Effect of intermetallic phases on corrosion resistance of superduplex and superaustenitic stainless steel weldments A J Leonard+, P Woollin+ and D C Buxton* + TWI Ltd Granta Park, Great Abington, Cambridge, CB1 6AL, UK * Now at Capcis Ltd Capcis House, 1 Echo Street, Manchester, M1 7DP, UK Stainless Steel World Conference The Hague, Netherlands, 1315 November 2001 Manuscript number : P0185 Abstract Superduplex, superaustenitic and nickel alloys have exceptional corrosion resistance and have proven performance in many aggressive environments. However, these highly alloyed materials are susceptible to the formation of intermetallic phases when exposed to thermal cycles in the range 6001000°C as may typically be experienced during welding. The study examined four alloys: UNS S31254, S32760, S32750 and N08825, which were welded, using an overmatching NiCrMo filler in the case of the superaustenitic steel, nickel overalloyed superduplex fillers for the superduplex stainless steels and AWS E/ERNiCrMo3 type consumables for alloy N08825. Welds were produced with different levels of arc energy representing conditions: (i) well within typical industrial practice, (ii) at the top end of typical industrial practice and (iii) above typical practice. The volume fraction of intermetallic phase was determined for each material and welding condition. The corrosion performance of the welds was assessed using short term critical pitting temperature tests, and also using long term tests in simulated service environments: chlorinated seawater, CO 2 /O 2 brine, CO 2 /H 2 S brine and a severe refinery environment. The results showed that using good industrial practice, welds containing less than 1% intermetallic were produced in superduplex stainless steel and between 1.4 and 2.4% weld metal intermetallic was formed in the highly alloyed nickel base weld metals. In the case of superduplex consumables, this increased to between 1.7 and 2.9% when using conditions above typical industrial practice. The increase in arc energy was associated with a reduction in critical pitting temperature. However, in simulated service environments, the welds showed no significant corrosion. Recommendations are made with regard to the setting of fitness for service criteria. 1. Introduction The excellent localised corrosion resistance of superduplex and superaustenitic stainless steels and NiFeCr alloys stems from the high levels of alloying elements, in particular Cr, Mo and N. However, these highly alloyed materials are susceptible to the formation of intermetallic phases, which form at temperatures between around 600°C and 1000°C. Such temperatures are experienced during a weld thermal cycle, giving the possibility of intermetallics being formed in weldments, especially for high alloy grades, where precipitation can begin within a few tens or hundreds of seconds [1,2] . Highly alloyed stainless steels and NiFeCr alloys are susceptible to localised corrosion, e.g. pitting, in aggressive chloride containing media and intermetallic precipitation may reduce the corrosion resistance significantly. Intermetallic phases are rich in alloying elements, notably Cr and Mo, compared to the surrounding matrix and therefore, as the intermetallic phase is formed, Cr and Mo depleted zone can be created around the particles. The extent of the alloy depletion will depend upon the thermal cycles experienced and will determine the reduction in corrosion resistance [1,2] . The aims of the current work were to quantify the variation of intermetallic phase formation over a range of welding conditions for superduplex, nickel and superaustenitic alloys, and to determine their effects on corrosion resistance in realistic service environments. 2. Experimental procedure 2.1. Materials Four corrosion resistant alloys were selected for testing, i.e. one superaustenitic grade: UNS S31254, two superduplex grades: S32760 and S32750, and one NiFeCr alloy: N08825. These are referred to as materials A to D respectively. The materials were in plate form, with a nominal thickness of 10mm. Table 1a shows their chemical compositions. 2.2.Welding The TIG process was utilised for the root and second passes, whilst fill and capping passes were deposited using the
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14/07/2015 Effect of intermetallic phases on corrosion resistance of superduplex and superaustenitic stainless steel weldments
Effect of intermetallic phases on corrosion resistance ofsuperduplex and superaustenitic stainless steel weldmentsA J Leonard+, P Woollin+ and D C Buxton*
+ TWI Ltd Granta Park, Great Abington, Cambridge, CB1 6AL, UK
* Now at Capcis Ltd Capcis House, 1 Echo Street, Manchester, M1 7DP, UK
Stainless Steel World Conference The Hague, Netherlands, 1315 November 2001 Manuscript number : P0185
AbstractSuperduplex, superaustenitic and nickel alloys have exceptional corrosion resistance and have proven performance inmany aggressive environments. However, these highly alloyed materials are susceptible to the formation ofintermetallic phases when exposed to thermal cycles in the range 6001000°C as may typically be experiencedduring welding.
