welding journal | february 2014

Introduction

Efficient production and consumptionof energy has become a critical issue in theUnited States and around the world dueto a number of factors, including increaseddemand due to rising population, deple-tion of fossil fuel sources, and detrimentalimpacts on the environment. Figure 1shows the relative amounts of U.S. energyproduction and consumption by fuelsource from 1949 to 2011 (Ref. 1), where itis apparent that fossil fuels and nuclearpower have been the primary fuel sources.It has been estimated (Ref. 2) that worldenergy consumption will increase by 56%between 2010 and 2040. Although renew-able energy and nuclear power are theworld’s fastest growing energy sources,fossil fuels will continue to supply almost80% of world energy use through this time

period. As a result of this increase in en-ergy consumption, carbon dioxide emis-sions from energy production are expectedto increase significantly by approximately46%, from 31 billion metric tons in 2010to 45 billion metric tons in 2040.

These factors point to the need to pro-duce and consume fossil and nuclear fuelswith increased efficiency. Materials playan important role in these applications.For example, the efficiency of coal-firedpower plants is most effectively improvedwith increases in the operating tempera-tures and pressures within the plant. Newso-called “ultra super critical” powerplants are currently being designed to op-erate at significantly higher pressures and

temperatures in order to realize thesegains in efficiency. When plant efficiencyis increased, more power can be generatedwith less consumption of coal and reducedcarbon dioxide emissions. Successful im-plementation of these new plant designshinges on the ability to develop materialswith improved creep strength and corro-sion resistance in order to provide ade-quate design lives out to about 30 years.Nickel-based alloys fill an important needin these new plants. In fact, in some casesNi alloys provide the only suitable candi-date material that can provide the re-quired level of high-temperature strengthand corrosion resistance.

There is a very wide range of Ni-basedalloys that have been developed for anequally impressive range of applications.The extensive development and use of Ni-based alloys can, at least in part, be attrib-uted to two unique characteristics. First,Ni is capable of dissolving high concentra-tions of alloying elements compared toother metals. The explanation for thisdates back to some of the early work onthe electron configuration of metals con-ducted by Pauling and has been attributedto the relatively full shell of d-band elec-trons in the Ni atom (Ref. 3). Second, theaddition of Cr (and/or Al) to Ni providesexcellent corrosion resistance resultingfrom the formation of a protective Cr2O3(or Al2O3) surface oxide layer. This per-mits use of Ni-based alloys in a wide vari-ety of applications that require protectiondue to various forms of degradation, suchas aqueous corrosion, oxidation, and sul-fidation. The subsequent discovery of im-provements in creep strength provided bythe addition of Ti and Al to promote pre-cipitation of the ordered γ ′-Ni3(Ti, Al)phase extended the use of these alloys toapplications requiring a combination ofhigh-temperature strength and corrosionresistance.

While these attributes are certainlybeneficial, the wide range and high con-centration of alloying elements used in Ni-based alloys can present a challenge for

SUPPLEMENT TO THE WELDING JOURNAL, FEBRUARY 2014Sponsored by the American Welding Society and the Welding Research Council

Welding of Nickel-Based Alloys for EnergyApplications

A summary of the Comfort A. Adams Lecture presented at the 2013 AWS Annual Meeting

BY JOHN N. DuPONT

KEYWORDS

Corrosion ResistanceEnergy ProductionFossil FuelsHigh TemperatureNickel AlloysNuclear Energy

JOHN N. DuPONT is R. D. Stout DistinguishedProfessor, Department of Materials Science andEngineering, Lehigh University, Bethlehem, Pa.

This is the 2013 Adams Lecture presented at the94th annual meeting of the American Welding So-ciety, Chicago, Ill.

ABSTRACTThe development of new engineering alloys with improved strength and corrosion

resistance is vital to the production of energy from a wide variety of sources. In manyapplications, critical components used in energy production will require welding dur-ing initial fabrication, maintenance, and repair. Nickel-based alloys are expected toplay a key role in this area since they can often provide levels of high-temperaturestrength and corrosion resistance that cannot be obtained with ferritic steels oraustenitic alloys. However, the thermal cycle associated with welding causes signif-icant microstructural modification to these alloys, usually in a manner that has adetrimental impact on properties. Thus, it is important to understand and controlthese modifications so that full-scale use of new Ni-based alloys is not limited whenwelding is required as a means of construction. This article summarizes some recentadvances made at Lehigh University in the last ~ 10 years toward understanding thewelding metallurgy and weldability of Ni-based alloys for energy applications. Theuse of new computational tools and microstructural characterization techniques havebeen important for obtaining a detailed understanding of welding metallurgy issues,and examples are provided in this review. The advances are described by several ex-amples for specific energy applications.

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microstructure control during welding.The primary reason for this can be under-stood by comparing typical thermal cyclesused for heat treatment during alloy pro-duction to those associated with welding.An example of this for the newly devel-oped Ni-based superalloy IN740H isshown in Fig. 2. During alloy production, aheat treatment (Fig. 2A) is applied thatconsists of an initial hold at 1135°C for oneh, followed by a water quench and re-heatat 800°C for four h. The initial high-temperature treatment is needed to ho-mogenize microsegregation and dissolvesecondary phases in the ingot. The quenchis required to produce a single-phaseaustenite matrix that is supersaturatedwith respect to γ ′-forming elements (Nb,Ti, Al). The final aging treatment at 800°Cis used to produce a fine distribution of γ ′precipitates that provide the best distri-bution and morphology for high-tempera-ture strength. Figure 2B shows a typicalthermal cycle for various locationsthroughout a fusion weld. Comparison ofthe weld thermal cycles to the original heattreatment thermal cycle reveals significantdifferences. The weld is exposed to a muchwider range of temperatures (from abovemelting to ambient temperature) and timeframes that are on the order of seconds(compared to hours for the original heattreatment).

The rapid weld thermal cycles associ-ated with fusion welding cause significantalteration to the microstructure of fusionwelds in Ni-based alloys, usually in a man-ner that has a negative effect on mechani-cal and corrosion properties. Whileprogress is being made to develop new Ni-based alloys to meet future energy de-mands, the joining technology has not kept

pace with the alloy development efforts. Inmost cases, the severe microstructural gra-dients in welds of Ni alloys lead to inferiorproperties that severely limit the overallperformance of the component. Thus, it iscritical that the joining technologyprogress in parallel with alloy develop-ment efforts so that these materials can beused in applications that require welding.This study describes several important ap-plications of fusion welds in Ni-based al-loys that are critical aspects to successfuloperation of various power plants. Exam-ples include use of Ni-based weld claddingin conventional fossil-fired plants operat-ing with low NOx burners, welding of newNi-based superalloys for advanced coal-fired power plants, and welding of Gd-enriched nickel alloys for spent nuclearfuel applications.

