Microstructural Evolution of INCONEL Alloy 740H Fusion ... · Microstructural Evolution of INCONEL Alloy 740H Fusion ... in compliance with the ASME code case for alloy 740H ... the

Microstructural Evolution of INCONEL� Alloy 740H� FusionWelds During Creep

DANIEL H. BECHETTI, JOHN N. DUPONT, JOHN J. DE BARBADILLO,BRIAN A. BAKER, and MASASHI WATANABE

Electron microscopy techniques have been used to investigate the cause of premature creepfailure in the fusion zone of INCONEL� Alloy 740H� (INCONEL and 740H are registeredtrademarks of Special Metals Corporation) welds. The reduced creep rupture lives of all-weld-metal and cross-weld creep specimens (relative to base metal specimens) have been attributed tothe presence of large grain boundary regions that were denuded in fine c¢ but contained coarse,elongated particles. Investigation of creep rupture specimens has revealed four factors thatinfluence the formation of these coarsened zones, and the large particles found within them havebeen identified as c¢. Comparisons of the microstructural characteristics of these zones to thecharacteristics that are typical of denuded zones formed by a variety of mechanisms identified inthe literature have been made. It is concluded that the mechanism of c¢-denuded zone formationin alloy 740H is discontinuous coarsening of the c¢ phase. The discontinuous reaction is cata-lyzed by the grain boundary migration and sliding which occur during creep and likely pro-moted by the inhomogeneous weld metal microstructure that results from solute segregationduring solidification. The increased susceptibility to the formation of the observed c¢-denudedzones in the weld metal as compared to the base metal is discussed in the context of differencesin the contributions to the driving force for the discontinuous coarsening reaction.

DOI: 10.1007/s11661-014-2682-6� The Minerals, Metals & Materials Society and ASM International 2014

I. INTRODUCTION

PREDICTIONS by the International Energy Agencywhich assert that the global demand for electricitygeneration by burning coal will double by the year2030, coupled with growing environmental concernssurrounding the release of greenhouse gases by theburning of fossil fuels, has sparked initiatives in theUnited States, Europe, and Asia aimed at developing anew generation of coal-fired power plants.[1–3] Theseplants, termed Advanced Ultra-Supercritical (A-USC),are slated to operate at steam temperatures of 973 K to1033 K (700 �C to 760 �C) and steam pressures of 35 to45 MPa, both of which are higher than the operatingconditions in current generation plants. By doing so,significant plant efficiency gains and reductions in CO2

emission can be achieved.[4,5] However, these harsherconditions will require header materials with creeprupture and corrosion properties that exceed what isattainable by ferritic steels, stainless steels, and mostsolid solution strengthened Ni-based superalloys. Atten-tion has therefore shifted to the c¢ strengthened Ni-basedalloys for this application. INCONEL� Alloy 740H�,

which was developed by modifying Nimonic 263 forbetter coal ash corrosion performance, is a leadingcandidate material for A-USC boilers and has shown theability to survive the aforementioned operating condi-tions in a laboratory setting.[2,4,6–10] However, applica-tion in coal-fired boilers will invariably require weldedjoints, and alloy 740H weldments have been shown tounderperform the base material by up to 30 pct in creeprupture tests across a wide range of applied stresses andtemperatures.[11] This weld strength reduction (WSR) isin compliance with the ASME code case for alloy 740Hin A-USC applications,[12] which requires a WSR factorof 0.7 or larger. However, improvement of the WSRwould allow alloy 740H weldments to be loaded closer tothe base metal design stress and therefore realize themaximum efficiency gains from the A-USC cycle. Inorder to make such improvements, it is necessary to firstidentify the mechanism behind the strength reduction.Therefore, the objective of this research is to understandthe reduced creep rupture behavior and associatedmicrostructural evolution of alloy 740H fusion welds.Through this understanding, the cause(s) of prematureweld creep failure can be established and a basis forfuture material and/or processing modifications to mit-igate this behavior can be established.

II. EXPERIMENTAL

Creep rupture tests on alloy 740H base metal andmultipass weld samples were performed by SpecialMetals Corporation. Multipass weld specimens were

DANIEL H. BECHETTI, Research Assistant, JOHN N.DUPONT, Professor, and MASASHI WATANABE, AssociateProfessor, are with Lehigh University, Bethlehem, PA 18015. Contacte-mail: [email protected] JOHN J. DE BARBADILLO, Manager,Product & Process Development, and BRIAN A. BAKER, Productand Application Development Engineer, are with the Special MetalsCorporation, Huntington, WV 25705.

Manuscript submitted April 2, 2014.Article published online December 6, 2014

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 46A, FEBRUARY 2015—739

fabricated from a 30 deg V-groove (60 deg includedangle) manual gas tungsten arc weld (GTAW) made ona 15 mm thick alloy 740H plate at 198 A and 11 Vunder 75/25 Ar/He shielding gas at 51 m3/h using a3.2 mm diameter W-2 pct Th electrode. The root pass(4.8 mm root opening) of this weld was deposited withan average filler metal feed rate of 26 mm/s and averagetorch travel speed of 0.9 mm/s, while the fill passes weredeposited at an average feed rate of 19 mm/s andaverage torch travel speed of 3.4 mm/s. A total ofapproximately 33 passes were used to fill the groove,and a schematic illustration of the weld is given inFigure 1. All base plates were given the following heattreatment prior to welding: 1394 K/1 hour (1121 �C/1 hour) solution anneal, water quench, 1073 K/4 hours(800 �C/4 hours) age, air cool. After fabrication, allweldments were given a 1073 K/4 hours (800 �C/4 hours) aging treatment and air cooled. All specimenswere in the form of smooth round bars machined andtested to failure in accordance with ASTM E139[13] inlever-arm type creep assemblies. Samples containingonly base metal were 6.35 mm in diameter with a31.75 mm gage length. Specimens containing weld metalwere 6.35 mm in diameter with a 57.15 mm gage length.Select base metal and weld metal specimens tested in thelongitudinal (along the weld axis) and transverse (cross-weld) orientations were chosen for microstructuralanalysis, and a summary of their creep rupture behavioris given in Table I. Specimens 1 to 3 were taken from theroot, center, and cap of the multipass weld shown inFigure 1, respectively. Also, because they are cross-weldspecimens, samples 4 and 5 contained both weld metaland base metal in the gage length. An all base metalspecimen tested under the same conditions as sample 5was also characterized. All specimens containing weldmetal failed within the weld metal. Creep rupturespecimens were cross-sectioned using wire electricaldischarge machining (EDM) for metallographic ana-lysis. The compositions of the base metal, filler metal,

