Failure of a super duplex stainless steel reaction vessel

Failure of a super duplex stainless steel reaction vessel

Engineering Failure Analysis 11 (2004) 243–256 www.elsevier.com/locate/engfailanal Failure of a super duplex stainless steel reaction vessel V.M. Lin...

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Engineering Failure Analysis 11 (2004) 243–256 www.elsevier.com/locate/engfailanal

Failure of a super duplex stainless steel reaction vessel V.M. Lintona,*, N.J. Laycockb, S.J. Thomsenb, A. Klumpersb a University of Adelaide, Adelaide, 5005, South Australia Materials Performance Technologies, Industrial Research Ltd, Lower Hutt, New Zealand

b

Received 1 July 2002; accepted 21 May 2003

Abstract Crevice corrosion and stress corrosion cracking (SCC) were recently discovered in a vessel used to strip vinyl chloride monomer from a water-based slurry of PVC granules. The vessel was manufactured from UNS S32750 super duplex stainless steel and the welds were produced using matching welding consumables. Although localised corrosion might have been expected, the occurrence of SCC was inconsistent with the majority of the published literature: in particular, the nominal operating temperature should have been too low for chloride-induced SCC of super-duplex stainless steel. However, damage was found mainly in the vicinity of the circumferential and longitudinal welds, and part of the subsequent failure investigation was therefore focused on the possibility of poor weld quality being the cause of failure. This task was approached primarily by measuring the influence of welding parameters on the value of the critical pitting temperature (CPT), and attempting to correlate the results with observed changes in the weld microstructure. CPT values were determined by a potentiodynamic method, using samples cut from the failed vessel and from a range of reference welds manufactured using known welding parameters. # 2003 Elsevier Ltd. All rights reserved. Keywords: Superduplex stainless steel; Stress corrosion cracking; Failure

1. Introduction A vessel, designed to strip vinyl chloride monomer (VCM), failed during service such that extensive offsite repairs were required. The vessel was constructed of UNS S32750 super duplex stainless steel, fabricated into a number of strakes, joined together with circumferential butt welds. A schematic of the vessel is given in Fig. 1. The vessel contained a number of trays and tray rings, most of which were also made of super duplex stainless steel. The remainder of the trays, at the top of the vessel, were fabricated from Incoloy 825. Visual inspection of the inside of the vessel during the investigation, revealed localised corrosion and cracking in many of the circumferential and longitudinal shell welds, circumferential butt welds 1, 3 and 5 being the most significantly affected. In all cases, the cracking originated from the inside of the vessel. Welds 3 and 5 contained at least 20 transverse cracks, many of which were associated with pits in the weld heat affected zone. Cracking, running along the weld toe parallel to the weld axis, was also seen associated * Corresponding author. E-mail address: [email protected] (V.M. Linton). 1350-6307/$ - see front matter # 2003 Elsevier Ltd. All rights reserved. doi:10.1016/j.engfailanal.2003.05.011

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Fig. 1. Schematic of the failed VCM vessel.

