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Mistakes easily made in welding duplex stainless steel

1. Preface

People have been producing and using duplex stainless steels for almost 80 years. These alloys are characterized by essentially 100% ferrite at solidification, and austenite is necessary to nucleate and grow in the solid state. Early alloys such as malleable alloy 329 and casting alloy CD4MCu contain far more ferrite than austenite.
In addition, people did not pay attention to the importance of nitrogen, many alloys contain very little nitrogen, so that in the cooling state, austenite nucleation and growth is too slow to obtain a balanced amount of austenite in the heat-affected zone of the weld when no heat treatment is performed after welding.
When superalloyed weld filler metals are used in order to promote austenite formation in the melt zone, the post-weld heat affected zone of these alloys usually becomes brittle and has poor corrosion resistance. This defect is also found in the weld metal of gas welding.
In the 1980s, the importance of nitrogen to duplex stainless steels was fully recognized, and minimum nitrogen content requirements are usually specified. By adding an appropriate amount of nitrogen to the base metal and using a weld filler metal with increased nickel content, the weld can be made to have approximately equal amounts of austenite and ferrite in the post-weld state, which significantly improves the mechanical properties and increases the corrosion resistance.
The next step is to make the heat input reasonable in order to achieve a reasonable austenite-ferrite balance in the heat affected zone through a suitable cooling rate. Too low heat input will result in too much ferrite and too high heat input will result in intermetallic phase precipitation. It is now conventional practice to weld 22% Cr duplex stainless steel with heat input between 0.5 and 2.5 kJ/mm, and 25% Cr duplex stainless steel with heat input between 0.5 and 1.5 kJ/mm.
Although the vast majority of duplex stainless steel weldments are used in the post-weld condition, there are at least two cases where post-weld heat treatment (annealing) is usually required. Duplex stainless steel castings almost without exception need to be annealed, and, if the casting defects are repaired by welding, it is necessary to anneal the weld. Large welded heads manufactured by welding, whether cold-formed or hot-formed, need to be annealed.

2. Mistakes made

1.1 Unreasonable base metal specification

In the past 25 years, the most common duplex stainless steel is the alloy called 2205. The literature almost always uses the composition range of UNS S31803 to describe this alloy. However, the UNS S31803 nitrogen content is as low as 0.08%, and this level of nitrogen has proven to be too low to maintain good performance in the heat affected zone and melt zone under post-weld conditions.
After realizing this problem, ASTM has used UNS S32205 to define 2205 since 2000. the composition ranges for UNS S31803 and UNS S32205 are shown in Table 1. It is worth noting that S32205 has increased minimum chromium and molybdenum content in addition to increased minimum nitrogen content compared to S31803. The importance of nitrogen in controlling the ferrite/austenite phase balance during welding has been well illustrated by Ogawa and Koseki.
Figures 1 to 3 are taken from their report. Figure 1 shows the microstructure and the distribution of alloying elements between its ferrite and austenite phases for a forging steel that conforms to the composition of UNS S31803 but not to that of UNS S32205. The nitrogen content of 0.12% is too low for S32205. In Figure 1(a) the ferrite is darker gray, while the austenite is almost white. It can be seen from the figure that within the ferrite phase, it is rich in chromium and molybdenum, and within the austenite phase it is rich in nickel and nitrogen.
In particular, chromium is about 25% within ferrite and only 20% within austenite, see Figure 1(b), molybdenum is about 3.5% within ferrite __ and only 2.5% within austenite, see Figure 1(d), while nickel is 7.5% within austenite and about 5% within ferrite, see Figure 1(c), and nitrogen is about 0.3% within austenite and ferrite is zero, see Fig. 1(e). The distribution of phases is basically balanced. Since the steel is a hot rolled product, the organization is striped.
In contrast to Figure 1, the microstructure and distribution of alloying elements in the melt zone of a self-fusing GTA weld of UNS S31803, which is the same material as within Figure 1, is shown in Figure 2. Austenite is not well distributed within the ferrite grains and is mainly located at thin layers of the original ferrite grain boundaries.
It is difficult to separate the distribution of chromium, nickel and molybdenum between ferrite and austenite. One can hardly see the pattern of alloying element distribution in Figures 2(b),2(c) or 2(d) similar to the phase distribution in Figure 2(a). However, the nitrogen distribution shows that the concentration of nitrogen within the austenite lamellae around the ferrite grain boundaries is higher than the concentration of nitrogen within the ferrite grains.
This may be due to the fact that the nitrogen within these regions has enough time to diffuse into the austenite on the ferrite grain boundaries, causing the regions next to the austenite lamellae to be practically free of nitrogen, as shown in Figure 2(e). Inside the ferrite grain, the nitrogen is blocked and has no chance to enter the austenite and precipitate as chromium nitride, which is clearly visible, see Fig. (2a). For the melting zone, the combination of large ferrite grains and chromium nitride precipitates is very harmful for the toughness and corrosion resistance of the melting zone.

