Stress corrosion cracking of duplex and super duplex stainless steels in sour environments

Corrosion Science\ Vol[ 39\ No[ 5\ pp[ 898Ð829\ 0887Þ 0887 Elsevier Science Ltd[ All rights reserved[\ Pergamon Printed in Great Britain[

9909Ð827X:87 ,*see front matter

PII] S9909Ð827X"87#99911Ð4

STRESS CORROSION CRACKING OF DUPLEX AND SUPERDUPLEX STAINLESS STEELS IN SOUR ENVIRONMENTS

A[ A[ EL!YAZGI� and D[ HARDIE

Dept[ of Mechanical\ Materials and Manufacturing Engineering\ University of Newcastle upon Tyne\ Newcastleupon Tyne NE0 6RU\ U[K[

Abstract*The resistance of two commercial duplex stainless steels to sulphide stress corrosion cracking in sourenvironments has been assessed by means of slow straining in a variety of solutions saturated with hydrogensulphide over the temperature range from 1 to 84>C[ The results suggest that\ although one is not susceptible toany loss in ductility\ the other su}ers a signi_cant loss in ductility and a ductility minimum occurs at sometemperature between 14 and 79>C that depends upon the environment[ The embrittlement at lower temperaturesis attributed to absorption of the hydrogen generated at the steel surface\ but this is mitigated by the barrier topermeation presented by the passive _lm also generated[

Associated polarisation measurements and appropriate pitting experiments clearly indicate a reduction of therange of potential in which passivation occurs with increasing temperature and severe pitting above a temperaturethat depends upon the alloy content of the steel[ The prevalence of severe pitting corrosion provides an explanationof the attenuation of the hydrogen embrittlement e}ects at the higher temperatures\ and the resultant occurrenceof a minimum ductility[ Þ 0887 Elsevier Science Ltd[ All rights reserved

Keywords] A[ stainless steel\ C[ stress corrosion cracking

INTRODUCTION

Sulphide stress corrosion cracking "sscc# has long been recognised as a serious problem forthe petroleum and petrochemical industries\ particularly in high strength and stainlesssteels[0Ð7 In recent years\ the rising energy demand\ increasing oil and gas prices\ anddwindling reserves have led to the development of deeper oil and gas wells "often ×5999m#[The environments often encountered in such deep wells contain substantial amounts of saltwater\ hydrogen sulphide\ and carbon dioxide "sour wells#[ This hostile environment\depending upon its composition and temperature\ may cause general corrosion\ localisedcorrosion\ and stress corrosion cracking of the materials used[ One means of combatingsuch problems is to inject inhibiting agents\ but it was soon realised that these were neitherwholly e}ective nor economically viable[ Another solution relied on the use of highlyalloyed materials that can withstand the corrosiveness of the well ~uids[ Among the variousspecial materials that can be utilised in such hostile environments\ duplex stainless steelsrepresent an optimum technical!economic choice[

�Dr A[ A[ El!Yazgi has now returned to the Department of Materials Science\ Tajura Nuclear ResearchCentre\ Tripoli\ Libya[

Manuscript received 04 July 0886^ in amended form 09 September 0886[

898

A[ A[ El!Yazgi and D[ Hardie809

When\ as in a sour environment\ H1S comes into contact with water\ it dissociates toform a weak acid according to the reactions]

H1S:HS−¦H¦ "0#

H1S:1H¦¦S1− "1#

and hence\ a corrosion reaction takes place between these ions and steel resulting in theformation of ferrous ion at the anoxic sites and the reduction of hydrogen ion at thecathodic sites on the steel surface]

Fe:Fe¦¦¦1e− "at the anoxic site# "2#

1H¦¦1e−:1Hand 1H:H1 "at the cathodic site# "3#

Thus the overall reaction can be represented as]

Fe¦H1S"aq#:FeS¦1H "4#

The recombination of atomic to produce molecular hydrogen is inhibited by the presenceof the dissociated hydrogen sulphide\ and this results in a greater hydrogen fugacity\ whichfacilitates penetration of hydrogen into the steel\ hence intensifying any embrittlemente}ects[

