its effect on temperature distributions at the welded zone of thin sheets of duplex stainless steel have been evaluated. The duplex stainless steels have many features due to unique microstructural combination of austenite and ferrite grains. The phase balance can be easily shifted depending on the welding parameters. Two sheets were welded using pure argon and pure argon plus 2% of nitrogen as shielding gas. The thermal profile had shown that N₂-supplemented shielding gas lead to higher peaks of temperature using similar welding parameters. Microstructural examination showed that the austenite phase in the weld increased with the presence of nitrogen in the shielding gas. The added nitrogen promoted primary austenite formation and slightly increases the microhardness at the solidified zone. Microhardness mapping and metallographic imaging presented information about microstructures, confirming the formation of secondary phases during thermal cycle in the temperature range 850 °C and 950 °C. Control of ferrite amounts in the welds is essential mainly to improve mechanical properties and corrosion resistance of welding zones.

**Keywords:** duplex stainless steel, cooling rates, nitrogen, GTAW, thermal cycle, microhardness.

## 1. Introduction

Duplex stainless steels (DSS) have basically a microstructure consisting of austenite and ferrite. In general, DSS has around 50% of each phase in the base material but typically a variation of 35 to 60% in ferrite content results in acceptable properties¹,².

The main commercial factor pushing the advance, the development and use of duplex stainless alloys was the increase of activity in the offshore oil industry which demanded a stainless steel material to handle aggressive environments and new steel production techniques². The new technique using addition of nitrogen and lowering of carbon content made it possible to produce much cleaner steels improving high temperature stability as shown by duplex structure¹,⁵. Thus, the DSS alloys have become a suitable alternative to austenitic stainless steels due to their special combination of good mechanical properties and corrosion resistance and lower costs¹,²,⁶. The nuclear industry, especially requests alloys with high thermal stability in pipes for power generation systems and safe transportation equipment's for radioactive material⁷,⁹. DSS also presents satisfactory weldability, however it is possible to occur the formation of excessive ferrite, nitrides and intermetallic phases in the solidifying zone (SZ) and in the heat affected zone (HAZ)¹,²,¹⁰. These formations of deleterious phases need to be avoided and a balanced ferrite / austenite microstructure must be maintained¹¹.

The GTAW (gas tungsten arc welding) processes are generally preferred because they produce very high-quality welds²,³,¹². In many cases distortion is the major problem in welding of thin sections. However, the distortion is controlled in pulsed and magnetic arc oscillation GTAW process¹². Weld root quality of stainless-steel pipe and tubes can be ensured by using an inert purging gas³,⁵,¹¹. The pure Ar is the most commonly used purging gas in GTA welding of several grades of steels⁷,¹¹. However, 2% nitrogen is recommended in pure argon mixture for shielding gas and optimizing the phase balance, impact toughness and corrosion resistance of DSS¹³. Taking all aspects into account, Ar + 2%N₂ is the preferred shielding gas for welding duplex and super duplex with the GTAW process⁶,⁹.

Several authors have been mentioning the important role of N₂ in the microstructure evolution by being the strongest austenite forming alloying element⁶,⁹,¹⁰,¹³⁻¹⁷. Moreover, N₂ is an economical and efficient substitute for nickel¹⁰,¹⁷. Generally, DSS has a low bulk nitrogen content standardized in the range 0.15% - 0.35%. Moreno et al.¹⁷ determined using electron probe microanalysis (EPMA) technique an average content of N₂ in the range 0.045% - 0.070%wt in ferrite and 0.520% - 0.770%wt in austenite for five different DSS alloys. The addition of C and N₂ strengthens both ferrite and austenite by dissolving at interstitial sites in the solid solution²,¹⁶,¹⁷,¹⁸.

<footnote>ªInstituto de Pesquisas Energéticas e Nucleares (IPEN), São Paulo, SP, Brasil</footnote>

<footnote>ᵇDepartamento de Engenharia Mecânica, Universidade Federal do Espírito Santo, Vitória, ES, Brasil</footnote>

