The NACE Ir ternational Annual Conference and Exposition

WELDING TECHNIQUES FOR HIGH ALLOY AUSTENITIC STAINLESS STEEL

T G Goocll and P Woollin‘IWI

Abin ~ton HallAbington

Cambrid g,e CB1 6ALUnitec Kingdom

ABSTRACT

Factors controlling corrosion resistant.e of wcklments in high alloy austenitic stainlesssteel are described, with emphasis on micro:;egregation, intermetallic phase precipitation andnitrogen loss from the molten pool. The application is considered of a range of welding,processes, both fusion and solid state.

Autogcnous fusion weldments have co-rosion resistance below that of the parent, but lowarc energy, high travel speed and use of Nz-.bcaring shielding gas are recommended for bestproperties. Conventional fusion welding practice is to use an overalloycd nickel-base filler rneta[to avoid preferential weld metal corrosion, and attention is given to the effects of consumablecomposition and level of wcldpool dilution by base steel. With non-matching consumables,overall joint corrosion resistance may bc limii cd by the prcscncc of a fusion boundary unmixedzone: better performance may bc obtained us .ng solid state friction welding, given appropriatecomponent geometry.

Overall, the effects of welding on supcraustcnitic steels are understood, and the materialshave given cxccllent service in welded fabrications. The paper summarises recommendations 011preferred welding proccdurc.

INTRODUCTION

Austenitic stainless steels offer an attrz ctivc combination of corrosion resistance and cas~of fabrication. Their corrosion resistance dcpimds upon the prcscncc of a passive surface layerl

the stability of which is controlled by the le~ds of several clcmcnts, especially Cr, Mo and NHence, there has been a continuous trend of increasing alloy lCVCIS(Mo and N in particular]relative to the common 300 series alloys] in order to produce steels with improved corrosionresistance (Table 1). The resultant “supcraustc nitic” steels offer cxccllcnt performance in a widerange of environments and redox potential, including chloride media, notably seawater. Matcria15with about 6%M0 and 0.2%N have rcachcd w idcsprcad acccptancc, and attention is being given;!to the application of newer grades with even higher nitrogen contents (Table 1).

Copyright

01996 by NACEInternational. Requests for permission to publish this m: nuscript in any form, in part or in whole must be made In writing to NACEInternational. Conferences Division, P.O. Box 218340, Houston, Texas 77218-8340. The matertal presented and the views expressed in tl~ispaper are solely those of the author(s) and are not necessarily endorsefl by the Association. Printed in the U.S.A.

Welding plays an essential part in modern fabrication practice, mainly by fusion, but alsoemploying solid state processes. However, ~vclding leads to changes in the local materialcomposition and microstructure which may have a pronounced effect on the passive layer andhence on resistance to environmental attack3-”. This paper describes the effects of a weldingoperation on high alloy austcnitic steels, with particular reference to the corrosion resistance ofcomplctcd joints.

METALLURGICAL EFFECTS OF WELDING

Alloy Element Segregation

Parent wrought high alloy austenitic stc el typically shows a homogeneous structure, withsome “banding” reflecting through–thickness compositional variation (Fig. la). When suchmaterial is melted and allowed to cool as in a cc nventional fusion weld, a quite different situationarises. These highly alloyed steels generally s ~lidify with a fully austenitic structure (Fig.lb),which causes marked segregation of alloying cl ~rnentsg. Chromium and Mo arc rejected from thesolid phase into the liquid, although nitrogen n-ay segregate in the reverse sense. Because of theaustenitic structure, diffusion coefficients are low, certainly for substitutional alloy clcmcnts, andconsequently segregation from solidification remains when the weld is cooled to roomtemperature. It is found that Mo displays pc~haps the most pronounced segregation into theinterdcndritic liquid, and microanalysis has shcwn that the difference between the dendrite coreand the interdendritic region may approach a f: ctor of two in a conventional fusion weld (Table

~).

This segregation leads to the formation of Me–lean areas which have significantly lowerpassive film stability than the nominal bulk composition suggests (Fig.2). Consequently, fusedmaterial has much poorer overall corrosion resistance than the parent (Fig.3)5’b unless anappropriate homogenizing heat treatment is employed, which requires high temperatures(>1OOO”C a d “) n 1s seldom practical, apart pcrh: ps from products such as scam-welded tube.

Segregation across dendrites (so–called rnicroscgrcgation) depends on the conditions underwhich solidification occurs, i ,c. the local tcmp~:raturc gradient G and the growth rate R. At lowgrowth rates, a planar solid/liquid intcrfdcc nlay be stable so that rnicroscgrcgation is not aproblem, but equilibrium segregation will o;cur ahead of the intcrface9. This leads to acompositional gradient along the solidified it~ml. At higher growth rate, the planar interfacebecomes energetically unstable due to constitu [ional supercooling and is replaced by a cellularor a dendritic interface. With increasing solid/liquid interface speed, the scale of the dendritesbecomes generally finer and the lCVC1of micro segregation dccrcascs due to solute trapping, i.e.insufficient time for equilibrium partitionin~! at the intcrfacelO, Eventually, at very highsolidification rates, morphological stability the[@l indicates that reversion to a planar interfacewill occur without segregation (Fig.4). This variation of alloy microscgregation indicates scopefor controlling corrosion resistance in as–solid ficd high alloy austenitic steels via modificationof the conditions under which solidification and segregation occur. The effect of thermal gradientis minor in relation to the solidification speed, but high thermal gradients also tend to promotefine structures with low lCVCISof microsegreg:~tion.