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with weld 3. Comparison of radiographs taken of the vessel after fabrication and during the failure investigation showed that all of the cracking had occurred in service. Transverse cracking was detected in the fillet welds attaching the trays to the vessel walls and delaminations were seen in the trays. These transverse cracks were similar to the cracks associated with the longitudinal and circumferential welds. The selection of S32750 for the vessel was based on the satisfactory performance of Incoloy 825 in a parallel process stream of similar fluid composition. Incoloy 825 generally has a lower corrosion resistance than does S32750 and, consequently, the selection of S32750 for the vessel under investigation was thought to be a conservative one. However, the preliminary failure investigation [1] indicated that the vessel had failed due to localised corrosion and stress corrosion cracking (SCC), and some key differences between the two process streams were identified. SCC of S32750 VCM strippers has also been reported by Notten [2], whilst Davies and Potgieter [3,4] have described similar failures of S31803 in this application. The significance of these failures is that they apparently involve the chloride-SCC of S32750 at about 105  C, whilst it is reported elsewhere that S32750 is effectively immune to chloride-SCC at temperatures up to 300  C [5,6]. Critical temperatures for chloride-SCC are only an empirical measure of resistance, but are a popular method of comparing different alloy grades. In particular, it is often stated that chloride-induced stress corrosion cracking (SCC) of UNS S30400 and S31600 stainless steels does not occur below about 50– 60  C [5–8], whilst duplex grades can withstand higher temperatures. For example, in 3% NaCl, S31803 does not suffer SCC at up to 150  C [9], and in hot MgCl2 solutions the common duplex grades (such as S31803) will resist SCC at up to 130  C [5]. However, the oxygen concentration in these solutions is low, which reduces the risk of SCC. The drop evaporation test (DET) allows easier oxygen access to the sample surface and so creates a more aggressive test [10]. Using MgCl2 in the DET, the common duplex grades suffer SCC at 100  C [11–13]. Steinsmo and Drugli also report service failure of S31803 at 4110  C due to seawater trapped beneath vessel insulation [11,12]. In the reported stripper failures, cracking was found to initiate at welds, and it is of course well-known that poor welding can lead to large reductions in SCC resistance. Although duplex stainless steels are very resistant to classical sensitisation due to precipitation of chromium carbides [14,15], sensitisation is now commonly used as a catch-all term to describe precipitation of intermetallics, such as s (sigma) or x (chi), and other particles or phases such as Cr2N (chromium nitride) or a0 (alpha prime). Of all these, s formation is by far the most detrimental factor in terms of reducing SCC resistance [10]. Welding can also reduce SCC resistance when the ferrite content in the weld metal is too high [9]. However, many authors find that the effect of welding is small when it is done properly [9,13]. Since the failure of the vessel under consideration in this paper was unexpected, further investigation was required to determine the reasons for the failure and offer advice on the susceptibility of the repaired vessel to failure. Detailed metallography of the microstructure of the welds was performed to determine the relationship between the microstructure and the cracking. The influence of welding parameters on the susceptibility of different microstructures to cracking was assessed through the use of critical pitting temperature (CPT) measurement on samples taken from the failed vessel and from test welds made under known conditions.

2. Vessel fabrication The vessel was designed and constructed to AS1210-1989 Class 1, with some additional requirements covering the fabrication and finishing of welds. One of these requirements was that the internal surface of the shell welds should be ground smooth and flush with the vessel wall. This was performed for all of the welds except circumferential welds 3 and 5, which both had up to 4 mm of weld reinforcement on the internal surface. Examination of the welds during the failure investigation indicated that the circumferential welds (except welds 3 and 5) and all of the longitudinal welds had been welded by submerged arc welding (SAW) from the outside of the vessel. Circumferential welds 3 and 5 appeared to have been fabricated using

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shielded or gas metal arc welding (SMAW or GMAW) for the first pass on the inside of the vessel and completed using SAW from the outside of the vessel. These welds were completed in five and four passes, respectively. The tray support rings were welded into place using a continuous fillet weld on the upper side and fillet stitch welds on the under side of the ring.

3. Service environment The process in the stripper involves a slurry of polyvinyl chloride resin entering the top of the vessel at 90  C, passing down through the trays, and finally leaving the bottom of the stripper at 110  C. VCM is stripped from the slurry by an upward flow of steam. The supporting medium in the slurry is potable water, containing trace amounts of hydrochloric acid, which is produced by hydrolysis of VCM during polymerisation, with the amount of chloride increasing if the residual oxygen in the reactor is increased. In addition to chloride produced from hydrolysis of VCM, it is also produced to an estimated concentration of 110 ppm from the catalyst formation reaction of ethyl chloroformate with hydrogen peroxide and calcium hydroxide. Overall, the worst case estimate of chloride concentration is 220 ppm, whilst the expected pH is 3.5–4. The oxidants present include oxygen, nitrite, and organic peroxides used to initiate polymerisation of the VCM. The oxygen content of the starting ‘‘batch’’ of water is minimised by application of a vacuum followed by back flushing with nitrogen, and there is no ingress of oxygen during the process. The organic peroxide is destroyed by passage through a spiral heat exchanger upstream of the stripper, but the degradation products are likely to include another oxidising agent, possibly oxygen. Nevertheless, the primary oxidising agent in the stripper is probably nitrite, which is added upstream of the spiral heat exchanger, at a target level of 30 ppm, to prevent polymer deposition on the plates of the spiral heat exchanger.

