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Use of supermartensitic stainless steel pipe for offshore flowline applications (June 2006)

   
P Woollin and A Kostrivas

Proceedings of OMAE 2006 25th International Conference on Offshore Mechanics and Arctic Engineering
June 4-9, 2006, Hamburg, Germany. Paper OMAE2006-92351

Abstract

Supermartensitic stainless steels (SMSS) have high strength and good resistance to corrosion in produced fluids containing CO 2 and are cheaper than other competing corrosion resistant alloys. Hence, they are attractive flowline materials and they have been successfully used in a number of offshore applications. Nevertheless, service failures have occurred and two failure mechanisms in particular have caused difficulties at welds: (i) hydrogen embrittlement/ hydrogen induced stress cracking resulting from hydrogen picked-up under cathodic protection and (ii)intergranular stress corrosion cracking (IGSCC). This paper presents experimental data on each of these two failure phenomena and gives details of the currently available ways of avoiding these problems, highlighting where further information is required.

Introduction

Supermartensitic stainless steels typically contain approximately 12%Cr, 1.5-7%Ni, up to 2.5%Mo and about 0.01%C. Two classes of SMSS may be identified: 'lean' grades, with typically <4%Ni and no molybdenum addition, and 'high' grades, with >4%Ni and 2-2.5%Mo. The microstructure of such steels is predominantly tempered martensite with some austenite but delta ferrite and untempered martensite may form in weld HAZs.

Subsea SMSS components are subjected to cathodic protection in practice due to concerns over pitting corrosion resistance on the seawater side. Typically, in-service potentials may be of the order of -1000 to -1100mV SCE . At such potentials, generation of atomic hydrogen at the steel surface may lead to significant hydrogen pick-up. The ferrite and martensite phases are generally susceptible to hydrogen embrittlement, whilst austenite is much less susceptible.

Hydrogen induced stress cracking (HISC) has caused failures in subsea SMSS pipelines as a result of embrittlement of HAZs and matching composition weld metal. [1,2]

Recent failures of lean grade material in service and cracking of both lean and high-alloy grades during qualification testing in hot acidic brine have revealed sensitivity to intergranular stress corrosion cracking (IGSCC) in the HAZ of girth welds, [3-5] although no problems have been reported to date in high alloy SMSS flowlines.

Experimental programme

HISC Tests

Materials

Testing was performed on parent material taken from 275mm OD x 15mm wt supermartensitic UNS S41426 pipe, with nominal composition 12%Cr, 6%Ni, 2.5%Mo with about 0.1%Ti, Table 1. Further testing was performed on cross-weld specimens from the cap side of two girth welds in the pipe: (i) an automatic pulsed GMA weld (one half as-welded and the other given brief postweld heat treatment (PWHT))and (ii) a manual GTA weld, which was given brief PWHT. Postweld heat treatment of the welds was carried out by induction heating to within the temperature range 630°C-650°C for 5 minutes.

Table 1 Chemical analyses of the supermartensitic steel pipes used

Pipe nominal compositionElement, wt%
CNSiMnCrNiMoCuTi
12Cr6Ni2.5MoTi* 0.009 0.005 0.20 0.43 12.2 6.4 2.51 0.03 0.12
12Cr5Ni2Mo 0.013 0.009 0.12 0.54 11.8 5.1 2.03 0.04 <0.005
12Cr6Ni2Mo 0.010 0.011 0.17 0.18 12.4 5.8 2.18 0.03 0.020
13Cr5Ni1Mo 0.013 0.006 0.16 0.65 13.5 5.1 0.78 0.03 0.088
11Cr1.5Ni0.5Cu 0.010 0.006 0.18 1.14 10.9 1.6 <0.01 0.49 0.01

* UNS S41426

Welding of the pipe was carried out with superduplex filler wire and following normal superduplex heat input and interpass temperature limits (i.e. 0.5-1.5kJ/mm and <150°C).

HISC Test details

Parent material

Constant load tensile tests were performed on smooth round cross-section specimens machined from the parent supermartensitic pipe (P1-P5). The gauge section dimensions were 25mm long x 3.8mm diameter. Testing took place at ambient pressure and temperature (approximately 15-20°C), in seawater, with polarisation to -1100mV SCE , following a 14-day pre-charge under similar conditions but without stress. Specimens were loaded to give stresses between 0.96 and 1.15 times the measured parent steel 0.2% proof stress (735MPa). Testing was for nominally 30 days.