and as-deposited weld metal were determined using wetchemical and OES techniques and are given in Table II.Secondary electron and backscattered electron imag-

ing were performed in a Hitachi 4300SE/N Schottkyfield emission scanning electron microscope (SEM) at anoperating voltage of 15 keV. Given the composition ofthe alloy, Monte Carlo simulations using the CASINOprogram[14] predict an electron beam interaction depthof approximately 650 nm with lateral spatial resolutionof approximately 850 nm under these conditions. Elec-tron backscatter diffraction (EBSD) was performed inthe same instrument at 70 deg tilt and an acceleratingvoltage of 7 to 8 keV, using the EDAX-TSL OIM 5software package for data acquisition and analysis.After SEM characterization, a coarsened zone was

chosen for extraction and thinning to electron transpar-ency using focused ion beam (FIB) techniques in an FEIStrata DB235 with a Ga ion source operated at 30 keV.After FIB thinning, the TEM specimen was Ar ionmilled in a Fischione NanoMill 1040 at an acceleratingvoltage of 900 eV for 20 minutes to remove Ga iondamage. Final specimen thickness was approximately150 to 200 nm. Microstructural analysis of the thinnedspecimen was carried out in both the transmissionelectron microscope (TEM) and SEM. Imaging, selectedarea electron diffraction (SAED), and X-ray energydispersive spectrometry (XEDS) were performed using aJEOL 2000-FX TEM at an operating voltage of200 keV. The predicted lateral spatial resolutions forimaging/diffraction and X-ray analysis in this setup wereapproximately 15 and 40 nm, respectively.[14] AfterTEM analysis, the aforementioned Hitachi SEM wasused to acquire additional compositional informationfrom the FIB sample via XEDS line scans across themicrostructure. The specimen was held perpendicular tothe electron beam using a scanning transmission elec-tron microscopy (STEM) in SEM sample holder, asdescribed in the work of Patel and Watanabe.[15] Forthis analysis, an accelerating voltage of 30 keV wasused, resulting in a predicted X-ray spatial resolution ofapproximately 100 nm.Scheil solidification calculations were performed using

the ThermoCalc software package[16] and the TTNi7thermodynamic database[17] in order to assess thecompositional gradients present within the matrix (aus-tenite) immediately after weld metal solidification. TheScheil solidification simulation assumes a planar solid/liquid interface, no diffusion of substitutional elementsin the solid, complete mixing in the liquid, and equilib-rium at the solid/liquid interface. It therefore representsthe worst case of microsegregation for substitutional

Fig. 1—Schematic illustration of multipass gas tungsten arc weldused for creep tests and stress-free aging treatments.

Table I. Summary of Creep Specimens Analyzed in This Study

Specimen No. Specimen Type Test Temperature [K (�C)] Stress (MPa) Rupture Life (h)

1 all weld metal (root) 1023 (750) 280 270.92 all weld metal (midweld) 1023 (750) 250 546.13 all weld metal (cap) 1023 (750) 200 1312.04 cross-weld 1023 (750) 160 2434.15 cross-weld 1123 (850) 100 224.26 base metal 1123 (850) 100 744.5

740—VOLUME 46A, FEBRUARY 2015 METALLURGICAL AND MATERIALS TRANSACTIONS A

elements. In actual fusion welds, there can be reductionsin the extent of microsegregation due to solute enrich-ment at the dendrite tip, tip undercooling, and backdiffusion toward the dendrite core. However, backdiffusion of substitutional elements is generally notexpected under cooling rate conditions associated withwelding, and previous work has shown that tip enrich-ment and undercooling are small under the typicalwelding conditions considered here.[18] During solidifi-cation simulations, all phases included in the relevantdatabases were allowed to be active. P and S, which arenot included in the TTNi7 database, were excluded.Carbon was entered as a fast-diffusing component,which is justified based on previous results.[18] Followingsegregation profile calculations, equilibrium calculationswere performed to assess predicted changes in phasestability as a result of compositional differences frommicrosegregation. Only phases that have been ob-served[19,20] in alloy 740H (c, c¢, g, MC, M23C6) wereactive during equilibrium calculations to speed simula-tion time.In addition to the multipass weld described above,

Special Metals Corporation also fabricated a single passbead on plate GTAW on a 7.3 mm thick plate of alloy740H using 1.14 mm diameter alloy 740H filler metalwire. The weld was made using the same amperage,voltage, and shielding gas conditions as the multipassweld described above. This weld was fabricated usingthe same heat of material as the aforementioned creepspecimens and was given the same direct age treatmentafter welding. Stress-free aging was performed onsections of the single and multipass fusion welds at1023 K (750 �C) in air in a Lindberg tube furnace.Aging times were chosen to correspond to the creeprupture lives of four of the creep specimens that weretested at 1023 K (750 �C): 270, 540, 1312, and2434 hours. Samples were air cooled upon completionof aging.All metallographic specimens were prepared using

standard metallographic techniques and electrolyticallyetched at 6 V in a solution comprised of 20 mLH3PO4 and 150 mL H2SO4 saturated with CrO3.