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Even though equilibrium conditions are not obtained during welding, the extent ofsegregation in a given alloy depends on the 1iquid/solid partition coefficients of the individualelementsg. It is well established that segregati{m of Cr and Mo is fairly low in alloys solidifyingto primary ferrite, in part because of rapid t iffusion in the solid state. Available cvidencex’l~suggests that segregation of these elements is also reduced in nickel-base alloys relative to thesuperaustenitic steels (Fig.5 and Table 2), although still pronounced compared to ferriticmaterials. Nevertheless, it is unlikely that significant change in the degree of segregation can beobtained in high alloy austenitic steel weld metals from optimisation of bulk composition.

Intermetallic Phase Precipitation

The formation of intcrmetallic comp rends such as sigma phase is most commonlyassociated with materials containing ferrite, due to the higher diffusion rates in ferrite than inaustenite. However, at sufficiently high alloy levels (especially Cr, Me), fully austenitic steelsbegin to show precipitation of Cr and Mo rich phases over relatively short timcscalcs4’]3comparable to the duration of a weld tlm-mal c YClC(Fig.6). The formation of intermetallic phasesis controlled by diffusion of substitutional elenlents. Hence, a particle is surrounded by an alloy–lean zone and the adjacent passive layer is ICSSprotective than for the homogcnous parent.Precipitation of interrnctallic phases involves complex interaction of several alloying elements,principally Cr, Mo, Cu and W. The effect is generally more pronounced as the total alloying isincreased, although study of high alloy aust .mitic steels has indicated a beneficial effect ofnitrogen, with Icvels of approximately ().4% acting to rcducc intermctallic formationappreciably ]2’14.

As a precipitation process, direct analc gy can be drawn between intcrmctallic formationand the problem of “weld decay” from chron lium carbide formation in the heat affcctcd zone(HAZ) of conventional austcnitic steels. Both phenomena cause loss of corrosion resistance vialocal depletion in alloying elements promoting passivity, immediately adjacent to the precipitatedparticle. The extent of alloy depletion depends on the size of the particle, and hcncc on the timein the precipitating temperature range availabl; for nucleation and growth, i.e. on the total weldthermal cycle. Noting that precipitation also depends on the local alloy content, intermctallicphases in alloy-rich intcrdcndritic weld metal regions have been variously rcported,7’]5’16as wellas in the HAZS of fusion wclds4.

ARC FUSION WELDING

Autogenous Welding

For many applications, the usc of auto[;enous welding procedures (i.e. without deliberatefiller metal addition) is desirable for producirg root passes. Root runs arc commonly made bythe gas tungsten arc (GTA) process. Problems of controlling root gap reliably are avoided byusing C1OSCfitting root lands in joint prepara ions, while control of the molten weld pool andhence bead profile is facilitated. Further, back purging with inert gas to prevent local oxidationis easier, while autogcnous welding removes uncertainty as to whether or not adequate filler hasbeen added. Unfortunately, autogcnous GTA welding gives heavily segregated dendritic weldmetal, with typical secondary arm spacings of the order of 5–10pnl, and, for example, variation

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of Mo across a dendrite arm by a factor of apploximatcly 1.5 to 2.07’8.This leads to a significantloss of corrosion resistance relative to the part nt for high alloy austcnitic weldments.

Trials to alter solidification rate have indicated that some improvement in corrosionresistance may be obtained at higher travel spc eds17/lower arc energies15 (Fig.7). This has beenfound for both established steel to UNS S31254, and also for more recent, higher nitrogen alloys(Fig.7a). The principle has been employed in laser surface treatment as a method of improvingthe corrosion resistance of autogcnous weldmellts and as–cast structurcs17’18.At sufficiently highbeam scan speeds and with appropriate surface preparation, it is possible to create reproduciblyvery fine dendritic structures (secondary arm spacings of the order of l–3~m) with rcduccdmicroscgrcgation (c .g. only 20% difference in 110 content between dendrite cores and edges) andcorrosion resistance more equivalent to that of the original homogeneous plate (Fig.7b).

These studies have shown that some control is possible over the degree of segregation inautogenous high alloy austcnitic weld metal and hence on its corrosion resistance. Specifically,low heat input and rapid solidification is to bc preferred. At the same time, the extreme case oflaser treatment will seldom be of practical application, while restriction of GTA I-oot runs to verylow arc energy will rcducc overall productivit:f. Further, even with these measures, it must berecognised that some segregation remains and that the corrosion resistance of the autogcnousweld bead is still inferior to that of the base rr,ctal.

