Hydrogen Induced Stress Cracking of Duplex Stainless Steel Subsea Components

8/10/2019 Hydrogen Induced Stress Cracking of Duplex Stainless Steel Subsea Components

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Hydrogen induced stress cracking of duplex stainless steel subseacomponents

Amir Bahrami and Paul WoollinTWI LtdCambridge, UK

Paper presented at 29th International Conference on Offshore Mechanics and Arctic Engineering (OMAE 2010), Shanghai, China, 6-11 June 2010.

Abstract

A small number of duplex and superduplex stainless steel components have failed in subsea service due to hydrogen induced stresscracking (HISC). The significance of these failures has led to research to define critical loading conditions for HISC, to allowconfident design of components in future. Data relating to the Foinaven superduplex hub failures were published at OTC in 1999 and

NACE Corrosion conference in 2001 and data from TWI Group Sponsored Projects were published at OMAE in 2004. DNV RP F112has been based on these and other data, to provide a conservative approach to design.

There are a number of gaps in the published literature and in the data available when DNV RP F112 was prepared, related todifferences between small scale specimens tested at typical seabed temperature and actual components in operation, ie theoperating temperature and pressure, and data from full-scale tests on pipe material with fine austenite spacing and hence goodresistance to HISC.

The paper presents new data on these issues and indicates where technology gaps remain.

Introduction

A small number of duplex and superduplex stainless steel components have failed in subsea service due to hydrogen induced stresscracking (HISC). The significance of these failures has led to research aimed at defining critical conditions for HISC, to allow

confident design of components in future. Data relating to the Foinaven superduplex hub failures were published at OTC in 1999[1]

and the NACE corrosion conference in 2001

[2]

and data from other twi investigations were published at OMAE in 2004.

[3]

The susceptibility to HISC shows a strong correlation with austenite spacing and tests performed on smooth samples have shown

that coarse-grained microstructures are more susceptible to HISC.[3] Small-scale constant load tests in seawater with cathodic

protection (CP) are generally used to characterise the HISC susceptibility of duplex and superduplex materials. [2] Constant load,tensile HISC tests on hub materials from the Foinaven field, at -1100mVSCE showed that if the superduplex material was loaded to

an initial strain of 0.5%, creep and crack initiation would ensue very quickly and failure would eventually occur at a substantially

higher total strain, as a consequence of low temperature creep.[1] Tests on full scale hubs have been reported,[1,2] which indicatedthat once cracks had initiated, they could propagate through thickness in about 10 days without further increase in applied load.Hence it was concluded that initiation of HISC must be avoided if failure is to be prevented, and the assessment criterion should be

related to the threshold initial stress or strain for initiation of HISC.[1]

DNV RP F112[4,5] has been based on these and other data, to provide a conservative approach to design. However data areavailable to show that some product forms including pipes are substantially more resistant to HISC than coarse grained forgings,due to finer austenite spacing. RP F112 does not allow advantage of this finer spacing to be taken reliably, as it is based on a

measurement of austenite spacing, which is not the subject of a recognised standard with anecdotal evidence from industry that it isnot reproducible and hence open to mis-interpretation. There were also a number of other gaps in the available data available whenDNV RP F112 was prepared, related to (i) differences between small scale specimens tested at typical seabed temperature andambient pressure and actual components in operation, ie at elevated operating temperature and pressure, and (ii) full-scale tests forfine grained pipe material. This paper presents data on these issues and indicates where technology gaps remain.

Experimental procedure

Introduction

Small-scale HISC tests were performed on tensile specimens taken from as fairly coarse grained superduplex stainless steel. Testswere of 30-day duration and performed at a potential of -1100mVSCE. Comparative tests were performed at 20 and 80C (1 bara),

and 1 and 100bara (20C).

Large-scale four-point bend HISC testing was carried out on fine-grained girth-welded and fillet-welded seamless pipes for a

maximum duration exceeding six months. Residual stress measurements were taken prior to testing. Testing was performed inseawater under cathodic protection at -1100mV and strain was recorded during testing at different locations on the welded pipe.Visual inspection, dye-penetrant examination, metallographic and fractographic studies were performed.