<footnote>ᶜUnité Matériaux et Transformations, Université de Lille, Cité Scientifique, Villeneuve d'Ascq, França</footnote>
During welding process, the austenite dissolves while being heated, it is followed by grain growth in the delta ferrite region, and finally restoring to the austenitic phase during cooling¹⁹,²⁰. However, the loss of nitrogen and a decrease of austenite content during GTAW welding would lead to toughness reduction¹⁰,¹⁷. Hosseini et al.¹⁰ monitored economical content and microstructure for each weld pass in super duplex stainless steel, by using wavelength dispersive X-ray spectrometry; It was observed that after four passes a loss of nitrogen from 0.28% in the base metal - BM to 0.17% (in mass) for low heat input. Karlsson et al.²¹ using low heat input autogenously laser welding on 2507 SDSS observed highly ferritic microstructures with nitride precipitates due to nitrogen loss at high cooling rates. During the thermal cycle, the alloy is heated to a very high temperature above 1350 °C for a short time, held for a few seconds and subsequently cooled to room temperature²¹,²². Later on, in the cooling stage, the austenite re-precipitates around ferrite grain boundaries²³. In this way, adjustment of a low cooling rate and proportion of N₂ allows adequate austenite formation, but not too low cooling rate in order to avoid precipitation of intermetallic phases¹⁸,²²⁻²⁴. Westin²⁴ reported that high ferrite contents resulted in Cr-rich nitrides formation in HAZ and SZ due to large ferrite grain supersaturated in nitrogen. If the gas protection is insufficient during welding or if the nitrogen is added to the shielding gas, the formation of chromium nitrides caused by an uptake of nitrogen from the atmosphere may occur²³.
Nitrogen is added to argon shielding gas in duplex stainless-steel welds to preserve weld metal nitrogen levels¹⁵,¹⁶. According to Zhang et al.¹⁵, there are two ways to compensate for the nitrogen loss during welding: one is directly adding the nitrogen to the BM and filler wires in the originally producing supplies¹³,¹⁷,²³ or a proper proportion of nitrogen gas is added into shielding gas and backing gas during welding¹¹,¹⁴,¹⁶,²⁴⁻²⁸.

The main goal in the present work was to assess the effects of the nature of the nitrogen amount in the shielding gas (2% N₂ in Ar) on cooling rates, microstructural properties, microhardness and grain size of 1.8 mm thickness sheets of UNS S32304 duplex stainless-steel autogenously welded by pulsed GTAW process. The amounts of phases were evaluated by microstructural image phase analysis and electron probe microanalysis (EPMA) technique was used to measure the N₂ contents in transversal profile of welds.

## 2. Materials and Methods

The specimens were obtained with dimensions of 72 x 72 x 1.8 mm³. Table 1 shows the chemical composition of used UNS S32304 duplex stainless steel. The alloy's chemical composition was certified by the commercial supplier Aperam South America S/A.

The microstructure of the ferrite and austenite was observed by optical microscopy (OM). The specimens for microstructure investigation were ground successively from P300 to P2000 grit emery paper and polished using diamond pastes of 6 µm, 3 µm and 1µm. With the purpose of characterization of various types of austenites in the welding zones, the as-polished specimens were etched in Behara I solution for 15-20 s. The optical microstructures of the SZ and HAZ were observed by means of the Image J® software with bright light micrographs according to ASTM E562²⁹ standard to determine the phases ratios. The free application 3D Builder® was used to better illustrate dual phase proportion in base metal and solidified zone.

Thin samples were clamped in a welding table in order to reproduce the conventional set welding process. The DSS sheets were welded using GTAW process with pulsed current and direct polarities were used with automatic drive systems. The batch of samples was welded using two commercial kinds of pure argon and a mixture of argon and 2% of nitrogen shielding gas. The gas flow rate in both cases was 10 L.min⁻¹. The schematic diagram of the welding process is shown in Figure 1.

<figure><img src="image_3.png" /><figcaption>**Figure 1.** Schematic impression of the GTAW process on thin plates of duplex stainless steel UNS S32304.</figcaption></figure>

<table><caption>**Table 1.** Chemical composition (% mass and ppm) of the duplex stainless steel UNS S32304 used in the present study.</caption><thead><tr><th>Cr</th><th>Ni</th><th>Mo</th><th>Mn</th><th>Si</th><th>C</th><th>P</th><th>S</th><th>Ti</th><th>Cu</th><th>Co</th><th>N<sub>ppm</sub></th><th>Fe</th></tr></thead><tbody><tr><td>22.210</td><td>3.521</td><td>0.255</td><td>1.402</td><td>0.250</td><td>0.016</td><td>0.023</td><td>0.001</td><td>0.004</td><td>0.417</td><td>0.091</td><td>1030</td><td>Bal.</td></tr></tbody></table>
The AWS Class EWTh-2 electrode was kept negative and was located at 2 mm with a 60° angle from the sheet, according to direct polarity process. The electrical current and voltage values were measured for each welding step. In the case of the pulsed current, the wave was maintained in balance with background time and pulse values equal to 0.9 s. Table 2, indicates the welding and thermal cycle parameters.