Segregation of alloying elements affe(;ts passive film stability in a range of media.Particular attention has been paid to the influer ce of segregation on chloride pitting resistancel’zsince the superaustcnitic steels are attractive candidates for seawater service, being able towithstand a wide range of flow rates. However, segregation reduces passive film stability alsoin chloride–free nlcdia8, serving, for example, to increase the active/passive transition potentialin sulphuric acid environments (Fig.8). Given that segregation is inherent in autogenous weldsand reduces corrosion resistance in a range of environments, it has become normal practice toweld the supcraustcnitic steels with non-match ng ovcralloyed fillers as considered bclow~. Thisis not to say that autogcnous runs are not of practical USC,since they may be perfectly acceptablein conditions where supcraustcnitic alloys are :;pccificd bccausc materials such as 316 steel aremarginally inadequate.

Autogcnous welding may further bc viable if attention is paid to the nitrogen level of thedeposit. It is not always recognised that nitrogc n loss is an important factor in limiting corrosionresistance of autogcnous deposits rnadc using conventional argon shielding gas: nitrogen loss ofGTA runs is roughly 20% (Table 3) leading to corresponding reduction of corrosion resistance.With base steel dcvclopmcnt giving increased TJZIcvcls, such an effect may become increasinglyimportant in the newer alloys. Study of both 13TA7’15and plasma7 processes has indicated thebeneficial effect of nitrogen addition to the shi:lding gas on col-rosion resistance of the deposit.The use of Ar+3’%Nz shielding gas for GTA wdding can largely negate nitrogen loss in S31254steel, and even incrcasc the final nitrogen content, and gives approximate] y 10“C improvementin critical pitting temperature (CPT) (Fig.7a). Full nitrogen recovery in GTA welds may be moredifficult in high nitrogen alloys such as S32654, although better results have been reported forplasma welding.7 It should be noted for both proccsscs that excessive nitrogen gas addition canlead to electrode erosion and process instability. Usc of nitrogen-containing gases has not

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reached general industrial practice, but autogerlous weld deposits can be produced by this meanswith pitting resistance very similar to those o:’ welds made using non-matching fillers.

Welding with Filler Addition

The consequences of segregation in au togenous deposits were recognised fairly early inthe development and application of superaustcnitic stainless steels. The potential advantages ofusing an ovcrallo yed consumable were manifest, and the vast majority of fabrication in S3 1254has utilised a nickel–base filler obtaining about 9%M0, such as AWS ENiCrMo–3 orERNiCrMo-3 typcsb (Table 1), selected primarily on the basis of ready commercial availabilitycoupled with acceptable consumable costs. The recent, more highly alloyed steels may demandconsumables with increased Mo content7, say 16%M0 as in ENiCrMo–4 or ERNiCrMo–4, toobtain commensurate improvement in weld metal corrosion resistance, but the principle of“ovcrrnatching” remains valid (Fig.9)19. With o~cralloycd fillers, segregation still occurs but evena]loy-depleted regions contain sufficiently hig I Cr and Mo levels for the corrosion resistance ofbulk weld metal to approach that of the base :;tccl.

It must bc emphasised that there have been very fcw failures reported of welds madeusing non-matching consumables. Ncvcrthcless, even using such fillers, a number of points mustbc appreciated in designing a welding proccdurc. First, cxccssivc dilution of the molten pool bybase metal should be avoided. This is clearly a prc–requisite to maintain sufficient alloy contentthat all regions of the weld have high passive film stability, but is ncccssary also to avoid weldmetal compositions which are particular y sensitive to solidification cracking.b

Second, heat input to the joint should not be unduly high. This is advisable to reduce therisk of solidification cracking and possibly to limit the degree of segregation, but more especiallyto avoid any damaging effects of intcrmetallif: formation in the heat affcctcd zone around thejoint (Fig. 1())(q).A maximum arc energy of abollt 1 to 1.5kJhnm is in principle desirable, althoughthis will depend on the joint thickness and cffectivc heat sink, and a rather lower limit would berequired, for, say, 3mnl material. Together witl[ restriction of heat input, it is necessary to controlinterpass temperature and a maximum of 100--1500C is to be advised(b).

During a welding operation, it may be {Iifficult to ensure maintenance of a fixed root gapand thus to achicvc entirely consistent filler zddition. Some regions may therefore have fairlyhigh dilution by base steel and reduction in alloy content: as illustrated by T~ble 4, dilution andsegregation can result in regions of weld metal from 97oMo fillers containing only about 670M0.It has therefore been rccommcndcd that even m ore highly alloyed consumables should bc utilised,such as nickel-base fillers with 1670M0 for 6~$Mo S31 254 steel. This approach has not rcachcdwidespread use, and in f~ct may not bc particl]larly beneficial in terms of the overall corrosionresistance of the completed joint.20 Most data 01 the effect of composition on corrosion resistanceof high alloy materials have been derived from either base steel or weld pads made from a rangeof consumable types. The situation in a com~ letcd joint, however, is somewhat different, and,although filler addition is to bc preferred, increasing the total alloy content of the weld metaldots not necessarily improve the performance of an actual joint. This is illustrated in Fig. 1lN,showing that butt welds in 670Mo S31Z54 stcc display a CPT of 55–60°C, WC1lbelow the parentsteel lCVC1of 70”C, even with significantly ovcralloycd weld metal. Moreover, dilution so high

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that 6070 of the weld metal is constituted by fu wd base steel dots not greatly diminish corrosionresistance.