The study examined four alloys: UNS S31254, S32760, S32750 and N08825, which were welded, using anovermatching NiCrMo filler in the case of the superaustenitic steel, nickel overalloyed superduplex fillers for thesuperduplex stainless steels and AWS E/ERNiCrMo3 type consumables for alloy N08825. Welds were produced withdifferent levels of arc energy representing conditions: (i) well within typical industrial practice, (ii) at the top end oftypical industrial practice and (iii) above typical practice. The volume fraction of intermetallic phase was determinedfor each material and welding condition. The corrosion performance of the welds was assessed using short termcritical pitting temperature tests, and also using long term tests in simulated service environments: chlorinatedseawater, CO2/O2 brine, CO2/H2S brine and a severe refinery environment.
The results showed that using good industrial practice, welds containing less than 1% intermetallic were produced insuperduplex stainless steel and between 1.4 and 2.4% weld metal intermetallic was formed in the highly alloyednickel base weld metals. In the case of superduplex consumables, this increased to between 1.7 and 2.9% whenusing conditions above typical industrial practice. The increase in arc energy was associated with a reduction in criticalpitting temperature. However, in simulated service environments, the welds showed no significant corrosion.Recommendations are made with regard to the setting of fitness for service criteria.
1. IntroductionThe excellent localised corrosion resistance of superduplex and superaustenitic stainless steels and NiFeCr alloysstems from the high levels of alloying elements, in particular Cr, Mo and N. However, these highly alloyed materialsare susceptible to the formation of intermetallic phases, which form at temperatures between around 600°C and1000°C. Such temperatures are experienced during a weld thermal cycle, giving the possibility of intermetallics beingformed in weldments, especially for high alloy grades, where precipitation can begin within a few tens or hundreds ofseconds[1,2] .
Highly alloyed stainless steels and NiFeCr alloys are susceptible to localised corrosion, e.g. pitting, in aggressivechloride containing media and intermetallic precipitation may reduce the corrosion resistance significantly.Intermetallic phases are rich in alloying elements, notably Cr and Mo, compared to the surrounding matrix andtherefore, as the intermetallic phase is formed, Cr and Mo depleted zone can be created around the particles. Theextent of the alloy depletion will depend upon the thermal cycles experienced and will determine the reduction incorrosion resistance [1,2] . The aims of the current work were to quantify the variation of intermetallic phaseformation over a range of welding conditions for superduplex, nickel and superaustenitic alloys, and to determinetheir effects on corrosion resistance in realistic service environments.
2. Experimental procedure2.1. MaterialsFour corrosion resistant alloys were selected for testing, i.e. one superaustenitic grade: UNS S31254, twosuperduplex grades: S32760 and S32750, and one NiFeCr alloy: N08825. These are referred to as materials A to Drespectively. The materials were in plate form, with a nominal thickness of 10mm. Table 1a shows their chemicalcompositions.
2.2.WeldingThe TIG process was utilised for the root and second passes, whilst fill and capping passes were deposited using the
MMA process. No preheating was used and interpass temperatures were restricted to below 150°C in all cases. Theshielding and backing gases were both argon. Typical weld metal compositions for the consumables used are listedinTable 1b . An overmatching NiCrMo filler was used in the case of S31254, nickel overalloyed superduplexconsumables were used for the two superduplex grades, and AWS E/ERNiCrMo3 type consumables were used foralloy N08825.
Welding current and travel speed were varied to produce welds with three different levels of arc energy, i.e. low (L):from 0.9 to 1.6kJ/mm, medium (M): between 1.1 and 2.0kJ/mm and high (H): between 1.9 and 3.2kJ/mm,representing conditions (i) well within typical industrial practice, (ii) at the top end of typical industrial practice and(iii) above typical industrial practice.
The welds were coded with a 'W', followed by 'A, B, C or D' to represent the base steel and then 'L, M or H'. Wheremore than one weld of each type was made, these were distinguished by a number, e.g. 1 or 2.
