fabrication and characterization of graded transition joints...

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Introduction

Dissimilar metal welds (DMWs) be-tween ferritic low-alloy steels andaustenitic alloys are commonly used inpower-generation plants. The less-expen-sive low-alloy steels are used in the low-temperature regions of the plant, whilethe higher temperatures in the super-heater regions require the superior corro-sion resistance and greater creep strengthof more-expensive austenitic alloys. In atypical power plant, there can be thou-sands of DMWs, which are frequently cy-cled between room temperature and650°C (Ref. 1). The DMWs are prone topremature failure due to differences inchemistry, thermal expansion, and creepstrength between the two alloys (Refs. 1,2). Premature failure of these DMWs can

result in forced plant outages that can costa power company $250,000–$850,000 perday in lost revenue (Ref. 3). The failuresoccur from simultaneous metallurgicaland mechanical factors, including devel-opment of strength gradients during bothwelding and service, formation of a weakcarbon-depleted zone due to carbon mi-gration, and concentration of stress fromthermal cycling that causes acceleratedcreep failure.

The microstructure of DMWs in theas-welded condition consists of a sharpconcentration gradient across the weld in-terface that separates the ferritic and

austenitic alloys. During fusion welding,the combination of the high alloy contentof the austenitic filler metal and fast cool-ing rate produces a hard martensite bandin the partially mixed zone (PMZ) (Refs.2, 4, 5). High temperatures encounteredduring either postweld heat treatment(PWHT) or service provide the activationenergy for carbon diffusion to occur downthe chemical potential gradient from theferritic steel toward the austenitic alloy(Refs. 1, 6–11). This can lead to formationof a soft carbon denuded zone near the in-terface on the ferritic steel, and nucleationand growth of carbides on the austeniticside that have very high hardness. Theselarge differences in microstructure andhardness occur over very short distancesacross the weld interface (~ 50–100 μm)(Refs. 2, 6, 12). A band of carbides alsoforms along the weld interface in the fer-ritic side of the joint (Ref. 8). The differ-ence in hardness across the weld interfaceincreases with increasing aging time due tonucleation and growth of the interfacialcarbides (Ref. 6). At the same time, strainis localized along the carbon denudedzone due to differences in creep strength(Refs. 13–15) while localized stresses de-velop due to differences in thermal expan-sion coefficients of the steels (Refs. 16,17). As a consequence of the hardness andstrength gradients, these stresses are con-centrated in the weak carbon-depletedzone, generating creep voids around car-bides that lead to eventual creep rupture(Refs. 16, 18).

There are two distinct morphologies ofthese carbides that can evolve in DMWsduring aging (Refs. 1, 18–22). Type I car-bides are the ones most frequently ob-served. These carbides form very close tothe weld interface (~ 1 μm) in the HAZ ofthe ferritic steel. These carbides initiallyform with a spherical shape, but graduallyacquire a lenticular morphology and caneventually form regions of continuous orsemicontinuous carbides as they grow andcoalesce. The Type II carbides generallyform as a wide band and are associated

Fabrication and Characterizationof Graded Transition Joints for

Welding Dissimilar Alloys

A GTAW system employing dual wire feeders was used to fabricate graded jointsbetween 2.25Cr-1Mo steel and three austenitic alloys, and all cases showed a