III. RESULTS AND DISCUSSION

Figure 2 shows a subset of the available creep data foralloy 740H.[21] It presents the results for base metal andweld metal specimens tested at 1023 K (750 �C). Con-tained within this figure are the four creep specimens inTable I that were tested at 1023 K (750 �C). As shown,the alloy 740H weld metal displays a large reduction increep rupture life as compared to the base metal acrossall applied stresses.

A. Morphology of Grain Boundary c¢ Denuded Zones inAlloy 740H

Microstructural investigation of the specimens givenin Table I revealed significant differences in the mor-phology and distribution of the strengthening c¢ phase inthe base metal and weld metal, especially near the grain

Table

II.

INCONEL

�Alloy740H

�Compositions

Technique

Material

Ni

Cr

Co

Nb

Ti

Al

Mo

Fe

Si

CMn

VW

Zr

Ta

PCu

S

Wet

chem

ical/OESanalysis

base

metal

49.17

24.35

20.08

1.53

1.45

1.28

0.53

1.07

0.20

0.05

0.30

0.007

0.008

0.020

<0.001

0.002

<0.001

—filler

metal

50.23

23.88

19.52

1.50

1.31

1.36

0.52

1.04

0.21

0.05

0.30

——

0.029

<0.0001

0.005

—<0.01

as-deposited

weldmetal

50.20

23.90

19.40

1.52

1.28

1.31

0.54

1.10

0.22

0.05

0.29

0.006

0.013

0.018

<0.001

0.0022

0.105

0.0013

Values

are

inwtpct.


boundaries. Figures 3(a) and (b) show the as-agedmicrostructures of the base metal and weld metalinvestigated in this study, and Figures 3(c) through (f)demonstrate the evolution of these microstructuresduring creep. The as-aged micrograph from the weldmetal was taken mid-cross-section and is representativeof the microstructures in the cap and root passes, exceptfor some variations in precipitate size due to thereheating caused by multiple weld passes. As shown,there exists an even distribution of fine (<200 nm) c¢ upto and including the grain boundaries in the base metal(Figures 3(a), (c), and (e)) in both the as-aged and creptconditions. The weld metal, however, displays aninhomogeneous distribution of fine c¢ in the as-agedcondition (Figure 3(b)) and a large number of grainboundaries characterized by 1 to 3 lm wide regions thatare devoid of the fine c¢ but contain coarse, typicallyelongated second phase particles after creep(Figures 3(d) and (f)). Grain boundary regions ofsimilar morphology have been identified in the literatureas precipitate-free zones (PFZs).[22–64] For comparativepurposes, and because PFZs are often not truly free ofsecond phases, ‘‘denuded zone’’ will be used here as ageneral term for these regions, and ‘‘coarsened zone’’(CZ) will be used to specifically distinguish the grainboundary features in alloy 740H. Figure 3(f) alsodemonstrates the typical accumulation of creep damagealong and within the CZs in alloy 740H. Creep voidsappear to nucleate preferentially in the vicinity of thelarge second phase particles, and all of the observedcreep cracks in the series of specimens examined in thisstudy were propagated along grain boundaries contain-ing coarsened zones. To confirm that the voiding at theweld metal grain boundaries is in fact due to creepdamage accumulation rather than effects from metallo-graphic preparation and etching, micrographs of creepspecimen 4 in the as-polished and etched conditionswere acquired and compared. Figure 4(a) shows atypical large weld metal creep crack in the gage lengthof this specimen in the as-polished condition.Figure 4(b) shows the same region after etching. Note

that only minor changes in the shape and size of thevoids have occurred. These are most likely due to eithermatrix removal exposing wider, subsurface portions ofthe defects or an edge effect whereby the electrolyticreaction is intensified at the perimeter of the voids.Further evidence that the presence of the large voids isdue exclusively to the CZs is shown in the highermagnification micrograph of Figure 4(c), which demon-strates the impending convergence of two cracks alongthe 2 lm wide c¢ denuded region. It is unlikely thatetching would preferentially attack only a portion of thegrain boundary that was denuded in c¢ and leave such along section unaffected. Also, since the narrower coars-ened region in the lower left of Figure 4(c) does notexhibit evidence of attack, it can be concluded that whilethe electrolytic etching process used in this study mayslightly increase the size of the voids/cracks within thesample for the reasons given above, it does not createthe voids.The clear preference for creep cracks to run along

boundaries containing CZs indicates that they aresignificant damage accumulation zones, which is inagreement with similar behavior associated with grainboundary denuded zones as reported by otherauthors.[23–27] As such, it is believed that the reductionin creep performance observed in alloy 740H weld metalis due to the presence of these micron-scale grainboundary CZs.

B. Factors That Influence the Formation of CoarsenedZones in Alloy 740H

The creep specimens given in Table I, as well assamples that were aged under stress-free conditions atidentical times and temperatures to the creep tests, wereexamined in order to understand the influence of startingmicrostructure (i.e., base metal vs fusion zone), stress,time, and temperature on the development of grainboundary coarsened zones in alloy 740H. First, fromFigure 3, which compares base metal and weld metalfrom the same cross-weld sample, it is shown that alloy740H weld metal is significantly more susceptible tocoarsened zone formation than the base metal. In cross-weld creep specimens, the weld metal has been observedto contain many CZs, while the base metal within thegage length of the same specimens shows little evidenceof CZ formation. Secondly, it appears that applied stressplays a significant role in the formation of CZs in alloy740H. As shown in Figures 5(a) and (b), weld metaltested in creep with an applied stress of 200 MPa(specimen 3 from Table I) contains large coarsenedzones, while a section of the same multipass weldexposed to a stress-free age at the same temperatureand exposure time shows little evidence of PZ formation.Next, exposure time during creep also influences theformation of CZs in alloy 740H. As given in Figures 5(c)and (d), in the base metal of a cross-weld specimen testedat 1123 K (850 �C) with 100 MPa applied stress (spec-imen 5 from Table I), no evidence of CZ formation isseen after 224.2 hours. However, some micron-scale c¢denuded zones have developed in an all base metalspecimen after 744.5 hours of exposure to the same