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Figure 1. Wrought microstructure and element partitioning in UNS S31803

Table 1 typical chemical composition of duplex stainless steel

ASTM EN-Nr. C max. Cr Ni Mo N Mn Cu W PRE*
S31803 1.4462 0,03 21,0-23,0 4,5-6,5 2,5-3,5 0,08-0,20 2,0 33
S32205 1.4462 0,03 22,0-23,0 4,5-6,5 3,0-3,5 0,14-0,20 2,0 35

In fact, as mentioned before, 100% of the melting zone solidification is ferrite. Then, diffusion is required as the transformation to austenite begins. Since chromium, nickel and molybdenum are replacement elements that diffuse relatively slowly in the solid state, they cannot achieve a balanced distribution between ferrite and austenite under normal weld cooling conditions. However, nitrogen is an interstitial element, which diffuses about 100 times faster than the replacement element. As a result, it has the ability to enter austenite more often, although not completely as the composition shown in Figure 2.
In materials containing 0.12% nitrogen, the behavior of the hottest part of the heat affected zone is similar to that of the weld metal. In particular, it is composed of 100% ferrite, which is then partially transformed into austenite in the solid state. Therefore, it is more inclined to form large ferrite grains and austenitic strips and sheets along the original ferrite grain boundaries. Although the composition of the melting zone can be controlled by accelerating the formation of austenite by selecting a filler metal with a high nickel content, there is little control over the heat affected zone. In this way, it is best to avoid this composition in the post-weld structure in order to achieve optimum performance.

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Figure 2

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Figure 3
Figure 3 shows the microstructure and distribution of alloying elements in the GTA melt zone when the nitrogen content of alloy 2205 is higher than that shown in Figure 2, meeting the compositional limits of UNS S31803 and S32205. Due to the higher nitrogen content of the weld metal in Figure 3 (0.18% nitrogen compared to 0.12% nitrogen in Figure 2), the post-weld organization is significantly altered.
In particular, it can be clearly seen in Fig. 3(a) that more austenite is formed than that formed in Fig. 2(a), and the austenite is dispersed through the entire ferrite grain instead of being mainly confined to the ferrite grain boundaries. In contrast to Fig. 2(a), no precipitation of chromium nitride is seen in Fig. 3(a). Instead, all the nitrogen enters the austenite, as shown in Fig. 3(e). And clearly in Fig. 3(b) and (d) a small distribution of chromium and molybdenum can be seen in this region, with less chromium and molybdenum in the austenite first formed along the original ferrite grain boundaries than in the original ferrite grains.
The higher nitrogen content of this specimen may cause austenite formation to begin at a higher temperature than the 0.12% N alloy, and, because diffusion and transformation begin at higher temperatures, molybdenum and chromium may diffuse more rapidly and take longer to diffuse. It can also be seen in Figure 3(a) that the austenite bars formed on the entire original ferrite grain break up the original coarse ferrite grains.
As the original ferrite grains are broken into small grains by the internal austenite bars, the toughness is improved. And because there is no chromium nitride precipitation, the corrosion resistance is improved. This is also true in the high temperature part of the heat affected zone. In this way, it can be concluded that UNS S32205 with higher nitrogen content in Figure 3 is significantly better than UNS S31803 with lower nitrogen content in Figure 2.
Therefore, it is clear that UNS S31803 is not suitable as a base metal for post-weld applications, UNS S32205 should be used. in duplex stainless steel, UNS S31803 is not the only unsuitable base metal for welded structures. This is also the case for Alloy 255, which is part of the UNS S32550 composition.
Table 2 compares the composition of UNS S32550 and UNS S32520 used for overlay welding, the composition is largely the same, however, the minimum nitrogen content of UNS S32520 is higher than that of UNS S32550, so it is clear that UNS S32520 is more suitable for use in the post-weld condition. Alternatively, UNS S32550 can be used, but only if the nitrogen content of this steel composition is controlled to the upper limit.