relatively smooth transition in composition

BY G. J. BRENTRUP and J. N. DUPONT

ABSTRACT

Functionally graded materials have potential for joining dissimilar materials inmany applications. In this work, graded transition joints were fabricated for joiningferritic and austenitic alloys. A gas tungsten arc welding system employing dual wirefeeders was used to fabricate graded joints between 2.25Cr-1Mo steel and threeaustenitic alloys, including IN800, IN82, and 347 stainless steel. All three joint com-binations were characterized to determine compositional variations, microhardnessprofiles, and microstructural evolution. Tensile tests were also performed to evaluatethe high-temperature mechanical properties. In all cases, a relatively smooth transi-tion in composition was achieved over ~ 50 mm, which represents an increase inlength of approximately three orders of magnitude compared to composition gradi-ents observed in traditional dissimilar metal welds. Measured composition data fromthe transition joints were converted to Ni and Cr equivalents and used with theSchaeffler diagram to predict the phase distribution along the joints, and good agree-ment was obtained between the observed and predicted results. Peaks in hardnesswere observed along the joints and attributed to the formation of martensite. The ten-sile properties from 20° to 650°C were generally within the range of those expectedfor the end member alloys except for the T22-IN82 joint, which exhibited slightlylower yield and tensile strengths. Solidification cracking was observed in the T22-800joints in the region that exhibited a fully austenitic solidification mode.

KEYWORDS

Cr-Mo SteelsStainless SteelsFractureGraded Transition Joints

G. J. BRENTRUP and J. N. DUPONT([email protected]) are with the Dept. of Materi-als Science & Engineering, Lehigh University,Bethlehem, Pa.

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with the martensite region that formswithin the PMZ. For DMWs preparedwith nickel-based filler metals, the Type Icarbides that form along the interface pro-vide sites for nucleation and growth ofcreep cavities that eventually lead to pre-mature cracking. For older DMWs madewith stainless steel filler metals, crackingtypically occurs along the prior austenitegrain boundaries (PAGBs) in the ferriticHAZ at a location of about one or twograins away from the weld interface.

Several different approaches have beentaken to extend the service life of DMWs.Nickel-based filler metals reduce the coef-ficient of thermal expansion (CTE) mis-match between the austenitic and ferriticsteels. Due to the low solubility and low dif-fusivity of carbon in Ni-based alloys, carbonmigration is reduced over the service life ofthe weld (Refs. 6, 9, 19, 23, 24). The Ni-based fillers can improve the weld lifetimeby a factor of five compared to austeniticfiller metals (Ref. 23).

Another approach is to separate thecarbon and stainless steels with a materialof intermediate thermal expansion,known as a trimetallic joint. Severalnickel-based alloys, notably Inco Alloy800/800H, have a CTE intermediate be-tween the two steels and have been usedas suitable candidates. Studies ontrimetallic joints (Refs. 19, 25, 26) haveshown a fourfold increase in the lifetimeand a 38% reduction in stress compared totraditional DMWs. However, failure hasstill been observed, indicating that thesteep CTE, microstructure, and propertygradients are still problems.

Building on the idea of the trimetallicjoint, graded transition joints could be de-

veloped to replaceDMWs with an inter-mediate section inwhich the compositionvaries continuouslyalong its length. Thegraded joint smoothlytransitions from theferritic steel composi-tion to the stainlesssteel composition, allowing two similarwelds to replace the one dissimilar weld.By continuously grading the composition,the sharp changes in both compositionand properties of traditional DMWswould be extended over the whole lengthof the component, potentially eliminatingmany of the factors that promote failure.

Functionally graded steels have beenproduced by a number of different re-search groups. Mohandesi et al. (Refs. 27,28) have used transport processes, specif-ically diffusion of alloying elements, toproduce functionally graded steels withlayers of ferrite, austenite, bainite, andmartensite. Coco et al. (Ref. 29) createdgraded Fe-C and Fe-Mn-C steels by par-tial decarburization. While these tech-niques appear to work well for low-alloyand carbon steels, compositional controlof high-alloy steels and other alloys re-quires other methods. A preliminary studyby Farren et al. (Ref. 30) demonstratedthat a functionally graded material(FGM) that transitioned from carbon tostainless steel could be fabricated by theLaser Engineered Net Shaping (LENS)process. The FGM exhibited a gradualchange in microstructure and hardness,indicating the sharp gradients in DMWscan be eliminated.

This paper describes the fabrication,characterization, and testing of gradedjoints with a smooth transition in composi-tion, microstructure, and properties thatcan be used for joining ferritic alloys toaustenitic alloys. The transition joints werefabricated by a dual-wire gas tungsten arcwelding process and characterized by opti-cal and electron microscopy, hardness test-ing, and high-temperature tensile testing.Future applications of the graded joints arejoining ferritic and austenitic alloys in thepower generation industry.