Fig. 2—Select creep data for alloy 740H base metal and weld metaltested at 1023 K (750 �C).


temperature and stress (specimen 6 from Table I). Notonly does this illustrate the effect of time on CZformation, but it also indicates that alloy 740H basemetal is not immune to their formation. Finally, thecreep temperature was identified as a key factor in theformation of CZs in alloy 740H. As shown inFigures 5(e) and (f), the base metal in a cross-weldspecimen (specimen 4 from Table I) tested at 1023 K(750 �C) does not exhibit significant evidence of CZformation, while an all base metal sample (specimen 6

from Table I) tested 1123 K (850 �C) clearly containsthem. This occurs despite the lower temperature speci-men being subject to a higher applied stress and longerexposure time.These observations qualitatively support the idea that the

operativemechanismofCZ formation in alloy 740H is likelydriven by stress-assisted diffusion. The temperature and timedependence of CZ formation, coupled with the influence ofstress and apparent exacerbation by the inhomogeneousweld metal microstructure, support this assertion.

Fig. 3—SEM micrographs showing the typical morphology and distribution of c¢ in alloy 740H: (a, b) as-aged, (c through f) after creep rupture.All images after creep rupture acquired from creep specimen 4 in Table I.


C. Mechanisms of c¢ Denuded Zone Formation

Five different formation mechanisms for regions withsimilar morphologies to the CZs in alloy 740H havebeen observed and reported for a wide variety of alloysystems over the past 60 years.[22–64] Each one producesa specific set of microstructural characteristics that canbe used to identify the operable mechanism of CZformation in alloy 740H.

1. Nabarro–Herring creepFormation of grain boundary denuded zones by N–H

creep follows the traditional description of N–H masstransfer. It involves diffusion from grain boundariesoriented parallel to the maximum tensile stress throughthe grain interiors to grain boundaries oriented perpen-dicular to that stress. This mass flux has been linked tothe dissolution of precipitates at the perpendicular grainboundaries and a concomitant increase in precipitatevolume fraction at the parallel boundaries in a variety ofalloy systems.[28–32] Although the validity of this mech-anism has been debated,[32–38] it has received enoughattention in the literature to warrant a discussion of itspossible applicability to the formation of CZs in alloy740H. With regard to Ni-based alloys, Tien and Gam-ble[31] proposed that for Ni-16Cr-5Al-4Ta, movement ofCr toward boundaries oriented perpendicular to themaximum tensile stress during N–H creep can cause localenrichment of Cr around these boundaries, leading todissolution of the c¢ phase. Unfortunately, this explana-tion is not entirely convincing because Cr usually acts toincrease the stability of c¢.[39,40] It is more likely that someother mechanism was contributing to the observedreduction in c¢ at the grain boundaries (for example,grain boundary sliding, as described below and arguedby Burton and Reynolds[37]). In any case, proponents ofN–H creep as a mechanism of denuded zone formationgenerally present that their microstructural characteris-tics should include symmetry of the denuded zone aboutthe grain boundary and a clear orientation dependence,whereby they are all aligned perpendicular to the appliedtensile creep stress.

Perhaps the most important consideration when eval-uating N–H creep as a possible source of the CZs in alloy

740H is the fact that N–H creep is typically onlyconsidered to be an active deformation mechanism attemperatures greater than approximately 90 pct of analloy’s absolute melting temperature. Considering thatthe low end of the melting range for alloy 740H is 1561 K(1288 �C) and the highest experimentally applied creeptemperature was 1123 K (850 �C), it can be shown thatthe most severe creep conditions occur at only 72 pct ofthe melting temperature of alloy 740H. In addition,Figures 6 and 7 demonstrate that the microstructuralfeatures of the CZs observed in alloy 740H are notconsistent with N–H creep. The SEM micrographs ofFigure 6 typify the lack of CZ orientation dependencewith respect to the tensile loading axis (horizontal, asindicated) that was observed in all the creep samples.While some CZ containing boundaries are orientedalmost normal to the loading direction, the majority arenot, and many such boundaries are nearly parallel to thetensile axis. It is recognized that the true geometry ofthese boundaries cannot be assessed by a two-dimen-sional section alone, but evidence shown to date by otherresearchers shows that when N–H creep is the activedenuded zone formation mechanism, there is a very clearpreference for formation on boundaries oriented perpen-dicular to the applied tensile load, even in planarmicrographs.[28,29,36] Figure 7 demonstrates that CZs inalloy 740H are asymmetric about the grain boundaries,i.e., they grow preferentially into only one of the grains.The BSE image of a grain boundary with a CZ demon-strates the coarse second phase within the CZ and thesurrounding fine distribution of c¢, both of which appeardark due to a lower weighted average atomic numberthan the matrix. As shown, EBSD patterns acquired atpoints B (within the CZ) and C (in the interior of grain 2)in the BSE image are identical, implying that they are ofthe same orientation and therefore share the same grain.The EBSD pattern acquired at point A (in the interior ofgrain 1) reflects a markedly different orientation, whichindicates that the actual location of the grain boundary inthe BSE image shown in Figure 7 is at the upperboundary of the CZ, between points A and B. Thus, itis asymmetric about the boundary. Given this evidence, itis concluded that N–H creep is not the active mechanismof CZ formation in alloy 740H.