4. Experimental 4.1. Metallurgical examination Dye penetrant examination and radiography of the failed vessel revealed a number of defects associated with the welds. These included surface undercut, crater cracks and over roll. Some of these defects would not have been acceptable to the original fabrication standard. Approximately half of the circumferential and longitudinal butt welds in the vessel contained traverse cracking. Some of these cracks extended into the heat affected zone and some into the parent material. Long traverse cracks were often associated with pitting in the heat affected zone. These pits were jagged and had flat bottoms and therefore cannot be considered as classic chloride induced pits. Similar traverse cracks and pits were seen in the tray support ring fillet welds. Circumferential butt weld 3 additionally contained cracking at the toe of the weld running parallel to the weld direction. A section of longitudinal weld adjacent to circumferential weld 5 and circumferential welds 3 and 5 were cut out of the vessel and therefore were available for metallurgical examination. Samples of the shell material, circumferential weld 3 and the longitudinal shell weld adjacent to circumferential weld 5 were chemically analysed using atomic emission spectroscopy and wet LECO measurements. The results of these analyses are given in Table 1. These results indicate that the composition of the shell material complies with the requirements of UNS S32750. The weld metal compositions are in line with what would be expected from a AWS25.10.4L type electrode or similar. A tray support ring had been welded to the vessel wall just above circumferential weld 3, the distance between the adjacent toes of the butt and fillet welds being approximately 5 mm. The weld reinforcement of

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circumferential weld 3 on the inside of the vessel had been ground to flatten the weld reinforcement, leaving an untidy pattern of grinding marks. However, the reinforcement had not been fully removed and averaged about 4 mm across the weld. These grinding marks were covered with an adherent pink deposit, as was the rest of the inside of the vessel and so it can be assumed that the grinding marks had been introduced by the OEM. Weld spatter, which was to be removed according to the fabrication specification, was visible around the weld region (Fig. 2). Circumferential weld 5 was close to a tray support ring, the distance between the adjacent toes of the fillet and butt welding being 45 mm. The surface of circumferential weld 5 on the inside of the vessel was in the as-welded condition. It had a single capping pass and approximately 4 mm of weld reinforcement. Some remedial grinding had been performed to remove spatter. Metallographic samples were prepared from sections removed from the longitudinal weld adjacent to circumferential weld 5 and circumferential welds 3 and 5 to examine the microstructure in the cracked regions. The samples were etched in a mixture of nitric and hydrochloric acids in methanol and examined using an Olympus BM60X optical microscope and a Philips XL20 scanning electron microscope. 4.2. Corrosion testing Sample materials for this work were taken from various locations in the failed VCM stripper as follows: S32750 parent material; circumferential weld 3, circumferential weld 5, and a longitudinal weld. In addition, three reference butt welds were manufactured in S32750 with shielded metal arc weld runs, using heat input values of approximately 0.8, 1.6 and 2.4 kJ/mm, respectively.

Table 1 Chemical analyses of the shell material, circumferential weld 3 and a longitudinal weld, wt.%

Parent Long Circ

Cr

Ni

Mo

C

S

P

Mn

Si

Cu

Al

Sn

V

Ti

N

24.2 24.2 24.1

6.5 8.7 8.2

3.71 3.82 3.83

0.020 0.019 0.020

0.001 0.001 0.001

0.016 0.013 0.014

0.92 0.38 0.52

0.32 0.5 0.45

005 0.1 0.08

0.016 0.036 0.030

0.005 0.011 0.010

0.14 0.088 0.1

0.007 0.006 0.007

0.272 0.216 0.248

Fig. 2. A section of circumferential weld 3 with the tray support ring fillet weld at the top of the picture.