Procedure

Three types of graded transition jointswere fabricated: T22 steel to IN82, T22steel to IN800, and T22 steel to 347 stain-less steel. These combinations were se-lected based on previous experience andtheir use in fossil-fired power plants. MostDMWs are currently welded with IN82filler metal, and IN800 is often used as anintermediate section between ferritic andaustenitic steels to reduce the CTE gradi-ent. The T22 - 347 grade was explored tounderstand the behavior of a direct transi-tion between a ferritic and austenitic steel.

Although the LENS system is usefulfor making graded transition joints (Ref.

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Fig. 1 — Picture of a T22 plate that had 50 layers of T22-Alloy 800 gradedmaterial added through the GTAW process to create graded joints, which weresubsequently machined into the tensile bar geometry shown on the right.

Fig. 2 — EDS composition (top) and hardness traces of traditional DMWdemonstrating the sharp gradients in composition and properties (Ref. 4).


30), several significant drawbacks of theLENS process have been recognized dur-ing the course of this work. First, the sys-tem is only capable of producing alloysfrom powders. Unfortunately, many alloysare not readily available in powder form,particularly for the size ranges required bythe LENS powder feeders. Thus, fabrica-tion of custom powders is typically re-quired, which can be quite expensive. Sec-ond, previous research with the LENSprocess (Ref. 31) has shown that onlyabout 5–14% of the powder is actually de-posited, with the remaining being wasted.As a result, it is difficult to control the heatflow and deposit composition with pow-ders because the precise amount of pow-der that enters into the liquid pool is un-

known. Last, thedeposition rate ofthe LENS processis extremely slow,making fabrica-tion of large partsimpractical. Thus,a gas tungsten arcwelding (GTAW)process using dualwire feeders wasdeveloped. In thissystem, wiresrather than pow-ders are used, soall of the materialis fed directly intothe melt pool andthe compositioncan be more pre-cisely controlled.In addition, thewires are morereadily availableand generally lessexpensive. Theuse of wire feeders

also allows for much higher depositionrates of ~3 kg/h compared to the LENSsystem with a deposition rate of only ~0.2kg/h.

All wire compositions are listed inTable 1. The diameter of the wires was0.035 in. (0.9 mm). A 6.35 × 76.2 × 610mm (0.5 × 3 × 24 in.) plate of T22(2.25Cr-1Mo) steel (composition pro-vided in Table 1) was clamped betweentwo water-cooled copper plates. Materialwas deposited onto the edge of the T22plate, and after each layer the copperplates were raised to ensure repeatablewelds were made for each layer. The totalfeed rate of the two wires was maintainedat a constant 1.27 m/min (50 in./min), and

individual wire speeds were varied witheach new layer to adjust the compositionand manufacture the grade. The first layerwas deposited with the T22 and austeniticwire feed rates at 1.25 and 0.025 m/min (49and 1 in./min), respectively. The secondlayer was deposited with the T22 fillermetal at 1.22 m/min (48 in./min) andaustenitic filler metal at 0.05 m/min (2in./min). These changes continued until 50layers were deposited. Once the grade wascomplete, two inches of additional “pure”austenitic material was deposited so thattensile bars could be machined with theentire grade contained in the gauge lengthof the test sample. A wire brush andethanol were used to clean each layer aftercooling. The welding parameters were atravel speed of 1 mm/s, a total wire feedrate of 1.27 m/min (50 in./min), a currentof 250 A, a voltage of 13 ± 1 V, and a rootopening of 6.35 mm (0.25 in.).