GMA Weld, as-welded

Similar constant load tensile tests were performed on cross-weld specimens machined from the cap side of the GMA weld, in the as-welded condition (GMAW [1-3] ). Testing was performed at ambient pressure, at 4°C, in 3.5wt% NaCl solution, with polarisation at -1100mV SCE . Three specimens were tested following 14, 21 and 25 days pre-charging, respectively. The stress levels were 1, 1.08 and 1.08 times the parent steel 0.2% proof stress.

GMA Weld, after PWHT

Further constant load tensile tests were performed on three cross-weld specimens machined from the cap side of the GMA weld, in the PWHT condition (GMAW HT1-3). Testing took place at ambient pressure, 4°C, in 3.5wt% NaCl solution, with polarisation to -1100mV SCE . The stress level for all tests was the parent steel 0.2% proof stress. Further testing took place at 30bara in an autoclave pressurised using a pre-mixed N 2 /O 2 mixture to obtain 0.21 bar O 2 partial pressure, at 4°C, in 3.5wt% NaCl aqueous solution, at -1100mV SCE . Two specimens (GMAW HT [4,5] ) were tested, following a 14-day pre-charge. The stress levels were 0.91 and 1.00 times the 0.2% proof stress of the parent steel.

GTA Weld, after PWHT

Three cross-weld specimens were machined from the GTA weld cap (GTAW HT [1-3] ) and examined by ultrasonic testing and penetrant testing to ensure freedom from detectable flaws. The specimens were tested at the parent steel 0.2% proof stress. Testing took place at ambient pressure, at 4°C, in 3.5wt%NaCl solution, with polarisation to -1100mV SCE . The samples received a brief pre-charge of approximately one-hour, prior to loading. Further testing on three specimens (GTAW HT [4-6] ) took place at 30bara, at 4°C, in 3.5wt% NaCl aqueous solution, at -1100mV SCE .

Intergranular SCC tests

Materials

Five low carbon SMSS pipes were used, Table 1. Compositions ranged from 10.9-13.5%Cr, 1.5-6.4%Ni and 0-2.5%Mo for the steels examined. Two types of automatic pulsed GMA girth weld were examined, (i) three welds made with superduplex solid filler wire throughout(W1-W3) and (ii) two made with approximately-matching composition metal cored filler wires (W4 and W5). Table 2 summarises the welding matrix.

Table 2 Pulsed GMA girth welding matrix for IGSCC tests

Weld codePipeWelding wireShielding gasPWHT
  
W1 12Cr6Ni2.5MoTi 25%Cr Ar/He/CO 2 /N 2 None
W1P 12Cr6Ni2.5MoTi 25%Cr Ar/He/CO 2 /N 2 650°C/5min*
W2 12Cr5Ni2Mo 25%Cr Ar/He/CO 2 /N 2 None
W2P 12Cr5Ni2Mo 25%Cr Ar/He/CO 2 /N 2 650°C/5min*
W3 12Cr6Ni2Mo 25%Cr Ar/He/CO 2 /N 2 None
W3P 12Cr6Ni2Mo 25%Cr Ar/He/CO 2 /N 2 650°C/5min*
W4 13Cr5Ni1Mo 1.5%Mo SMSS Ar+0.5%CO 2 None
W4P 13Cr5Ni1Mo 1.5%Mo SMSS Ar+0.5%CO 2 650°C/5min*
W5 12Cr6Ni2Mo 2.5%Mo SMSS Ar+0.5%CO 2 None
W5P 12Cr6Ni2Mo 2.5%Mo SMSS Ar+0.5%CO 2 650°C/5min*

Examples of each weld type (which were given a 'P' designation) were subjected to brief PWHT by induction heating. The specified heat treatment cycle was 5 minutes at 640-660°C at the weld cap. The root temperature was typically 15-35°C less than the cap temperature, i.e. 620-640°C, depending on the pipe wall thickness. The effect of PWHT on hardness was examined, using 5kg Vickers hardness measurements on transverse weld sections.