Fig. 4—SEM micrographs of creep voids in the weld metal of specimen 5. (a) As-polished, (b) etched, and (c) higher magnification micrographof boxed region in (b).


2. Grain boundary carbide formationThe second mechanism relevant to the formation of

denuded zones involves c¢ former depletion by MCcarbide formation. This mechanism involves growth of(Nb,Ti)C carbides at the grain boundaries at the expenseof the local c¢. It can occur because these MC-typecarbides can have greater high temperature stabilitythan c¢. The removal of Nb and Ti, which are directlyinvolved in c¢ formation, from the grain boundary as thecarbides grow results in the development of a denudedzone. This mechanism has been suggested for the nickel-based alloy PE16 by Maldonado and Nembach.[41]

Scheil solidification simulations and thermodynamiccalculations were conducted to determine the stability ofthe MC phase within the alloy and to assess whether thismechanism is possible in the fusion zone of alloy 740Hwelds. Figure 8 shows the composition profile of thematrix (austenite) across a half dendrite after non-equilibrium solidification as predicted by the Thermo-Calc Scheil solidification module[16] using the TTNi7thermodynamic database.[17] As shown, significant mi-crosegregation of the major alloying elements, as well asthe c¢ formers Nb and Ti, is predicted. Recent results[65]

have verified the accuracy of these calculations via

Fig. 5—Influence of applied stress, time at temperature, and exposure temperature on evolution of CZs in alloy 740H.


experimental XEDS measurements across a series ofparallel dendrites in as-welded alloy 740H. This micro-segregation will affect the stability of phases at the creeptemperatures, so equilibrium calculations were con-ducted at various locations along the simulated halfdendrite to determine the phases that are likely to bestable after long-term high temperature exposure. Asgiven in Table III, the microsegregation induced duringsolidification should promote the stability of MC, c¢,and g at the interdendritic regions during exposure totemperatures between 1023 K to 1123 K (750 �C to850 �C). The predicted MC carbides are of the(Nb,Ti)C-type with an Nb:Ti ratio that is dependentupon temperature and the local composition (whichvaries as in Figure 8). Calculated composition ranges for

the MC phase are given in Table IV. As shown, it ispossible for carbides that are rich in the c¢ formers toevolve at the grain boundaries within alloy 740H weldmetal, which indicates that the mechanism of CZformation by c¢ former depletion is also possible. Notethat, while g is also predicted to be thermodynamicallystable at the interdendritic regions and has been linkedto denuded zone formation,[8] it has not been experi-mentally observed in alloy 740H due to the low Nbcontent and the slow kinetics of g formation in this alloysystem.[20] It is also worth noting that the microsegre-gation discussed here a significant effect on the generalprecipitation behavior of c¢. Earlier experimental workhas shown that microsegregation during solidification ofalloy 740H can cause the interdendritic regions to be

Fig. 6—SEM micrographs demonstrating the lack of an observed orientation relationship between CZs and direction of applied tensile stress increep specimens 2 (left) and 5 (right).

Fig. 7—Backscattered electron image of a CZ in creep specimen 4 (top) and EBSD patterns taken from the indicated locations demonstratingasymmetric nature of CZs (bottom).


enriched in the c¢ forming elements by up to 5 times thenominal composition before aging.[65] As shown inFigures 4(b) and (c), this results in inhomogeneousprecipitation of c¢ during aging of the weld metal.

Denuded zones formed due to (Nb,Ti)C are expectedto be symmetric about their grain boundaries (contraryto the experimental evidence given in Figure 7) and tocontain (Nb,Ti)C particles. Figure 9 shows the locationof extraction (using secondary electron imaging in theFIB) and the resulting cross-section (using the TEM) ofa specimen extracted from a CZ in specimen 4 fromTable I. A typical XEDS spectrum acquired from thelarge particles within the CZ of this specimen, as shownin Figure 10(a), indicates that they are not only enrichedin Nb and Ti, but are also enriched in Ni and Al. Thesurrounding matrix is depleted in Nb, Ti, and Alcompared to the particles, as shown in Figure 10(b).The calculated compositions of the MC phase(Table IV) indicate negligible solid solubility of Al andNi in MC. Therefore, experimentally observed compo-sition of the particles within the CZ is not consistentwith that of (Nb,Ti)C. In fact, these compositional dataare more consistent with c¢, Ni3(Ti,Al,Nb). This pointwill be discussed in detail with regard to discontinuousreactions as a mechanism of denuded zone formation ina subsequent section. Given these calculated and exper-imental results, it is concluded that the primary mech-anism of CZ formation in alloy 740H weld metal is notMC carbide formation.

Table III. Calculated Mole Fraction of Stable Phases Along a Half Dendrite Calculated Using the Composition Sets Indicated inFig. 8

Temperature [K (�C)]

Fraction Solid

0 0.50 0.80 0.95 0.99

1023 (750) c—0.907 c—0.871 c—0.773 c—0.669 c—0.457c¢—0.087 c¢—0.123 c¢—0.221 c¢—0.248 c¢—0.366M23C6—0.006 M23C6—0.006 M23C6—0.006 g—0.077 g—0.174

M23C6—0.006 MC—0.003

1073 (800) c—0.925 c—0.887 c—0.785 c—0.691 c—0.466c¢—0.069 c¢—0.107 c¢—0.209 c¢—0.215 c¢—0.329M23C6—0.006 M23C6—0.006 M23C6—0.006 g—0.092 g—0.020

MC—0.002 MC—0.003

1123 (850) c—0.949 c—0.909 c—0.812 c—0.707 c—0.479c¢—0.045 c¢—0.084 c¢—0.185 c¢—0.173 c¢—0.292M23C6—0.005 M23C6—0.005 MC–0.003 g—0.117 g—0.226MC—0.0001 MC—0.0002 MC—0.003 MC—0.003

Table IV. Calculated MC Carbide Composition Ranges in alloy 740H, Generated from the Simulation Results Given in Table III

1023 K (750 �C) 1073 K (800 �C) 1123 K (850 �C)

Nb 89 86 to 89 64 to 88Ti 0.9 1 to 4 0 to 22C 10 9 to 10 10 to 13Cr 9 9 10�5 0.001 to 0.03 0.01 to 0.03Ni 5 9 10�7 9 9 10�7 to 2 9 10�6 1 9 10�4 to 1 9 10�3

Co 7 9 10�7 1 9 10�4 to 3 9 10�4 4 9 10�6 to 2 9 10�5

Al 3 9 10�13 2.6 9 10�13 to 2.7 9 10�13 5 9 10�13 to 1 9 10�12

Values are in wt pct.

