134567205 103677583 Corrosion of DSS Weldments

CHAPTER 5

Corrosion of Duplex Stainless SteelWeldments

DUPLEX STAINLESS STEELS are two-phase alloys based on the Fe-Cr-Ni system.These materials typically comprise approxi-mately equal amounts of body-centered cubic(bcc) ferrite and face-centered cubic (fcc)austenite in their microstructures and are char-acterized by their low carbon content (

100 / Corrosion of Weldments

Table 1 Composition and pitting resistance equivalent numbers (PREN) of wrought duplex stainlesssteels covered by UNS designations

Composition(a), wt%

UNS No.Common

designation C Mn S P Si Cr Ni Mo Cu W N2 PREN

Low-alloy grades (PREN

Chapter 5: Corrosion of Duplex Stainless Steel Weldments / 101

Fig. 1 The development of hot-rolled duplex stainless steelsin terms of their nitrogen versus chromium plus molyb-

denum contents. Source: Ref 1

Fig. 2 Effect of orientation plane on the microstructure of Fe-22Cr-5.5Ni-3Mo-0.15N wrought duplex stainless steel base materialelectrolytically etched in 40% NaOH. (a) Parallel to rolling direction. (b) Transverse to rolling direction. (c) Plan view. 100

show a typical microstructure representing theweld metal in the as-welded and the solution-annealed conditions. The darker areas, as forexample in Fig. 3(d), represent the ferrite phase,while the lighter areas represent the austenitephase. The austenite islands are coarser in thesolution-annealed microstructures as comparedto the as-welded microstructures. The coarsegrain HAZ of the as-welded condition wasremoved by solution-annealing heat-treatment,as shown in Fig. 3(b) and Fig. 3(e). Thus, solution-annealing has the benecial effect ofeliminating the coarse grain HAZ which isusually detrimental due to carbide precipita-tion. The weld metal microstructures (Fig. 3c and 3f ) revealed coarser austenite regions in thesolution-annealed weld metal as compared to theweld metal in the as-welded condition.

Mechanical and Physical Properties. Du-plex stainless steels characteristically arestronger than either of their two phases consid-ered separately. The duplex grades have yield

strengths twice those of the common austeniticgrades while retaining good ductility (Table 2).In the annealed condition, the duplex gradeshave outstanding toughness. With the morerecently developed intermediate and high-alloygrades, it is possible to retain toughness and cor-rosion resistance after welding.

The coefcient of thermal expansion and theheat-transfer characteristics of the duplex stain-less steels fall between those of the ferritic andthe austenitic stainless steels.

When installing a duplex stainless-steel com-ponent in an existing austenitic stainless-steelstructure, consideration should be given to therelative strengths and expansion coefcients ofthe materials. The high strength of the duplexgrade and its relatively low expansion coef-cient may impose high stresses on the transitionwelds or the host structure.

Elevated-Temperature Properties. Thehigh alloy content and the presence of a ferriticmatrix render duplex stainless steels susceptibleto embrittlement and loss of mechanical proper-ties, particularly toughness, through prolongedexposure to elevated temperatures. This iscaused by the precipitation of intermetallicphases, most notably alpha prime () sigma (),chi (), and Laves () phases. For this reason,duplex stainless steels are generally not used attemperatures above 300 C (570 F). Figure 4shows the phases that can be formed in duplexstainless steels over the temperature range of 300to approximately 1000 C (5701830 F).

Corrosion Resistance. Duplex stainlesssteels comprise a family of grades with a widerange of corrosion resistance. They are typicallyhigher in chromium than the corrosion-resistant, austenitic stainless steels and havemolybdenum contents as high as 4.0%. Thehigher chromium plus molybdenum combina-tion is a cost-efcient way to achieve good chlo-

(a) (b) (c)


Fig. 3 Solidication morphologies of fusion welded alloy 2205. (a) As-welded base metal. (b) As-welded composite region. (c) As-welded weld metal. (d) Postweld heat treated solution-annealed base metal. (e) Solution annealed composite region. (f) Solu-

tion annealed weld metal. Source: Ref 2

ride pitting and crevice corrosion resistance.Many duplex stainless steels exceed the chlo-ride resistance of the common austenitic stain-less-steel grades and also alloy 904L (UNSN08904) (Table 3). SAF 2507 (UNS S32750)has chloride resistance comparable to the 6%molybdenum austenitic stainless steels.

The constraints of achieving the desired bal-ance of phases dene the amount of nickel in aduplex stainless steel. The resulting nickel con-tents, however, are sufcient to provide signi-cant benet in many chemical environments. Asshown in Table 4, alloy 2205 and Ferralium 255(UNS S32550) compare favorably with type

317L (UNS S31703) and alloy 20 (UNSN08020) in a variety of chemical environments.

One of the primary reasons for using duplexstainless steels is their excellent resistance to ClSCC. Compared with conventional austenitics,they are clearly superior (Fig. 5). The morehighly alloyed superduplex grades are moreresistant to Cl SCC than those with lower alloy-ing contents. The SCC resistance in the an-nealed condition of the superduplex grades iscomparable to that observed with highly alloyedaustenitic grades like 20Cb-3 (UNS N08020)and the 6% molybdenum superaustenitics likeAL-6XN (UNS N08367).