Many variables are known to in~uence sscc of iron base alloys and these may beclassi_ed into two broad categories]

i# Environmental e}ects that include hydrogen sulphide concentration "pressure#\ chlor!ide concentration\ temperature and applied potential and

ii# Mechanical and metallurgical variables such as alloy composition\ microstructure\heat treatment\ cold work\ hardness or strength level\ and surface condition[

The aim of this work was to assess the contribution of some of these factors to embrittlementinduced in commercial duplex stainless steels[

EXPERIMENTAL

The investigation was facilitated by the availability of two commercial duplex stainlesssteels] types 1294 and 1496 "Table 0#[ A three!dimensional assessment of the microstructuresof these materials after polishing and etching in an aqueous solution containing 19)hydrochloric acid and 2g potassium metabisulphite "K1S1O4# reveals "Fig[ 0# that bothconsist of elongated islands of austenite "light!etching phase# in a matrix of ferrite "dark!etching phase#[ The proportions of the two phases were assessed by optical image and X!ray di}raction analysis as well as by means of a ferrometer\ as previously described[8 All

Table 0[ Chemical Analysis "wt)# of the Duplex Stainless Steel Pipes Used

Cr Ni Mo Mn Si N Cu Co V C P S Ti

Steel A 12[99 4[99 1[88 9[87 9[37 9[03 9[92 9[91 9[94 9[904 9[905 9[990 ³9[90Steel B 14[29 5[69 2[72 9[18 9[13 9[18 9[09 9[04 9[93 9[913 9[904 9[991 9[992

Cracking of stainless steels in sour environments 800

gave consistent results with an average value of 24[4 and 49[1) austenite for steels A andB respectively[

Smooth tensile specimens having an 00[24mm gauge length of 2[19mm diameter andan overall length of 069mm were machined with their axis parallel to the longitudinaldirection of the pipes and the gauge lengths were polished to 599mm before straining insolution[ The straining was achieved in a hard tensile machine at a cross head speed thatproduced a strain rate "as referred to the original gauge length# of 1[9×09−5:s[

The straining was conducted in three di}erent solutions\ all saturated with H1S at 0 bar]

"i# 2[4) sodium chloride solution with initial pH"pHi �5[4#\"ii# NACE standard solution\ 49 g sodium chloride ¦4g glacial acetic acid in 834ml

deionized water "pHi �1[6#\ and"iii# Deionized water with varying amounts of sodium chloride\ from 9 to 599wppm

Cl−"pHi �5[4#[

The test cell was a specially designed cylindrical PTFE cell closed by a screw!in top lid thatcould be sealed to the main body of the cell by an O!ring seal "Fig[ 1#[ It incorporates athermometer\ a control thermocouple\ a reference electrode "saturated calomel electrodefor solution "i# and "ii# but sulphate reference electrode for solution "iii##\ a platinum counterelectrode for testing under controlled potentials\ and glass inlet and outlet tubes to allowcontinuous bubbling of the hydrogen sulphide gas into the solution during the whole of thetest[ The outlet from the cell is connected to a 1[9 litre container of 0M NaOH solution\through a 0[9 litre plastic container to prevent back ~ow of the caustic solution into thetest cell\ in order to absorb hydrogen sulphide waste before venting to atmosphere[ The

Fig[ 0[ Three!dimensional representation of the microstructure of the as!received materials[


Fig[ 1[ Arrangement of the test cell for straining in solutions saturated with hydrogen sulphide[

H1S saturation was initially achieved by continuous bubbling through the test solutions ata rate of 299ml:min for 2 hours\ at which time the open circuit potential\ which wasmonitored continuously\ stabilised at about −499mV "sce#\ which was taken as the criterionfor hydrogen sulphide saturation[ The bubbling rate is then lowered to 099ml:min\ whichproved to be su.cient to maintain saturation throughout the entire test period[

The test solution is heated or cooled to the desired test temperature "from 1 to 84>C#before tensile straining is applied[ Sub!ambient temperatures were achieved by using animmersion cooler[ A copper block\ into which the cooling coil was inserted\ was fastenednext to the testing cell and the whole arrangement was wrapped with glass wool forinsulation[ The arrangement was then left for 5 hours to cool and stabilise at the desired


lower temperature before the solution was saturated with hydrogen sulphide[ The lowesttemperature used was 1>C\ to avoid partial solidi_cation of the solution[