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For critical pitting temperature (CPT) measurements, the test samples were made from sections cut from the (welded) plate material such that the total surface area exposed to the test environment was 10–20 cm2. All faces of the samples were exposed to the test solutions, so that all orientations of the microstructure were tested simultaneously. Weld samples were prepared so that parent, HAZ and weld metal were all included in the tested area. All surfaces of the test electrodes were wet abraded to a 220-grit finish and rinsed with distilled water prior to testing. The CPT was measured by obtaining anodic polarisation curves in de-aerated 1M NaCl at different temperatures (from 40 to 80  C) using a sweep rate of 0.5 mV s1. Breakdown of passivity in each test was indicated by a sharply increasing anodic current, and the ‘‘breakdown potential’’, Eb, can be defined as the potential where the anodic current density reached 100 mA cm2. By plotting Eb as a function of temperature, the CPT can be easily identified to within 5 or 10  C [2,3]. Another series of experiments was then carried out, in which the test electrodes were mounted in epoxy resin so that only one face of the sample was exposed and the different orientations of the microstructure could be tested individually. The test surface was polished to a 1 mm diamond paste finish, and nail polish was used to mask some of the surface so that only  1 cm2 was exposed to the test solution in each experiment. A temperature of 86  2  C was used in all tests, and breakdown potentials were measured as above. After each of these tests, the sample surface was etched using a two-step electrolytic process: 20 s in 10 wt.% oxalic acid solution at  2 A cm2, followed by 30 s in 20 wt.% sodium hydroxide solution at  2 A cm-2 . The surface was then examined using an optical microscope at up to 400 magnification, and the position of pits, with respect to the microstructure, was recorded.

5. Results 5.1. Metallurgical examination The metallographic examination of the circumferential welds showed that the weld metal and HAZ had a ferrite/austenite phase balance within the range considered acceptable for fit for purpose welds. The same was true for the longitudinal weld, although this weld contained more austenite than ferrite in the weld metal. The HAZ transformed microstructure was made up of large grains of ferrite with grain boundary and secondary austenite. A small percentage, no more than 2%, of third phase was noted in the weld metal and HAZ of the root run in both circumferential welds. In areas away from significant cracking, there was a small amount of preferential corrosion at the toe of the weld on the inside surface of the vessel, extending only a few micrometres into the material. Both circumferential welds contained transverse cracks in the weld metal, starting from the inside surface of the weld and propagating all of the way through the vessel wall in some places. The majority of these cracks initiated at the toe of the weld and some were associated with large pits at the weld/HAZ boundary. The cracks were open at the surface, filled with corrosion product and branched. The cracks ran along the ferrite/austenite phase boundaries with some preferential corrosion of the austenite phase. The crack crossed the continuous ferrite phase, taking the shortest route from one austenite/ferrite boundary to the next as it progressed through the plate wall thickness. In some regions the transverse cracks propagated through the HAZ and parent material, particularly in the area between the wide root pass and the narrow second pass. In these regions, although the general direction of the crack was through wall, the crack branched significantly to run parallel to the surface of the plate, following the ferrite/austenite phase boundaries in the parent material. On the inside of the vessel there were large pits at the toe of circumferential weld 3. Transverse sections through these pits showed that they had a depth/width ratio < 1 and that the pits were full of corrosion product. Preferential corrosion of austenite had occurred at the periphery of each pit, Fig. 3. A number of

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Fig. 3. Preferential corrosion adjacent to a pit at the toe of circumferential weld 3, 200.