Table 2 Comparison of mechanical properties of commonly used stainless steels in the annealedcondition

Tensile strength

UNS No. Common designation MPa ksi MPa ksi Elongation, %

Austenitic grades

S30400 Type 304 515 75 205 30 40 88N08020 20Cb-3 585 85 275 40 30 95Duplex grades

S31803 2205 620 90 450 65 25 30.5 HRC(b)S32750 2507 800 116 550 80 15 32 HRC(b)Ferritic grades

S40900 Type 409 415 60 205 30 22(c) 80S44625 E-Brite 26-1 450 65 275 40 22(c) 90(a) At 0.2% offset. (b) Typical values. (c) 20% elongation for thicknesses of 1.3 mm (0.050 in.) or less. Source : Ref 1

Yield strength(a)

Hardness (max), HRB

General Welding Considerations

The performance of duplex stainless steelscan be signicantly affected by welding. Due tothe importance of maintaining a balancedmicrostructure and avoiding the formation ofundesirable metallurgical phases, the weldingparameters and ller metals employed must beaccurately specied and closely monitored. Thebalanced microstructure of the base material(that is, equal proportions of austenite and fer-rite) will be affected by the welding thermalcycle. If the balance is signicantly altered andthe two phases are no longer in similar propor-tions, the loss of material properties can beacute. Because the steels derive properties fromboth austenitic and ferritic portions of the struc-ture, many of the single-phase base materialcharacteristics are also evident in duplex mate-rials. Austenitic stainless steels have excellentweldability and low-temperature toughness,whereas their Cl SCC resistance and strengthare comparatively poor. Ferritic stainless steelshave high resistance to Cl SCC but have poortoughness, especially in the welded condition.A duplex microstructure with high ferrite con-tent can therefore have poor low-temperaturenotch toughness, whereas a structure with highaustenite content can possess low strength andreduced resistance to Cl SCC (Ref 3). The highalloy content of duplex stainless steels also ren-ders them susceptible to formation of inter-metallic phases from extended exposure to hightemperatures. Extensive intermetallic precipita-tion may lead to a loss of corrosion resistanceand sometimes to a loss of toughness (Ref 4).

Duplex stainless steels weldability is gener-ally good, although they are not as forgiving asaustenitic stainless steels or as prone to degrada-tion of properties as fully ferritic stainless steels.The current commercial grades are low in carbon(less than 0.03 wt%), thereby essentially elimi-nating the risk of sensitization and intergranularcorrosion from carbide precipitation. The basematerial and ller metals also have low sulfur andphosphorus levels (less than 0.03 wt%), which incombination with the ferritic solidicationreduce the likelihood of solidication cracking(hot cracking). Hydrogen cracking (cold crack-ing) resistance is also good due to the high hydro-gen solubility in the austenite and the high per-centage of austenite in the matrix. Nevertheless,hydrogen cracking can occur in duplex alloys,and is discussed later in the section CorrosionBehavior of Weldments.

Fusion Welding. Nearly all of the arc weld-ing processes that are employed for other stain-less steels can be used with duplex alloys,except where the process characteristic is toweld autogenously, such as electron-beamwelding and laser-beam welding. In such cir-cumstances a PWHT is nearly always necessaryto restore the correct phase balance to the weldmetal and remove any undesirable precipitates.There are few reported differences in corrosionresistance between welding processes, but thenonmetallic inclusion distribution would beanticipated to have an effect. In most instancesof pipe welding where access is from one sideonly, gas-tungsten arc welding (GTAW) isalmost exclusively employed for the root pass (Ref 5). This provides a controllable, high-


Fig. 4 Time-temperature-transformation (TTT) curve for alloy2205 (UNS 31803) showing the effect of alloying ele-

ments on precipitation reactions. These phases can negativelyaffect the corrosion resistance and the ductility of the materialand are the most serious threats to the successful applications ofduplex grades. Source: Ref 4

Table 3 Critical crevice corrosion temperatures

Critical crevice temperature in 10%FeCl3 6H2O; pH = 1;24 h exposure

UNS No. Common name C F

Duplex grades

S32900 Type 329 5 41S31200 44LN 5 41S31260 DP-3 10 50S32950 7-Mo PLUS 15 60S31803 2205 17.5 63.5S32250 Ferralium 255 22.5 72.5Austenitic grades

S30400 Type 304


Fig. 5 Stress-corrosion cracking (SCC) resistance of selectedduplex stainless steels (S31803, S32304, and S32750)

relative to austenitic stainless steels (S30400, S30403, S31600,and S31603) as a function of temperature and chloride concen-tration in neutral O2-bearing solutions (approximately 8 ppm).Test duration was 1000 h. Applied stress was equal to yieldstrength. Source: Ref 1

the as-welded phase balance and increase aus-tenite content. The ferrite content of a weld madewith a nickel-enriched consumable would de-crease signicantly if it underwent a PWHT. Itmay suffer from slightly reduced weld metalstrength and could also be more susceptible to phase formation during heat treatment (Ref 6).The nickel level in the enriched weld metal willbe approximately 2.5 to 3.5% greater than in thebase material (for example, for the Fe-22Cr-5.5Ni-3Mo-0.15N duplex stainless steel basematerial containing 5.5% Ni, the ller metal willcomprise 8.0 to 9.0% Ni, depending on consum-able manufacturer and form).