Corrosion test details

Four point bend tests were performed in two hot 25%NaCl solutions, acidified with 10bar CO 2 . The first test had a temperature of 110°C and calculated pH of 3.3, whilst the second test was at 120°C and had 500ppm NaHCO 3 added to raise the calculated pH to 4.5. Rectangular test specimens, 100x15x3mm, were taken cross-weld from each of the girth weld types, in both as-welded and PWHT conditions. Specimens had the root intact and were deflected to give a strain equivalent to the parent steel 0.2% proof stress in the HAZ.

Results

HISC Tests

None of the parent steel HISC test samples failed or cracked within 30 days, Fig.1 and Table 3. However, six of the eight GMA weld samples and two of the six GTA weld samples failed. The two GMA samples that did not fail were GMAW1, which was as-welded and loaded at the 0.2% proof stress of the parent steel(738MPa), and GMAW HT4, which was postweld heat treated and loaded to 0.92 of the proof stress (679MPa). The four GTA samples that did not fail included all three tested at 30 bara and one tested at 1 bara, all loaded to the parent steel 0.2% proof stress. Examination under a binocular microscope, of the specimens that did not fail, did not reveal any surface-breaking flaws or small cracks, Table 4.

Fig.1. Comparison of HISC test results for the various supermartensitic weldments, together with an indication of the pre-existing flaw type and size. No flaws were found in specimens that did not fail
Fig.1. Comparison of HISC test results for the various supermartensitic weldments, together with an indication of the pre-existing flaw type and size. No flaws were found in specimens that did not fail

Table 3 HISC test results for parent steel in seawater at ambient pressure and temperature, at -1100mV SCE .

SampleStress (MPa)Strain (%)Time to fail (hours)Result
P1 708 0.8 >720 DNF, no cracks or defects
P2 735 1.0 >720 DNF, no cracks or defects
P3 771 2.0 >720 DNF, no cracks or defects
P4 813 4.0 >720 DNF, no cracks or defects
P5 846 5.4 >720 DNF, no cracks or defects

DNF = did not fail

Table 4 HISC test results from cross-weld specimens in 3.5%NaCl at 4°C under cathodic protection at -1100mV SCE

Sample (test pressure, bara)Stress, MPa (fraction of 0.2% proof stress)Time to fail (hours)Result (max. flaw size, mm)
GMAW1 (1) 738 (1.00) >720 DNF, no cracks or surface defects
GMAW2 (1) 797 (1.08) 82 Cracking from small weld pore (0.05)
GMAW3 (1) 798 (1.08) 23 Cracking from small weld pore (0.05)
GMAW HT1 (1) 738 (1.00) 0.1 Cracking from SWLOF (1.5)
GMAW HT2 (1) 738 (1.00) 0.5 Cracking from 2 x SWLOF (0.3)
GMAW HT3 (1) 739 (1.00) 0.01 Cracking from SWLOF (2.0)
GMAW HT4 (30) 679 (0.92) >720 DNF, no cracks or surface defects
GMAW HT5 (30) 742 (1.01) 0.7 Cracking from 2 x SWLOF (0.3)
GTAW HT1 (1) 735 (1.00) 763 Cracking from weld pore/IRLOF (0.4)
GTAW HT2 (1) 738 (1.00) 148 Cracking from IRLOF (1.0)
GTAW HT3 (1) 738 (1.00) >720 DNF, no cracks or surface defects
GTAW HT4 (30) 739 (1.00) >720 DNF, no cracks or surface defects
GTAW HT5 (30) 739 (1.00) >720 DNF, no cracks or surface defects
GTAW HT6 (30) 723 (0.98) >720 DNF, no cracks or surface defects

SWLOF = side wall lack of fusion, IRLOF = inter-run lack of fusion.

Examination of the fracture surfaces of the GMA weld specimens revealed the presence of small pre-existing sidewall lack of fusion flaws (0.3-2.0mm long) and an adjacent brittle fracture morphology where the HISC crack had initiated, Fig.2, changing to a dimpled ductile appearance further away from the flaws. Cross-sections taken along the axis of the broken halves of the specimens showed that the brittle part of the fracture coincided with weld metal very close to and parallel to the weld fusion boundary, Fig.3 and 4. This implies HISC crack development in superduplex weld metal close to the fusion boundary, from small sidewall lack of fusion flaws. No effect of pressure or PWHT was apparent from the tests undertaken.