Fig. 8—Calculated variation in the concentration of alloying ele-ments in austenite across a half dendrite during solidification.Dashed lines represent composition sets chosen for equilibriumphase stability calculations.


A third mechanism of denuded zone formationidentified in the literature is that of localized depletionof Cr near the grain boundaries due to the growth of Cr-rich grain boundary M23C6-type carbides. This mecha-nism is similar to the previously described mechanism of(Nb,Ti)C formation, but involves the formation ofdenuded zones by removal of Cr, which plays a role instabilizing c¢ but is not directly involved in its formation.It has been proposed that such Cr depletion increasesthe solubility of Al, Ti, and Nb in austenite, which inturn leads to dissolution of the c¢ phase. This mechanismhas been proposed for the Nimonic PE16 system, amongothers.[42,43]

To assess whether this mechanism is possible in thealloy 740H system, a series of thermodynamic simula-tions were performed. Starting with the nominal alloy740H composition given in Table II, a series ofisopleths with Cr content ranging from the nominal25 wt pct down to 5 wt pct were generated. An

increase in Ni content accommodated for the changein Cr content. Figure 11 shows the influence of Cr onthe solubility of Al, Ti, and Nb in austenite. Theshaded regions correspond to the composition andtemperature space within which c and c¢ are stabletogether. Correspondingly, the lines that define theupper boundaries of the shaded regions are thecalculated c¢ solvus lines for the modeled systems.Note that the stability of c¢ in this system is predictedto decrease, which implies that the solubility of the c¢formers in austenite is expected to increase, withdecreasing Cr concentration. Thus, localized Cr deple-tion should promote c¢ dissolution in this alloy.Microstructural evidence that this mechanism is activeincludes a symmetric denuded zone that is depleted inCr and a distribution of Cr carbides along the grainboundary itself.Again, review of the microstructural observations in

Figure 7 demonstrates that the CZs in this alloy are

Fig. 9—(a) Region of interest for TEM sample extraction from creep specimen 4 and (b) bright-field image of resulting TEM specimen.

Fig. 10—Typical XEDS spectra of (a) large, elongated particle within CZ in alloy 740H weld metal and (b) the CZ matrix surrounding the largeparticle in (a), acquired in the TEM.


asymmetric about their associated grain boundaries. Inaddition, depletion of Cr in the matrix of the CZcompared with the matrix in a grain interior can beassessed by analysis of XEDS line scans taken across theCZ boundary. Two such line scans were acquired viaSTEM in SEM techniques, at the locations indicated inFigure 12(a). Care was taken to avoid sampling from thefinely dispersed c¢ particles, but they could not be avoidedcompletely. A qualitative comparison of the Cr contentin the two regions can be made by taking the FWHM CrKa/Ni Ka intensity ratio at each point, as given inFigure 12(b). The reported values are average intensityratios at each point, with the 95 pct confidence intervalprovided. From the plot, it is shown that the matrix Crconcentrations in the CZ and two-phase grain are verysimilar. Note that there is also no significant change in Ti(a c¢ former) content across the boundary. This will beaddressed in a following section. Finally, as shown inFigure 10(a), the elongated particles within the CZs aredepleted in Cr and enriched in the elements involved in c¢formation. All of these observations are inconsistent withthe microstructural characteristics that are expected to bepresent if Cr-rich M23C6 formation is the active mech-anism of CZ development in alloy 740H.

3. Grain boundary sliding/grain boundary migrationGrain boundary sliding (GBS) and grain boundary

migration (GBM) during creep are characterized by themovement of grain boundaries away from their originallocations because of local shearing and dislocationpileup. This occurs because it is more energeticallyfavorable for one grain to maintain strain compatibilitywith its neighbor by migration of the grain boundarythan by the formation of a creep void.[28] The mechan-ical shearing of this grain boundary sliding/migratingand dislocation pileup at the grain boundaries duringcreep can both contribute to an increased grain bound-ary vacancy concentration, which can lead to localizedprecipitate dissolution. Denuded zone formation byGBS/GBM has been proposed by a variety ofauthors for magnesium- and nickel-based alloysystems.[29,35,44–48] Based on their findings, microstruc-tural expectations for denuded zones formed by theGBS/GBM mechanism include a lack of orientationdependence with respect to the axis of applied stress,asymmetry about their associated grain boundary, and,if all of the precipitates in the wake of the migratingboundaries do not dissolve, a distribution of coarse,equiaxed precipitates should be found along them.

Fig. 12—(a) STEM in SEM image of the TEM sample shown in Fig. 9, with regions of interest for XEDS line scans across the CZ reactionfront indicated and (b) XEDS line scan data.

Fig. 11—Calculated variation in c¢ former solubility within austenite matrix and c¢ stability with changing alloy 740H Cr content. Values associ-ated with shaded regions are in units of wt pct Cr. Ni content is increased to accommodate changing Cr contents below the nominal value of25 wt pct Cr.