The higher alloy ller metals are sometimesused for welding a less alloyed base material (forexample, a duplex stainless steel ller metal with25% Cr could be used for the root run in a Fe-22Cr-5.5Ni-3Mo-0.15N base metal). This is usu-ally done to improve root weld metal corrosionresistance and thereby pass the qualication testrequirements. In most cases, this does not lead toloss of mechanical properties; indeed, the morehighly alloyed ller metal in the case previouslyis likely to have greater strength.

To avoid all the requirements for weld metalphase balance and microstructural control nec-essary with duplex ller metals, nickel-baseconsumables (for example, AWS A 5.14 ERNi-CrMo-3) have been used. The yield strength,however, is slightly below that of the morehighly alloyed grades, and the lack of nitrogen

and the presence of niobium in the ller metalmay contribute to unfavorable metallurgicalreactions and the formation of intermetallic pre-cipitates and areas of high ferrite content in theHAZ (Ref 7, 8).

Preheat is generally not recommended forduplex stainless steels, but may sometimes bespecied in low-nitrogen grades, because thicksections and low heat input welding processesmay, in combination, develop highly ferriticHAZs (Ref 9). For the more highly alloyedduplex stainless steels, a preheat can be highlydetrimental and reduce corrosion resistance andmechanical properties.

Postweld heat treatment is not commonlyused except in autogenous welds or welds witha ller metal composition that exactly matchesthe base steel. Although not always necessary,particularly if a nickel-enriched ller metal isused, it is common to PWHT duplex stainlesssteel welded pipe after longitudinal seam weld-ing. The PWHT will largely be for the purposeof restoring the correct phase balance and redis-solving unwanted precipitates. Postweld heattreatment temperatures of approximately 1050to 1100 C (1920 to 2010 F) are used, depend-ing on grade, followed by the same heat treat-ment applied to the base material during solu-tion annealingusually water quenching. Theheat treatments commonly used for structuralsteels (for example, 550 to 600 C, or 1020 to1110 F) are totally inappropriate for duplexalloys and should never be considered.

Corrosion Behavior of Weldments

Corrosion characteristics of duplex stainlesssteel weldments are complex. The HAZ suffersmore corrosion attack than either the base metalor the weld metal because of the unbalance inaustenite/ferrite fractions in the HAZ (Ref 2).Pitting corrosion resistance of wrought duplexstainless steels is superior to the cast version.Less austenite is typically present in the caststructure. Thus, the duplex stainless steel weldconsumables are enriched with nickel toachieve higher levels of austenite in their as-welded microstructures, since the weld metal isessentially a cast material.

The welding heat input affects the pitting cor-rosion resistance of the duplex stainless steelweldments. As shown in Fig. 6, the best pittingcorrosion resistance is achieved when the weld-


Fig. 7 Effect of ferrite-austenite balance on pitting resistanceof Fe-22Cr-5.5Ni-3.0Mo-0.12N gas tungsten arc stain-

less steel welds. Source: Ref 3Fig. 6 Inuence of heat input on corrosion of welded S31803

steel in ferric chloride. Source: Ref 3

ing practice involves higher heat inputs (Ref 3).In addition, cooling rates also affect the pittingcorrosion resistance. The slower the coolingrate, the better is the pitting corrosion resistance(Ref 2). Best corrosion resistance and mechani-cal properties are achieved when approximatelyequal amounts of austenite and ferrite are pres-ent in both the weld metal and the HAZ (Ref 2).A balanced austenite/ferrite content can beachieved by slowing the cooling rate, throughhigh heat input, preheating in multipass weldingoperations, and controlled interpass tempera-tures. The interpass temperature is usually keptbetween 150 and 200 C (300 and 400 F).

Alloying elements in duplex stainless steelsplay a key role in determining the mechanicaland corrosion properties. Due to the highchromium content, duplex stainless steels haveexcellent high-temperature oxidation resis-tance. However, they are prone to carbide pre-cipitation and to phase and chromium nitride(e.g., CrN, Cr2N) formation (Ref 2). The car-bide precipitation and other problems related tothe high chromium content can be resolvedthrough solution annealing and controlled weld-ing practices such as slower cooling rates andcontrolled-interpass temperatures (Ref 2).Nitrogen reduces the partitioning of chromiumbetween the austenite and the ferrite phases. Italso improves the pitting and crevice corrosionresistance of the duplex stainless steels (Ref 2).Very high nickel contents (e.g., 10 wt%) induplex stainless steel weld metal degrades pit-ting corrosion resistance by diluting the nitro-gen content in the austenite (Ref 2).

Inuence of Ferrite-Austenite Balance onCorrosion Resistance. The distribution ofaustenite and ferrite in the weld and HAZ isknown to affect the corrosion properties and themechanical properties of duplex stainless steels.Figure 7 shows the effect of the ferrite-austenitebalance on the pitting resistance of a duplexstainless steel. To achieve a satisfactory balancein properties, it is essential that both base metaland weld metal be of the proper composition.For example, without nickel enrichment in theller rod, welds can be produced with ferritelevels in excess of 80%. Such microstructureshave very poor ductility and inferior corrosionresistance. For this reason, autogenous welding(without the addition of ller metal) is not rec-ommended unless postweld solution annealingis performed, which is not always practical. Toachieve a balanced weld microstructure, a lowcarbon content (approximately 0.02%) and theaddition of nitrogen (0.1 to 0.2%) should bespecied for the base metal. Low carbon helpsto minimize the effects of sensitization, and thenitrogen slows the precipitation kinetics associ-ated with the segregation of chromium andmolybdenum during the welding operation. Ni-trogen also enhances the reformation of austen-ite in the HAZ and weld metal during cooling.