Fig. 4. Crack running from a pit at the toe of circumferential weld 3, 500.

cracks radiated from each pit with one crack running through the HAZ, parallel to the fusion boundary. Other cracks ran out from the pits, turning to run away from the weld, parallel to the surface of the steel. These cracks were branched, ran along the austenite/ferrite phase boundary and had areas of preferential austenite corrosion associated with them, Fig. 4. Some of the cracks, initiated at the toe of the weld, ran through the parent material parallel to both the plate surface and the weld along the weld toe. In one instance, a second crack ran in the opposite direction

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so that the cracks almost met. When these cracks meet, there is a chance that the material bounded by them and the plate surface could be removed, forming a large flat-bottomed pit adjacent to the toe of the weld. Microsections through fillet welds joining the tray support ring to the vessel wall revealed that the welds had a small throat thickness and failed to penetrate into the parent materials where they intersected. The weld metal microstructure of these fillet welds had an.approximately equal ratio of ferrite and austenite. No deleterious phases were observed in either the weld metal or the HAZ of these welds. The reference butt welds had a weld structure similar to, but with narrower and deeper passes than, circumferential welds 3 and 5. This difference can be explained by the fact that the reference welds were produced using SMAW whereas the vessel welds were fabricated using a combination of SMAW/GMAW and SAW. The reference welds had microstructures composed of approximately equal proportions of ferrite and austenite, with no deleterious phases. 5.2. Corrosion testing The Eb values for all samples are shown as a function of temperature in Fig. 5. From these data, the approximate CPT values can be obtained (Table 2). An extensive series of breakdown potential measurements was also made at 86  C, which is well above the CPT for all the materials considered. Although the samples in these experiments had a very smooth surface finish, anodic breakdown in all cases was due to pitting or, in relatively few cases, crevice corrosion at the sample edges. When substantial crevice corrosion was observed, the electrochemical data typically also showed a relatively slow increase of the anodic current around the breakdown potential. The results from these tests were discarded, such that the data plotted in Figs. 6 and 7 are considered pitting potentials, Epit. N samples were successfully tested for each experimental condition, and the results are presented here as the cumulative probability (n/N) of failure, where n is the number of failed samples at any given potential. Thus, each data point represents the Epit of an individual sample. Fig. 6 shows the probability of failure vs. potential for parent S32750, in comparison with results for samples from the longitudinal and circumferential welds. Clearly, the weld samples show significantly lower pitting potentials than the parent material, but there is little difference in the results for the two different welds. This is consistent with the CPT results in Fig. 5. In Fig. 6, the microstructural orientation of the sample surface has been ignored. However, the different faces of the samples were tested individually, and the results for the parent samples are shown in Fig. 7. This figure suggests a small influence of orientation, but there is not enough data to justify detailed interpretation of this result. Post-test examination of welded samples revealed that most large pits were in the weld metal or at the boundary between the weld and the parent material. It was generally not possible to positively identify the initiation sites of pits in these areas. However, in the surrounding parent metal, and on parent samples, it was usually possible to identify the initiation sites of many small pits, with the results shown in Fig. 8. In a few tests, crevice corrosion initiated at the sample edges beneath the nail polish used to mask off the surrounding surface area. In a couple of these cases, where the extent of corrosion was limited, there was clearly preferential attack along the grain boundaries, and possibly a slight tendency to preferential corrosion of the ferrite phase—see Fig. 9.

6. Discussion The cracking of the stripper was clearly initiated at weldments, and it is well-known that duplex stainless steels are susceptible to formation of intermetallic compounds during welding or heat treatment, and that these compounds can be detrimental to corrosion and cracking resistance. ASTM A923 [16] describes

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Fig. 5. Breakdown potentials in 1 M NaCl as a function of temperature for welded samples and parent S32750. Top: Weld samples cut from the failed vessel. Bottom: Reference welds with varied heat input.

Table 2 Measured CPT values for various samples of S32750 Origin of weld test plate

Details

Microstructure

CPT ( C)

Samples cut from failed vessel

Parent circ weld 3 circ weld 5 long weld 0.8 kJ/mm 1.6 kJ/mm 2.4 kJ/mm

Even a/g. No s phase. Even a/g s phase. Even a/g. s phase. Even a/g. No s phase. Even a/g. No s phase. Even a/g. No s phase. Even a/g. No s phase.

70–75 50–60 50–60 50–60 40–50 40–50 40–50

Reference butt SMAW welds

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Fig. 6. Comparison of pitting potentials measured at 862  C for S32750 samples (cut from the VCM stripper) with a 1 mm diamond finish and an exposed surface area of about 1 cm2. The microstructural orientation in the exposed sample surface has been ignored.