Fig.2. Fracture surface in superduplex weld metal of sample GTAW HT2, which failed from a weld metal flaw
Fig.2. Fracture surface in superduplex weld metal of sample GTAW HT2, which failed from a weld metal flaw
Fig.3. Fracture path in sample GMAW3
Fig.3. Fracture path in sample GMAW3
Fig.4. Brittle fracture surface in superduplex weld metal close to fusion line of sample GMAW3
Fig.4. Brittle fracture surface in superduplex weld metal close to fusion line of sample GMAW3

Examination of the GTA weld specimen fracture surfaces in an SEM revealed small pores (~0.05mm across) and interrun lack of fusion flaws (~0.4-1.0mm across), lying parallel to the sample axis. Flaws with this orientation are difficult to detect by UT due to their small cross-sectional area. The fracture surface again had a brittle morphology close to the flaws, changing to a dimpled ductile appearance further away from the flaws. Examination of cross-sections showed that fracture had initiated from flaws in the weld metal and then progressed to the fusion line/HAZ and into the parent steel, Fig.5

Fig.5. Section along axis of specimen GTAW HT2, showing superduplex weld metal, with inter-run flaw
Fig.5. Section along axis of specimen GTAW HT2, showing superduplex weld metal, with inter-run flaw

ICSCC Tests and Effect of PWHT

Corrosion testing

Table 5 lists the results of the SCC tests. Most of the as-welded specimens showed intergranular cracking in the HAZ in environment A, at a variety of locations ranging from immediately adjacent to the fusion line (e.g. W1,12Cr6Ni2.5MoTi steel welded with superduplex wire) to about 0.5mm from the fusion line (W2, 12Cr5Ni2Mo steel welded with superduplex wire), Fig.6 and 7. No such cracking was found in any weld after PWHT. Weld W5 (12Cr6Ni2Mo steel welded with SMSS wire) showed no cracking in the as-welded condition. In the second environment, with pH raised to 4.5, similar trends were observed, i.e. as-welded specimens tended to crack and PWHT specimens did not. Crack location was similar to that in the environment A. Weld W3 (12Cr6Ni2Mo steel welded with superduplex wire) and W4 (13Cr5Ni1Mo steel, welded with matching composition wire) showed no cracks in the as-welded condition.

Table 5 Results of intergranular stress corrosion cracking tests.

Weld codeParent pipeMax root HAZ HV10PWHTA (pH 3.3, 110°C)B (pH 4.5, 120°C)
W1 12Cr6Ni2.5MoTi 327 No HAZ cracks HAZ cracks
W1P 12Cr6Ni2.5MoTi 313 Yes No cracks No cracks
W2 12Cr5Ni2Mo 345 No HAZ cracks HAZ cracks
W2P 12Cr5Ni2Mo 345 Yes No cracks No cracks
W3 12Cr6Ni2Mo 332 No Small HAZ crack No cracks
W3P 12Cr6Ni2Mo 319 Yes No cracks No cracks
W4 13Cr5Ni1Mo 351 No HAZ cracks No cracks
W4P 13Cr5Ni1Mo 327 Yes No cracks No cracks
W5 12Cr6Ni2Mo 347 No No cracks Shallow HAZ cracks
W5P 12Cr6Ni2Mo 312 Yes No cracks No cracks
Fig.6. Intergranular cracking in the HAZ of weld W1 (pipe A, 12Cr5Ni2Mo), environment A
Fig.6. Intergranular cracking in the HAZ of weld W1 (pipe A, 12Cr5Ni2Mo), environment A
Fig.7. Detail of cracking close to the fusion boundary of weld W3 (pipe C, 12Cr6Ni2.5MoTi), environment A
Fig.7. Detail of cracking close to the fusion boundary of weld W3 (pipe C, 12Cr6Ni2.5MoTi), environment A

Effect of PWHT on hardness

In general, the as-welded weld root/mid thickness position had higher maximum HAZ hardness than the weld cap, presumably reflecting effects of reheating/straining. Maximum HAZ values, as-welded, were in the range 332-351 HV5, with the highest hardness always being about 2mm from the fusion line. After brief induction PWHT, peak HAZ hardness was typically reduced at the root position, by up to 54 HV5 for lean grade steel but more typically by 10-15 HV5 for the higher alloy grades. The weld cap HAZs showed a mixed response with hardening observed in some cases (up to +12 HV5) and softening (up to -26 HV5) in others.