Because these duplex alloys have been usedfor many years, guidelines relating to austenite-ferrite phase distribution are available. It hasbeen shown that to ensure resistance to Cl SCC,


welds should contain at least 25% ferrite. Tomaintain a good phase balance for corrosionresistance and mechanical properties (espe-cially ductility and notch toughness) compara-ble to those of the base metal, the average ferritecontent of the weld should not exceed 60%.This means using welding techniques that min-imize weld dilution, especially in the root pass.Conditions that encourage mixing of the lower-nickel base metal with the weld metal reduce theoverall nickel content. Weld metal with a lowernickel content will have a higher ferrite content,with reduced mechanical and corrosion proper-ties. Once duplex base metal and welding con-sumables have been selected, it is then neces-sary to select joint designs and weld parametersthat will produce welding heat inputs and cool-ing rates so as to produce a favorable balance ofaustenite and ferrite in the weld and HAZ.

Researchers have shown that the high-ferritemicrostructures that develop during welding inlean (low-nickel) base metal and weld metalcompositions can be altered by adjusting weld-ing heat input and cooling rate. In these cases, ahigher heat input that produces a slower coolingrate can be used to advantage by allowing moretime for ferrite to transform to austenite. Thereare, however, some practical aspects to considerbefore applying higher heat inputs indiscrimi-nately. For example, as heat input is increased,base metal dilution increases. As the amount oflower-nickel base metal in the weld increases,the overall nickel content of the deposit de-creases. This increases the potential for moreferrite, with a resultant loss in impact toughness,ductility, and corrosion resistance. This wouldbe another case for using an enriched ller metalcontaining more nickel than the base metal.Grain growth and the formation of embrittlingphases are two other negative effects of highheat inputs. When there is uncertainty regardingthe effect that welding conditions will have oncorrosion performance and mechanical proper-ties, a corrosion test is advisable.

Effect of Welding on Pitting and SCCResistance. The weld is usually the part of asystem with reduced corrosion resistance andlow-temperature toughness, and therefore inmany cases it is the limiting factor in materialapplication. From a corrosion standpoint, weld-ing primarily affects pitting corrosion and ClSCC.

Pitting corrosion resistance can be affectedby many features of the welding operation,including:

Localized segregation of alloying elements tothe different constituent phases in the micro-structure, producing areas lean in molybde-num and chromium

Incorrect ferrite/austenite phase balance Formation of nitrides or intermetallic phases Loss of nitrogen from the root pass Presence of an oxidized surface on the under-

side of the root beadThe extent to which the reduction of corrosionresistance occurs depends on which of thesefactors are active and to what degree. Partition-ing of alloying elements between the austeniteand ferrite occurs in the weld metal, withchromium, molybdenum, and silicon partition-ing to the ferrite and carbon, nickel, and nitro-gen to the austenite (Ref 10, 11). The effect isnot so apparent in as-deposited weld metals, butit becomes more signicant as a result of reheat-ing a previously deposited weld pass.

Weld metal and HAZ microstructures withvery high ferrite contents are also less resistant topitting attack than are balanced structures. This islargely because predominantly ferritic structuresare more prone to chromium nitride precipita-tion, which locally denudes the chromium con-centration and lowers resistance to pitting attack.

Nitrogen loss in the root pass may reduce weldmetal corrosion resistance. Up to 20% loss ofnitrogen has been reported for GTA welds (Ref12), and nitrogen-bearing backing gases havebeen explored and used in limited applications.

Cleanliness of the root side purge gas mayalso affect pitting resistance. Figure 8 shows theeffect of reducing oxygen content in an other-wise pure argon purge gas and its benecialeffect on pitting resistance. Also shown is theapparent benet of using a reducing gas (NH10),which would signicantly reduce the tendency

Fig. 8 Plot of pitting temperature versus oxygen content ofbacking gas for Fe-22Cr-5.5Ni-3Mo-0.15N and Fe-

23Cr-4Ni-0.1N duplex stainless steels tested in 3% NaCl and0.1% NaCl solutions, respectively, both at anodic potential of+300 mV. Source: Ref 13


Fig. 9 Pitting corrosion resistance of base metal relative to weld metal placed in 6 wt % FeCl3 solution for 24 h duration per ASTM648 (method A). Source: Ref 14

for oxide formation and leave the underbeadappearance very shiny (Ref 13).

The net effect on pitting corrosion resistancemay be observed by applying the ASTM G 48pitting corrosion test to welds and base materialwith the same PREN value, then assessing thereduction in critical pitting temperature (that is,the temperature at which pitting in the ferricchloride solution is rst observed). The differ-ence is approximately 20 C (35 F), as reportedin Fig. 9, thereby quantifying the effect ofreduced weld metal properties. Figure 9 alsoshows that the use of a superduplex stainlesssteel ller metal with a PREN value of about 40on an Fe-22Cr-5.5Ni-3Mo-0.15N parent steel(which typically has a PREN value of about 33to 35) will improve the weld metal pitting cor-rosion resistance, as assessed by the ASTM G48 test, to approximately match that of the basematerial.