2.3. CharacterisationTwo different point counting methodologies were adopted for determining the localised peak and average volumefraction of intermetallic phases in each weld: (i) To determine 'peak' localised precipitate levels: 16 measurementfields were counted on a prepared metallographic specimen, in a square 4x4 field array, in areas identified by priormicroscopic examination as containing the highest levels, at a magnification of either X1250 or X1400. (ii) Todetermine average precipitate levels: 16 measurement fields were counted, using a 100 point (10x10) grid, atrandom locations within the root and second weld runs. The magnification used was either X1250 or X1400. Thephase balance was also determined in each of the following areas of the superduplex materials: parent steel, weldmetal cap, weld metal root, cap HAZ and root HAZ.
2.4. Short term ferric chloride corrosion testsCritical pitting temperature (CPT) ferric chloride tests were conducted generally following ASTM G48 Method C [3] .The weld surfaces were left aswelded except for degreasing, and the specimens sides and ends were prepared to 120and 1200 grit finish respectively. Exposure commenced at 27.5°C. The temperature increment was 2.5°C. Thisprocedure was repeated until either metal loss was detected by visual inspection or weight loss exceeding 20mg wasrecorded. The first temperature at which pitting occurred was taken as the CPT.
2.5. Long term corrosion tests in simulated service environments2.5.1. Test samplesThe specimens were ~25 by 50mm by full thickness, so that the exposed surfaces included both weld cap and root,and cut edges. The weld surfaces were degreased, whilst the specimen edges and ends were prepared to 120 and1200 grit finish respectively. Specimen mass was recorded before and after test. Following completion of thecorrosion test, samples were examined visually and if signs of corrosion were observed, examination in a scanningelectron microscope (SEM) was undertaken.
2.5.2. Chlorinated seawater testDuplicate weld samples of the superaustenitic and both superduplex grades were immersed in natural seawater heldat a constant temperature of 40°C. The test vessel was open to the atmosphere, allowing diffusion of oxygen intothe test solution. A diluted solution of sodium hypochlorite (approximately 1%) was continuously added to the testsolution, to achieve a free chlorine level between 0.5 and 1.0ppm throughout the test. Electrochemical potential wasrecorded during the exposure period.After 36 and 62 days exposure, the specimens were removed and inspected visually under a binocular microscope.Mass loss measurements were taken. As no evidence of corrosion was found, the test temperature was increasedfrom 40°C to 53°C on reimmersion after the second inspection until the end of the test on day 91. Followingremoval from the test solution, the specimens were cleaned in water, inspected under a binocular light and reweighed.
2.5.3. CO2/O2 Brine testThe CO2/O2containing brine was designed to simulate produced water systems where temperatures can be high (80100°C) and small amounts of O2 can be present. Specimens of superaustenitic steel (UNS S31254), one of thesuperduplex grades (UNS S32750) and the NiFeCr alloy N08825 were suspended in deaerated 3%NaCl solution at80°C, a gas mixture of CO2 and 0.4%O2 was introduced and bubbled through the solution throughout the 90 daytest period. The total test pressure was 2.5 barg, which gave a partial pressure of CO2 of 2 bar. Following completionof the test, the specimens were cleaned in water, inspected visually and reweighed. The mass loss for each specimenwas also determined after a second cleaning operation, using 40% hydrochloric acid to remove tarnishing [4] .
2.5.4. CO2/H2S brine testThe H2Scontaining environment was chosen to match the NACE MR017599 recommended service limits for UNS
S32760 [5] . The test solution comprised deaerated 12%NaCl, with an addition of 0.34 g/l of sodium bicarbonate togive a pH value of 4, at test temperature and pressure. Four point bend test specimens were taken from the twosuperduplex grades. The specimens had the root intact and were strained to give 0.2% plastic strain at the weld toe.After heating to test temperature, i.e. 100°C, the CO2/H2S/N2 test gas mix was introduced to the vessel, providingCO2 and H2S partial pressures of 2 and 0.2 bar, respectively and the test lasted 90 days. At the end of the testperiod, the specimens were inspected as described in section 2.5.2 above.
2.5.5. Severe refinery testThe test involved preparation of NH4Cl paste (160ml of distilled water and 330g of NH4Cl), which was placed in aglass container inside an autoclave. Approximately one litre of distilled water was poured into the autoclave. Aftersealing, the system was purged with N2 gas to remove oxygen prior to testing. Two sets of specimens wereemployed: one set was placed in the paste, the other was suspended in the vapour. All four materials were includedin the test. The autoclave was heated to a temperature of 148°C and was pressurised to 18 barg with N2 gas. Thetest duration was one month. The water in the bottom of the vessel provided approximately 20% humidity in thevapour phase, assuming a water vapour partial pressure of 4.5 bar. Following unloading of the autoclave, thespecimens were washed in water, visually inspected for evidence of corrosion under a binocular microscope and reweighed.