Fig. 7. Pitting potentials of parent material as shown in Fig. 6, but showing the influence of the microstructural orientation of the tested surface: Face A is the internal vessel wall surface; Face B is a cross-section parallel to the circumferential weld run; Face C is a cross-section parallel to the longitudinal weld run.

various methods for detecting the presence of these detrimental phases, one of which involves measurement of the CPT. This standard specifies minimum acceptable values for S31803 and S32205, but not for S32750. In this work, we have used CPT measurements in an attempt to detect and quantify the influence of detrimental phases in and around the welds in the stripper, and we have compared our results against those published by others. The CPT we obtain for parent S32750 is slightly lower than the general literature value of about 80  C, suggesting that our test method is conservative, but reasonable. The measured CPT values for all welded samples are between 40 and 60  C, which is similar to the highest values reported in

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Fig. 8. Location of pit initiation sites in samples tested at 86  C in 1 M NaCl (all orientations): s indicates a pit that initiated within a ferrite grain; g indicates a pit initiated within an austenite grain; gs, ss and gg indicate pits initiated at grain boundaries; and ‘‘BIG’’ represents those cases where the pits were too large for their initiation site to be identified.

Fig. 9. Preferential grain boundary attack at crevice corrosion site on parent S32750 sample tested in 1 MNaCl at 86  C, 100.

the literature. For example, Wallen et al. [17] found 45–50  C for gas tungsten arc welds and 45–62  C for SMAW welds. On the other hand, Steinsmo et al. [18] report values as low as 32  C. We conclude that the failed welds in the VCM Stripper have a lower resistance to localised corrosion than the parent material, but that the localised corrosion resistance was as good as could reasonably be expected for welded S32750. It is also important to note that the CPT test does not directly assess SCC resistance. However, we do know that s phase formation is detrimental to both pitting and SCC resistance. Nilsson and Wilson [14]

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carried out carefully controlled heat treatments and detailed microstructural studies to show the effect of s phase on the CPT of S32750. A brief summary of their results is shown in Table 3. Comparing these data with our results, it is apparent that only a small amount ( 2 vol.%) of s phase (or possibly Cr2N) is required to cause the observed reductions in pitting resistance at welds. In early duplex steels, the segregation of Cr and Mo to the ferrite phase ensured preferential pitting of the austenite. More recently, it has been claimed [14,15] that the pitting resistance of each phase in S32750 is the same, such that preferential pitting of one phase should not normally occur. However, modern 25Cr duplex steels with high N levels, similar to S32750, have generally been found to suffer preferential pit initiation in the ferrite phase [19–21]. In this investigation, our corrosion tests revealed a clear tendency for preferential pit initiation at sg grain boundaries, although some pits did initiate within both s and g grains, Fig. 8. These results are consistent with the idea that tertiary phases, such as s and Cr2N, are the most likely pit initiation sites in this material since these phases are generally formed at grain boundaries. There was also some evidence of preferential intergranular attack at crevice corrosion sites, as shown in Fig. 9, which is consistent with localised chromium depletion of the matrix due to formation of chromium rich tertiary phases. The relatively high SCC resistance of duplex steels is primarily due to the different partitioning of alloy elements such that one phase usually cathodically protects the other in the aggressive local environment at the crack tip [10,22]. In early duplex stainless steels, e.g. those made before 1970, the austenitic phase had a lower corrosion potential at the crack tip and so cathodically protected the ferrite. Hence, SCC in these materials propagated through the austenite and was blocked by the ferrite [10]. However, modern duplex stainless steels are highly alloyed in ways that have tended to ennoble the austenite phase. Consequently, it is now more usual for cracks to run through the ferrite and be blocked by protected austenite [10,22,23]. The maximum in SCC resistance at 50% ferrite is probably due to this structure requiring the most tortuous path for a planar crack to progress through one phase alone. Our metallurgical examination of the crack path in the failed vessel showed predominantly intergranular crack propagation. However, there was a strong tendency for preferential corrosion of the austenite, with the cracks generally making only very short jumps across the continuous ferrite phase. These results are very similar to those reported by Notten [2], who found that cracks originated in welds and propagated almost exclusively in the austenitic phase in both the weld metal and the parent material. Both cases are contrary to the expectation, based on recent literature, that cracks in super-duplex stainless steels would tend to propagate mainly in the ferrite. There were several significant differences in weld fabrication technique between circumferential welds 3 and 5 and the other circumferential welds. Welds 3 and 5 were completed in four or five passes, with the first pass being deposited from the inside of the vessel. This contrasts with the other circumferential welds, which had the last pass deposited from the inside of the vessel. Therefore, the second pass in welds 3 and 5 would have reheated the weld and HAZ of the first pass and, depending on the interpass temperature and the time taken to complete the weld, provided the potential for formation of small amounts of third phases in the region of the weld pass on the vessel interior. The quantity of third phase formed, less than 5%, generally would not be considered to be sufficient to degrade the corrosion resistance of the welds. However, the corrosion testing has shown that even small amounts of third phase in welds can have a Table 3 Influence of tertiary phases on the CPT of S32750 [14] Volume fraction of s phase (%)