The effect arises because, given filler addition, the corrosion resistance of the weld islimited by the fusion boundary area where an Immixed zone (UMZ) exists.4>b’z0This is a narrowband of melted and resolidified parent steel which has not become mixed with the added filler(Fig.lZ). The region displays similar segregat .on of Cr and Mo to that in an autogenous weld(Fig.13) and thus loss of corrosion resistance relative to base steel (Fig.14), although the shortexistence time of the UMZ in the liquid state probably means that nitrogen loss is restricted.Weldrncnt pitting resistance in terms of the CPT can be directly related to the minimum alloycontent of dendrite or CC1lccntrcs resulting fr~)m segregation on solidification, as indicated bythe dotted line in Fig. 11. It has been found t~at the UMZ width depends on welding process,presumably via solidification conditions and pool liquid flow rate, and it tends to be morepronounced with GTA than shielded metal arc (SMA) welding for examplcm. The UMZ may alsobc narrower at the free surface than in the ccm rc of the bead depth, but formation of a UMZ tosome degree is inevitable with a fusion welding operation employing non-matching filler. Indeed,if autogenous welds are made with argon/nitrogen gas to enhance corrosion resistance of the weldmetal, it is again the UMZ which is limiting on corrosion pcrforrnance. The UMZ is by no meansspecific to supcraustcnitic steels but arises in z range of materials, including corrosion resistantnickel–base systems j(21~and it is difficult to sce how its formation can bc reliably avoided.Fortunately the practical conscqucnccs have b~:cn very few, and, even with an UMZ, corrosionresistance of welds in supcraustcnitic steels aml similar alloys can bc cxpcctcd to bc far superiorto that of joints in lCSShighly alloyed conventional austcnitic steels.

The interdendritic segregation of alloying clcmcnts can promote formation of intermetallicphases in weld metal, even under “normal” coaling conditions which cause only negligible HAZprecipitation. This is the case in both autogcn ms welds and in Ni–base deposits .7’15’lbA rangeof phases is formed, sigma, chi ctc, as indicat(:d by the EDX analysis data in Table 4. In ferricchloride testing at high tcmpcraturcs, attack [:an take place in association with second phaseparticles in the weld metal (Fig.15). However, they do not seem to have been associated withpractical problems, and it is likely that any tffcct of alloy depletion around the particles issecondary to those of bulk segregation and of the UMZ.

OTHER WELD[NG PROCESSES

Resistance Welding

Given their high resistance to passive film breakdown, the supcraustenitic steels arcattractive materials for plate heat exchangers, resisting attack at the internal crcviccs. Such unitshave given satisfactory service in a range of conditions and with diverse cooling media.

Plate heat exchangers arc commonly nxidc using resistance spot or scam welding to formthe necessary channels, and cxpericncc has shown that the supcraustcnitic steels can bc reliablyresistance welded for this purpose (Fig.16). Ai the same time, it should be noted that the alloyshave higher mechanical strength than conventional austcnitic grades, and this affects theirresponse to a resistance welding operation. Because of the rapid solidification and cooling of the

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nugget, resistance welds commonly display a :endcncy for centreline shrinkage defects to form,and the increased strength of superaustenitic gladcs means that high loads are required during theforge stage of the welding cycle to consolidate the nugget and minimise “porosit y“. Furthermore,it is nornlal in resistance welding to set condil ions such that a simple peel test results in failurearound the nugget. However, the high strength of supcraustenitic grades relative to 300 seriesalloys reduces the ability to deform locally at the nugget periphery, and it may be ncccssary toincrease the nugget size to obtain pull out fail[lre on peel testing with the higher alloy materials.

A further characteristic of superaust{:nitic steels, and of other high alloy materialsincluding nickel–base systems, is that grain b(nmdary segregation can occur in the HAZ duringthe welding cycle. This can be quite pronount cd in resistance welding (Fig. 16), and conditionsmust be controlled such that the scgrcgatcd b~mndaries do not extend to the free surface wherethey might promote some reduction in corrosion resistance. Hcncc, while superaustenitic steelsarc certainly arncnablc to resistance wcldin ~, conditions must be carefully controlled anddemonstrated to bc reproducible by appropriac prc-production trials.