3. Results3.1. Metallography3.1.1. Material A: superaustentic steel UNS S31254Each of the weld metals of material A contained widespread interdentritic intermetallic phases, the volume fraction ofwhich generally increased with increasing arc energy, Tables 2 and 3 . The peak volume fraction increased from 1.8%to 2.4%. Most dense precipitation was in reheated areas. The weld cap and bands in the lower parts of reheatedpasses showed less precipitation. The average volume fraction of intermetallics was lower, but showed a similar trendto the peak volume fraction results increasing from low arc energy (1.4%) to high arc energy (1.8%). Figure 1 showsa typical microstructure of the weld root in the high arc energy weld. The intermetallic particles were irregular inshape and typically 1µm to 4µm in all dimensions. No HAZ precipitation could be quantified, although, there wasevidence of preferential etching of the HAZ grain boundaries and at higher arc energies some 'spottiness' andthickening of the boundaries when etched.
Fig.1. Typical light micrograph of the root microstructure ofUNS S31254 weld WAH1
3.1.2. Material B: superduplex steel UNS S32760All welds in material B showed a typical ferriteaustenite structure, with islands of austenite in a ferrite matrix. Somesmall patches of secondary austenite were present, typically lying between weld passes, and most numerous for thedeliberately 'abusive' weld procedures, i.e. the medium and high arc energy. The ferrite volume fractions for each ofthe welds are shown in Table 4 . The reported values were within the expected range [6] .
Welds of medium and high arc energy both contained intermetallic phases in a region encompassing the root and 2ndweld passes. In these samples, the peak volume fractions were 0.8% and 1.7% respectively in the middle of theareas of precipitation ( Table 2 ). Random point counting gave average volume fractions in the root and second passof 0.6 % and 0.5% respectively ( Table 3 ). The low arc energy weld contained only a few particles in the weld root,
possibly intermetallics. Figure 2 shows the root bead microstructure in sample WBH1. The intermetallic particles wereirregular in size and morphology, Fig.2b, and tended to be elongated on the austeniteferrite interfaces. Typicalexamples were 12µm wide and up to around 10µm in length.
a) light photomicrograph
b) backscattered electron image of intermetallic phase
Fig.2. Microstructure of the root of WBH1 UNS S32760 welds
The HAZ contained precipitates, with morphology suggesting that some were nitrides and some intermetallics. All ofthe different arc energy samples showed similar high temperature HAZ features with groups of etch pits withinferritic grains and decoration of ferrite subgrain boundaries. These are consistent with chromium nitrides.
3.1.3. Material C: UNS S32750Each weld contained a small fraction of intermetallic phases, predominantly in the second weld run and some smallregions of secondary austenite towards the weld cap. The high temperature HAZ showed limited fine intragranularnitrides, but no intermetallic phases were found. The ferrite volume fraction in each weld is shown in Table 4 and waswithin the expected range [6] .
The peak level of intermetallic phases was highest in. the medium arc energy weld. This was determined to be 2.6%c.f. 0.8% in the low arc energy sample and 1.3% in high arc energy sample ( Table 2 ). However, random pointcounting revealed lower average volume fractions, which increased with increasing weld arc energy, Table 3 . As forS32760, typical precipitates were 12µm wide and up to around 10µm in length.
3.1.4. Material D: NiFeCr alloy UNS N08825All of the weld metals contained second phase interdendritic particles, which was fairly uniform throughout the weldmetal. The volume fractions of the intermetallic phase particles are displayed in Tables 2 and 3 , and increase slightlywith increasing arc energy.
Examination in the SEM confirmed the presence of second phase particles of irregular morphology and with
dimensions up to 5µm. The HAZ precipitation could not be quantified in terms of size or volume fraction, but thegrain boundaries in this zone showed preferential etching when compared to parent material, becoming increasinglypronounced at higher arc energy.
3.2. Ferric chloride testingThe results are displayed in Fig.3. The superaustenitic steel samples gave the highest CPT values, up to 55°C, whilstpitting was first observed on the superduplex alloys at temperatures between 35 and 40°C. Generally, as the weldarc energy increased, the CPT was reduced, although there was little difference between the medium and high arcenergy samples. The parent material critical pitting temperatures calculated from published data [3] were found to beapproximately 60°C. In all cases, the measured weld values were below the parent material value, as would beexpected.