Weld Cladding for Corrosion Control in Low-NOx Burners

In an effort to reduce boiler emissions,many coal-fired power plant operators havemoved toward a staged combustion process.By delaying the mixing of fuel and oxygen,the amount of nitrous oxides (NOx) that arereleased as a by-product of combustion isreduced (Refs. 4, 5). Prior to this, mostboiler atmospheres were oxidizing, allowingfor formation of protective metal oxides onwaterwall tubes made out of carbon- or low-alloy steels (Refs. 4, 6). Under those condi-tions, failure due to accelerated waterwallwastage was generally not a major problem.Staged combustion boilers, on the otherhand, create a reducing atmosphere in theboiler due to the lack of oxygen. Sulfur com-pounds from the coal are transformed intohighly corrosive gaseous H2S (Ref.7). In ad-

dition, corrosive deposits may form on thewaterwall tubes due to the accumulation ofsolid particles in the combustion environ-ment, such as ash and unburnt coal. As a re-sult, low-alloy steels are often susceptible toexcessive wastage rates and unsatisfactoryservice lifetimes (Refs. 4, 7). Thermal sprayand chromium diffusion coatings were ini-tially evaluated for protection in these envi-ronments, but generally do not provideadequate corrosion resistance. Commer-cially available nickel-based weld claddingis currently the industry standard for corro-sion protection of waterwalls in boilers op-erating with low-NOx burners (Ref. 8).These alloys provide excellent resistance togeneral corrosion and can extend the serv-ice life of waterwalls relative to bare tubes.However, recent experience has shown thatthese coatings are susceptible to prematurefailure due to corrosion-fatigue cracking(Ref. 9). In fact, waterwall failures are theleading cause of forced outages of coal-firedpower plants and can cost a utility companyfrom $250,000 to $850,000 a day in down-time and lost revenue (Ref. 10).

Figure 3A is a photograph of an IN625cladding that was applied with the gasmetal arc welding (GMAW) process andremoved from service due to the presenceof extensive corrosion-fatigue cracking(Ref. 9). Figure 3B shows a cross-sectionalphotomicrograph of several small cracksthat were examined early in the crackingstage, and Fig. 3C shows the distributionof alloying elements across the dendritic

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Fig. 1 — Relative amounts of U.S. energy production (A) and consumption (B) by fuel source from 1949to 2011.

Fig. 2 — A — Typical heat treatment of the basemetal for newly developed IN740 nickel-based su-peralloy; B — typical thermal cycle at various lo-cations throughout a weld that are associated withfusion welding.

AA B

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substructure of the cladding. Figure 3Dprovides a lower magnification view thatdemonstrates the cracks initiate at the val-leys of the weld ripples. The dendrite coresin the cladding exhibit a minimum in alloyconcentration due to the relatively rapidsolidification conditions associated withwelding. As a result, the corrosion rate isaccelerated in these regions and localizedattack occurs at the dendrite cores. Theselocalized penetrations form stress concen-trations, which eventually grow into full-size fatigue cracks under the influence ofresidual and service-applied stresses,where the service-applied stresses ariseprimarily through thermal cycling. Asshown in Fig. 3D, most cracks initiate at aregion in the valley of surface weld ripplewhere an additional stress concentrationexists. The high residual stress that resultsfrom welding also probably contributes tothe cracking problem. In addition, dilutionfrom the underlying steel tube substrate,which results in reduced alloy content ofthe cladding, compromises the corrosionresistance of the cladding.

The primary metallurgical factors thatcontribute to corrosion-fatigue cracking(weld ripples, microsegregation, highresidual welding stresses, dilution) are allassociated with the localized heating,melting, and solidification of the weldingprocess. Thus, use of a coating that can beapplied with more uniform heating in thesolid state should help mitigate theseproblems and improve the cracking resist-ance. Work is in progress (Ref. 11) to eval-uate coatings made by the coextrusionprocess in order to eliminate or reduce theinherent problems of weld cladding. Withthis process, a cylindrical shell of a corro-sion-resistant alloy is first joined by explo-sive welding to a steel substrate, and thebimetallic billet is then coextruded at anelevated temperature to produce a tubewith an outer coating. Since there is nomelting/solidification involved, the coatingmicrostructure consists of equiaxed grainswith no microsegregation, similar to thatexpected for a wrought alloy.

The corrosion resistance of Alloys 600(Ni-16Cr-8Fe) and 622 (Ni-22Cr-13Mo-2Fe) weld cladding and coextruded coat-ings have recently been compared usingthermogravimetric testing in a simulatedcombustion gas (10%CO-5%CO2-2%H2O-0.12%H2S-N2, in vol-% at 600°C,and the results are shown in Fig. 4 (Ref.11). A sample of wrought Alloy 600 wasalso tested. The coextruded coatings ex-hibit significantly better corrosion resist-ance compared to the weld cladding of thematching alloy. Figure 5 shows a weldcladding sample that was corrosion testedunder solid-state conditions and thenetched to reveal the dendritic substruc-ture. Note that preferential corrosion hasoccurred at the dendrite cores (arrows).

Figure 6 provides an EDS line scan thatwas acquired across the dendritic sub-structure of the weld cladding. As ex-pected (Ref. 12), the dendrite cores aredepleted in Mo, with Mo concentrationlevels down to ~ 11 wt-% (the nominalMo concentration of the filler metal is ~ 13 wt-%). Figure 7 shows the mi-crostructure of the coextruded coating,and an EDS line scan acquired across sev-eral grains of the coating is shown in Fig. 8.The coextruded coating exhibits a uni-form, equiaxed grain structure and a uni-form distribution of alloying elements.

For Alloy 600, the corrosion resistanceof the wrought alloy and coextruded coat-ing are comparable. This indicates that thecoextrusion coating process has no adverseeffect on the inherent corrosion resistanceof the alloy. The improved corrosion re-sistance of Alloy 622 relative to 600 is at-tributed to the higher Cr and Mo contentof Alloy 622. The difference in corrosionperformance between the coating typescan be attributed to two factors. First, theweld claddings evaluated in this work con-tain an additional 10 wt-% Fe from dilu-tion with the steel substrate. Thisadditional Fe stems from dilution with the

underlying steel substrate and is alwayspresent in any weld cladding. The additionof Fe results in a corresponding decreasein the Cr and Mo content and often has adetrimental effect on corrosion resistance.The 10% dilution value used to preparethe sample for these tests represents alower limit on the dilution level for com-mercially applied coatings. Such dilutioneffects do not occur with the coextrudedcoating. Second, the weld cladding exhibitsmicrosegregation of alloying elements, re-sulting in preferential corrosion of thealloy-depleted dendrite core regions. Thecoextruded coatings also have a uniformcoating thickness and smooth surface fin-ish that should help reduce the localizedstress concentrations that can aggravatethe corrosion-fatigue problem. In addi-tion, the heating and cooling cycles expe-rienced during coextrusion are much lesssevere and more uniform compared to fu-sion welding, so the residual stressesshould be significantly reduced.