Resistance to Cl SCC does not appear to beaffected signicantly by welding per se (Ref15). Nevertheless, welds are likely regions ofattack for Cl SCC due to the presence of highstresses and the structural inhomogeneity pres-ent at the weld. If localized pitting is a necessaryprecursor for Cl SCC, the effects describedabove will also ultimately affect Cl SCC resis-tance. See also the discussion on SCC of alloy2205 in the section Corrosion Behavior ofAlloy 2205 Weldments.

Hydrogen-Induced Cracking. Duplexstainless steels can suffer from weld metalhydrogen cracking, but HAZ cracking has not

been reported in practice and is consideredhighly unlikely to develop. Hydrogen crackingfrom welding and in-service hydrogen pickuphas been observed (Ref 16). The duplexmicrostructure provides a combination of a fer-ritic matrix, where hydrogen diffusion can befairly rapid, with intergranular and intragranularaustenite, where the hydrogen diffusion is sig-nicantly slower, thereby acting as a barrier tohydrogen diffusion. The net effect appears to be that hydrogen can be trapped within fer-rite grains by the surrounding austenite, particu-larly where it decorates the prior ferrite grainboundaries. Due to these characteristics, low-temperature hydrogen release treatments are noteffective, and the hydrogen is likely to remain inthe structure for a long period (Ref 16). Whethercracking actually develops will depend upon anumber of factors, including the total amount oftrapped hydrogen, the applied strain, and theamount of ferrite and austenite in the structure.Weld metal hydrogen content from coveredelectrodes can be relatively high, and levels up to25 ppm have been reported (Ref 17).

The problem of weld metal hydrogen crack-ing in practice must not be overstated. Thereported incidences of hydrogen cracking induplex stainless steels have been restricted tocases in which the alloy has been heavily coldworked or weld metals have seen high levels ofrestraint or possessed very high ferrite contentsin combination with very high hydrogen levelsas a result of poor control of covered electrodesor the use of hydrogen-containing shielding gas.


Table 5 Chemical compositions of alloy 2205 specimens tested and ller metals used in Ref 19Element, %

Specimen size and conguration C Si Mn P S Cr Ni Mo Cu N

Parent metals

48.1 mm (1.89 in.) OD, 3.8 mm(0.149 in.) wall tube

0.015 0.37 1.54 0.024 0.003 21.84 5.63 2.95 0.09 0.15

88.9 mm (3.5 in.) OD, 3.6 mm(0.142 in.) wall tube

0.017 0.28 1.51 0.025 0.003 21.90 5.17 2.97 0.09 0.15

110 mm (4.3 in.) OD, 8 mm (0.31in.) wall tube

0.027 0.34 1.57 0.027 0.003 21.96 5.62 2.98 0.09 0.13

213 mm (8.4 in.) OD, 18 mm (0.7in.) wall tube

0.017 0.28 1.50 0.026 0.003 21.85 5.77 2.98 0.10 0.15

20 mm (3/4 in.) plate 0.019 0.39 1.80 0.032 0.003 22.62 5.81 2.84 . . . 0.13Filler metals1.2 mm (0.047 in.) diam wire 0.011 0.48 1.61 0.016 0.003 22.50 8.00 2.95 0.07 0.131.6 mm (0.063 in.) diam rod 0.011 0.48 1.61 0.016 0.003 22.50 8.00 2.95 0.07 0.133.2 mm (0.125 in.) diam wire 0.011 0.48 1.61 0.016 0.003 22.50 8.00 2.95 0.07 0.133.25 mm (0.127 in.) diam covered

electrode0.020 1.01 0.82 0.024 0.011 23.1 10.4 3.06 . . . 0.13

4.0 mm (0.16 in.) diam coveredelectrode

0.016 0.94 0.78 0.015 0.011 23.0 10.5 3.13 . . . 0.11

Indeed, other studies have shown how resistantduplex stainless steel weld metals are to hydro-gen cracking, even with consumables intention-ally humidied (Ref 18), and that hydrogen-containing backing gases can be employedwithout producing cracking. There is no doubtan effect of hydrogen on the ductility of duplexstainless steel, and to avoid fabrication-relatedcracking problems, high-hydrogen-potentialwelding processes, such as SMAW, should becontrolled by careful storage and use of elec-trodes, and by ensuring that the weld metalphase balance is within acceptable limits.

Corrosion Behavior of Alloy 2205Weldments

The inuence of different welding conditionson various material properties of alloy 2205(UNS S31803, Fe-22Cr-5.5Ni-3.0Mo-0.15N)has been studied (Ref 19). Chemical composi-tions of test materials are given in Table 5, andthe results of the investigation are detailed in thefollowing sections.

Intergranular Corrosion. Despite the useof very high arc energies (0.5 to 6 kJ/mm, or 13to 150 kJ/in.) in combination with multipasswelding, the Strauss test (ASTM A 262, prac-tice E) failed to uncover any signs of sensitiza-tion after bending through 180. The results ofHuey tests (ASTM A 262, practice C) on sub-

merged-arc welds showed that the corrosionrate increased slightly with arc energy in thestudied range of 0.5 to 6.0 kJ/mm (13 to 150kJ/in.). For comparison, the corrosion rate forparent metal typically varies between 0.15 and1.0 mm/yr (6 and 40 mils/yr), depending on sur-face nish and heat treatment cycle.