Long term pitting tests were conducted on 09mm×09mm specimens cut from bothmaterials A and B[ Specimens were polished to 0mm _nish\ after which they were slightlyetched in order to identify the sites of pit initiation[ They were then carefully washed withacetone\ to prevent crevice e}ects that might be caused by dirt or oil spots\ before beingimmersed in test solution "ii#[ Tests were conducted at two di}erent temperatures *29 and89>C* in glass ~asks[ After the solution in the ~ask had been saturated with hydrogensulphide\ the ~ask was sealed and left at the desired temperature for the duration of thetest "619 hours#[ Specimens were removed and washed successively with distilled water andalcohol before examination by optical and scanning electron microscopy[

After soaking at the higher temperature\ where corrosion product is expected to form\the fracture surfaces were subjected to cathodic polarisation in alkaline sodium cyanidesolution\ for a few minutes\ with simultaneous ultrasonic cleaning\ and this produced veryclean fracture surfaces[ Polished sections perpendicular to the fracture surface were alsoexamined after etching in 119ml HCl¦79ml H1O¦2g K1S1O4\ which attacks the ferrite"dark# and leaves the austenite unattacked "bright#[

Potentiodynamic polarisation measurements were conducted\ on 1[9 cm1 standard cyl!indrical specimens of both materials\ using a potentiogalvanic scan coupled with a log!arithmic output current sink[ The specimens were _rst dry polished to 599 grit paperfollowed by a wet polish[ The test cell was _lled with the test solution and heated or cooledto the desired temperature\ before purging the solution with hydrogen sulphide for 09minutes[ The bubbling rate is lowered to 099ml:min until the open circuit potential ofthe specimen stabilised "1 hours#\ which was the criteria adapted for hydrogen sulphidesaturation\ before commencing with the potential scan which was conducted from −0999to ¦0999mV at a rate of 09mV:s[ Some polarisation tests were conducted in a solutionpurged with nitrogen for one hour before saturation with hydrogen sulphide\ in order toinvestigate the e}ect of excluding oxygen from the environment\ but the curves subsequentlyobtained were identical to those done without such purging\ which indicates that adequatesaturation with hydrogen sulphide is as e}ective in the expulsion of oxygen from the testsolution[

DISCUSSION OF RESULTS

The variation in percent reduction in area with straining temperature reveals thatembrittlement of steel A _rst increases with temperature up to a certain intermediatetemperature and then decreases at higher temperatures "Fig[ 2#[ Maximum embrittlementoccurs at 59Ð69>C in NACE!90!66 solution but at 29Ð39>C in non!acidi_ed sodium chloridesolution[ Many investigators have observed similar trends in results for both high strengthand duplex steels1Ð7\09Ð02 using di}erent loading methods[ The signi_cance of this relation!ship\ however\ is not so much the particular temperature at which maximum loss in ductilityoccurs\ but rather the fact that the increasing amount of embrittlement as temperatureincreases is curtailed by recovery of ductility at a higher temperature\ which suggests thattwo competing processes are involved[ One possibility would involve hydrogen embrit!tlement at lower temperatures and increasing corrosion at higher temperatures[ The decreasein the susceptibility at lower temperatures could be attributed to a decrease in the rate ofhydrogen arrival at the critical site within the bulk of the specimen due to a decrease in the


Fig[ 2[ E}ect of test temperature upon the ductility of duplex stainless steel A strained at 1×09−5:sin either NACE "solid symbols# or 2[4) sodium chloride "open symbols# solution\ saturated with

hydrogen sulphide[

thermally!activated hydrogen di}usion process and the prevention of hydrogen entry bythe formation of some form of protective _lm[ On the other hand\ it is feasible to assumethat the recovery of ductility at the higher temperatures is due to increased dissolution\


which is also thermally activated\ that causes blunting of any cracks initiated and alsohinders the entry of hydrogen\ by the corrosion product[ Steel B\ however\ showed no

susceptibility to either anoxic corrosion or hydrogen embrittlement over the whole tem!