While the mechanism of corrosion fa-tigue cracking is generally understood,there is a need to understand the corro-sion fatigue behavior on a more funda-mental basis in order to assess the relative

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Fig. 3 — A — Photograph of a weld cladding with extensive circumferential cracks; B — cross-sectionalscanning electron photomicrograph of several small cracks early in the cracking stage; C — distributionof alloying elements across the dendritic substructure of the cladding; D — photograph showing crack ini-tiation at the valley of the weld ripple.

A B

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cracking susceptibility of candidate alloysand coating processes. In order to accom-plish this, an experimental techniqueneeds to be developed that accurately sim-ulates the corrosion fatigue mechanism ofweld cladding in service. Typically, fatiguetests involve the use of standard specimen

configurations such ascompact tension (C(T))or single edge-notched(SEN). These types oftests involve fatiguecrack propagation usinga single crack design.While this type of ap-proach offers the abilityto measure propagationrates of isolated cracks,it is not entirely ade-quate for studying thecorrosion fatigue behav-ior of weld cladding incombustion conditionsfor several reasons.First, these approachesdo not provide informa-tion on the corrosion fa-tigue crack initiationbehavior of multiplecracks. Since crack initi-ation can comprise alarge portion of the fa-tigue life, it is impera-tive that the crackinitiation behavior becharacterized. Addi-tionally, single crack ex-periments do not takeinto account the effectof crack interactions onthe crack propagationbehavior. Numerouscircumferential cracksform on the surface ofNi-based weld claddingduring service (Refs. 9,13). A series of crackson the surface can alter

the crack propagation behavior by reduc-ing the stress intensity factors to a levelwell below that of a single isolated crack(Refs. 14, 15). Therefore, in order to un-derstand the corrosion fatigue resistance,the effects of crack interactions need to beconsidered. This includes understanding

the crack initiation behavior that affectsthe crack depths and distribution.

Figure 9 shows a recently developedcorrosion fatigue cracking test apparatusinvolving a Gleeble thermomechanicalsimulator. With this test, a retort is posi-tioned around the sample to allow the ap-plication of simulated combustion gases.The test sample is resistively heated to aconstant temperature of 600°C, which is atypical surface temperature of Ni-basedcladdings in service. A representative(Refs. 16, 17) sulfidizing gas of N2-10%CO-5%CO2-0.12%H2S circulatedthrough the retort during the test. Corro-sion fatigue tests to date have been con-ducted with an alternating stress profileinvolving a minimum tensile stress of 0MPa and a maximum tensile stress of 300MPa. The minimum and maximumstresses were alternated every five min(ten-min fatigue cycles). A maximumstress of 300 MPa was chosen because it isabove the 200 MPa yield strength of Alloy622 at 600°C. This was done to simulatethe residual tensile stresses that develop inthe waterwall tubes that cause significantyielding.

The validity of the experimental ap-proach was first examined by comparingthe laboratory-induced cracking mecha-nism to the established mechanism fromfield samples (Ref. 9). Figure 10A demon-strates the layered multiphasic scales thatdeveloped on the surface of the samples.An embryonic corrosion fatigue crack isshown, which formed within the multilay-ered corrosion scale. Figure 10B shows theinitiation of a corrosion fatigue crack at apreferentially corroded dendrite core (thecorrosion scale was removed by the etch-ing process). These results are consistentwith the corrosion fatigue mechanism ob-served on field samples. Figure 10C and Dillustrates that the mature corrosion fa-tigue cracks propagated down the mainaxis of the dendrite cores and exhibit a sec-ondary or “spinal” phase along the length


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Fig. 4 — Thermogravimetric results from the gaseous corrosion testing for Alloy 600 (A) and Alloy 622 (B).

Fig. 5 — Light optical photomicrographs of the 300-h corrosion sample ofthe weld cladding after it was etched to reveal the dendritic substructure. Notethat preferential corrosion has occurred at the dendrite cores (arrows).

A B


of the crack. This is also consistent withthe corrosion fatigue mechanism observedfrom examination of samples removedfrom the field. Figure 11 shows represen-tative images of corrosion fatigue cracks inweld cladding samples tested out to 440cycles. Figure 11A shows the presence of asmall corrosion fatigue crack after 25 cy-cles. This observation indicates that cor-rosion fatigue cracks initiated relativelyquickly as a result of the accelerated testconditions. Thus, the experimental ap-proach was able to accurately reproducethe corrosion fatigue mechanism in a rel-atively short amount of time. Additionally,multiple corrosion fatigue cracks were ap-parent within each of the samples after 25cycles, indicating that cracks continuouslyformed throughout the test. The variationin the corrosion fatigue crack depths withincreasing number of corrosion fatigue cy-cles further supports this conclusion. Se-rial sectioning and quantitative imageanalysis techniques have recently beenused to measure the frequency and depths

of corrosion fatigue cracks that developduring this test. As an example, Fig. 12compares the maximum crack depths of acladding applied with the GMAW processto that of a laser weld cladding. Note thatthe laser weld cladding exhibits better per-formance than the GMAW cladding interms of both time to crack initiation andcrack depth for a given number of cycles.Work is currently in progress to under-stand the microstructural differences thataccount for this improvement in corro-sion-fatigue cracking resistance and to testother coating systems, such as the coex-truded coating described above.

Welding of IN740H for Ultrasupercritical Power Plants

As described above, advanced ultra-supercritical (AUSC) power plants arecurrently being designed to operate athigher pressures and temperatures for in-creased efficiency. These more aggressiveoperating conditions place heavy demands

on tubing and piping components withinthe plants, particularly with regard tocreep resistance. Figure 13 (Ref. 18) showsthe maximum allowable stress as a func-tion of temperature for a wide range ofstainless steels and nickel alloys. The tem-peratures and pressures associated withAUSC conditions are expected to be ~ 1300° –1400°F (704° –760°C) and 4500lb/in.2 (31 MPa), respectively. Note thatthese operating conditions lie on theupper bound of allowable stresses andtemperatures even for the commonly usedsolid solution strengthened Ni alloys. Thissituation represents a significant challengeto the successful implementation of thesenewer, more efficient plants. In responseto this need, a new precipitation-strength-ened nickel-based superalloy has recentlybeen developed known as IN740H. Thisalloy has significant additions of Nb, Ti,and Al in order to form the γ ′ precipitatefor high-temperature creep strength. Ex-tensive fabrication by fusion welding willbe required on this alloy during plant con-

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Fig. 6 — EDS line scan. A — Acquired across the dendritic substructure of the weld cladding showing the composition profiles for Fe, Ni, and Cr; B — acquiredacross the dendritic substructure of the weld cladding showing Mo depletion at dendrite cores.