Materials

The microstructural characteristics of the five materials are summarised in Table 1. Material A had a fairly coarse aligned austenitestructure, with fairly consistent austenite island size. Material B had a 'primary' coarse, aligned austenite structure and finer, random

Oil & Gas, Published Papers


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equiaxed 'secondary' austenite islands in between the coarse units. The measurements of austenite spacing were made in thismaterial (i) including all austenite, and (ii) to reflect only the coarse primary austenite spacing, ie ignoring the fine secondaryaustenite.

Table 1 Materials examined

Material type Average austenite spacing (m) Ferrite (%) Hardness (HV5) Third phases

A: 25%Cr (UNS S32760) bar 20 (transverse)

39 (longitudinal) 55 5 261 None

B: 22%Cr (UNS S32205) pipeTransverse: 8 (11 ignoring fine austenite)Longitudinal: 11 (43)

56 5 244 None

Effect of temperature and pressure on small scale HISC testsTwo series of constant load tensile HISC tests were performed on material A (25%Cr superduplex bar) with specimens machined inthe longitudinal direction pre-charging and testing were at -1100mV SCE and 1 bara pressure at 20 and 80C, in natural seawater. A

salt bridge was used, so that the Ag/AgCl reference electrode could be kept cool.

Additional constant load HISC tests were performed at -1100mV SCE, 20C in natural seawater on material A (specimens machined

longitudinally) with pre-charging and testing at 100 bara. The autoclave was pressurised with nitrogen. An Ag/AgCl referenceelectrode was used. Specimens were stressed to around the threshold stresses for crack propagation and initiation in 30 days, asestablished at room temperature and pressure (749 and 553 MPa respectively). Allowance for the internal autoclave pressure was

made to the measured applied stress on the specimens, as described in EFC 17.[6]

After test the specimen hydrogen contents were measured by vacuum hot extraction and crack numbers and depths were measuredon metallographic sections through specimens that had not failed at the end of test.

Full scale welded pipe testsPipe samples in material B (22%Cr duplex), which were 4m long, with 15mm wall thickness and 168mm outer diameter were usedfor full scale testing in a purpose-built, four point bend load frame. Two specimen weld geometries were examined:

(i) a pipe with a girth weld at mid-length.(ii) a pipe with a fillet weld to a circular patch at mid-length, simulating an anode attachment pad. The girth-welded pipes,designated GW1 and GW2, were welded by mechanised TIG (with the pipe horizontal and rotated) employing a Zeron 100Xsuperduplex filler wire. The fillet-welded pipe, designated FW1, was welded by manual TIG, employing similar wire.

Residual stresses were measured prior to testing using the centre-hole drilling technique in the weld toe/HAZ areas. The locationswere chosen to minimise impact on subsequent test. The welded pipes were tested with a cell mounted around the weld area,containing natural, flowing seawater at a temperature of 10C with a potential of -1100mVSCE applied by potentiostat. The pipes

were pre-charged, without load applied, for seven days, prior to test. The load applied to pipe GW1 and FW1 was then increasedincrementally to identify the threshold load for macroscopic crack development. Each loading step was maintained for seven daysand the pipe was examined for the onset of cracking employing a binocular microscope. After seven days exposure the applied loadwas increased and the procedure repeated until cracking initiated. In order to record the strain during testing, the pipes were straingauged. Following determination of an approximate threshold load from the first two tests, the girth-welded pipe GW2 was testedwith loading to give 0.5% total strain, as measured on strain gauges away from the weld and held for a period of 6 months. Thewelded pipes were examined by dye-penetrant inspection (DPI) at the end of the test to identify any fine cracking at the weld andsections were taken through relevant areas identified by DPI.

Stress concentration factors (SCFs) for the weld toes were estimated based upon comparison of the weld geometries and previous

finite element analysis performed at TWI.[7] The two weld geometries examined had the following estimated SCFs:

(i) girth-welded pipes. The SCF was about 2.6 0.2. The errors quoted allow for the variability of geometry in real welds comparedwith FE models. The main variables determining the SCF at the toe of a butt weld are the angle at the weld toe and the overallprofile of the weld. For the girth welds, the toe angle was measured to be about 45degrees and the weld overfill was considered tobe in the form of a circular arc.