Friction Welding

Joining supcraustcnitic steels by a solid state process is attractive since segregation effectsassociated with fusion should bc greatly rcduccJ or even eliminated. Under production conditions,solid state joining is most frequently achicvc[l by friction welding in which one component ofthe joint is moved relative to the other to generate frictional heat: sufficiently high metaltcmpcraturcs arc rcachcd for local softening to occur so that joining can then be effected bystopping the relative movement and applying a forge stage. Most commonly, components ofcircular symmetry arc joined by relative rota ing movement, but in principle both orbital andlinear oscillation movement can be utilised.

It has been shown that supcraustenit ic steels can be successfully joined by frictionwclding22 (Fig. 17). The high mechanical strclgth at elevated temperature requires fairly highfriction and forge pressures to induce sufficic]lt heating and plastic flow at the faying surfaces,but the materials arc, with this proviso, tc lcrant to variation in welding conditions. The“conditioning” time of frictional heating is higllcr than for 300 series steels: this has no pal-titulareffect on quality, but for mass production appl cations it may be desirable to increase the frictionload appropriately to maximise production rate. The friction process leads to some localhardening from the effective cold work in tht: bond area although this does not seem to be ofparticular significance in most practical purposes.

By avoiding significant fusion and associated segregation, friction welds in supemustcniticsteels of a given pitting index value display appreciably higher corrosion resistance than dofusion welds (Fig.18). Indeed, the corrosion rcsistance of a friction weld may be very similar tothat of the base metal. The full performance (If the parent steel is not nomlally achieved, sincethe scvcrc local strain associated with friction welding seems to enhance precipitation ofintcrmctallic phases, and this has an adverse effect on corrosion rcsiskmcc (Fig.19). Frictionwelding leads to the formation of a flash at the joint (Fig. 17) which must be removedmechanically for optimum properties, but the total joining time may bc much lCSSthan that ofan arc weld, and the process can offer economi~ advantage for components of suitable geometry.

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Power Beam Welding

Because of the high travel speed and thus joint completion rate achievable, industrialinterest is increasing in laser and electron b(;anl welding. Little information is available onapplication of the Iattcr process to superausten itic alloys, but, even if weld metal segregation islimited by the inherent rapid solidification, st)me reduction in nitrogen content and corrosionproperties must be expected, bearing in mind that welding is conventionally carried out in avacuum.

Trials with the laser process, however, have indicated that this is viable for supcraustcniticsteel s.17’23Autogenous laser welding with n[~mlal gas shielding reduces base metal pittingrcsistancc, although to a lesser extent than GTA welding. However, addition of nitrogen to thelaser gas shroud above the molten metal significantly enhances pitting pcrfornlancc,17 to a lCVC1equivalent to that of arc welds with added no~.-rnatching filler.

GENERAL COMMENTS

The superaustcnitic steels should be wel(led with similar precautions regarding clcanlincss,backing gas etc as employed for conventional zustcnitic materials.b Indeed, with the proviso thatopen root gaps are essential to ensure adcquatf: filler addition, welding practice rcrnains largely

b9~-~@It nlust bc rccogniscd thatunaltered, and all the common arc proccsscs art of application. ( ‘-welding normally entails nickel-base consumables. The specific handling characteristics of nickelalloy SMA clcctrodcs can differ slightly frorr austcnitic fillers, while, unlike stainless steels,Ar/02 shielding gases would not bc used for GldA wclding,2q’25since they tend to give weld pooloxidation and pmccss instabilityy. For this process, a range of gas can be considered: pure Ar isprobably the most common, but Ar/He, He/AriC02 ctc have been rcconm~endcd.25 Appropriatewelder training is therefore essential, particularly perhaps in root run technique.

From the corrosion viewpoint, the kc~r point in filler selection is that the weld metalshould bc sufficiently ovcralloycd relative to thl; base steel to accommodate dilution and maintainhigh alloy content in all regions of the solidified structure. For steels such as S31Z54, W6M0fillers secm to bc generally adequate, while l(l%Mo consumables are required for recent morehighly alloyed grades. It has been observed that bend ductility can bc fairly low using sornccommercial consumables containing Nblb: tk is dots not necessarily mean that weld metalmechanical properties arc inadequate for scrvi~:e, but the view has been expressed that Nb–freefillers arc to be preferred.

Even with filler addition (or argon/nitrogen shielding gases), fusion welds may not havecorrosion resistance fully equivalent to the ba: e metal, most especially bccausc of the effect ofthe fusion boundary unrnixcd zone. This is thl; case both for established 670Mo steels such asS31Q5420 and also for newer more highly alloyed grades.7 It is rcitcratcd that the scrvicc history

of welded supcraustenitic steels has been cxc(:llcnt, but it is clearly prudent to include weldedjoints in any material evaluation study for sc-vice under conditions where there is little priorexperience.

42C/8

Recognizing the high corrosion rcsistanc~; of superaustcnitic alloys, even autogenous weldswith segregation may well be adequate for service where the supcraustcnitic steels are specifiedto replace common 300 series alloys. Autogcr ous welding can give joint corrosion resistanceequivalent to that of welds with filler if low arc energy is utilised in combination withargon/nitrogen shielding gases. This approach is not yet common industrial practice, but certainlyhas merit for future fabrications.