Fig. 7 — Light optical photomicrograph showing the microstructure of thecoextruded coating.

Fig. 8 — EDS line scan acquired across several grains of the coating.


struction and subsequent maintenance.Thus, retention of the creep properties ofthe welds is of primary importance.

Figure 14 shows recent results fromOak Ridge National Laboratory (Ref. 19)that compares the stress rupture proper-ties of wrought IN740H tube material tothat of cross-weld samples, where a signif-icant reduction in the weld metal creep lifeis apparent. These results have requiredthe use of a weld strength reduction factorfor the welds of 0.70, which can placerather appreciable restrictions on the load-bearing capacity of welded components.Thus, it is important to understand thecause of this reduced creep strength so

that a viable solution can be developed.Electron microscopy techniques have

recently been used (Refs. 20, 21) to char-acterize both base metal and weld samplesexposed to a wide range of creep test con-ditions in order to identify the cause forthe reduced creep life in the welds. Figure15 compares scanning electron microscopy(SEM) photomicrographs of base metaland weld samples after creep testing. Thebase metal sample exhibits a uniform dis-tribution of γ ′ precipitates directly up tothe grain boundary. In contrast, the weldmetal has precipitate-free zones (PFZs) atthe grain boundaries, and the creep voidsthat initiate premature failure in the weld

are commonly observed to initiate at thePFZs. The localized creep-failure alongthese PFZs is generally attributed to theirlower strength due to the lack of a densepopulation of fine γ ′ precipitates. Thepresence of PFZs in creep-resistant alloyshas been observed in many alloy systemsand been attributed to a number of differ-ent formation mechanisms, includingstress-assisted diffusion associated withNabarro-Herring (NH) creep conditions,grain boundary carbide formation, anddiscontinuous precipitation/coarsening(DPC) (Refs. 22–34). Each mechanism isassociated with a set of unique mi-crostructural features that can be used toidentify the responsible mechanism, whichis critical for formulating a solution. Whilea complete review of these mechanisms isbeyond the scope of this paper, several dis-tinguishing characteristics are worthy ofdiscussion.

Nabarro-Herring creep and grainboundary carbide formation will each pro-duce PFZs that are symmetric about thegrain boundary. In addition, the coarsesecondary phases within the PFZs typicallyexhibit a globular morphology. Precipi-tate-free zones formed under NH creepconditions have the added feature of beingdependent on the direction of appliedstress, where the PFZs will form normal tothe applied stress direction. In contrast,PFZs that form due to DPC exhibit alamellar morphology and are asymmetricabout the grain boundary (i.e., the PFZwill be confined to one grain). These fea-tures are the result of precipitation (orcoarsening) in the presence of a movingboundary. Figure 16 shows the lamellarmorphology that is commonly observedwith the PFZs in welds of IN740H thathave been creep tested. Transmission elec-tron miscroscope techniques have beenused to confirm that the coarse secondary


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Fig. 9 — A — A corrosion fatigue sample that has been clamped into the grips of the Gleeble; B — the retort that has been designed to seal around a corrosionfatigue sample and allow for the application of corrosive gas.

Fig. 10 — A — Multilayer corrosion scale that develops under the simulated combustion conditions; B— corrosion fatigue cracks initiating at the alloy-depleted dendrite cores. (The etching process removedthe corrosion scale on the surface of the sample); C, D — a mature corrosion fatigue crack that propa-gated down a dendrite core. A secondary phase was observed along the length of crack.

A B

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phase within the PFZ is γ ′. Figure 17shows electron backscattered diffraction(EBSD) patterns acquired across theboundary that contains a PFZ. The EBSDpatterns demonstrate that the PFZ is lo-cated entirely in Grain 2. Detailed com-parisons of creep samples from a widerange of test conditions also showed thatthere was no relation between the PFZorientation to that of the applied stress. Itis worth noting that similar PFZs are oc-casionally observed in the base metal ofIN740H after creep. However, the basemetal requires significantly longer times(or higher temperatures) for PFZ forma-tion. Thus, the reduced creep strength inwelds of IN740H can be attributed to en-hanced susceptibility to PFZ formation as-sociated with DPC.

The enhanced driving force for PFZsin the weld metal can be understood withreference to Figs. 18 and 19. Figure 18shows results from EDS traces acquired

across the dendritic substructure of theweld. The data points reflect the averageand standard deviation from multiplemeasurements, while the lines representresults from Scheil solidification simula-tions (Ref. 20). Note there is good agree-ment between the model andexperimental results. In this case, the cal-culations are probably more reflective ofthe actual compositions within the veryedge of the interdendritic regions due tothe spatial resolution limitations of theelectron beam (which is about 1 μm3) andthe large concentration gradient within theinterdendritic region. The results alsoshow that, due to microsegregation, theconcentrations of Ti and Nb are higher atthe grain boundary and interdendritic re-gions (the Al concentration is also slightlyhigher). The increased localized concen-tration of these γ ′ forming elements is im-portant, since it causes an increasedsupersaturation beyond the solubility

limit, thus leading to the enhanced grainboundary precipitation shown in Fig. 19.Classical solutions to the diffusion equa-tions associated with precipitation (Ref.35) reveal that the growth rate of the pre-cipitate/matrix interface is directly pro-portional to the degree of supersaturationin the matrix, ΔCo. Here, ΔCo can beviewed as the driving force for precipita-tion and is given by the difference betweenthe actual matrix concentration and thesolubility limit at a given temperature. Fig-ure 20 shows the ΔCo values (at 800°C,the typical operating temperature) associ-ated with the γ ′ forming elements Nb, Ti,and Al that demonstrate the significantlyenhanced driving force for precipitationassociated with interdendritic microsegre-gation. Also note from Fig. 19 that thereis significant grain boundary curvature.This could be associated with the original(nonequilibrium) grain boundary curva-ture typical in the as-solidified condition,

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Fig. 11 — Corrosion fatigue cracks observed in the GMAW cladding. A— 25cycles; B — 50 cycles; C — 100 cycles; D — 442 cycles.

Fig. 13 — Maximum allowable stress as a function of temperature for a widerange of stainless steels and nickel alloys.