The microstructural evidence gathered in this studyseems to be moderately consistent with the predictionsof the GBS/GBM mechanism. As shown in Figures 6and 7, the CZs in alloy 740H are asymmetric and do nothave a preferred orientation with respect to the appliedstress. However, the elongated morphology of theparticles within the CZs of alloy 740H is generally notconsistent with the equiaxed morphology of particlesthat remain within denuded zones formed purely byGBS/GBM.[29] Also, the mechanism of GBS/GBMwould not explain the apparently greater propensityfor CZ formation in the weld metal of alloy 740H. Asshown above, the weld metal in a cross-weld creepspecimen is significantly more prone to CZ developmentthan the base metal in the same specimen. Since bothbase metal and weld metal are contained within the gagelength of that specimen, they would both be subject tothe action of grain boundary sliding/migration as adeformation mechanism and would therefore displaycomparable levels of CZ formation if GBS or GBM wasthe primary mechanism of their evolution. For thesereasons, it is believed that GBS/GBM is not the primarymechanism of CZ formation in alloy 740H fusion welds.However, it is important to note that the movement ofgrain boundaries is essential for the activation of thefinal possible mechanism of denuded zone formation,described in the following section. It is thereforeproposed that GBS/GBM may play an important butnon-primary role in CZ formation in alloy 740H.

4. Discontinuous precipitation (DP) and discontinuouscoarsening (DC)

These mechanisms rely on the enhanced diffusivity ata moving grain boundary to promote acceleratedgrowth of a particle which is coherent with one of thegrains and incoherent with the other. Discontinuousreactions must involve a moving grain boundary,although whether the moving boundary causes thediscontinuous reaction or vice versa is situation andsystem dependent.[49] It has been reported that thesereactions are often forced to occur when precipitationand/or coarsening are thermodynamically favored, butbulk diffusion is too slow to produce the desiredmicrostructural change.[50] Thus, the reactions takeadvantage of the enhanced grain boundary diffusivityto accelerate precipitation and/or coarsening. DPinvolves the transformation of a single-phase into atwo-phase lamellar structure, while DC involves thedirectional coarsening of a phase that has alreadyprecipitated (either continuously or discontinuously).[51]

It is important to note here that DC is not dependentupon prior DP, and that it can occur from a micro-structure subject only to prior uniform precipitation.[52]

DP and DC have been identified in a wide range of alloysystems, and discontinuous reactions involving of c¢ orgrain boundary carbides (such as M23C6) have beenpresented as the operative denuded zone formationmechanism in nickel-based alloy systems such as alloy800 and alloy 738.[53–64] Also, reactions involving astrengthening phase (e.g., d¢ in Al-Li and c¢ in superal-loys) have been identified as one of the main types ofdiscontinuous transformations.[49] Typical microstruc-

tural characteristics of this phenomenon include adenuded zone that does not maintain orientationdependence with applied stress and contains a distribu-tion of coarse particles that are elongated in the growthdirection, typically forming a lamellar morphology withthe matrix. Discontinuous reaction products can beeither symmetric or asymmetric about the grain bound-aries, but the symmetric (so-called ‘‘double seam’’)morphology typically only arises when large fractionsof the microstructure are subject to the discontinuousreaction. Interfacial energy, coherency strain at theparticle–matrix interface, local chemistry, and bulkstrain can contribute to the driving force for increasingparticle size.Figure 13(a) shows an SAED pattern collected from

the CZ shown in the bright-field TEM image ofFigure 9, acquired with the specimen tilted to a h100imatrix zone axis. This pattern contains primary diffrac-tion spots from the austenite matrix as well as a largenumber of superlattice reflections. In order to elucidatewhich phase(s) were generating the superlattice reflec-tions, SAED patterns for a variety of different phasesassociated with the alloy 740H system were simulatedusing the Web Electron Microscopy Applications Soft-ware (WebEMAPS).[66] Figure 13(b) presents the simu-lated selected area diffraction pattern generated bycombining the h100i zone axis patterns from austenite(lattice parameter a = 3.60 A),[19] c¢ (a = 3.57 A),[19]

and Cr23C6 (a = 10.6 A).[20] As shown, Figure 13(b) isidentical to the experimentally acquired diffractionpattern within the alloy 740H CZ (Figure 13(a)). Thus,the observed morphology of the CZs in alloy 740H weldmetal creep specimens is characterized by an austenitematrix that contains elongated particles that are iden-tified as c¢ through both diffraction analysis (Figure 13)and compositional analysis (Figure 10). In the case ofthe particular CZ analyzed in Figure 9, a primaryCr23C6 carbide was captured within the selected areaaperture, thus giving rise to the superlattice reflections inFigure 13 that cannot be attributed to c¢. This primarycarbide, whose Cr-enriched XEDS spectrum is given inFigure 14, would be expected to form along the originalgrain boundary. Therefore, its location in the lowerright corner of the bright-field image in Figure 9(b)coincides with the original position of the highlightedgrain boundary and gives an indication (though only atwo-dimensional one) of the growth direction of the CZ.Finally, Figure 4 demonstrates the continuous (albeit

inhomogeneous) c¢ precipitation in as-heat treated alloy740H, which has transformed all of the single-phaseaustenite parent material into two-phase austenite +c¢.The completion of continuous precipitation within thealloy, combined with the lack of experimentally ob-served grain boundary CZs in this as-heat treatedcondition, implies that they did not form upon initialprecipitation of the c¢ phase and therefore discountsdiscontinuous precipitation as the mechanism of CZformation in alloy 740H.The combination of these experimental observations

with the CZ asymmetry and lack of an orientationdependence points to discontinuous coarsening as theprimary mechanism of CZ formation in alloy 740H weld


metal during creep. Given the conditions to which thespecimens investigated here were exposed, it is mostlikely that the grain boundary c¢ formed during contin-uous precipitation were driven to their elongated shapevia a DC reaction, which was catalyzed by grainboundary movement during creep.