Finally, highest joint corrosion rcsistancl; is likely to be achieved by solid state proccsscs,most especially friction welding. Such an approach should be considered for components havingsuitable geometry.

CONCI.,USIONS

The superaustcnitic steels have now been used for welded fabrications for over a decade,and have given an excellent scrvicc history. In large part, the guidelines for successful weldingarc WC1lunderstood. Arc welding remains dominant in fabrication, and the following aspects inparticular should be noted.

i.

ii.

. ..111

iv

v.

vi

Welds should be made using ovcralloycd non-matching filler: nickel-base consumableswith 9%Mo arc normally employed for established grades such as S31254, while fillerswith 16YoM0 are required for more highly alloyed steels.

The joint preparation and welding proce~ure should be controlled to ensure that adequatefiller addition is made, most especially ill root runs exposed to the corrosive environment.From the corrosion standpoint, dilution of up to 6(E% may be tolerable, but lower levelsare prcfcrablc to rcducc the risk of solidification cracking.

Arc energy and intcrpass temperature should be fi~irly low. For most applications, amaximum arc energy of lkJ/nml should be used, although this would be dcpcndcnt on thecomponent thickness. Interpass tcrnpcrzturcs should normally be below 100–150°C.

Welding practice follows that for conv(,ntional 300 series austcnitic steels, although theparticular rcquircmcnts of Ni-base fillers must bc recognised.

Autogcnous welding nornlally causes lCSSof base metal corrosion resistance, because ofalloy element segregation and nitrogc n loss. However, addition of nitrogen to theshielding gas can give joint corrosio;1 resistance equivalent to that of welds with

overalloycd filler.

Joint corrosion resistance approaching Ime metal behaviour can bc obtained using solidstate friction welding.

ACKNOWLEDGEMENTS

The authors thank colleagues at TWI for advice and assistance regarding this paper.

Grateful acknowledgement is made to Avcsta- Sheffield AB for permission to use Figure 16.

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REFERENCES

1.

9Q.

3.

4.

5.

6.

7.

8.

9.

10.

11.

1~.

13.

14.

15.

16.

17.

Davison R M, et al, CORROSION’86, Paper 185 (Houston, NACE, 1986).

Wal16n B, Liljas M, Stenvall P, Applications of Stainless Steels ’92, Stockholm Sweden,(The Institute of Metals, London, 199~ p. 23.

Garner, A: “Pitting corrosion of higk alloy stainless steel weldments in oxidisingenvironment”, Weld. J. 62, 1(1983), p. 27.

Liljas M, Holmberg B, Ulander A: “Welding of a high molybdenum stainless steel”, Procconf stainless steels ’84, (London, The Metals Society, 1984) p. 232.

Suutala N, Kurkcla M: vidc rcf 4, p. 240.

Rabensteincr, G: “The welding of fully imstenitic stainless steels with high molybdenumContents “, Welding in the World 27, IL! (1989), p. 2

Liljas M, Stenvall P, “Welding of UNS 32654- corrosion properties and rnctallurgicalaspects”, 12th International Corrosion Congress, Houston 1993 (NACE International,Houston, 1993) p. 2882.

Marshall P I, Gooch T G, Corrosion 49, 6 (1993), p. 514.

David S A, Vitck J M, Int Mat Rev, 34, 5 (1989), p. 213.

Aziz M J, J Appl Phys, 53, 2 (1982), p. 1158.

Kurz W, Giovanola B, Trivcdi R, Acts Mctall, 34, 5 (1986), p. 823.

Ogawa T, Koscki T, J Jap Weld Sot, 9, 1 (1991), p. 154.

Salbu H, “The influence of microstructure on corrosion resistance and impact strength of6Mo castings”. Corrosion and Matcriak Offshore, Oslo 1994 (NITO, Oslo, 1994).

Charles J et al, “A new high nitrogen : ustcnitic stainless steel with irnprovcd structurestability and corrosion resistance propc]tics”, vidc ref.13.

Ginn B J, Gooch T G, “Pitting rcsistanc{: of autogenous welds in UNS S31254 high alloyaustenitic stainless steel”, vidc rcf 7, p. 2895.

Karjalaincn L P, Jiirvenpaa S A, Lcinon~:n J I, “Welding of Cr–Ni–6Mo–N Type Steels”.

JOM5 Helsinki 1991 ,(JOM-Institute, F.elsingor, Dennlark 1991) p. 462.

Woollin P, “Laser surface melting o:’ high alloy austcnitic stainless steel”, Fourth

International Conference on Trends in \Vclding Research, Gatlinburg, USA, 1995 (ASM

420’10

International, to be published).

18.

19.

’70.

Nakao Y, Nishimoto K, “Effects of laser surface melting on corrosion resistance instainless steel and nickel-base alloy clal~ layers of cast hi–metallic pipes”, IIW Doc.IX–1666–92, 1992.