Fig.3. Ferric chloride CPT data
The location of attack was found to vary with the UNS S31254 specimens mainly suffering corrosion on the specimensides at the fusion line. For the two superduplex grades, attack was again on the specimen sides, generally in theroot or second/third pass weld metal region.
3.3. Long term corrosion tests in simulated service environments3.3.1. Chlorinated seawater testThe weight change results, displayed in Fig.4, were small, which indicates that no significant corrosion occurredduring the test period. In addition, there was no significant difference between the low and high arc energy weldresults. The mass loss increased with time and the maximum recorded mass loss for the 91 day exposure was5.7mg. This represents a very low uniform corrosion rate of <0.001mm/yr. The visual examination results supportthe weight change data, as the inspection after 36 days exposure only revealed some weldment discolouration andlight etching, which revealed the weld metal.
Fig.4. Mass loss results from the chlorinated seawater tests
The corrosion potentials that were recorded during the test period are plotted in Fig.5a, b and c. The initial potentialreadings were approximately 200 to 300mV SCE, but during the first few days of the test, the corrosion potentialsof the specimens increased to values above 0 mV. During the three month test, potentials of +500 to +600mV wererecorded on all the alloys. However, the potentials were not steady and significant fluctuations were observed. Typicalbehaviour is shown in the plot of time versus corrosion potential, Fig.5d, which was recorded after 19 days exposure.
a) UNS S31254
b) UNS S32760
14/07/2015 Effect of intermetallic phases on corrosion resistance of superduplex and superaustenitic stainless steel weldments
Fig.5. Corrosion potentials during the chlorinated seawater tests:
Potential fluctuation in UNS S31254
3.3.2. CO2/O2 brine testAfter exposure, all of the test specimens appeared tarnished brown, with some black discoloration around the weldcap. There was no visual evidence of pitting corrosion or of any crevice corrosion where the specimens had beensupported, but some corrosion at weld spatter on the cap had occurred.The mass loss results are given in Table 5 and are generally low, indicating no corrosion attack with the exception ofWCL1 and WCH2F2. The results after light brushing showed a small weight gain, probably reflecting the buildup ofthe Fe/Cr oxide on the weldment surface. Further oxidation of the weld oxides during the corrosion test could lead toan increase in weight and a brown discoloration. The mass losses recorded after cleaning in HCl suggested that theseoxides were removed. The losses recorded for specimens WCL1 and WCH2F2 were not accompanied by any visualevidence of corrosion and are considered unlikely to be the result of corrosion attack, but more likely to be caused bythe removal of weld oxides and weld spatter. A similar magnitude mass loss was observed in a ferric chloride test,during initial exposure.
3.3.3. CO2/H2S brine testMass loss data are given in Table 6 . The visual examination performed after the test did not reveal any evidence ofcorrosion attack and all the specimens were similar in appearance. All the specimens suffered some mass loss duringthe test period (Table 6). The maximum value was 34mg on specimen WBH1F4, but as for the CO2/O2 brine test, asthere was no visual evidence of corrosion, the mass loss is believed to arise from loss of weld oxide. Although the
differences were small, high arc energy welds gave higher weight loss than the low arc energy welds. The S32760steel welds showed greater weight loss than the S32750 steel.3.3.4. Severe refinery water testThe mass loss results, which are displayed in Table 7 , are very low. Visual examination confirmed that minimalcorrosion had occurred.
4. Discussion4.1. Precipitation characteristicsOver the arc energy range examined, the volume fraction of intermetallic phase in nickel based weld metals increasedby only up to 50%, whilst superduplex weld metals showed an increase from less than 0.5% to around 2.6%maximum. It should be noted that the highest arc energies employed were well beyond the manufacturers'recommended levels and equally beyond the maximum levels that would typically be expected for industrial welding.Perhaps of greater significance is the fact that individual particle sizes varied little between the different arc energylevels. For superduplex weld metal, particles tended to grow along austenite/ferrite boundaries and become morenumerous but the width of particles varied little with arc energy. Nickel based weld metals generally showedincreased numbers of particles rather than larger particles at higher arc energy. Single phase weld deposits transformto intermetallic phase in alloy enriched interdendritic areas and dendrite size was relatively constant for the weldsexamined, typically requiring order of magnitude changes in cooling rate for significant changes to occur. Insuperduplex deposits, intermetallic tends to form in thin areas of ferrite between two islands of austenite. Asaustenite essentially does not transform to intermetallic phase, the particle size is limited to some extent by thespacing of austenite units.