Regarding the parameters such as voltage (U), pulsed and background current ($I_p, I_b$), pulsed and background time ($t_p, t_b$), the welding speed (v) along with 60% arc efficiency for the pulsed GTAW welding, it was possible to calculate the welding heat input (H) per mm using equation 1¹². The welding heat input shown in Table 2 was determined using the average voltage, current and welding speed.

$$H = \frac{60.U.(I_p.t_p + I_b.t_b)}{(t_p + t_b).v} \quad (1)$$

The thermal cycles were measured during welding by inserting thermocouples type K and S in the UNS S 32304 DSS sheet, in small drilled holes and inserting the thermocouple tip in this hole. The temperature was measured and recorded using a data acquisition system. Four thermocouples were positioned and fixed to the sample sheet's surface at different distances, along transversal and longitudinal lines of the weld bead. The thermocouples were located at 0 mm, 3 mm and 4.5 mm from the joint line.

The signals from the thermocouples were acquired by 8-channel universal data acquisition system amplifier using MX board - PT1000 for room temperature automatic conditioning. The measured total error limit at 300 K ambient temperature was ± 1 K and the temperature drift (K type) was used K/10K ratio where the uncertainty was $\le \pm 0.5$.

An electron probe micro-analyzer (EPMA) was used to quantitatively evaluate the nitrogen in the welded regions. This electronic microprobe is equipped with four WDS (wavelength dispersive spectrometer) spectrometers. For the point quantifications it was used voltage acceleration of 10 kV and beam current of 100 nA.

The crystal used to detect the $K\alpha$ of N was PC2. The $N_2$ profile was composed of EPMA-WDS point analyses distributed over a 4300 µm line for sample welded using

pure Ar and a 4500 µm line for welding using Ar+2%$N_2$ shielding gas, crossing the base metal, heat-affected zone till center of solidified zone.

Vickers microhardness was performed for both welds using 0.5 N of load for 15 s (HV₀.₀₅). A microhardness tester was used to determine microhardness maps on the cross sections of the welded sheets. The thickness of the sheets is 1.8 mm. In the total, five lines with 140 hardness points for 60 µm steps were performed for each sample passing through the base metal, ZAC and SZ.

## 3. Results and Discussion

In order to study thermal cycles, the temperature distributions were plotted for both specimens as shown in Figure 2. First, the sheet welded using pure Ar as shielding gas is shown in Figure 2a. The thermocouples TK1 and TK2 (4.5 mm away from the welding line) were placed closer to the weld bead where the peaks of temperature and cooling rate were close to 543 °C with 10.6 °C.s⁻¹ and 556 °C with 12.1 °C.s⁻¹, respectively. In this region of the sheet (HAZ), the holding time were maintained between 500 °C and 600 °C for 6 s, approximately, for both TK1 and TK2 thermocouples that leads to precipitations of secondary phases as $\alpha'$, $\sigma$, $\epsilon$ and $\pi^{15,22,25}$. The TK3 thermocouple presented a peak of temperature around 760 °C at 3 mm away from the center of welding line. This value is higher than that obtained in other thermocouples for the HAZ which was measured a cooling rate of 53.2 °C.s⁻¹. The welding process with high cooling rate lead to ferritization with coarse grains of ferrites in the welding zone³⁰.

The Figure 2b shows the thermal cycles for welding with Ar+2%$N_2$ as shielding gas. The thermocouples TK1 and TK2 (4.5 mm distant from the melting pool) presented peaks of temperature and cooling rates near to 759 °C - 25.2 °C.s⁻¹ and 773 °C - 26.5 °C.s⁻¹, respectively. The holding time were maintained above 700 °C for 4 s which would lead to possible intermetallic phases as $M_{23}C_{6}$, $\gamma_2$, $\chi$ and $\sigma^{22,30}$. For TK1 and TK2 also were subjected to holding time above 500 °C for 15 s approximately that would lead to higher precipitation of secondary phases such as $\alpha'$, $\sigma$ and $\pi^{15,22,30}$.