Bonnefois B, Gagncpain J C, Dupoiron I;, Charles J, “Welding of supcraustenitic stainlesssteels with very high nitrogen contents”, 34 ieme Journee du Cercle d’Etudes des M6taux,“Nitrogen as alloying or strengthening ~:lcmcnt”, Ecolc Nationale Superieure des Minesde Saint Etienne, France, 1995.

Gooch T G, Elbro AC, “Welding corrosion resistant high alloy austenitic stainless steel”,

Corrosion in natural and industrial env ronments: problems and solutions, Grade, Italy

1995 (NACE International, Italia, Monj:a), p. 351.

Fkischc L, “Localised corrosion of the unmixed zone in nickel-base alloy weldments”vidc rcf 7, p. 2907.

Gooch T G, Gunn R N, Dunkerton S 13. “Friction welding and corrosion resistance ofhigh alloy austcnitic stainless steel”, VJcldability of Materials, (Materials Park, Ohio,ASM International, 1990), p. 81.

Arlt N, Heuser H, Gross V, Rcvista de Soldadura 24, 1 (1994), p. 14.

Agarwal D C, vidc rcf 22, p. 107.

Holrnbcrg B, Stcnbacka N, Svctscn 49, Issue lE (1990) p. 24.

Davison R M, Redmond J D, van Dcellm B J, vide ref 22, p. 91.

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T~ble 1 Compositions of typical high alloy au:;tenitic stainless steels and welding consumables

Element, wt% (a)UNS 01 AWS MinimumNo. c Mn Fe Cr Ni Mo cll Nb N PRFW(b)

S31254 0.02 1.0 bal 19.5- 17.5- 6.0- 0.5- - 0.18- 44.720.5 lE.5 6.5 1.0 0.Z2

N08367 0.03 2.0 bal 20.0– 2! .5- 6.0- - - 0.18- 45.222.0 25.5 7.0 0.25

S34565 0.03 5.0- bal 23.0- 16.0- 4.0- - 0.1 0.4-0.6 48.27.0 25.0 18.0 5.0

s3~fj54 0.02 2.0– bal 24.0– 21.o– 7.0- 0.3- - 0.45- 60.64.0 25.0 ‘2:.0 80 0.6 0.55

ENiCrMo–3 0.10 1.0 7.0 20.0– >;5.0 8.0- 0.5 3.15- - 46.423.0 10.0 4.15

ERNiCrMo–3 0.10 0.5 5.0 20.0– >:8.0 8.0- 0.5 3,15- - 46.423.0 10.0 4.15

ENiCrMo–4 0.02 1.0 4.0- 14.5- baI 15.0- 0.5 — — 64.0(c) 7.0 16.5 17.0

131{NiClMo-4 0.02 1.0 4.0 14.5- ba1 15.0- 0.5 — — 64.0

(c) 16.5 17.0

3 All figures quoted arc maximum value:; except where a mngc is specified:

(b)

(c)

T~ble

bal = balance: - = not specified.PRE~ = ‘%Cr + 3.3%M0 + 30%N<2.5%C0, SO.3570V, 3.O–4.5TOW

2 Results of electron microprobe analysis on alloy weld metals.(~)

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Table 3 Nitrogen analyses on root runs in supc raustenitic steels.(’)

Nitrogen, wt%Steel type Welding Filler type

process Measured Calculated*

S31254 GTA Auto gcnous 0.15 0.21GTA ERNi(;rMo–3 0.09 0.11SMA ENiCrMo-3 0.13 o.1~

S3~654 GTA Auto/ ynous 0.36 0.50GTAN Autoi;enous 0.41 0.50

,

GTA: Ar shieldingGTAN: Ar/lO%Nz shielding* Calculated from base steel and filler nitrogc]- levels and weld dilution.

Table 4 EDX analyses on root runs of butt wc .ds in S31254 base metal.

Process/Consumable Element

GTA/9Mo(b) MoCrNb

sMA/9Mo(b) MoCrNb

Region analys

‘ulk E

6.1 4.?

20.2 19.3

7.6 5.8~1,4 20.7~.7 0.9

79 59‘70.7 ~~.3

1.6 0.5

clement wt%

lntcrdendriti Secondc phase

10.5 15.9~().9 1~.~

5.3 13.7

9..3 9.621.2 19.71.1 8.6

GTA/16Mo(c) M() 11.3 9.8 14.4 38.5Cr 17.3 172 18.1 17.6

SMA/16Mo(c) Mo 11.6 9.0 1~3 —

Cr 18.7 18.4 195

(a) Autogcnous (b) E/ER NiCrMo-3 (c) E/ER NiCrMo-4

420,13

.“ ,,. ” ../“-”\. -+ ,.

.,‘.,

--

\. . ‘-’ ,“” ‘1,,

\\. ~,, ●

(a) . -1-

AC2492.- ‘!, \,

Fig.1:

a) Optical micrograph of typical austenitic base steel microstructure, x200;b) Scanning electron micrograph showing segregation and coring in fully austenitic, high alloy

weld metal, nominal x2000.