Discussion

HISC

Parent supermartensitic stainless steel (12Cr6.5Ni2.5Mo) pipe did not show any sensitivity to HISC in 30 day constant load tensile tests in seawater at -1100mV SCE , at stresses up to 1.15 times the 0.2% proof stress, which is approaching the UTS. Similar testing on superduplex stainless steel hub material has previously shown that it is susceptible to HISC at stresses equal to or above 96% of the 0.2% proof stress in similar tests. [6] It is noted however that HISC is reported to require dynamic plastic strain, which ruptures the passive film. [7] In duplex and superduplex stainless steel, low temperature creep processes induce dynamic plastic strain under constant load conditions but supermartensitic steels showed little low temperature creep in these tests. This probably explains why the parent steel samples did not crack here, as previous work has shown that parent SMSS is susceptible to HISC under dynamic test conditions. [7]

Cross-weld specimens taken from girth welds in supermartensitic pipe, made with superduplex wire, were sensitive to HISC from weld flaws when stressed to around the 0.2% proof stress of the parent pipe. Various flaw types were found including weld metal pores, sidewall lack of fusion (GMA weld) and inter-run lack of fusion (GTA weld). The HISC failure path in smooth specimens was initially through superduplex weld metal. The failure path was in weld metal very close to the fusion line for sidewall lack of fusion flaws. Cracks then ran into the supermartensitic HAZ and parent steel. Specimens that did not contain flaws survived 30-day tests without any evidence of HISC.

Cracking in the weld metal is perhaps not surprising given that the applied stresses were based on the proof stress of the supermartensitic pipe, and were in the range 679-798MPa, which is probably of the order of, or above, the proof stress of the weld metal. Proof strength for undiluted weld metal is typically in the range 700-750MPa. Other authors have previously observed HISC fractures initiated in the HAZ during SSRT testing under cathodic protection of cross-weld specimens from supermartensitic welds, perhaps reflecting that the weld metal strength was over-matching compared to the parent pipe in that instance and probably an effect of dynamic strain on the susceptibility to cracking of the SMSS. [7]

When flaws were present, a brittle fracture face characteristic of a ferritic-austenitic microstructure was found in weld metal, with cleavage through ferrite grains and evidence of where the crack has passed through the austenite (Fig.2). However, for cracking from lack of sidewall fusion, when crack propagation was along the interface, a different fracture morphology was found, which was not typical of a HISC fracture in a ferritic-austeniticmicrostructure ( Fig.4). This is consistent with the fact that the fusion line region, where a partially mixed zone (PMZ) exists, does not have the same composition and microstructure as the bulk weld metal. The fusion boundary area/PMZ has a different distribution of ferrite and austenite to the bulk weld metal and may contain martensite.

This programme did not examine in detail how cracking initiated in the weld metal or at the fusion line and then progressed into the HAZ and parent steel. However, practical experience of failures in service [6,7] indicates that initiation at a weld toe (i.e. in the HAZ, on the fusion boundary or in weld metal close to the fusion line) followed by propagation predominantly in the parent steel is a failure mode of particular concern. Consequently, there is a need to define the parameters contributing to susceptibility to HISC in this region and whether PWHT is of benefit. It was not possible to draw a definitive conclusion on the effect of PWHT here due to the over-riding effect of weld flaws but the PWHT examined clearly did not prevent HISC from running into a HAZ. Also, the fact that only a modest change in HAZ hardness occurred suggests that PWHT is unlikely to have a strong effect on HISC. Similarly, no clear effect of test pressure could be discerned as a result of the influence of weld flaws.