<table><caption>**Table 2.** Welding parameters used for the pulsed GTAW process.</caption><thead><tr><th>Sample</th><th>#6</th><th>#7</th></tr></thead><tbody><tr><td>Shielding gas (10 L.min<sup>-1</sup>)</td><td>Pure Ar</td><td>Ar+2%N<sub>2</sub></td></tr><tr><td>Voltage (V)</td><td>11</td><td>11</td></tr><tr><td>Pulse current - I<sub>p</sub> (A)</td><td>150</td><td>140</td></tr><tr><td>Background current - I<sub>b</sub> (A)</td><td>80</td><td>70</td></tr><tr><td>Pulse time - t<sub>p</sub> (ms)</td><td>0.9</td><td>0.9</td></tr><tr><td>Background time - t<sub>b</sub> (ms)</td><td>0.9</td><td>0.9</td></tr><tr><td>Welding speed (cm.min<sup>-1</sup>)</td><td>35</td><td>35</td></tr><tr><td>Arc efficiency (%)</td><td>60</td><td>60</td></tr><tr><td>Heat input (kJ.mm<sup>-1</sup>)</td><td>0.20</td><td>0.17</td></tr></tbody></table>
Through the TK3 thermocouples it is shown a temperature peak and cooling rate around 947 °C and 27.5 °C·s⁻¹ at 3 mm from the center of the solidified zone. Through the thermal cycle it is also possible to suggest that the beneficial effect of the addition of 2% N₂ on the stabilization of the temperature distribution during cooling process. This stabilization can be seen from 56 s at time on the graph shown in Figure 2b, and it can be observed that all curves are dropping at the same slope. This curve behavior could not be observed for the sample welded using pure Ar (Figure 2a). Lin et al.³¹, also showed that the peak temperature of a thermal cycle increased with increasing N₂ content as more heat is carried into shielding gas. Thus, increasing the N₂ content more heat is carried into the work piece and increases the weld metal area³². In addition, according to the study by Tseng and Chou¹², using pulsed GTAW process higher ratio amplitude can reduce the temperature difference between the melting zone and the unaffected base metal in welding and therefore the welding residual stress can be reduced.

The Figure 3 shows the temperatures obtained on the bottom of the center of welding zone (0 mm) and HAZ (3 mm and 4.5 mm) during passage of the welding. In these points it was obtained values in the order of 1173 °C for welding with pure Ar and 1221 °C for welding with Ar + 2% N₂, which suggests an influence of the shielding gas on the temperature distribution of welding. In addition, all welding points measured with the mixture of Ar + 2% N₂ presented higher temperature than welding with pure Ar.

Figure 4 shows an optical micrograph of the microstructure that consisted of austenite and ferrite for the as received material. The ferrite phase is shown as dark, while the light region is austenite phase (see Figure 4a). In the continuous ferrite matrix is embedded island-like austenite phase without apparent secondary phase precipitation. Both grains are elongated along the rolling direction. Figure 4b shows the quantitative evaluation of the content of both ferrite and austenite phases. Based on microstructural investigations and using image analyzer software the amounts of phases in different zones of the joints were estimated. Up to five OM images were collected for each zone (SZ, HAZ and BM) and then were processed.

<figure><img src="image_4.png" /><figcaption>Figure 3. Peaks of temperature for thermal cycles of welding.</figcaption></figure>

In order to detail the modifications in the biphasic microstructures and content variations of phases due to the welding conditions, optical micrographs were obtained, and the separation of the ferrite and austenite phases was attained with image editing effect.

Micrographs for both samples are shown in Figure 4 using the modified Behara I reagent to reveal the etched microstructure. The chemical etching resulted in a blue coloration for ferrite and a yellow coloration for austenite as reported in the literature¹⁵,³³,³⁴.

The Figure 4 b shows the micrograph of the solidified zone for welded sheet using pure Ar as shielding gas. The Behara I reagent also resulted in dual effect of dark colors (blue) for ferrite and clear (yellow) for the austenite microstructure.

In this way, through the chemical etching it was possible to observe the ferrite morphology separately from the austenite as well as the austenite fraction phase resulting from the welding process with pure Ar as shielding gas. Besides that, secondary austenite was observed quite dispersed inside the ferrite coarse grains.

The microstructure of the solidified zone of the sample welded with Ar + 2% N₂ as the shielding gas using purge gas is presented Figure 4c. For the identification of ferrite and austenite phases, the same modified Behara I etching was used.