CPT ( C)

<0.2 1.3 1.9 7 to 9 0—but Cr2N formed instead

80 65 60 <40 <65

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deleterious effect on the performance of materials in marginal environments. The size of the weld passes in welds 3 and 5, and comparisons with the reference butt welds, indicate that these welds were fabricated using a heat input in excess of 2 kJ/mm, higher than the recommended maximum heat input of 1.5 kJ/ mm generally cited for welding super duplex stainless steel. Weld over roll defects were found associated with welds 3 and 5, exceeding the limits permitted by the standard. Finally, welds 3 and 5 had up to 4mm of weld reinforcement on the inside of the vessel, with a pronounced stress concentration at the toe of the weld. These factors, namely the existence of a deleterious third phase, weld defects, and the presence of stress concentrators, combined with residual stresses inherently associated with the welds, made welds 3 and 5 particularly susceptible to SCC. Nevertheless, the failure of S32750 in the VCM stripper environment remains difficult to reconcile with published laboratory testing results. Cracking at lower than expected temperatures could possibly be explained by the presence of very high chloride concentrations, but the nominal chloride concentration was low in the failed stripper and localised evaporative concentration was unlikely. The cracking might also have been explained by the presence of poor weld microstructures, but the welds in this case would have passed most reasonable acceptance criteria in terms of both their microstructure and their corrosion resistance. We have previously argued that the relatively high oxidising power of the stripper environment may explain the failure, and this remains, in our opinion, the most plausible explanation. However, further work is clearly required to more accurately define the range of safe working conditions for welded S32750, together with appropriate welding guidelines and acceptance criteria.

7. Conclusions

1. The VCM stripper vessel failed due to a combination of localised corrosion and stress corrosion cracking. 2. The volume fraction of third phase present in the failed welds was within the limits generally considered insignificant for super duplex stainless steel in service. 3. In corrosion tests, all welded samples performed worse than the parent material. However, based on comparison with the literature, the CPT values of the welded samples were as high as could reasonably be expected. 4. Corrosion testing did not reveal any significant difference between results for the different welds, including those that showed substantially different performance in-service. 5. The microstructural variations between the welds in the vessel that cracked and those that did not are small, indicating that factors such as fabrication methodology and weld finishing, in addition to the presence of a deleterious third phase in the microstructure, contributed to the failure. 6. The SCC of duplex S32750 in this environment would not have been expected based on the published results of SCC tests, and the microstructure of the failed welds would have passed most reasonable acceptance criteria. Therefore, further work is required to more accurately define the safe-use limits for this material, with a particular focus on the condition of welds.

Acknowledgements We would like to thank P.T. Wilson and C. Jones for their substantial contributions to the original failure investigation, and Y.L. Siew and M. Sinathuraja for assistance with metallography. The original investigation was funded by Australian Vinyls Corporation, whilst follow-up work at MPT was funded by the New Zealand Foundation for Research, Science and Technology under contract CO8X0012.

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