1.2 Unreasonable welding heat input

The traditional view on welding heat input is that when welding duplex stainless steel with 22% chromium content, heat input should be limited to 0.5 to 2.5 kJ/mm, and when welding super duplex stainless steel with 25% chromium content, heat input should be limited to 0.5 to 1.5 kJ/mm.
When using lower (here means <0.5kJ/mm, small heat input, small heat, fast cooling rate!) ) heat input, even for duplex steels containing high nitrogen, austenite formation at very fast cooling rates is not sufficient. When using higher (here >1.5kJ/mm, high heat input, high heat and slow cooling rate!) of heat input, there is a tendency for intermetallic compounds to precipitate within the ferrite under slow cooling conditions.
The trend is more pronounced for 25% Cr super duplex stainless steel compared to 22% Cr duplex stainless steel.
Karlsson et al. pointed out that the tendency to form precipitates during welding of higher nitrogen containing, 22% Cr duplex stainless steels such as UNS S32205 is quite low and there is no risk during welding as long as the above mentioned welding heat input limits are observed.
However, he further pointed out that when welding 25% Cr super duplex stainless steel, even if the welding heat input is limited to 0.5 to 1.5 kJ/mm, there is no guarantee that multiple welds will be free of precipitates. In these high-alloyed steels, multiple heating cycles of the weld can lead to precipitation of chromium nitride, secondary austenite and various intermetallic compounds, including the σ phase. Table 3 lists the composition ranges for base metals and welding rods.
Note that in addition to the high nickel content of the filler metal, which by convention promotes austenite formation in the post-weld state, small amounts of copper and tungsten are added to the filler metal, in order to match the base metal. Many filler metal manufacturers recommend a combination of filler metal/base metal materials.
The thickness of the plate used for the process evaluation test was 9.5 mm, and the joint bevel was a single-sided V-bevel with a bevel angle of 60 degrees, a root gap of 1.5 mm, and a blunt edge of 3 mm. 3.2 mm electrode was used for the initial process evaluation test. After ten welds in the V-bevel, the root was purged to expose the intact metal and then two more passes were made to complete the weld.
The average weld heat input for all passes was 0.7 kJ/mm. Small size (8 mm thick) Charpy V-notch specimens were cut from the weld metal and heat affected zone at -40°C and tested. The impact test requirement was 27J and the heat affected zone far exceeded that requirement. However, two of the three Charpy V-notch specimens did not reach 27 J during the initial and repeat tests on the weld metal.

Table 2 composition limits of alloy 255

ASTM EN-Nr. C max. Cr Ni Mo N Mn Cu W PRE*
S32520 1.4507 0,03 24,0-26,0 5,5-8,0 3,0-4,0 0,20-0,35 1,5 0,5-2,0 >40
S32550 1.4507 0,04 24,0-27,0 4,5-6,5 2,9-3,9 0,10-0,25 1,5 1,5-2,5 >40

Table 3 UNS S32760 and E2595- 15 composition ranges

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In order to find out the reason for the poor impact test results of the weld metal, the weld specimens for the process evaluation test were examined using a scanning electron microscope. The microstructure of the weld metal near the middle thickness of the specimen is shown in Figure 4. There was a large amount of angular precipitates within the ferrite alone. However, it was not determined exactly what the precipitates were. Our conclusion is that the precipitates were produced by repeated heating during the twelve passes of the test specimen.
Therefore, a new process evaluation test was performed using the same joint design and electrode. In the new process test, the welding speed was reduced in order to keep the welding heat input between 1.2 and 1.3 kJ/mm, and the welding was accomplished by performing four passes in the upper part and one pass after root clearing. At a temperature of -40°C, small size Charpy V-notch impact specimens of the same size completely exceeded the requirement of 27J. There were also no various precipitates within the microstructure.
There is a possibility that the root weld channel within the pipe presents a special case of improper heat input. When training welders of carbon steel pipes, they are required to weld root passes at a fairly high speed, usually with a cellulose electrode for downward standing welds, and then use a high heat input of “hot passes” to prevent hydrogen cracking of carbon steel.
However, after the “hot pass” with high heat input, the welding of the root pass with low heat input can overheat the root pass and cause the precipitation of intermetallic compounds in the root pass of super duplex stainless steel.
This is a very dangerous situation because the root weld surface is usually in contact with corrosive media during use. Although intermetallic compounds are detrimental to toughness, intermetallic compounds buried within the welded joint away from the exposed surface are less hazardous than intermetallic compounds within the root weld channel because intermetallic compounds buried within the welded joint are generally not in contact with the corrosive medium, while intermetallic compounds within the root weld channel are in contact with the corrosive medium.
The standard operation in welding duplex stainless steel, especially super duplex stainless steel piping is that the heat input of the root weld channel is greater than the initial filler weld channel. The thickness of the root weld channel of about 6mm is used quite well.