Similar results were obtained in Huey tests ofspecimens from bead-on-tube welds producedby GTAW welding. In this case, the corrosionrate had a tendency to increase slightly with arcenergy up to 3 kJ/mm (75 kJ/in.).

Pitting tests were conducted in 10% ferricchloride (Fe Cl3) at 25 and 30 C (75 and 85 F) inaccordance with ASTM G 48. Results of tests onsubmerged-arc test welds did not indicate anysignicant change in pitting resistance when thearc energy was increased from 1.5 to 6 kJ/mm (38to 150 kJ/in.). Pitting occurred along the bound-ary between two adjacent weld beads. Attackwas caused by slag entrapment in the weld; there-fore, removal of slag is important.

Gas tungsten arc weld test specimens (arcenergies from 0.5 to 3 kJ/mm, or 13 to 75 kJ/in.)showed a marked improvement in pitting resis-tance with increasing arc energy. In order forduplicate specimens to pass the FeCl3 test at 30C (85 F), 3 kJ/mm (75 kJ/in.) of arc energywas required. At 25 C (75 F), at least 2 kJ/mm(50 kJ/in.) was required to achieve immunity.Welds made autogenously (no nickel enrich-ment) were somewhat inferior; improvementswere achieved by using higher arc energies.


Fig. 12 Preferential attack of the continuous austenite phasein an autogenous gas tungsten arc weld in Ferralium

alloy 255. Crevice corrosion test was performed in synthetic sea-water according to ASTM D 1141 at 100 C (212 F). Etched with50% HNO3. 100

Fig. 11 Preferential corrosion of the ferrite phase in the weldmetal of Ferralium alloy 255 gas tungsten arc welds

in 10% FeCl3 at room temperature. Base metal was 3.2 mm (1/8

in.) thick.

For comparison with a different alloy, Fig. 10shows the effect of heat input on the corrosionresistance of Ferralium alloy 255 (UNS S32550,Fe-25.5Cr-5.5Ni-3.0Mo-0.17N) welds madeautogenously and tested on FeCl3 at 15 C (60F). Preferential corrosion of the ferrite phase isshown in Fig. 11. In a different test, Ferralium

alloy 255 was welded autogenously and testedin a neutral chloride solution according toASTM D 1141 at 60 to 100 C (140 to 212 F).In this case, preferential attack of the austenitephase was observed. An example is shown inFig. 12. Similar results would be expected foralloy 2205.

A study of the alloy 2205 weld microstructures(Ref 19) revealed why high arc energies werefound to be benecial to pitting resistance. Manyinvestigations have indicated that the presence ofchromium nitrides in the ferrite phase lowers theresistance to pitting of the weld metal and theHAZ in duplex stainless steels. In this study, bothweld metal and HAZ produced by low arc ener-gies contained an appreciable amount ofchromium nitride (Cr2N) (Fig. 13). The nitrideprecipitation vanished when an arc energy of 3 kJ/mm (75 kJ/in.) was used (Fig. 14).

The results of FeCl3 tests on submerged-arcwelds showed that all top weld surfaces passedthe test at 30 C (85 F) without pitting attack,irrespective of arc energy in the range of 2 to 6kJ/mm (50 to 150 kJ/in.). Surprisingly, the weldmetal on the root side, which was the rst to bedeposited, did not pass the same test temperature.

The deteriorating effect of high arc energieson the pitting resistance of the weld metal on theroot side was unexpected. Potentiostatic testscarried out in 3% sodium chloride (NaCl) at 400mV versus saturated calomel electrode (SCE)conrmed these ndings. Microexamination ofthe entire joint disclosed the presence ofextremely ne austenite precipitates, particu-larly in the second weld bead (Fig. 15) but also

Fig. 10 Effect of welding heat input on the corrosion resis-tance of autogenous gas tungsten arc welds in Fer-

ralium alloy 255 in 10% FeCl3 at 10 C (40 F). The base metalwas 25 mm (1 in.) thick. Source: Ref 20


Fig. 14 Microstructure of bead-on-tube weld made by auto-genous GTAW with an arc energy of 3 kJ/mm (76

kJ/in.). Virtually no chromium nitrides are present, which resultsin adequate pitting resistance. 200. Source: Ref 19

Fig. 15 Microstructure of the second weld bead of a sub-merged-arc weld joint in 200 mm (3/4 in.) duplex

stainless steel plate. The extremely ne austenite precipitate wasformed as a result of reheating from the subsequent weld pass,which used an arc energy of 6 kJ/mm (150 kJ/in.). 1000. Source:Ref 19

Fig. 13 Microstructure of bead-on-tube weld made by auto-genous GTAW with an arc energy of 0.5 kJ/mm (13

kJ/in.). Note the abundance of chromium nitrides in the ferritephase. See also Fig. 14. 200. Source: Ref 19

in the rst or root side bead. The higher the arcenergy, the more austenite of this kind was pres-ent in the rst two weld beads. Thus, nitridesgive rise to negative effects on the pitting resis-tance, as do ne austenite precipitates that werepresumably reformed at as low a temperature asapproximately 800 C (1470 F).