perature range studied[The polarisation curves for steel A "Fig[ 3# revealed that a decrease in temperature

promotes a dramatic decrease in corrosion rate\ as indicated by the measured currentdensity[ In addition\ no passive region occurs on the polarisation curve at 74>C\ and thussmall increases in the potential at this temperature will understandably result in a signi_cantincrease in the corrosion rate[ However\ as temperature is lowered\ the metal tries topassivate until\ when the temperature reaches 1>C\ some form of protective _lm\ believedto be a type of iron sulphide\ occurs over a substantial range of potential between −399and −199mV "sce#[ The measured open circuit potential during tensile straining testsperformed in both NACE!90!66 solution and non!acidi_ed sodium chloride solution satu!rated with H1S at 1>C\ was −399mV "sce# and this appears to correlate well with the regionof passivation[ This protective _lm may hinder the entry of atomic hydrogen into the metaland so prevent embrittlement\ but if the tensile specimens were polarised to slightly morenoble potential this would be expected to shift the metal away from conditions that favouredformation of protecting _lms and thus facilitate hydrogen entry[ At 1>C the ductility ")RA#almost regains its value of 67) in air\ but when two specimens were strained to failure inNACE!90!66 solution at 1>C whilst polarised to a slightly more negative potential than theopen circuit i[e[ −499 and −449mV "sce#\ both failed in a brittle manner\ with )RA of20 and 22 respectively "Fig[ 2#[ Thus\ this small cathodic shift in potential causes the steelto be embrittled by hydrogen[ This can be attributed to the expected evolution of hydrogenat such a potential and pH\ and should result in 34 to 44)RA[03 However\ the more severe

Fig[ 3[ Potentiodynamic polarisation curves for duplex stainless steel A obtained at a sweep rateof 09 mV:s in NACE!90!66 solution at various temperatures[


loss in ductility "20 to 22) RA# may perhaps be due to the more e.cient delivery ofhydrogen to the critical sites because the applied potential shifts the metal from the passiveregion "−199 to −399mV"sce# in Fig[ 3# to a more active potential where the _lm thatpresents a barrier to hydrogen becomes unstable or does not form at all and this clearlyfacilitates the entry of atomic hydrogen into the metal[ Furthermore\ when two specimenswere strained to failure in the same solution "NACE!90!66# at 79>C\ which is near thetemperature of maximum embrittlement\ whilst polarised to −249 and −399mV "sce#\ i[e[respectively 049 and 099mV more noble than the measured open circuit potential at 79>C"−499mV sce#\ extensive pitting became visible "Fig[ 4# and the straining had to beterminated before _nal failure "separation#Ðtests were stopped after about 07 h in each case[It is obvious that the imposed potential caused the metal to be in the active state and henceextensive pitting occurred[ Thus\ a small shift in potential at a temperature where hydrogenembrittlement might be expected to play a greater part in the failure may obviouslyintroduce excessive dissolution instead[

After establishing that specimens from steel B did not su}er any detectable damage\

under any of the test conditions used\ polarisation curves for this material were deemednecessary to shed light on possible explanations[ These revealed that\ even under the mostsevere embrittlement conditions\ i[e[ in NACE solution near 74>C\ a passive _lm still formson the surface\ in marked contrast to steel A under the same conditions\ where no pas!sivation is possible[ This can be attributed to the fact that material B contains higher Cr\Mo\ Ni\ and N\ all of which are known to stabilise the passivity of stainless steels over awider range of temperature and pressure[04 This high alloying content gives rise to aPitting Resistance Equivalent number "PREn# of 32 "calculated05 from PREn � ) Cr ¦2[2")Mo#¦29 ")N## as compared with steel A where the PREn �24[ Actually\ it has beenrecommended that duplex stainless steels to be used in sour environments should have aminimum pitting resistance equivalent number of 39[ The corrosion potential "for steel B#appears to be 099mV more active than that of steel A\ which is in agreement with thereported increases in corrosion potential with increasing PREn of the alloy[06

The initial pH "pHi# of both neutral solutions "i# and "iii# were 5[4 before saturatingwith hydrogen sulphide[ After saturation\ however\ the pH was reduced to about 3[1 forboth solutions[ But for NACE!90!66 solution\ with pHi\ of 1[6\ saturation with hydrogensulphide increased the pH to about 2[1[ Therefore the actual "e}ective# pH to which thespecimens were subjected were 3[1 and 2[1 for neutral sodium chloride and NACE solutionsrespectively[ The degree of embrittlement was always higher in NACE!solution "pH�2[1#than in sodium chloride "pH�3[1# "Fig[ 2# and this may be attributed to the lower pH andgreater hydrogen activity of the NACE solution[1