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Figure 4 Angular precipitates in the ferrite of E2595-15 reheated weld metal

1.3 Unreasonable post-weld heat treatment

If welded castings or welded formed heads require post-weld heat treatment, then the use of commonly used nickel-rich filler metals, coupled with the use of base metal specifications to meet but unreasonable annealing temperatures will make another mistake for duplex stainless steel processing and manufacturing plants (for duplex stainless steel, typically 9% nickel content, other similar to filler metals, such as E2595 shown in Table 3 -(15 filler metal).
The general requirement is to anneal at a minimum temperature of 1040°C and then water quench from the annealing temperature. Since it is not well understood that σ-phase is almost always formed in duplex stainless steel during heating to annealing temperature, and higher nickel content increases the solid solution phase line temperature of the σ-phase. In this case, the nickel-rich weld metal is dangerous. the effect of nickel on the solid solution phase line temperature of the σ-phase of the 25% Cr-3.5% Mo alloy is shown in Figure 5, produced by Grobner.
Although the alloy used to draw this figure does not contain alloying elements such as manganese, silicon and nitrogen, it is convenient to understand the effect of nickel in essence. It clearly shows that the σ-phase solid solution phase line temperature increases with increasing nickel content. In particular, it shows that the σ-phase solid solution phase line temperature of a weld metal containing 9% Ni will be at least 50°C higher than the σ-phase solid solution phase line temperature of a matched base metal containing 5% Ni.
In essence, Figure 5 can also be applied to alloys with 22% Cr content, such as 2205 welded using a nickel-rich filler metal. the weld metal contains 8.3% Ni. it was annealed at 1040°C for 96 hours due to concerns that the σ-phase formed during heating to annealing temperature would dissolve too slowly at the annealing temperature. A large amount of σ-phase can be clearly seen after water quenching from the annealing temperature.
It can be concluded that the σ-phase within this composition is stable at the temperature of 1040°C. Note that due to the long annealing time, the microstructure is much coarser compared to the one shown later. Figure 6 shows the weld metal exactly matched to the alloy 255 composition (5.8% Ni). The weld metal is water quenched before annealing at a temperature of 1040°C for 4 hours, it contains no σ-phase and is quite ductile (34% elongation in a 4:1 scale length to diameter ratio tensile test).
When alloy 255 with weld metal of otherwise similar composition but with 9% nickel content was annealed at a temperature of 1040°C and water quenched, the result was a large amount of σ-phase throughout the microstructure, see Figure 7. In Figure 7, the σ-phase appears gray, the austenite white, and the ferrite black. The weld metal is quite brittle, with an elongation of only 7% (compared to 26% in the post-weld condition).
It is noteworthy that the post-weld ferrite content of this 9% Ni weld metal is 54FN, however, after annealing at a temperature of 1040°C, the FN drops to 28, indicating that about half of the original ferrite becomes σ-phase due to this heat treatment. This is substantially consistent with the microstructure shown in Figure 7.
The same 9% Ni weld metal, annealed at a temperature of 1150°C, cooled in the furnace to 1040°C and held at 1040°C for 30 minutes, then water quenched before nucleation of the σ phase, yielded a measured 45 FN with 35% elongation and no σ phase within the microstructure.
The “graded annealing” allows for near-equilibrium distribution of nitrogen, which is where the Alloy 255 supplier claims the base metal has excellent corrosion resistance. It is clear from this paper that the nickel-rich filler metal requires a higher annealing temperature than the base metal in order to avoid the hazards associated with the σ-phase.
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3. Conclusion

Duplex stainless steels, including super duplex stainless steels, have proven to have good welding properties and are very important engineering materials. However, due to carelessness and lack of knowledge, some mistakes can be made. In order to have good properties in the heat-affected zone and weld metal, it is appropriate to specify the use of base metals with a high nitrogen content of at least 0.14%. Otherwise, at least in the heat-affected zone will have a large amount of ferrite, chromium nitride precipitation, damage corrosion resistance and mechanical properties.

In order to avoid precipitation of super duplex stainless steel in the multiple repeated heating zones, the use of a large number of low heat input, small welding passes should be avoided as an option. Also, when welding duplex and super duplex stainless steel pipes, to avoid precipitation on the inner surface of the pipe, the root weld channel should be welded with a higher heat input (greater than 1 kJ/mm) than the heat input of the first few consecutive weld channels. Post-weld annealing of welded duplex stainless steels requires consideration of the case where the temperature of the σ-phase of the weld metal with dissolved nickel-rich content is higher than that of the base metal.

By Demian J. kotecki

Source: China Large Diameter Flange Manufacturer – Yaang Pipe Industry (

(Yaang Pipe Industry is a leading manufacturer and supplier of nickel alloy and stainless steel products, including Super Duplex Stainless Steel Flanges, Stainless Steel Flanges, Stainless Steel Pipe Fittings, Stainless Steel Pipe. Yaang products are widely used in Shipbuilding, Nuclear power, Marine engineering, Petroleum, Chemical, Mining, Sewage treatment, Natural gas and Pressure vessels and other industries.)

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