<figure><img src="image_5.png" /><figcaption>Figure 2. Distribution of temperatures for samples (a) pure Ar and (b) Ar + 2% N₂ with pure argon as purging gas. It shows the temperature variation measured with each thermocouple for the UNS 32304 duplex stainless steel, 1.8 mm long stripe, while being welded.</figcaption></figure>
<figure><img src="image_6.png" /><figcaption>Figure 4. Microstructure of the transversal profiles of the both samples welded with a) pure Ar and b) Ar+2%N₂ as shielding gas.</figcaption></figure>

For ferrite resulting from the nitrogen addition in shielding gas, it is shown as a dark (brown) color and the voids that are occupied by the austenite. For austenite shown in bright light (yellow) color, this phase is more evident in the regions at the grain boundary of ferrite and less dispersed than the sample welded with pure Ar (see Figure 4b). Thus, the separation of the phases as presented in Figure 4, shows the relationship with the changes of the austenite fractions as described in the diagram showed in Figure 5.

Figure 5 shows the evolution of the austenite volumetric fraction regarding the composition of the shielding gas in the SZ, HAZ and BM. The percentages of austenite are significantly affected by the addition of nitrogen in the shielding gas and this leads to changes in the total austenite content after the welding process. About N₂ supplementation in the shielding gas, the increase of the austenite content in the SZ was shown: 22% to 37% and HAZ: 23% to 25.3%, which agrees with the results of Tahaei et al.³⁵ and Zhang et al.,¹⁵.

The addition of N₂ in the shielding gas significantly increases the amount of austenite for the DSS and GTAW welds. Zhang et al.¹⁵ reported that N₂ in the shielding gas mainly promotes the formation of primary austenite, suppressing Cr₂N precipitation at the root of the weld. According to Westin et al.²⁸, nitrogen can increase the temperature of ferrite transformation into austenite. Consequently, the higher content of nitrogen makes the alloy less sensitive to the fast cooling rates due to more efficient austenite reform and delays the ferritization²⁸. As seen in the present work, the microstructural changes in DSS that occur during welding can lead to the variation of the phase balance in SZ and HAZ.

<figure><img src="image_7.png" /><figcaption>Figure 5. The effects of shielding gas composition on the austenite content of the HAZ and solidified zone in DSS.</figcaption></figure>

In Figure 6, the microstructures of the UNS S32304 duplex alloy are shown in three different orthogonal directions. The first image, Figure 6a, shows the base metal in the three directions with indication of the rolling directions. No intermetallic phase was observed, both for samples as received and for the base metal.

The evolution of the microstructure of ZTA and SZ is strongly dependent on the base metal microstructure, peak temperature, residence time at high temperatures and cooling rates¹³. DSS typically have δ ferrite, primary austenite γ₁, secondary austenite γ₂, σ, χ, carbides and nitrites. For the solidified zone of DSS, the austenite phases are generally formed from ferrite in different modes: allotriomorphic austenite in the borders of ferrite grains (GBA), intragranular austenite (IGA) and austenite Widmanstätten (WA) within the grains³⁶.
For Figure 6b and Figure 6c, the effects of GTAW welding with 99.9% Ar as shielding gas and a mixture of Ar + 2% N₂ in three orthogonal directions are presented. The microstructure observed in both welded samples consists of large ferrite grains with intragranular austenite, grain border austenite and Widmanstätten austenite (γ<sub>w</sub>). For the solidified zone, with Ar + 2% N₂ (Figure 6c), elongated and thicker austenite grains are evident.

Figure 7 shows the cross-section profile of nitrogen amounts in mass%, for sheet welded with pure Ar. As reported in the literature, the loss of nitrogen in ZTA and SZ are unavoidable in GTAW welding using pure Ar as shielding gas<sup>13,15,37</sup>. Hosseini et al.<sup>37</sup> obtained a cross-section profile of nitrogen in superduplex stainless steel subjected to a GTAW arc by WDS analysis. The author still considers the amount of nitrogen in the solidified zone as virtually zero. As seen in Figure 7, a large number of point measurements of nitrogen within the ZTA and SZ are zero. The loss of nitrogen is accompanied by the reduction of the amount of austenitic phase leading to a lower corrosion resistance and loss of mechanical properties<sup>10</sup>.