Several specimens without flaws did not fail in constant load tests but, as described previously, this does not imply that defect-free welds will be immune to HISC under dynamic strain conditions. It does however indicate that the presence of stress concentrators, whether weld flaws or discontinuities at a weld toe, will encourage HISC. Hence, minimising the presence of such features is recommended and if practical, all welds should be coated to prevent hydrogen pick-up.

IGSCC

The sensitisation of lean grade SMSS HAZs has been linked to the formation of Cr-carbides on prior-austenite grain boundaries and adjacent Cr-depleted zones but this link has not been established for the high alloy grades. [8,9] However, formation of Cr-depleted zones on prior-austenite boundaries immediately underneath the welding oxide has been observed in high alloy grades. [10] Hence some uncertainty remains over the mechanism of IGSCC of high grade supermartensitic stainless steel and the effect of PWHT. Nevertheless, there is a substantial body of information supporting a consistent beneficial effect of brief PWHT for a broad range of supermartensitic grades. [8,9,11]

Based on the available data, it is recommended that PWHT should be applied to welds in supermartensitic stainless steel where there is a risk of intergranular SCC in service, i.e. in hot acidic environments. However, more data are required to fully confirm this beneficial effect and the control required. A PWHT temperature of about 650°C at the root is recommended based on current data and the heat treated zone should encompass the whole of the weld metaland HAZ. The maximum and minimum allowable temperatures have not been established but the current work extended from 620°C up to 660°C. Heating and cooling should be fairly rapid. The most appropriate PWHT duration has not been established but there is fairly common agreement that 5 minutes is an appropriate duration, as used here. Whilst the beneficial effect of PWHT with respect to IGSCC has been demonstrated for 30 day exposure tests, longer term data are required to confirm the applicability of the effect to long term service.

Postweld heat treatment may also have detrimental effects if not adequately controlled, e.g. (i) thickening of weld area oxides and associated loss of general/pitting corrosion resistance, (ii) formation of virgin martensite in theHAZ and increased hardness leading to reduced toughness and resistance to sour environments, (iii) loss of toughness in superduplex stainless steel weld metal and (iv) tempering of HAZ at temperatures that could induce sensitisation to intergranular SCC if the heat treated area is not wide enough.

The precise response to PWHT is specific to each individual grade of supermartensitic steel, although the data indicate that all steels examined here were fairly similar and the beneficial effect of 5 minutes at 620-660°C, withrespect to IGSCC, is applicable to 'lean' grades, with <1%Mo and 'high' grades with >2%Mo both with and without Ti addition.

Due to the limited information available, the use of welded supermartensitic stainless steel in the PWHT condition will require qualification on a case by case basis. The qualification programme should consider the effects of PWHT on toughness and sour service performance, in addition to IGSCC and address the extremes of the range of PWHT thermal cycles that may be experienced in practice.

Conclusions

When developing welding procedures for supermartensitic stainless steel flowlines, efforts should be made to minimise the risk of hydrogen induced stress cracking due to hydrogen pick-up from cathodic protection and intergranular stress corrosion cracking of the HAZ in hot, acidic fluids.

Wherever possible, welds in supermartensitic steel should not be exposed to cathodic protection at around -1000 to -1100mV SCE , particularly where dynamic plastic strain may be experienced in service. Where this is not possible, the occurrence of weld flaws should be minimised and efforts made to limit the stress concentrating effect of the weld toe. Further work is required to define the limiting loading conditions for the onset of HISC. There is no indication that brief PWHT at 650°C for 5 minutes reduces susceptibility to HISC.

The current work and other published studies indicate that brief induction PWHT at about 650°C for 5 minutes will remove sensitivity to intergranular stress corrosion cracking of the HAZ. Further work is required to establish the required control of the PWHT process to ensure a consistent beneficial effect but, based on the current work, an allowable temperature range of 620-660°C is suggested.

Acknowledgements

The following companies, who sponsored the work, are thanked for their permission to publish the results: Advantica, BP, Chevron, ENI, ESAB, Industeel, JFE Steel, J Ray McDermott, Lincoln Electric, Nickel Institute, Serimer Dasa,Shell, Statoil, Sumitomo Metal Industries, Tenaris and Total.

References

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