Therefore, the resistance of alloy 2205 to pit-ting corrosion is dependent on several factors.First, Cr2N precipitation in the coarse ferritegrains upon rapid cooling from temperaturesabove approximately 1200 C (2190 F) causesthe most severe impairment to pitting resis-tance. This statement is supported by a greatnumber of FeCl3 tests as well as by potentiosta-tic pitting tests. Generally, it seems difcult toavoid Cr2N precipitation in welded joints com-pletely, particularly in the HAZ, the structure ofwhich can be controlled only by the weld ther-mal cycle. From this point of view, it appearsadvisable to employ as high an arc energy aspractical in each weld pass. In this way, thecooling rate will be slower (but not slow enough


Table 6 Corrosion resistance of Ferralium alloy 255 weldments using various nickel-base alloyllers and weld techniques3.2 mm (0.125 in.) plates tested in 10% FeCl3 for 120 h

Critical pitting temperature

Gas tungsten arc

Filler metal C F C F C F

Hastelloy alloy G-3 3035 8595(a) 30 85(a) 3035 8595(b)IN-112 30 85(a) . . . . . . 3540 95105(b)Hastelloy alloy C-276 . . . . . . . . . . . . 2530 7585(a)Hastelloy alloy C-22 30 85(a) . . . . . . 35-40 95105(a)(a) Haz. (b) HAZ plus weld metal

Submerged arcGas metal arc

to encounter 475 C, or 885 F, embrittlement),and the reformation of austenite will clearlydominate over the precipitation of Cr2N.

In addition, if there were no restriction onmaximum interpass temperature, the heat pro-duced by previous weld passes could be used todecrease the cooling rate further in the criticaltemperature range above approximately 1000C (1830 F). Preliminary tests with preheatedworkpieces have shown the signicance of tem-perature in suppressing Cr2N precipitation. Cur-rently, the maximum recommended interpasstemperature for alloy 2205 is 150 C (300 F).This temperature limit does not appear to becritical, and it is suggested that this limit couldbe increased to 300 C (570 F). The maximumrecommended interpass temperature for Ferral-ium alloy 255 is 200 C (390 F). Excessivegrain growth as a result of too much heat inputmust also be considered to avoid loss of ductil-ity and impact toughness.

Second, the ne austenite precipitates foundin the reheated ferrite when high arc energiesand multipass welding were combined are com-monly referred to as 2 the literature. The harm-ful inuence of 2 on the pitting resistance hasbeen noted with isothermally aged specimens,but as far as is known, it has never beenobserved in connection with welding. It is felthowever, that 2 is less detrimental to pittingthan Cr2N. Moreover, 2 formation is believedto be benecial to mechanical properties, suchas impact strength and ductility.

A third factor that lowers pitting resistance isoxide scale. Where possible, all surface oxidesshould be removed by mechanical means or,preferably, by pickling. Root surfaces (in pipe),however, are generally inaccessible, and pittingresistance must rely on the protection from thebacking gas during GTAW. It is therefore advis-

able to follow the current recommendation forstainless steels, which is a maximum of 25 ppmoxygen in the root backing gas.

Stress-Corrosion Cracking. The SCC resis-tance of alloy 2205 in aerated, concentratedchloride solutions is very good. The effect ofwelding on the SCC resistance is negligiblefrom a practical point of view. The thresholdstress for various welds, as well as for unweldedparent metal in the calcium chloride (CaCl2)test, is as high as 90% of the tensile strength atthe testing temperature. This is far above allconceivable design limits.

Also, in environments containing both hydro-gen sulde (H2S) and chlorides, the resistanceof welds is almost as high as for the parentmetal. In this type of environment, however, itis important to avoid too high a ferrite content inweld metal and the HAZ. For normal welding ofjoints, the resulting ferrite- contents should notcause any problems. For weld repair situations,however, care should be taken so that extremelyhigh ferrite contents (>75%) are avoided. Topreserve the high degree of resistance to SCC,the ferrite content should not be less than 25%.

Another reason to avoid coarse weld micro-structures (generated by excessive weldingheat) is the resultant nonuniform plastic ow,which can locally increase stresses and inducepreferential corrosion and cracking effects.

Use of High-Alloy Filler Metals. In criticalpitting or crevice corrosion applications, the pit-ting resistance of the weld metal can be en-hanced by the use of high Ni-Cr-Mo alloy llermetals. The corrosion resistance of such weld-ments in Ferralium alloy 255 is shown in Table6. For the same weld technique, it can be seenthat using high-alloy llers does improve corro-sion resistance. If high-alloy llers are used, theweld metal will have better corrosion resistance


than the HAZ and the fusion line. Therefore,again, proper selection of welding techniquecan improve the corrosion resistance of theweldments.