All graded joints were examined by ra-diographic X-ray analysis to detect anycracks or other defects. As shown in Fig. 1,tensile bars were machined from theplates so that the entire grade was con-tained in the reduced section. Tensile testswere performed at 20°, 250°, 350°, 450°,550°, and 650°C in air using an exten-someter at a strain rate of 0.127mm/mm/min (0.005 in./in./min) throughyield and then 1.27 to 2.54 mm/mm/min(0.05 to 0.10 in./in./min) to fracture. Twosamples were tested at each temperature.

Additional samples were cut from thegraded transition joints for metallo-graphic analysis. The samples weremounted in cold-setting epoxy, groundusing SiC papers, and polished using 6-and 1-μm diamond and 0.05-μm colloidalsilica. The T22 half of the grade wasetched using 2% Nital for 5–7 s. The alloyIN800 and IN82 halves were immersion

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Fig. 3 — Plot of EDS and microhardness data as a function of distance for anas-welded T22-347 graded joint. Vertical lines correspond to phase boundariesfrom the Schaeffler diagram (A – austenite, F – ferrite, M – martensite).

Fig. 4 — EDS composition data for each of the three graded joint combi-nations plotted as nickel and chromium equivalents on the Schaeffler dia-gram, which was used to determine the positions of the phase boundariesdrawn as dark vertical lines in Fig. 3.

Fig. 5 — Plot of EDS and microhardness data for as-welded T22-IN82 gradedjoint. Vertical lines correspond to phase boundaries from the Schaeffler diagram(A – austenite, M – martensite). Lettered arrows correspond to locations of pho-tomicrographs shown in Fig. 6.


etched with Marble’s reagent for 5–10 s,while the 347 half was electrolyticallyetched using Lucas’s reagent at 3 V for 15s. Photomicrographs were recorded usingan Olympus BH-2 optical microscope andPax-It software. Microhardness measure-ments were performed on the gradedjoints using a Leco M400-FT hardnesstester with a 10-g load, 15-s dwell time, and50-μm spacing between indents.

Energy-dispersive X-ray spectroscopy(EDS) composition analysis was con-ducted using a Hitachi 4300 FE scanningelectron microscope (SEM) with anEDAX X-ray detector. An acceleratingvoltage of 20 kV was used for line scanmeasurements with 100-μm spacing and

20-s dwell time. A set of standards of eachmaterial (T22, 347, 800, IN82) whose com-position had been confirmed by wet chem-ical analysis was used for calibration of theEDS measurements.

Results and Discussion

As described previously, the sharp gra-dient in the composition, microstructure,and properties across the weld interface ofDMWs between ferritic and austenitic al-loys is the primary cause of premature fail-ure. An example of the sharp concentra-tion gradients in a conventional DMW isshown in Fig. 2 as a basis for comparison,where composition data are plotted as a

function of distance across the weld (Ref.4). The primary alloying elements, nickeland chromium, both decrease sharplyfrom 12 and 17 wt-% in the stainless steelto <1 wt-% in the ferritic steel. Thesecomposition changes occur over a distanceof only ~80 μm. The abrupt compositionalchanges lead to steep property gradientsas well, as seen in the microhardness tra-verse in Fig. 2. The hardness increasesfrom 200 Knoop in the ferritic steel to amaximum of 600 Knoop in the partiallymixed zone before decreasing back to 250Knoop in the stainless steel. The peak inhardness is associated with a band ofmartensite that forms along the weld interface.

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Table 1 — Certified Compositions (wt-%) of the Alloy Wires Used to Construct the Graded Joints

Material Al C Cr Cu Mn Mo Nb Ni Si Ti Fe

2.25Cr-1Mo 0.01 0.07 2.39 0.14 0.59 0.94 0.00 0.05 0.47 0.00 95.33Alloy 800H 0.42 0.09 20.97 0.27 1.03 0.16 0.15 34.65 0.44 0.56 41.12

347 0.00 0.05 19.43 0.13 1.68 0.20 0.63 9.13 0.46 0.00 68.23Inconel 82 0.00 0.10 20.00 0.50 3.00 0.00 2.50 67.00 0.50 0.75 3.00

A

CD

B

Fig. 6 — Representative photomicrographs from the T22-IN82 graded joint. Letters correspond to arrows in Fig. 5: A — Tempered martensite; B — transitionbetween tempered martensite and lath martensite; C — transition between martensite and austenite; D — austenite.