(ii) fillet-welded pipe. The SCF was estimated as 2.8 0.2. The main variables here are the toe angle and the ratio between the weldleg length and plate (pipe) thickness. The toe angle was 20-25 and the leg length and plate (pipe) thicknesses were 8mm and

15mm, respectively.

Results

Effects of temperature and pressure

HISC testing at 20 and 80CFigure 1 presents the results of tensile HISC testing of 25%Cr superduplex material A at 20 and80C, 1 bara pressure and -1100mVSCE. In terms of applied stress, the threshold for specimen failure in 30 days was lower at 80C

than at 20C by around 3% but no reduction in the stress for crack initiation (ie formation of small cracks that did not propagatethrough thickness) in 30 days was found, Figure 1a. The equivalent data are plotted in terms of strain in Figure 1b, which shows asimilar small reduction in strain for cracking at 80C. However, the 0.2% proof stress of material E was around 13% lower at 80Cthan 20C (530 and 608MPa respectively), indicating that the HISC crack initiation and propagation thresholds at 20C were lowerthan at 80C, when considered in terms of normalised stress, ie applied stress divided by 0.2% proof stress.


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Fig.1. Results of small-scale tensile HISC testsat 20 and 80C (all at 1bara pressure) Fig.1a)Plotted in terms of applied stress

Fig.1b) Plotted in terms of measured strain

HISC testing at 1 and 100 bara Figure 2 present the results of small scale tensile HISC testing of 25%Cr superduplex material Aat 1 and 100bara, 20C, and -1100mV SCE in seawater. A reduction in the stress threshold for crack propagation, by about 4%, was

noted at 100bara but no reduction in threshold stress for crack initiation was noted. However, it should be noted that fewer testswere performed at 100bara, such that the threshold values were less precisely determined than at 1bara. Hence, it is possible that asmall shift in crack initiation threshold exists but it could be no more than 5%.

Fig.2. Results of small-scale HISC tests at 1 and100bar, 20CFig.2 a) Plotted in terms of

normalised applied stress, ie applied stress/0.2%proof stress

Fig.2b) Plotted in terms of measured strain

Post-test characterisation Substantially higher hydrogen pick-up was observed at 80C compared to 20C, by a factor of 5 to 8,whilst measurements indicated an approximate factor of two increase in hydrogen pick-up at 100bara compared to 1bara (both at20C). It was noted that there were many more cracks formed at 80C than at 20C but increasing temperature to 80C seemed tohave reduced crack depth. Raising pressure increased crack depth for a given strain.

Full scale HISC tests on welded pipes

Residual stress A maximum tensile residual stress of 453MPa was measured in the HAZ of girth-welded pipe GW1 about 2mm fromthe fusion line. Maximum circumferential and axial tensile residual stresses of 447MPa and 309MPa were observed in the HAZ of pipe

GW1. A maximum tensile residual stress of 392MPa was measured in the HAZ of pipe GW2 and maximum circumferential and axialtensile residual stresses were 389MPa and 279MPa. For the fillet welded FW1 pipe a maximum tensile residual stress of 404MPa wasmeasured in the HAZ about 1.5mm from the fusion line.

Threshold stress and strain for HISCFigure 3 shows the 'crack' and 'no crack' applied stress levels from girth-welded pipe GW1,as calculated from the longitudinal strain measurements away from the weld and the stress-strain curve for the 22%Cr parentmaterial B, indicating the approximate threshold load. The global strain, ie away from stress concentrations, during testing wasconstant and no straining due to low temperature creep was observed at the gauges.


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Fig.3. 'Crack' and 'no crack' stress levels in thefull-scale girth welded pipe GW1

No cracking was observed after seven days at 0.63% strain, corresponding to a nominal longitudinal stress of 584MPa. Theequivalent normalised stress (with respect to the 0.2% proof stress) was 1.05. Cracking occurred after less than 12 hours at a strainlevel of 0.85% (measured far from the weld toe) corresponding to a nominal stress of 601MPa (normalised stress = 1.08). Failure inpipe GW1 occurred at the weld toe and propagated through the HAZ (Figure 4) into the parent material.