Secondary cracks were con_ned to a very small necked region in specimens tested inthe range 09 to 19>C\ but as the temperature increased the number increases dramatically\

Table 1[ Mechanical Properties of the Duplex Stainless Steel Used "Strained at 1[9×09−5s−0#

9[1) Flow Stress Ultimate Tensile ) ) Reduction inMN:m1 Strength MN:m1 Plastic Elongation Area at Fracture

Steel A 512 633 33 67Steel B 690 729 30 79


Fig[ 4[ Tensile specimens of steel A strained in NACE!90!66 solution at 79>C whilst polarised at"a# −249 mV"sce# and "b# −399 mV"sce#!condition after about 07 h[

especially above about 59>C "Fig[ 5#[ An attempt was made to count the number ofsecondary cracks in the gauge length of the tensile specimens and combine the results withthose from the pitting tests to provide an insight into the e}ect of temperature on bothcracking and pitting "Fig[ 6#[ As temperature increases\ both the density of pit formationon polished specimens immersed in NACE solution "pH�1[6# for 37 h\ and the numberof secondary cracks in the gauge length of specimens strained in NACE!90!66 solution"pH�2[1# increased only slightly up to 59>C but then showed a dramatic increase[ Thisagrees well with the reported pitting temperature of 59>C for steel A "PREn �24#[01 In allcases\ the pits initiate at ferrite!austenite interfaces and grow into the austenite phase[

On some of the tensile specimens strained to failure at 04Ð19>C\ a thick black _lmformed that was analysed using an X!ray di}ractometer "CuKa radiation l�0[43940#^ thescan speed was set at 9[04 deg:min to ensure good trace resolution[ The _lm was found toconsist of multiple layers of iron sulphides^ FeS1 among others "Fig[ 7#[ However\ when a


Fig[ 5[ Tensile specimens of duplex stainless steel A strained to failure at 1×09−5:s in NACE!90!66 solution at various temperatures[

polished specimen of steel A was immersed in NACE!90!66 solution\ for 619 h at 29> and89>C they showed "in both cases# that the austenite was attacked preferentially[ At 29>Cthe corrosion rate calculated for the dissolution of austenite "from the depth of penetration#was 9[36mm:yr[ The austenite was uniformly attacked "Fig[ 8# and\ when this specimenwas subjected to X!rays for _lm analysis\ the austenite peaks had disappeared due to thefact that the austenite was below the general level of the surface\ and hence the ferrite[Careful scrutiny of this specimen revealed that the ferrite phase did not su}er any corrosionand\ in fact\ a thin layer of ferrite covering the austenite will provide protection[ The X!rayanalysis of the _lm also revealed the presence of elemental sulphur "Fig[ 09#\ which couldbe removed "dissolved# by immersion in carbon disulphide for a few seconds[ When sub!sequent X!ray examination was carried out\ the sulphur lines had disappeared\ leavingbehind just an indication of iron sulphide "FeS#[ At the higher temperature "89>C#\ however\


Fig[ 6[ E}ect of temperature on the density of initiation sites for pits and cracks in steel Aindicating the number of pits generated in an area of 0cm1 during immersion in NACE solution for37h and the number of cracks found in the specimen gauge length after straining to failure in

NACE!90!66 solution[

the austenite in the duplex structure was not uniformly attacked^ once initiated\ dissolutionappeared to be con_ned to certain preferential sites where it accelerated[

When straining is conducted at lower temperatures "19Ð59>C#\ the secondary cracks arevery small and con_ned to the small necked region\ with no evidence of any dissolution ineither of the phases "Fig[ 00a#[ The cracks initiate in the ferrite and generally propagateperpendicular to the applied stress\ although crack branching at about 34> and along theferrite:austenite interface "Fig[ 00b# is frequently noted[ Cracks generally propagate throughthe ferrite phase\ the phase sensitive to hydrogen\ and avoid the austenite "i[e[ many cracksare arrested by the austenite#[ Cracking along the ferrite:austenite grain boundary can beexplained by the fact that these are hydrogen traps and high concentrations of hydrogencan develop\ especially with respect to the high di}usivity of hydrogen in the ferrite ascompared to that of the austenite "several orders of magnitude at room temperature#[07