Fig. 12 — A — Average crack depths; B — maximum crack depths for theGMAW and laser weld cladding samples.

Fig. 14 — Recent results from Oak Ridge National Laboratory that comparesthe stress rupture properties of wrought IN740H tube material to that of cross-weld samples.

A B

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which would then provide the driving forcefor boundary movement. The observed cur-vature could also possibly be the result ofthe DPC process itself. Work is in progressto more clearly understand the factors lead-ing to boundary migration.Work is also in progress to explore newfiller metal compositions that should bemore resistant to the formation of PFZsduring creep. Meanwhile, the results pre-sented above demonstrate that a postweldheat treatment (PWHT) that homogenizesmicrosegregation within the weld may helpto minimize the problem of PFZ formation.Dissolution of secondary phases that formduring solidification may also be helpful. Inlight of this, it is important to establish theinfluence of time and temperature on thehomogenization kinetics of welds inIN740H. A study in this area has recentlybeen completed and is summarized below(Ref. 20).

The most convenient PWHT would per-mit dissolution of secondary phases and ho-mogenization in a single initial step,followed by a subsequent aging treatment.However, it must be recognized that thelocal increases in solute concentrationwithin the interdendritic regions will causea corresponding decrease in the localsolidus temperature. This, in turn, couldlead to localized melting within the solute-enriched interdendritic regions if the initialPWHT is too high. This is demonstrated inFig. 21, which shows the stability of various


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Fig. 15 — Scanning electron microscopy photomicrographs of base metal and weld samples of Alloy IN740H after creep testing.

Fig. 16 — Scanning electron microscope photomicrograph showing the lamellarmorphology that is commonly observed with the PFZs in welds of IN740H thathave been creep tested.

Fig. 17 — Electron backscattered diffraction (EBSD)patterns acquired across the boundary of a weld inIN740H that contains a PFZ. The EBSD patternsdemonstrate that the PFZ is located entirely in Grain 2.


phases as a function of temperature inIN740H for the base metal nominal com-position (Fig. 21A, B) and for a value of0.99 fraction solid that is representative ofthe interdendritic composition — Fig.21C, D. (Figure 21B, D shows the same re-sults as Figs. 21A and 22C, but over a nar-rower temperature range for clarification.)Note that, if the heat treatment was basedon the nominal composition of the basemetal, the results suggest that the initialPWHT step could be conducted slightlybeyond 1300°C before any problems fromlocalized melting (also note that the MCcarbide will still be stable at these hightemperatures). However, when dendriticsegregation is properly considered (Fig.21C, D), the results demonstrate that theinitial PWHT temperature should not ex-ceed about 1150°C to avoid localized melt-ing in the interdendritic regions. These

results emphasize the importance of de-signing PWHT schedules based on the ac-tual phase stability in the fusion zone asaffected by compositional gradients.

Figure 22 shows the results of DIC-TRA calculations for IN740H that demon-strate the homogenization kinetics.Results are shown for the concentrationgradient of Nb, since this is the slowest dif-fusing element in the system and thereforethe rate-limiting step. The calculationsshow that a PWHT at 1100°C for four hshould eliminate the microsegregation,and this has been confirmed by experi-mental measurements (Fig. 23). Once thehomogenization is complete, a higher tem-perature solution temperature can be uti-lized if needed for more completedissolution of secondary phases. Althoughit is recognized that the 1100°C/4-h treat-ment is not practical for a PWHT con-

ducted in the field, the results can be ap-plied to shop fabrication conditions. Asmentioned previously, the potential for amore practical solution through design oftailored filler metals is needed and is cur-rently being pursued.

Fusion Welding of Gd-Enriched NiAlloys for Spent Nuclear Fuel Applications

Nuclear fuel plays an important role inenergy production in the United Statesand throughout the world, and the use ofnuclear power is expected to rise over thecoming years. One of the major challengesassociated with nuclear fuel is the handlingof spent nuclear fuel (e.g., nuclear waste).Safe disposition of spent nuclear fuel re-quires the development of thermal neu-tron-absorbing structural materials for

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Fig. 18 — EDS traces acquired across the dendritic substructure of a weld in IN740H. The data points reflect the average and standard deviation from multi-ple measurements, while the lines represent results from Scheil solidification simulations.

A B

Fig. 19 — SEM photomicrograph showing how increased localized concen-tration of γ ′-forming elements produces enhanced grain boundary precipi-tation in fusion welds on Alloy IN740H.

Fig. 20 — ΔCo values at 800°C associated with the γ ′-forming elements Nb,Ti, and Al that demonstrate the significantly enhanced driving force for pre-cipitation in the interdendritic regions of welds in Alloy IN740H.


nuclear criticality control. These materialswill be used for the internal baskets thatseparate spent fuel assemblies and are re-quired for structural support, spent nu-clear fuel geometry control, and nuclearcriticality safety. Given the large quantityof material required for this application,the material should be producible withconventional fabrication methods such asingot casting and hot working. Ultimately,the material will be formed and weldedinto an internal structure that will cradlethe fuel and maintain a specified geome-try, so the material must also exhibit goodweldability.

Recent research (Refs. 36–38) has fo-cused on development of gadolinium-(Gd) enriched Ni-based alloys for this ap-plication. Gd serves as an effective alloyaddition for this application because it hasa very high neutron absorption cross sec-tion (Refs. 23, 39). The results of recentresearch has also shown that Gd additionsat the nominal 2 wt-% level can success-fully be added to the commercial Ni-basedC-4 alloy while maintaining adequate hotworkability, weldability, and mechanicalproperties.

The current plan for fusion welding ofthis alloy involves the use of a Gd-freecommercial filler metal such as Alloy C-4or 59. With this approach, the Gd content

in the fusion zone will vary with weld metaldilution, and the dilution level is stronglyaffected by the welding parameters (Refs.40, 41). Considering the case in which abase metal containing 2 wt-% Gd iswelded with a Gd-free filler metal, theconcentration of Gd in the fusion zone willvary linearly with dilution and can be closeto 0 wt-% Gd (at low dilution values) to 2wt-% Gd (for an autogenous weld inwhich the dilution is 100%). The weldingparameters are typically selected based onthe required weld size and joint design re-quirements, and a wide range of weldingparameters can be expected in practice.This, in turn, can produce a wide range offusion zone Gd concentration values. TheGd concentration can also vary through-out the fusion zone of a given weld in mul-tiplepass welding applications. Recentresearch has shown that Gd controls thesolidification behavior of this alloy. In par-ticular, the solidification temperaturerange and amount of terminal eutectic-type constituents that form at the end ofsolidification are essentially dominated bythe Gd concentration (Ref. 38). In addi-tion, it has also been well established thatthe solidification cracking susceptibility isalso strongly affected by the solidificationtemperature range and amount of termi-nal eutectic-type constituents that form

during solidification of the fusion zone(Refs. 42, 43). Thus, it is important to de-termine the influence of Gd concentrationon the solidification cracking susceptibil-ity of Alloy C-4 so that these structuralmaterials can be welded while maintainingtheir structural integrity.