4.2. Corrosion resistance4.2.1. Ferric chloride testsA general reduction in the CPT values was observed for the three alloys A, B and C as the arc energy increased.However, the reduction was not large; significant drops in CPT compared to parent plate material behaviour, up to40°C, have been reported [7] for volume fractions of 1 to 4% intermetallic phase produced in superduplex steel byisothermal heat treatment (Fig.6). However, the scatter in this published plot of drop in CPT versus volume fractionof intermetallic is considerable and this is consistent with the size of individual precipitates being more influential thanthe volume fraction [1,2] . The temperature at which the precipitate is formed will also have an effect. For pittingresistance to be reduced, the depleted layer must be of sufficient size to contain a stable pit nucleus and the width ofthe depleted layer and the extent of alloy depletion are a function of diffusion rates. Hence, depleted layer width hasbeen suggested to be the factor controlling corrosion resistance, rather than volume fraction. For example, the size ofa stable pit nucleus might be 0.1µm [8] . It has been shown that precipitates produced by isothermal ageing, whichare larger than those formed during welding, are more detrimental to CPT, for a given volume fraction [1] .
Fig.6. Effect of intermetallic content on CPT reduction [3]
The present data is included with the published data in Fig.6 and it falls within the scatter band. The data has been
plotted twice; firstly by using published parent material CPT values to calculate the reduction in CPT (the parentmaterial value was ~60°C for all these alloys S31254, S32760 and S32750) and the 'peak' intermetallic volumefraction results, and secondly, by taking the difference between the low and high arc energy weld ferric chlorideresults and using 'average' intermetallic volume fraction data. It must be recognised that when comparing the weldand parent material CPT values, the significant reduction is not entirely the result of intermetallic formation. The twomethods of depicting the data in Fig.6 represent the two extremes that might be considered.
Whilst intermetallic precipitation was more pronounced in the nickel weld deposits, they were sufficiently overalloyedthat substantial loss of corrosion resistance would be required to bring the weld metal down to the parent steel level.Hence, in the superaustenitic and NiFeCr alloy weldments, the corrosion resistance was much more likely to becontrolled by the HAZ or fusionline, where precipitation or formation of an unmixed zone (UMZ), with associated Mosegregation, may occur. The formation of a UMZ, which consists of melted parent material which has not mixed withthe consumable, and its potentially limiting effect on corrosion resistance for welds in UNS S31254 has been welldocumented [9] . In the superaustenitic alloy the attack was observed on the specimen side at the fusionline and isbelieved to be associated with the presence of a UMZ which has previously been found to be the area of lowest pittingresistance for welds in this material [9,10] . Furthermore, the reduction in corrosion resistance that is associated withthe presence of welding oxide and the chromium depleted layer on the root and cap surface appears to be lesssignificant than the effect of the UMZ in this instance.
For the two superduplex grades, attack was also on the specimen sides, generally in the root weld metal region oraround midthickness, equivalent to the second/third passes. These areas were associated with the highestconcentration of intermetallic phases. The results indicate that the root surface, which is normally the surface that isexposed to the environment, has a better pitting resistance in a G48 test than the weld metal on the specimen sides.This probably reflects the fact that intermetallic phases do not tend to form on the root surface but concentratefurther into the weld, although in the high arc energy S32760, weld intermetallic phase was observed adjacent tothe weld root toes. However, attack was not observed at that location.
The ferric chloride test is a very severe accelerated corrosion test and although it is used to rank corrosionperformance of alloys, the significance of any comparison with corrosion resistance in quite different serviceconditions, e.g. as employed in the long term corrosion tests, is open to question.
4.2.2. Effect of intermetallic phase on corrosion resistance in service environmentsNo severe corrosion was observed in the longterm tests conducted in simulated service environments. The testenvironments represented servicelike conditions and were considered a reasonably severe test of the weldmentperformance.
The chlorinated seawater test commenced at 40°C, 10°C above the NORSOK standard limit [11] and because nocorrosion was detected the temperature was increased to 53°C. It is recognised that the NORSOK limit is consideredsevere and that the test specimens may have been 'conditioned' to some degree by exposure to the lower testtemperature, but overall it is considered that the test has demonstrated the corrosion resistance of the weldedspecimens to chlorinated seawater.