That is to say that hydrogen travels much faster in ferrite and when it arrives at austenite\in which the di}usivity of hydrogen is much reduced\ a net accumulation of hydrogen inthe ferrite near the austenite boundary will develop and so provide a favourable crackpropagation path[ The cracking is very similar to that occurring during the straining oftensile specimens in a hydrogen atmosphere or whilst cathodically polarised\03\08\19 whichsuggests that hydrogen is also the embrittling species in the H1S!bearing environment[

At temperatures above 69>C\ in both solutions\ numerous secondary cracks appear


Fig[ 7[ X!ray di}raction Analysis of black deposit on the surface of tensile specimen of duplexstainless steel strained to failure at 1×09−5:s in NACE!90!66 solution at 04>C[

along the gauge length and microscopical examination of the gauge length indicates thatthese initiate at pits "Fig[ 01a# in marked contrast to the initiation of cracks at lowertemperatures "Fig[ 01b#[ After pitting occurs\ small cracks tend to initiate and propagateby cracking of the ferrite and simultaneous dissolution of the austenite "Fig[ 02a#[ Astemperature increases\ the cracks tend to become blunted by the accelerated corrosionreaction and this produces a recovery in the measured ductility[ When the temperaturereaches 84>C the dissolution of the austenite is very extensive "Fig[ 02b# and leads to thegreatly enhanced ductility[

SEM examination of fractured surfaces has indicated that the ferrite always fracturesin the cleavage mode and river markings are revealed[ The cleavage surfaces are invariablyclean and free from any corrosion product[ On the other hand\ although the austenite phaseshowed some ductility at temperatures up to 29>C "Fig[ 03#\ the austenite islands near thecrack initiation site were severely corroded above this temperature "i[e[ 39>C# and therewere deposits of corrosion product[ When the corrosion product was analysed using EDX\it was found to be rich in sulphur "Fig[ 04#\ indicating that some form of iron sulphide ispresent[ Removal of these corrosion products by means of ultrasonic cleaning duringcathodic polarisation in a sodium cyanide bath produced remarkable results[ The newcleaned surface "Fig[ 05# showed distinct crevicing of various orientations\ within theaustenite islands[ The fractured surface was subsequently polished slightly to produce asmooth mirror!like _nish\ and viewed under a microscope to con_rm that all crevices markshad disappeared[ When this surface was etched with 19) HCl 79)H1O¦9[4 g K1S1O4 thestructure revealed transformation of the austenite near the fracture surface to martensite"Fig[ 06#[ The crevicing that was present before polishing was the result of corrosive attackof martensite plates by the environment[ The martensite was distinguished by its brown!blue appearance\ the same colour as the ferrite\ depending on the etching time[0


Fig[ 8[ Dissolution of austenite from a specimen of A immersed in NACE!90!66 solution for 619 hat 29>C] "a#exposed surface and "b# a section perpendicular to "a#[


Fig[ 09[ X!ray di}raction Analysis of the _lm formed on the polished surface of duplex stainlesssteel immersed in NACE!90!66 solution for 619 h at 29>C[ Note the absence of austenite lines[

Results of tests conducted in neutral aqueous solution\ having various chloride ionconcentrations ranging from 9 to 699wppm and saturated with hydrogen sulphide\ indicatethat chloride ions play a signi_cant role in the embrittlement mechanism[ In distilled water"9) Cl# there was no signi_cant loss in ductility but\ as the concentration of chlorideincreased the sulphide scc susceptibility also increased until the chloride level reached ½299wppm[ Increasing the Cl! level beyond this limit seems to have little or no in~uence on theembrittlement "Fig[ 07#[ The chloride ion seems to play a signi_cant role in the repassivationof the steel[ Once the protective _lm is broken by straining\ the presence of a su.cientquantity of chloride ion will inhibit the reformation of the _lm and hence hydrogenmay enter the steel and produce a loss in ductility\ which is in agreement with reportedresults[03\04\06