Figure 24 shows the influence of Gd onthe solidification cracking susceptibility ofAlloy C-4. The cracking susceptibility islow for the Gd-free base alloy. The crack-ing susceptibility then increases with in-creasing Gd concentration up to ~1 wt-%Gd and then decreases to a level that issimilar to the base alloy at Gd concentra-tions of 1.5 to 2.5 wt-%. Typical mi-crostructures of the Varestraint samplesare shown in Fig. 25. No significant crackhealing due to backfilling of solute-richliquid occurred in the alloy with 1.01 wt-%Gd. Similar results were obtained with thealloy that had 0 and 0.46 wt-% Gd. How-ever, a significant amount of crack healingwas observed in the alloys that had addi-tions of Gd at the 1.49 wt-% level andabove.

The Varestraint weldability resultsshow that the cracking susceptibilityreaches a maximum at ~ 1 wt-% Gd, anddecreases with both higher and lower Gdadditions. Previous research (Refs. 42, 43)has shown that the solidification crackingsusceptibility of engineering alloys isstrongly affected by the solidification tem-perature range and amount of terminaleutectic-type constituent that forms at theend of solidification. Solidification ofthese alloys initiates at the liquidus tem-perature by the formation of primary γ−austenite. Essentially, no Gd is dissolvedin the austenite matrix. Thus, as solidifica-tion proceeds, the liquid becomes increas-ingly enriched in Gd until the Liquid → γ+ Ni5Gd eutectic-type reaction is reached,at which point solidification is terminated.


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Fig. 21 — Stability of various phases as a function of temperature in IN740H for the base metal nominalcomposition (Fig. 21A, B) and for a value of 0.99 fraction solid that is representative of the interdendriticcomposition (Fig. 21C, D) in fusion welds.

Fig. 22 — DICTRA calculations for IN740H thatdemonstrate the homogenization kinetics. Results areshown for the concentration gradient of Nb, since thisis the slowest diffusing element in the system andtherefore the rate-limiting step.

A

DC

B


This reaction sequence and temperaturerange is generally similar to that expectedin the binary Ni-Gd system. Simple binaryNi-Gd alloys with less than about 13 wt-%Gd exhibit a similar two-step solidificationsequence consisting of primary austeniteformation followed by a terminal eutecticreaction involving the Ni17Gd2 inter-metallic at 1275°C (Ref. 44). By compari-son, the multicomponent Ni-Cr-Mo-Gdalloys examined here complete solidifica-tion at ~ 1258°C by a terminal eutectic-type reaction involving the Ni5Gd inter-metallic. Thus, although the secondaryphase within the terminal eutectic con-stituent is different in each case, the ter-minal reaction temperatures are verysimilar. In fact, as shown in Fig. 26, apseudo binary solidification diagram hasrecently been developed (Ref. 38) for thisalloy that is similar to the phase diagramof a binary eutectic alloy. In this case, the“solvent” is represented by the Ni-Cr-Mosolid solution γ -austenite phase and Gdis treated as the solute element. Althoughthe diagram does not account for theminor variation in matrix Mo concentra-tion that occurs due to microsegregation,the similarity of this γ -Gd binary systemto a binary eutectic system is readily evi-dent in several ways, including the as-solidified microstructure consists of pri-mary γ dendrites surrounded by an inter-dendritic eutectic-type constituent inwhich the secondary phase in the eutecticis solute rich; the amount of eutectic-typeconstituent increases with increasingsolute content; and the proportionalamount of each phase within the eutecticconstituent is relatively insensitive to nom-inal solute content (Ref. 38). Also notethat the eutectic temperature is notstrongly dependent on the nominal Gd-concentration.

The pseudo binary diagram developedfor these alloys is useful for interpretingthe weldability results because it permitsdirect determination of the solidificationtemperature range and amount of termi-nal γ /Ni5Gd eutectic constituent as afunction of Gd concentration. The frac-tion of terminal γ /Ni5Gd eutectic con-stituent (fe) that forms duringsolidification can be calculated with theScheil equation (Ref. 45) via

(1)

where CGde is the concentration of Gd in

the liquid at the eutectic reaction (14.7 wt-% Gd, Fig. 26), CGd

o is the nominal Gdconcentration, and k is the distribution co-efficient for Gd. Equation 1 is valid forconditions in which the diffusivity of solute(Gd) in the solvent (γ) is insignificant. Pre-vious work (Ref. 37) has shown that es-sentially no Gd is dissolved in the γmatrix, which results in the equivalent con-dition of negligible solute diffusivity in thesolvent. The lack of Gd solubility inaustenite also indicates that k for Gd is 0,and Equation 3 reduces simply to

(2)

The solidification temperature range offusion welds is best represented by theseparation between the on-heating liq-uidus temperature (TL) and on-coolingeutectic temperature (Te) because solidi-fication initiates epitaxially at the weld in-terface without the need for undercooling(Refs. 46, 47). As mentioned previously

and discussed in more detail elsewhere(Ref. 38), the eutectic temperature (Te)does not vary significantly with Gd con-centration and can be represented by anaverage value of Te = 1258°C. The liq-uidus line in Fig. 26 can be expressed by asimple linear equation of the form

(3)

where To is the melting point of the Ni-Cr-Mo “solvent” and mL is the liquidus slope.Linear regression analysis of the phase di-agram leads to values of To = 1422°C andmL = –11.2°C/wt-% Gd. Thus, the solidi-fication temperature range (ΔT) can be di-rectly determined as a function of Gdconcentration via

(4)

The solidification temperature range isimportant from a weldability perspective be-cause it controls the size of the solid + liq-uid “mushy” zone that trails the fully moltenweld pool. Assuming a constant tempera-ture gradient in the solid + liquid region(i.e., fixed welding parameters and samplesize/geometry), the size of the mushy regionis given simply by the ratio of the solidifica-tion temperature range to temperature gra-dient. Thus, alloys with narrow solidificationtemperature ranges are generally crack re-sistant because the mushy zone is relativelysmall. The influence of the amount of ter-minal eutectic liquid is a bit more compli-cated. At very low amounts of terminalliquid, the cracking resistance is generallynot adversely affected because there is notenough liquid to cover the solidificationgrain boundaries and interdendritic regions.Thus, solid-solid boundaries can be easily

fC

Ce

eGd

oGd

k1

1=

⎡

⎣⎢⎢

⎤

⎦⎥⎥

−

fC

Ce

oGd

eGd=

T T m CL o L oGd= +

T T m C To l oGd

eΔ = + −

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Fig. 23 — Semiquantitative (standardless) XEDS line scan across dendrites in a single-passGTA weld of Alloy 740H after homogenization at 1100°C for 4 h. A — Light optical micro-graph of region of interest; B — SEM micrograph of region of interest; C — concentrationprofile for major alloying elements; D — concentration profile for γ ′-forming elements.