Previous tests in CO2/O2containing brine have been reported by Rogne and Johnsen [12] . They considered that at80°C and 200ppb O2, localised corrosion on S31254 might initiate and a maximum service temperature of 60°C wasrecommended for welded S31254, if 200ppb O2 is present. The test conditions used were therefore more severe thanthose recommended, yet no corrosion was observed.
The CO2/H2S brine and NH4Cl tests were based on conditions from NACE [5] and industrial experience respectivelyand were considered fairly demanding. The bend samples subjected to the sour test did not show any evidence ofcorrosion or cracking and the unstrained samples showed no significant corrosion. Similar behaviour was observed forthe NH4Cl tests. Modest mass losses, up to 65mg, were observed on a number of specimens. However, there was novisual evidence of localised corrosion on any of these specimens and it is believed that the mass loss arose largelyfrom removal of weld spatter either during the test or posttest cleaning.
In summary, both the low and high arc energy weld samples had sufficient corrosion resistance to withstand thechlorinated seawater, CO2/O2 brine, CO2/H2S and NH4Cl test environments. Therefore, materials with intermetallicphase of similar type and at comparable volume fractions would be considered suitable for service in theseenvironments. Corrosion attack of the high arc energy specimens might have been expected if the welding processhad produced intermetallic phase that significantly reduced corrosion resistance. However, the shortterm testsindicated that the effect of the intermetallic was fairly small.
The intermetallic phase volume fractions produced in this work were low when compared to levels generated byisothermal ageing, but the range of arc energy was broad and went well beyond current recommended weldingpractice.
4.3. Practical implicationsFor optimum superduplex weldment corrosion resistance, welding conditions should be restricted, e.g. to 0.5 to
1.5kJ/mm arc energy and <100°C interpass temperature, to minimise intermetallic phase formation. Superausteniticalloy UNS S31254 and NiFeCr alloy N08825, do not require such close control of welding conditions to minimiseweld metal intermetallic formation, but lower arc energy gives improved corrosion resistance, perhaps by minimisingHAZ precipitation and UMZ segregation.
Metallographic examination for intermetallic phase is typically a part of weld procedure qualification for superduplexstainless steels. When some particles are formed, point counting is often performed, and attempts made to set anallowable upper limit, typically around 12%. However, the size of the precipitates is perhaps of significance as largerparticles are expected to produce greater depleted zones. From this work, welds with 23% intermetallic phase maybe considered for corrosion resistant service. Demonstration of fitness for purpose should be on the basis of realisticsimulated worst case service tests, rather than shortterm quality control type tests.
5. Conclusions1. Generally the intermetallic volume fraction increased with increasing arc energy in superduplex,
superaustenitic stainless steels and nickel alloys.
2. Welds in superduplex steel (UNS S32760 and S32750) had weld metal intermetallic volume fractions lessthan 1% when welded within typical good industrial practice limits.
3. Welds produced with high arc energy, up to 3.2kJ/mm, i.e. above typical industrial practice, had maximumweld metal intermetallic volume fractions, between 1.7% and 2.9%.
4. Interdendritic intermetallic precipitates formed in highly alloyed nickel base weld metals. When typical goodindustrial practice was followed, the volume fraction of intermetallic was between 1.4 and 2.4%.
5. Little variation of intermetallic particle size with arc energy was observed.
6. The corrosion resistance of weld specimens containing intermetallic precipitates produced by welding with arcenergy well above typical industrial limits, was reduced by a modest margin in shortterm pitting tests, but insimulated service environments, no difference in corrosion performance was recorded, as none of the samplesshowed significant corrosion.
6. AcknowledgementsThis work was funded by Avesta Polarit, BP Exploration Operating Co Ltd, DERA, ESAB AB, Esso Engineering (Europe)Ltd, Fabrique de Fer de Charleroi SA and Shell UK Exploration and Production. The support of the sponsors isgratefully acknowledged. The authors also wish to thank all their colleagues at TWI who assisted with the practicalwork.
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14/07/2015 Effect of intermetallic phases on corrosion resistance of superduplex and superaustenitic stainless steel weldments
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Table 1a Chemical analyses of each of the parent materials.
Material Element, % (m/m) PRE N/PRE WC Si Mn P S Cr Cu Mo Ni W NS31254(A) 0.014 0.40 0.40 0.020 <0.001 20.1 0.67 6.23 18.0 0.05 0.210 44.0