ACKNOWLEDGMENTS

Portions of this chapter were adapted from:

D.N. Noble, Selection of Wrought DuplexStainless Steels, Welding, Brazing, and Sol-dering, Vol 6, ASM Handbook, ASM Interna-tional, 1993, p 471481

K. F. Krysiak, et al., Corrosion of Weld-ments, Corrosion, Vol 13, Metals Handbook,9th ed., ASM International, 1987, p 344368

REFERENCES

1. Wrought Duplex Stainless Steels, AlloyDigest Source Book: Stainless Steels, J.R.Davis, Ed., ASM International, 2000, p 317356

2. D.K. Aldun, K.M. Makhamreh, and D.J.Morrison, Corrosion Resistance of Second-Generation Duplex StainlessSteel Weldments in NACE Solution,International Trends in Welding Scienceand Technology, Proc. 3rd Int. Conf.Trends in Welding Research, S.A. Davidand J.M. Vitek, Ed., ASM International,1992, p 661666

3. T.G. Gooch, Corrosion Resistance ofWelds in Duplex Stainless Steels, DuplexStainless Steels 91, Vol 1, J. Charles andS. Bernhardsson, Ed., Les ditions dePhysique, Oct 1991, p 325346

4. J. Charles, Super Duplex Stainless Steels:Structure and Properties, Duplex Stain-less Steels 91, Vol 1, J. Charles and S.Bernhardsson, Ed., Les ditions dePhysique, Oct 1991, p 151168

5. J. Harston, E. Hutchins, and S. Sweeney,The Development and Construction ofDuplex Stainless Steel Pipelines for UseOffshore in the Southern North Sea, Proc.3rd Int. Conf. Welding and Performanceof Pipelines, The Welding Institute, 1986

6. D.J. Kotecki, Heat Treatment of DuplexStainless Steel Weld Metals, Weld. J.,Nov 1989, p 431s441s

7. L. Odegarg and S.-A. Fager, Choice ofRight Filler Metal for Joining SandvikSAF 2507 to a Super Austenitic Stainless

6 Mo Steel, ASM Materials Week 87,ASM International, Oct 1987, p 441449

8. R.N. Gunn, The Inuence of Compositionand Microstructure on the CorrosionBehaviour of Commercial Duplex Alloys,Recent Developments in the Joining ofStainless Steels and High Alloys, EdisonWelding Institute, Columbus, Ohio, Oct1992

9. How to Weld 2205, Avesta WeldingTrade Literature

10. T. Ogawa and T. Koseki, Effect of Com-position Proles on Metallurgy and Cor-rosion Behaviour of Duplex StainlessSteel Weld Metals, Weld. J., Vol 88 (No.5), May 1989, p 181s191s

11. R.A. Walker and T.G. Gooch, PittingResistance of Weld Metal for 22Cr-5NiFerritic-Austenitic Stainless Steel, Br.Corros. J., Vol 26 (No. 1), 1991, p 5159

12. T.G. Gooch, Weldability of Duplex Fer-ritic-Austenitic Stainless Steels, DuplexStainless Steels, R.A. Lula, Ed., AmericanSociety for Metals, 1983, p 573602

13. L. Odegard and S.-A. Fager, The RootSide Pitting Corrosion Resistance ofStainless Steel Welds, Sandvik SteelWeld. Rep., No. 1, 1990

14. R. Dolling, V. Neubert, and P. Knoll, TheCorrosion Behaviour of Super DuplexSteel Cast Alloys with a PREN>41,Duplex Stainless Steels Conf. Proc., Vol2, Les ditions de Physique, Les UlisCedex, France, Oct 1991, p 13411351

15. T.G. Gooch, Welding in the World, Vol24 (No. 7/8), 1986, p 148167

16. P. Sentance, The Brae Field, Stainl. SteelEur., Vol 4 (No. 20), Sept 1992, p 3841

17. C.D. Lundin, K. Kikuchi, and K. K. Kahn,Measurement of Diffusible HydrogenContent and Hydrogen Effects on theCracking Potential of Duplex StainlessSteel Weldments, Phase 1 Report, TheWelding Research Council, March 1991

18. R.A. Walker and T.G. Gooch, HydrogenCracking of Welds in Duplex StainlessSteel, Corrosion, Vol 47 (No. 8), Sept1991, p 10531063

19. B. Lundquist, P. Norberg, and K. Olsson,Inuence of Different Welding Condi-tions on Mechanical Properties and Cor-rosion Resistance of Sandvik SAF 2205(UNS S31803), Paper 10, presented atthe Duplex Stainless Steels 86 Confer-ence, the Hague, Netherlands, Oct 1986


20. N. Sridhar, L.H. Flasche, and J. Kolts,Effect of Welding Parameters on Local-ized Corrosion of a Duplex Stainless Steel,Mater. Perform., Dec 1984, p 52 55

SELECTED REFERENCES

R. Bradshaw, R.A. Cottis, and M.J. Scho-eld, Stress Corrosion Cracking of DuplexStainless Steel Weldments in Sour Condi-tions, Mater. Perform., April 1996, p 6570

R.M. Davison and J.D. Redmond, PracticalGuide to Using Duplex Stainless Steels,Mater. Perform., Jan 1990, p 5762

B. Gissler and B. Bouldin, Corrosion Perfor-mance of Welded Super Duplex StainlessSteel with Multiple Re-weld/Re-work WeldRepairs in Aqueous H2S and FeCl3 Environ-ments, Paper No. 04143, presented at Corro-sion 2004, NACE International

H. Inoue, T. Ogawa, and R. Matsuhashi, Cor-rosion Behavior of Duplex Stainless SteelWelds in Highly Concentrated Sulfuric Acid,International Trends in Welding Science andTechnology, Proc. 3rd Int. Conf. Trends inWelding Research, S.A. David and J.M.Vitek, Ed., ASM International, 1992, p667671

134567205 103677583 Corrosion of DSS Weldments

Documents

Transcript of 134567205 103677583 Corrosion of DSS Weldments