Fig.4. Crack at the weld toe in thegirth-welded pipe GW1

The applied load for the pipe GW2, tested at one load for six months without cracking, is indicated by the dashed line (measured125mm from weld cap toe - gauge 6). The maximum total strain measured far from the weld toe was a little over 0.5%,corresponding to a stress of 571MPa and a normalised stress of 1.03.

Fig.5. Fracture surface of the crackin girth-welded pipe GW1

Discussion

Operating conditions

Increasing temperature to 80C and increasing pressure to 100bara both tended to reduce the threshold stress for specimen failurein 30 days but little or no effect on small crack initiation was found. When the reduction of proof stress at 80C is taken intoaccount, cracking behaviour is improved in terms of normalised stress but it is concluded that sensitivity to HISC is not substantiallychanged at 80C. The mechanism of the effect of pressure on hydrogen pick-up is not immediately apparent, although one effectmight be to inhibit recombination of H atoms. It appears that conditions that act to increase hydrogen pick-up tend to enhancespecimen failure, ie crack propagation, but have little effect on crack initiation, at least for the durations studied. The fact thatincreased surface hydrogen content did not significantly affect HISC initiation, suggests that the surface hydrogen level is alreadyadequately high for easy crack initiation at 20C and 1bara. However, a greater but still quite small, effect of pressure on crackpropagation was noted. This is consistent with an increased hydrogen level subsurface, arising from a higher surface concentration(the hydrogen diffusion coefficient being unchanged).

The effect of increasing temperature is not a simple one. The surface hydrogen concentration at 80C is expected to be much higher

than at 20C, as higher charging current densities are found at higher temperatures,[8] and this might account for the greaternumber of surface cracks at 80C compared to 20C. Hydrogen diffusion is also faster at 80C, leading to higher subsurfacehydrogen contents. However, this did not lead to greater crack depths at 80C. One potential explanation of this could be thatpropagation is controlled, at least in part, by the austenite structure rather than the rate of hydrogen penetration ahead of the crackalone.

Full-scale pipe behaviourThe work showed that full scale welded duplex stainless steel pipes, with fine austenite spacing, cathodically protected at-1100mVSCE can tolerate a global stress of 1.03 times the 0.2% proof stress, equivalent to a total strain of 0.5%, both measured

well away from any stress concentrator prior to the onset of HISC at a weld with and SCF of 2.6-2.8. Only small differences wereobserved between fillet-welded and girth-welded pipes. The pipe test failure loads are substantially greater than for coarse grained

superduplex forgings, which failed under an applied strain of only 0.25% at a stress concentrator.[2] The initiation and subsequent


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propagation of HISC apparently occurred at very similar stress in the large scale samples. The stress and strain at the weld toe arenot known, as there is no simple relationship between the SCF and the local conditions at the weld toe, which will be affected bylocal low temperature creep. When the work was originally undertaken, there was no opportunity or strong need to analyse theresults further but there is now an opportunity to model the stress and strain at the weld toe and to compare with the RP F112allowable loading.

Remaining issues: materials with improved hisc resistance and sub-surface flaws

RP F112 represents a big step forward in design to avoid HISC. However, it may be more conservative than necessary for someproduct forms, e.g fine-grained wrought pipe, partly because it is primarily validated against data for very coarse grainedsuperduplex forging material and partly because it makes conservative assumptions regarding the effect of residual stress. This hasled to reports from industry that some components are being over-designed, leading to production of components that are thicker,

heavier and more expensive than necessary. However, some product forms, including pipes, are substantially more resistant toHISC than coarse grained forgings, at least partly due to finer austenite spacing. Whilst RP F112 allows advantage of this enhancedHISC resistance to be taken for material with austenite spacing


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conf 'Corrosion 99', NACE International, paper 148.

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Hydrogen Induced Stress Cracking of Duplex Stainless Steel Subsea Components