Fig. 24 — Varestraint results showing maximum crack length as afunction of Gd concentration in Alloy C-4.


established without interference from theliquid, and the cracking resistance is gener-ally good. (The presence of surface-activeelements that modify the solid/liquid sur-face energy and promote extensive wettingof the boundaries by the liquid at even smallliquid fractions can alter this general trend.)As the fraction of terminal eutectic liquidincreases, the liquid more extensively wetsthe boundaries, thus interfering with theformation of solid-solid boundaries andcausing an increase in the cracking suscep-tibility. At considerably higher fractions ofliquid, backfilling of cracks can occur by theliquid and heal the cracks as they form. Thisphenomenon is similar to liquid feeding byrisers used to control solidification shrink-age defects in castings.

With this background in mind, the in-fluence of Gd concentration on the crack-ing susceptibility, amount of terminaleutectic constituent, and solidificationtemperature range are summarized in Fig.27. The lines in the bottom two figuresrepresent the fraction eutectic and solidi-fication temperature range, respectively,calculated with the equations above, andthere is good agreement between the cal-culated and measured values. This sup-ports the use of a pseudo binary analog formodeling the solidification behavior ofthese alloys. More importantly, the resultsprovide a basis for developing a detailedunderstanding of the weldability results.At low Gd concentrations, the solidifica-tion temperature range is high, but the

fraction of terminal eutectic liquid is verylow. Thus, the cracking susceptibility islow. As the Gd concentration increases,there is only a slight decrease in the solid-ification temperature range, but a rathersignificant increase in the fraction of ter-minal eutectic liquid. This leads to the in-crease in cracking susceptibility shown inFig. 27 that reaches a maximum at ~ 1 wt-% Gd. With still increasing amounts ofGd, the amount of terminal liquid in-creases to the point where cracking sus-ceptibility decreases due to backfilling ofsolidification cracks. The reduction in thesolidification temperature range with in-creasing Gd concentration also assists indecreasing the cracking susceptibility. Theamount of terminal liquid required to pro-mote backfilling has been suggested to bein the range of 7–10 vol-% (Ref. 48). Theresults in Fig. 27 support this, where thecracking susceptibility is observed to de-crease when the fraction of terminal eu-tectic reaches this range. This occurs whenthe Gd concentration reaches ~1.5 wt-%.This phenomenon is also supported by di-rect observation of backfilling that becameappreciable when the Gd reached 1.5 wt-%, as shown in Fig. 25.

These effects can be evaluated in morequantitative detail by combining solidifi-cation theory with simple heat flow equa-tions for determining both the size of thecrack-susceptible region of the solid + liq-uid mushy zone and variation in fractionliquid with distance within the mushy re-

gion. The well-known Scheil equation canalso be used to calculate the fraction of liq-uid (fL) as a function of temperature via

(5)

where, as before, To is the melting point ofthe Ni-Cr-Mo solvent and TL is the liq-uidus temperature of the alloy (as affectedby Gd concentration). T is the actual tem-perature. Noting that k for Gd is 0, Equa-tion 6 reduces simply to

(6)

The variation in temperature within thesolid + liquid region can be estimatedusing the Rosenthal heat flow solution(Ref. 49) which, for three-dimensionalheat flow, is given by

(7)

where T is the actual temperature, Tp isthe preheat temperature, η is the heatsource transfer efficiency, P is the arcpower, h is the thermal conductivity of thebase metal, α is the thermal diffusivity ofthe base metal, S is the heat source travelspeed, r is the radial distance from theheat source, and x is the distance behindthe heat source. The weld centerline is ofthe most interest here because this is thelocation where the temperature gradientis the lowest and, as a result, where thecrack-susceptible solid + liquid region isthe largest. This accounts for the experi-mental observation of the maximum cracklength occurring in the weld centerline re-gion. At the weld centerline, r = x andEquation 7 reduces t

fT TT TL

o

o L

k1

1=

−−

⎡

⎣⎢

⎤

⎦⎥

−

forfT TT T

k 0Lo L

o=

−−

=

T Tphr

S r x2

exp2p

ηπ α

( )= + ⎛⎝⎜

⎞⎠⎟

− −⎡⎣⎢

⎤⎦⎥


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Fig. 25 — Microstructures of the Varestraint samples from Alloy C-4 alloyed with various concentrationsof Gd. A — 1.01 wt-% Gd; B — 1.49 wt-% Gd; C — 1.90 wt-% Gd; D — 2.45 wt-% Gd.

A B

C DFig. 26 — Pseudo-binary solidification diagram de-veloped for Gd-enriched Alloy C-4.


(8)

Equation 8 permits an estimate of thetemperature variation (T) with position(x) behind the heat source. Thus, theseequations can be combined to obtain anexpression between the fraction liquid (fL)and distance along the centerline of theweld x

(9)

This expression is useful because it per-mits direct estimation of both the variationin fraction liquid with location and the sizeof the crack-susceptible region of themushy zone as a function of alloy parame-ters (To, mL, h, k, Co

gd) and welding pa-rameters (To, P). The solid + liquid regionwill exist where the temperature is be-tween TL and Te, and the fraction liquidwill vary from 1 at T = TL to fe at T = Te.

Figure 28 shows calculated fL – x curvesfor the Gd-containing alloys evaluated

here. The curves have all beenshifted so that the referencepoint is taken as distance fromthe solid/liquid interface wherefL = 1 (instead of distance fromthe heat source). This shift per-mits more direct comparison be-tween curves of differentnominal Gd concentrations. Theterminal fraction eutectic values(fe) calculated to form at theedge of the mushy zone for eachalloy are also noted in the figure.It should be noted that the vari-

ation in temperature with position calcu-lated through Equation 8 is not expectedto be highly accurate due to the assump-tions invoked to arrive at the Rosenthalsolution, most notably that the thermalproperties are assumed constant with tem-perature and the heat flow via convectionin the melt pool is ignored. However, theobjective here is to evaluate the influenceof nominal composition on the solid + liq-uid zone characteristics and resultantweldability under identical heat flow con-ditions. In view of this, it is only requiredto have an estimate of the functional formof the temperature variation with themushy zone. As demonstrated by the goodcomparison between experimental andcalculated values of ΔT and fe, the influ-ence of nominal composition on mushyzone characteristics can be accurately cap-tured with these equations. This is alsoconfirmed by noting that the experimen-tally determined maximum crack lengthsshown in Fig. 24 and the calculated mushyzone sizes shown in Fig. 28 are of similarsize. Solidification cracks are not expectedto propagate through the high tempera-ture region of the mushy zone where theliquid fraction is high because solid-solid

boundaries have not yet begun to form inthese regions, but rather are constrainedto a region within the mushy zone. Thus,the maximum crack length values shouldbe less than the size of the mushy region,and this trend is certainly observed by themaximum crack lengths shown in Fig. 27and the mushy zone sizes shown in Fig. 28.