The goal of graded transition joints isto eliminate these sharp composition andproperty gradients. Figure 3 shows micro-hardness and composition data for a T22-347 graded joint as a function of distance.Note the composition changes graduallyover a distance of ~50 mm, compared to~50–100 μm that is typically observed forconventional DMWs. The microhardnessdata also show that the sharp changes inhardness have been elongated over amuch greater distance compared to tradi-tional DMWs.

The expected distribution of phasesalong the joint can be understood by con-verting the composition data to chromiumand nickel equivalents for use on theSchaeffler diagram (Ref. 32). For this, theconcentration of major alloying elementswas measured directly via EDS. Concen-trations of carbon and minor alloying ele-ments were determined with the approachpreviously described by DuPont andKusko (Ref. 4). The dilution within eachlocation was first calculated with themeasured composition data of the majoralloying elements. The concentration ofcarbon and minor alloying elements wasthen estimated by backcalculation fromknowledge of the known dilution and fillermetal compositions. Figure 4 shows the re-sults plotted on the Schaeffler diagram(composition data for the other grades arealso provided on this plot). Combinationof the results shown in Figs. 3 and 4 per-mit estimation of the phase distributionwithin the joints, and these locations are

shown on the compositional plot in Fig. 3.Note that the increase in hardness that be-gins at ~9 mm corresponds reasonablywell to the point in the grade wheremartensite is expected to form. The hard-ness then gradually decreases across themartensite region from ~400–450 HV at20 mm to ~200 HV at 55 mm, which is thenominal hardness for T22.

Similar trends in the composition andmicrohardness data were observed for theT22-IN82 and T22-800 graded joints asshown in Figs. 5 and 7, respectively. Thesegraded joints also have a gradual changein composition over ~50 mm. Light opti-cal photomicrographs are provided in Fig.6 for the T22-IN82 joint as an example ofthe microstructures observed in thegraded joints. Each photomicrograph cor-responds to one of the phase fields pre-dicted by the Schaeffler diagram. The mi-crostructure shown for the 50-mmlocation (Fig. 6A) consists of temperedmartensite, as expected for the T22 end ofthe joint. The photomicrograph at 43 mm(Fig. 6B) shows the transition between thetempered martensite of the T22 basemetal and the as-quenched lath marten-site that forms where the composition be-gins to change. Figure 6C, at 38 mm, showsthe transition between the austenite +martensite to the single-phase austeniteregion. Finally, the micrograph from 8 mm(Fig. 6D) corresponds to the austenitic mi-crostructure observed for IN82. Similarmicrostructures were observed for theT22-347 and T22-800 joints.

Hardness peaks correspond to regionsof the joints where the martensite phase isstable. The width of the martensite-containing regions varies for all threejoints, from ~35 mm for the T22-347 toonly ~10 mm for the T22-IN82. As de-scribed in previous work (Ref. 4), this isdue to the variations in the width of thecompositional gradients for each transi-tion joint. The steepest concentration gra-dient occurs for the T22-IN82 joint, wherethe nickel and chromium contents de-crease from 67 and 20 wt-%, respectively,in the IN82 down to <1 and 2 wt-% in theT22. This steep concentration gradient re-sults in a relatively narrow martensiteband, as seen in Fig. 5. Conversely, theT22-347 joint has a smaller gradient incomposition and a wider martensite band— Fig. 3. The variation in the width of themartensite band is due to the variation inmartensite start (Ms) temperature withcomposition (Ref. 4). It is known that al-loying elements such as Ni, Cr, Mn, andMo reduce the Ms temperature (Ref. 34).Thus, for the IN82, the large compositiongradient pushes the Ms temperaturebelow room temperature at a relativelyshort distance from the T22/graded jointinterface. This stabilizes the austenite at aconcomitantly short distance from the in-terface, producing a relatively thinmartensite layer. In contrast, the relativelysmall concentration gradient of theT22/347 grade causes the Ms temperatureto drop below room temperature at alarger distance from the interface, thusproducing a larger martensite layer.