Overall\ the results indicate a complex interplay of hydrogen embrittlement\ passivationand pitting corrosion in duplex stainless steels\ where the ferrite and austenite in themicrostructure have very di}erent strength and ductility\ as well as di}ering susceptibilitiesto hydrogen embrittlement[ The orientation of the two phases in the microstructure andthe formation of e!martensite in the austenite phase as a result of deformation19 both havea profound in~uence upon crack propagation[ Hydrogen embrittlement of steel A has beenwell characterised in a wide variety of environments producing hydrogen at di}erentfugacities[03 Whether this embrittlement originated from straining in gaseous hydrogen\ asa consequence of thermal charging under high pressure hydrogen\ or as a result of cathodicpolarisation\ it was obvious that essential preliminaries were the actual uptake of hydrogenand the necessary transport to critical sites in the steel[ The introduction of hydrogensulphide into the system combines the provision of a high hydrogen fugacity with chemicalreaction at the steel surface[ The latter reaction may\ under appropriate conditions\ result


Fig[ 00[ Secondary cracking in the gauge length of a duplex stainless steel specimen strained tofailure in NACE!90!66 solution at 19>C "a# unimpeded and "b# de~ected by austenite stringers[


Fig[ 01[ Secondary cracking in the gauge length of a duplex stainless steel specimen strained tofailure in NACE!90!66 solution at "a# 79>C and "b# 59>C[


Fig[ 02[ Secondary cracking in the gauge length of duplex stainless steel specimens strained tofailure in NACE!90!66 solution at "a# 59>C and "b# 84>C[


Fig[ 03[ Fracture surfaces of duplex stainless steel specimens strained to failure in NACE!90!66solution at "a# 29>C and "b# 39>C[


Fig[ 04[ EDX analysis of corrosion product on the surface of the duplex stainless steel specimenstrained to failure in NACE!90!66 at 39>C "see 03"b##[

in either formation of a passive barrier to hydrogen absorption or breakdown of this barrierto allow both ingress of hydrogen and localised pitting[ If the pitting becomes serious\ i[e[above the critical pitting temperature for the particular steel\ then the e}ects of hydrogenbecome inconsequential[

CONCLUSIONS

"0# Slow straining of two commercial duplex stainless steels in chloride solutions satu!rated with hydrogen sulphide between 1 and 84>C reveals that\ whilst type 1496retains its ductility over the whole temperature range\ type 1294 su}ers a ductilityloss that increases to a maximum between 14 and 79>C but depends upon theparticular environment concerned[

"1# The embrittlement at the lower temperatures is attributed to absorption of hydrogengenerated at the steel surface but becomes mitigated by the barrier layer to per!meation presented by the passive _lm formed concurrently[

"2# The range of embrittlement becomes attenuated because of the severe pitting thatoccurs at higher temperatures where the passive _lm breaks down and recovery ofductility is observed[


Fig[ 05[ Fracture surface of duplex stainless steel specimen strained to failure in NACE!90!66solution at 39>C "i[e[ surface shown in 03"b## after cleaning[

Fig[ 06[ Fracture surface of duplex stainless steel specimen strained to failure in NACE!90!66solution at 39>C\ after polishing and etching[


Fig[ 07[ E}ect of chloride ion concentration on the ductility of duplex stainless steel strained tofailure in solutions saturated with hydrogen sulphide[

REFERENCES

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04[ Fontana M[ G[ and Greene\ N[ D[\ Corrosion En`ineerin`\ McGraw Hill\ 0867[05[ Marshall\ P[ I[ and Gooch\ T[ G[ Corrosion\ 0882\ 38\ 403[06[ Saarinen\ K[\ Pro`ress in The Understandin` and Prevention of Corrosion\ pp[ 474Ð478[07[ Perng\ T[\ Altstetter\ C[ J[ Acta metall[\ 0875\ 23\ 0660[08[ Zheng\ W[ and Hardie\ D[ Corrosion Science\ 0880\ 21\ 12[19[ Zheng\ W[ and Hardie\ D[ Corrosion\ 0880\ 36\ 681[

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Stress corrosion cracking of duplex and super duplex stainless steels in sour environments