-.,, f’..,- -../. ‘

Fig.2 Section through laser melted S31254 steel, illustrating attack in FeCIAsolution on dendrite centres, x200.

80 I 1 1 1 1 I I

Fig.3 Critical pitting temperature data fromferric chloride tests on stainless steel basemetals and weld metals.5

01” I I 1 I I I

.?5 30 35 40 45 50 55“ACr+33x O/. Mo. 13x0AN

420/” 4

>’/

/’//I

0/

\, ,0/0

\ #~ — 0(11. k(T1. M(T)--- D,kO; M= &dant

1

70”2 10-’ 10° 701 102 ?8R, cm/sec

Fig.4 Effect of solidification rate (R) on dendritecore composition (C,~), with diffusion coefficients(D)r partition coefficients (k) and solidus slope (m)as constants or functions of temperature (T).10

‘0” ~

I____l1 23510 2050507m

lime, min

Fig.6 Effect of time and temperature onintermetallic precipitation in S31254 steel .13

100

80

60

–.-05 1.0 1,5 2.0 2.5

Molybdenum sgqmtron mtlo

Fig.8 Dependence of stainless steel weld metalpassivation on interdendritic Mo segregation:30’XOH2S04 with 0.1 g/litre NHd CNS at 25”C.14

~,;y~ MO,Ill ---a

‘o—

r‘I”l

80 ?0---0 _Cr

04 os oti 0.7 0.8 09 m 1.1r 1.6

Ni~ /[r ~

Fig.5 Relationship between element

segregation, partitioning and bulk

composition and solidification mode.8

100

------ ------ ______ ______ 180t It--------------------------------------------

pt=’u

u 40 -

--- S34565 parent----- S31254 parent

20 - — S34565 weld metal {Ar)— S31254 weld metal (Ar)-- S31254 weld metal (Ar+3%N2)

1 1 1 Jo 0.5 1.0 1.5 :?.o

(a) Arc energy, kJ/mm

80- I 1

To Base●

mdal //. ● /~ ‘1

.-

1 ● 0- ●/ I

3,-1ma?

(b) Laser veloclty,mmlmin

Fig.7:

a) Effect of arc energy and shielding gas oncritical pitting temperature (CPT) in ferricchloride of autogenous GTA welds in S31254and S34565 steels;15

b) Effect of beam travel speed on measured CF’Tof laser treated parent UNS S31254 steel.Squares represent use of N2 trail gaS.17

420/15

70

m

40

I I I

t, H

fWY- -

~&/

-/ J.// . Super austenitic BM

● ,0 0 Matching WM!

El Overalloyed WM

Fig 9 Effect of composition on CPT in FeC13

tesls for high alloy materials.19

20 I I I

30 40 SO 60 n 80

Cr0A3,3 {Mo %+ 0,5woA)+30N%

Fig.10 Attack in ferric chloride on HAZ of SMA weld in S31254 steel

with ENiCrMo-3 filler, xIO.~

~ A= Parent; V= Autogenous

weld metal/“ ❑ = UMZ

● = GTA + SMA welds9%M0 + 15yoMo fine 1

do ~

30 40 50 60%Cr+3.3%Mo+ 16%N

Fig.11 Relationship between FeC13pitting resistance of welded joints inUNS S31254 steel and composition.20

(Shaded area = weld metal bulk composition,Numbers = % dilution by base metal,Dotted line = minimum local alloying level in

solidification structure. )

../

.

.

Fig.12 Unmixed zone at fusion boundary of SMA weldin S31254 steel with ENiCrMo-3 filler, x320.

420f” 6

I 1 I 1 1 1 1 I

20 40 60 80Distance, microns

Fig.13 Molybdenum EDX scan across fusion

boundary of GTA weld in S31254 steel with

ERNiCrMo-3 filler.zo

d-. .. c....” . .

... . ,.. .. . .

r). ,,, . .

-be “ .-O. ‘c

“- %’” 4’

)

Fig.14 pitting attack initiated at UMZ in GTAweld in S31254 steel: ERNi CrMo-3 filler, x25,

Fig.15 Preferential corrosion in FeC13 associated with

sigma phase in autogenous S31254 GTA weld metal,X500.

Fig.16 Transverse section through resista!lce seam weld in S31254 steel, showingHAZgrain boundary Iiquation, x30.

. ‘j ‘“’. -,, I

,... ,.

●

b’,. ~

●

. . ..- .. ‘

— ❑ parent material

/

P– – o Friction welds

----- ● GTA welds /

/

/

/

1/

I

/

I

20 30 40 50

%Cr+3.3%Mo+ 73%N

Fig.18 Pitting potentials in oxygen-fresimulated seawater at 50”C for GTAand friction welds in austenitic stainlesteels: effect of base steel and weldmetal composition.zz

‘, ,. ,9

,

‘, .-,, .

Fig.19 Pit initiation in C02 s]turated synthetic seawaterat 60°C, at bond line and HAZ in friction welded S31254,X12.22

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