More importantly, the calculated fL – xcurves, combined with the experimentalweldability results, aid in the developmentof a more detailed understanding on theinfluence of Gd on the mushy zone char-acteristics and resultant cracking suscepti-bility. At the lowest Gd concentration(0.46 wt-%), the mushy zone is the largestbecause the solidification temperaturerange is the widest. In addition, much ofthe mushy zone is occupied by relativelylow fL values. For example, the fractionliquid only varies from 0.03 at the edge ofthe mushy zone (at x ≈ 1.4 mm) to 0.07 ata distance of 0.5 mm behind the solid/liq-uid interface, (i.e., ~ 65% of the mushyzone is occupied by 0.03 < fL < 0.07). Thisliquid fraction range is high enough tocause moderate cracking, but too low forappreciable backfilling and healing of thecracks. The mushy zone size for the 1.01wt-% Gd alloy is slightly reduced due tothe smaller solidification temperaturerange. However, the higher amount of liq-uid present in the trailing edge of themushy zone aggravates cracking by pre-venting the formation of solid/solidboundaries, and the terminal fraction liq-uid value of 0.07 is too low to permit heal-ing of cracks by backfilling. This suggeststhat, at a terminal value of fe ≈ 0.07 andbelow, the liquid from the molten pool iscutoff from cracks that form in the mushyzone because the solid dendritic morphol-ogy is well developed.

T Tphx2p

ηπ

= + ⎛⎝⎜

⎞⎠⎟

fm C

T Tphx2

Ll o

Gd

o pηπ

=−

− +

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Fig. 27 — Influence of Gd concentration on the cracking sus-ceptibility, amount of terminal eutectic constituent, and solidi-fication temperature range in Alloy C-4.

Fig. 28 — Calculated fL - x curves for various Gd- containing C-4 alloys.


At Gd concentrations of 1.49 wt-% andhigher, the fL values never drop below 0.10in the mushy zone. Also note that the fLvalues do not decrease as substantiallywith increasing distance for these higherGd alloys. Each of these factors wouldhelp promote a more continuous liquidphase between the fully molten weld pooland the solid + liquid mushy zone, thuspermitting the backfilling that was ob-served experimentally. The decrease in so-lidification temperature range that occursdue to a reduction in the liquidus temper-ature also contributes to the improvedcracking susceptibility of the alloys withGd concentration higher than 1.49 wt-%.

Summary

The ever-increasing demand to pro-duce energy from a variety of ever-de-creasing resources has created the need todevelop new plants that use existing fuelsources more efficiently. This, in turn,places a heavy demand on developing newengineering alloys that can be utilized inaggressive conditions that cannot be toler-ated by existing alloy systems. As the alloysbecome more complex, it is likely that theywill also be subjected to more complicatedchanges in microstructure during the weldthermal cycle. In view of this, it is impor-tant for alloy producers, end users, and re-search organizations to collaborate so thatwelding technology can be developed inparallel with the discovery of new materi-als. This approach will permit strategicchanges to new alloys during the designstage, thus avoiding the potential for lim-ited use of newly developed alloys in ap-plications where welding is required.

It is important to note that many of theinferior properties of fusion welds are re-lated to the steep gradients in chemicalcomposition and microstructure that formacross the heat-affected zone and fusionzone. Thus, emphasis in future researchshould be placed on understanding, con-trolling, and where required, minimizingthese gradients through alloy and processcontrol. Computational modeling can playa key role in this area by accelerating the un-derstanding of microstructure and propertychanges that occur during welding. Whilesignificant progress is being made on thefront of microstructural modeling, muchwork is still needed to predict mechanicalproperties of welds from knowledge of themicrostructure. Similarly, techniques areneeded to determine long-term creep prop-erties from short-term tests. This area isparticularly challenging because the long-term creep properties are invariably con-trolled by microstructural changes thatoccur at very long times, and these changescan be difficult to properly simulate withshort-term tests.

Lastly, training of graduate-level engi-

neers is also vitally important to ensuresafe and reliable operation of weldedstructures that will be used in energy ap-plications. While physical metallurgyforms the basic discipline for understand-ing the complex phase transformationsthat occur during welding, the number ofengineers with basic skills in physical met-allurgy has declined over the years as thisform of training gets replaced with other“modern” materials topics in areas such asnanotechnology, biotechnology, etc. Whilethese fields are certainly important, it isequally important to ensure a sufficientlevel of graduate engineers are trained tomeet the demands of industrial fabricationand maintenance required in energy andother areas that rely so heavily on manu-facturing. Industry can play an importantrole in this area through support of grad-uate research programs in order to fill thegap from government funding that has de-clined over the years. A successful exam-ple of industry/university collaboration iscurrently represented by the National Sci-ence Foundation Center on IntegratedMaterials Joining Science for Energy Ap-plications. This center is a joint effortthrough four universities (Ohio State,Lehigh, Colorado School of Mines, andUniversity of Wisconsin-Madison) that issupported by both NSF and 31 membercompanies. Research at the center is fo-cused on a wide range of issues related tojoining for energy applications while si-multaneously supporting ~ 30 graduatestudents who are often hired by membercompanies. This type of industry/univer-sity collaboration is essential for continu-ing to meet the future needs of weldingand metallurgical engineers.

Acknowledgments

The author gratefully acknowledges fi-nancial support of this work through theNSF I/UCRC Center for Integrative Ma-terials Joining Science for Energy Appli-cations (CIMJSEA) under contract#IIP-1034703 and the U.S. Department ofEnergy, Assistant Secretary for Environ-mental Management, under DOE IdahoOperations Office Contract No. DE-AC07-99ID13727. The author is alsograteful to collaborators who made signif-icant contributions to the research pre-sented in this study, including Drs. CharlesRobino and Ron Mizia and graduate stu-dents Michael Minicozzi, Andrew Stock-dale, and Daniel Bechetti.

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