The Figure 7b, presents a slightly more expressive nitrogen amount content profile for samples welded with the gas mixture Ar + 2% N₂ as protection gas. However, the reduction of the amount of nitrogen is relative to the base metal, which is evident. Zhang et al.<sup>15</sup> assessed the content of nitrogen for each phase in the welding zone through quantitative element mapping using EPMA.

As reported by Zhang, the 2% N₂ addition in the shielding gas also promoted the nitrogen solution in the austenite¹⁵. According to Lin et al.³¹, experimental results show that the heat input increases with the increased welding voltage caused by the increased nitrogen content of the shielding gas. The authors also observed steeply decreases of the ferrite amount using nitrogen in shielding gas.

For the present work, it was also observed a higher concentration of nitrogen in allotriomorphic austenite (γ<sub>A</sub>) grains than in austenite grains type Widmanstätten. This fact is reinforced by the low amount of nitrogen found for the measurements path as described in the nitrogen measurement profile. According to Eghlimi et al.³⁸, the lowest amount of Cr, Mo and N in the Widmanstätten austenite grains is due to their formation at low temperatures when most of the alloying elements have already been consumed by the allotriomorphic austenite grain boundaries.

The microhardness test was performed on the cross section of the welded samples. With the extracted data the microhardness maps were constructed and compared with the generated images from the microstructure of the welded region, as seen in Figure 8. The microhardness was measured in an area of 1.8 mm x 8 mm which surrounds the base metal with the position of the thermocouple, the ZAC and the SZ observed through the micrographs of the cross section from both samples.

<figure><img src="image_8.png" /><figcaption>**Figure 6.** Three-dimensional microstructure of duplex stainless steel UNS S32304 welds in base metal and solidified zone of the GTAW welded joints using b) pure Ar and c) Ar+2%N₂ as shielding gas. The polishing samples were etched for 30 seconds in modified Behara I reagent, composed of 20 mL hydrochloric acid (HCl), 80 mL distilled water and 1 g potassium metabissulfide (K₂S₂O₅).</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>**Figure 7.** Nitrogen concentration as a function of transversal position measured by EPMA for both GTAW process with a) pure Ar and b) Ar+2%N₂ as shielding gas.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>**Figure 8.** Microhardness maps on transversal direction of the GTAW welded joints using a) pure Ar and b) Ar+2%N₂ as shielding gas.</figcaption></figure>

Observing the Figure 8, the base metal region highlighted in green on the microhardness map (approximately 220 HV₀.₀₅) is in accordance with that informed by the steel supplier. The sample welded with pure Ar as protection gas, see Figure 8a, presented islands of hardness mainly at the center of the weld bead. The same phenomenon is observed for the sample welded with Ar + 2% N₂ (see Figure 8b), where the islands with high hardness are larger and more evident.

A large amount of allotriomorphic austenite (γ<sub>A</sub>) of columnar ferrite grain contours is also observed in the micrographs, and especially in the equiaxial ones, with this region having higher hardness. The high amount of γ<sub>A</sub> was favored by the high concentration of nickel and under the influence of N₂ in the shielding gas³⁰. According to Ruiz et al.³⁹, the microhardness is higher in the austenite than in the ferrite and the nitrogen, as already discussed, contributes strongly to the austenite reforming.
## 4. Conclusion

Through the welding thermal cycles, it was possible to determine the temperatures of the duplex stainless steel in the melt pool and in the heat-affected zones. The higher peak temperature (947 °C) for the sample welded with addition of $N_2$ in the shielding gas did not directly imply on the increase of the cooling rate. The highest cooling rate was found for welding with pure Ar.

Based on results of the cooling rates and quantitative phase analysis, it was possible to conclude that the fast cooling of the welding zone is straightly related to the amount of austenitic phase. The amount of austenite decreased considerably for the sample welded using pure Ar. The intragranular austenite and grain edge for welded sample using pure Ar for three directions observed are smaller than ferrite for welding using the mixture Ar + 2% $N_2$ as the shielding gas, can suggest a direct relationship with the increase of microhardness in the solidified zone for this sample.

The addition of $N_2$ in the shielding gas increased the hardness of the solidified zone. By the nitrogen amount profile measurements in the cross section, it is shown that it is possible to achieve a good level of nitrogen in the microstructure with adequate shielding gas aiming the increase of austenite phase.

## 5. Acknowledgments

The authors want to thank Erasmus Mundus (EBW+), CAPES and CNPq for the financial support.

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