The hardness gradients observed in thetransition joints occur over distances of~5–10 mm and are considerably less se-vere than those observed in conventionalDMWs in which the same hardness gradi-ents occur over ~2–40 μm (Refs. 4, 6). Theeffect of the local hardness changes needsto be determined with long-term creeptesting. Any possible detrimental effectmay be minimized by more gradualchanges in composition across the gradedjoint in the regions where martensite isknown to form. This will be evaluated infuture work.

Figures 8–10 show the tensile proper-ties as a function of temperature for eachof the graded joints. Also shown on theplots are minimum values (shown as solidlines) for the T22, 347, 800, and IN82 ma-terials (elongation data are not availablefor the IN82 as deposited filler metal). Aschematic tensile bar shows the failure lo-cations as dotted lines for each test tem-perature, and the phase boundaries basedon the variation in composition from theSchaeffler diagram are also indicated assolid vertical lines.

The ductility and yield strength resultsfor the T22-347 and T22-IN800 grades fallwithin the range for the end member al-

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Fig. 7 — Plot of EDS and microhardness data as a function of distance for as-welded T22-800 gradedjoint. Vertical lines correspond to phase boundaries from the Schaeffler diagram (A – austenite, M –martensite).


loys, while the tensile strength is close tothat of the lower strength material at eachtemperature. In general, yielding followedby eventual fracture is expected to occur inthe weakest region of the joint. However,there are other factors to consider wheninterpreting tensile data from graded ma-terials. For example, the overall yieldstrength can be increased due to con-straint of deformation in a soft regionfrom a hard neighboring region that re-stricts plastic flow. This could account forthe yield strength of the graded joint beingabove that of the weaker alloy for theT22/347 and T22/800 transition joints. Inaddition, work hardening of the soft re-gion can produce a shift of the fracturefrom the region of initial deformation toanother region.

The data in Fig. 9 for the T22-IN82joint demonstrate that this alloy combina-tion exhibits yield and tensile strengthsslightly below the nominal alloy proper-ties. All of the failures occurred close tothe T22 side of the joint, which may be at-tributed to the reduced tensile strength ofT22 relative to IN82. The 20°C resultsshown in Fig. 9 are similar to those ob-served by Slaughter (Ref. 35) in DMWsbetween T22 and 304 stainless steel madewith IN82 filler metal. The conventionalT22-IN82-304 DMW exhibited yield andtensile strengths of 298 and 503 MP andfailed in the T22 base metal. This is simi-lar to the values of 300 and 490 MPa forthe T22-IN82 graded joint that also failedvery close to the T22 side of the joint.

The T22-800 graded joint exhibitedsolidification cracking as shown in the ra-diographs in Fig. 11 and light opticalphotomicrographs in Fig. 12. (Samplesfor tensile testing were extracted fromcrack-free locations.) The cracks areconcentrated in an area of the X-ray thatappears the lightest, indicating higherdensity associated with the fcc austenitephase. This light (high-density) regionextends from 11 to 22 mm along thegrade and, as shown in Fig. 7, occurs inthe fully austenitic region. Also note thatthe cracks are located along the grainboundaries, which is typical for solidifi-cation cracks. It is well known that alloysthat solidify as primary austenite aremore prone to solidification cracking.Alloy 800 is not commonly used in fillermetal form as done in this research. Theresults presented here indicate that fur-ther work would be required to identifythe cause of cracking. Composition mod-ifications may be needed before thisalloy is used in a transition joint. At thispoint, it appears more appropriate to uti-lize the T22-IN82 transition joint in favorof the T22-800 combination in order toavoid this problem, since the IN82 end ofthe transition is compatible with manyother austenitic alloys.

Conclusions

Functionally graded transition jointswere fabricated using a dual-wire GTAWprocess and used for microstructural andmechanical property characterization.The following conclusions can be drawnfrom this work.

1. The transition joints exhibited asmooth transition in composition over adistance of ~50 mm, which is significantlylonger than the concentration differencestypically observed in traditional dissimilarmetal welds that occur over distances of~50–100 μm.

2. Measured composition data from thetransition joints were converted to Ni andCr equivalents and used with the Schaefflerdiagram to predict the phase distributionalong the joints, and good agreement wasobtained between the observed and pre-dicted results.

3. Peaks in hardness were observedalong the joints and were attributed tothe formation of martensite.

4. The tensile properties from 20° to650°C were generally within the range ofthose expected for the end member alloysexcept for the T22-IN82 joint, which ex-hibited slightly lower yield and tensilestrengths.

5. Solidification cracking was ob-

served in the T22-800 joints in the regionthat exhibited a fully austenitic solidifica-tion mode.

Work is in progress to evaluate theaging and creep-rupture behavior ofthese joints and will be reported in futurepapers.

Acknowledgments

The authors gratefully acknowledge fi-nancial support through National ScienceFoundation Grant No. CMMI-0758622,the NSF Center for Integrated MaterialsJoining Science for Energy Applications,Grant IIP-1034703, PPL Corp., Contract00474836, The Idaho Engineering Labo-ratory through Contract IIP-1034703, andthe Pennsylvania Infrastructure Technol-ogy Alliance. Useful technical discussionswith Ruben Choug and Robert Schneiderof PPL Corp. and Ron Mizia and MikePatterson of INL are also gratefully ap-preciated.

References

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2. Klueh, R. L., and King, J. F. 1982.

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Fig. 8 — Plots of percent elongation, yield strength, and tensile strength as a function of temperature for agraded transition joint between T22 and 347. The minimum nominal alloy data (Refs. 13, 14, 36) are shownfor comparison. The failure locations for each temperature are shown by the dotted lines on the schematictensile bar above, along with the phase boundaries from the Schaeffler diagram.


Austenitic stainless steel-ferritic steel weldedjoint failures. Welding Journal 61(9): 302-s to311-s.

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5. Bala Srinivasan, P., and Satish Kumar, M.2008. Characterization of thin section dissimi-lar weld joint comprising austenitic and ferriticstainless steels. Materials Science and Technol-ogy 24: 392–398.

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13. Data Sheets on the Elevated-TemperatureProperties of Quenched and Tempered 2.25Cr-1Mo Steel Plates for Pressure Vessels (ASTMA542/A542M). 2003. 1-3 (National ResearchInstitute for Metals, Japan).

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15. Nath, B. 1982. Creep rupture and creepcrack growth behavior of transition joints. In-ternational Conference on Welding Technologyfor Energy Applications pp. 597–621.

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Fig. 9 — Plots of percent elongation, yield strength, and tensile strength as a function of temperature for gradedtransition joints between T22 and IN82. The minimum nominal alloy data (Refs. 13, 37) are shown for com-parison (except percent elongation data for IN82). All failures occurred between the dotted lines on theschematic tensile bar above. Also shown are the phase boundaries from the Schaeffler diagram.

Fig. 10 — Plots of percent elongation, yield strength, and tensile strength as a function of temperature fora graded transition joint between T22 and Alloy 800. The minimum nominal alloy data (Refs. 13, 38–40)are shown for comparison. The failure locations for each temperature are shown by the dotted lines onthe schematic tensile bar above, along with the phase boundaries from the Schaeffler diagram.

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B

Fig. 11 — X-ray radiographs of as-welded T22-800 graded joint showing evidence of so-lidification cracking: A —Light-colored (more dense) region showing cracks through-out that region of the grade; B — close-up view of the cracks.

A

Fig. 12 — Light optical photomicrographs of solidi-fication cracks in T22-800 graded joint.

A BC


fabrication and characterization of graded transition joints...

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