In addition to the microstructural characteristicsdescribed above, discontinuous reactions are oftenfound to leave discontinuous concentration profiles intheir wake, i.e., they create an abrupt change in matrixcomposition as the reaction front is crossed.[49,61,62]

However, the literature typically presents such concen-tration profiles in the as-reacted case immediately after adiscontinuous reaction has taken place, without consid-ering the effect of diffusion that may take place once adiscontinuously formed colony has reached an equilib-rium size during high temperature exposure (duringcreep, for example).[62] Furthermore, most studies thathave assessed discontinuous concentration profiles after

thermal exposure have done so on Al or Mg alloysexposed to low temperature annealing in the range of400 K to 500 K (127 �C to 227 �C) for short durations(20 minutes or less).[63,64] One study performed on a Ni-Cu system was annealed at 1023 K (750 �C) for 6 hours,and while this system was reported to maintain asignificant concentration gradient across the discontin-uous reaction front, compositional scans indicated thatthis was in the form of a peak rather than a disconti-nuity and that the maximum concentration was subjectto reduction with time due to redistribution by latticediffusion.[67] Neither of these cases is completely repre-sentative of thermal exposure of the alloy 740H systemat 1023 K to 1123 K (750 �C to 850 �C) for up to2434 hours, so the DICTRA software package,[16] inconjunction with the TTNi7 thermodynamic data-base[17] and the MOB2 mobility database,[68] was usedto assess the likelihood of a discontinuous concentrationprofile in the matrix persisting after such exposure. Theinitial discontinuous concentration profile (at t = 0)that was set up for this simulation is given inFigure 15(a). The composition for the ‘‘two-phasegrain’’ was chosen by simulating the standard alloy740H aging treatment on the composition associatedwith the core of a dendrite in the weld metal (shown inFigure 8). The composition for the ‘‘CZ’’ was chosen byrepeating the same calculation on the composition setassociated with an interdendritic region in the weldmetal (also shown in Figure 8). The region over whichthe simulation was performed has a total width of 4 lm(2 lm of CZ adjacent to 2 lm of a two-phase grain) andis intended to reflect homogenization of one-half of thewidth of a CZ. Note that the chosen width is larger thanthe typical half-width of an experimentally observed CZ(0.5 to 1 lm), so the model represents a conservativeprediction of the ability of the associated concentrationprofile to homogenize. Figure 15(b) shows the matrixcomposition in equilibrium with c¢ after a simulated2434 hours isothermal hold at 1023 K (750 �C). As

Fig. 14—XEDS spectrum of primary carbide located in the bottomright corner of the bright-field TEM micrograph presented inFig. 9(b).

Fig. 13—(a) Selected area diffraction pattern of the CZ shown in Fig. 9, taken along a h100i zone axis. Primary reflections from the austenitematrix, superlattice reflections from cross-sectioned elongated c¢ particles, and superlattice reflections from a Cr carbide in the field of view areconfirmed by (b) the simulated diffraction pattern overlay from these three phases aligned to a h100i zone axis.


shown, there is predicted to be significant homogeniza-tion of the initial concentration gradient, such that thelargest concentration variation for any element is lessthan 1 wt pct. This demonstrates that homogenizationof a discontinuous concentration profile created by theformation of a denuded zone by DC is possible at thetemperature and duration associated with the creep testsanalyzed in this study.

Experimental validation of this assertion is given bythe XEDS line scans shown in Figure 12(b). Asdescribed above, this evidence suggests that there islittle variation in the matrix composition across the CZboundary with respect to both a major alloying elementthat can influence c¢ and carbide stability (Cr) and a c¢former (Ti). This seemingly validates the calculatedresults described above and lending credence to the ideathat discontinuous concentration profiles may homog-enize during long-term thermal exposure.

5. Driving force for DC in alloy 740H weld metalWith the mechanism of DC in mind, it is pertinent to

consider the major driving forces for the development ofCZs in alloy 740H. The components for the thermody-namic driving force for discontinuous coarsening and itsassociated boundary migration have been identified asfollows[49,52,61]:

DGDC ¼ DGcDC þ DGe

DC þ DGchemDC þ DGc

DC; ½1�

where DGDC is the Gibbs free energy change per unitvolume associated with the discontinuous coarsening(DC) reaction, and the superscripts on each of itscomponents refer to the precipitate-matrix coherencystrain (c), bulk strain (e), chemical (chem), and interfa-cial (c) energy contributions to the driving force.

It is well known that the c/c¢ lattice mismatch in Ni-based superalloys is small,[22,40] and it has been reportedelsewhere that the mismatch in alloy 740H is as low as[19]

0.70 pct. The low coherency strain due to near matchingof the matrix and precipitate lattice parameters, incombination with the minimal evidence in the literaturethat DGc

DC provides anything other than a minorcontribution to the driving force for DC, leads toassumption here that the free energy contribution toDC from precipitate-matrix coherency strain isnegligible.[52,61,69]

The bulk strain component of the driving force forDC may not be neglected from consideration in thespecimens examined in this study, as is often the case inthe literature where annealed samples or bicrystals areinvestigated. This is because there are significant sourcesof plastic strain in this work, namely creep deformationand residual stress from welding.[70–72] Residual plasticstrain from welding extends into the base metal,[71,72]

and cross-weld creep specimens which subjected bothweld metal and base metal to plastic deformation underthe same conditions were analyzed here. Thus, consid-ering the minimal evidence of CZ formation in the basemetal of the cross-weld specimen shown in Figure 3, it isunlikely that the bulk strain itself is the main drivingforce for CZ formation in alloy 740H. However, theeventual development of CZs in the base metal at longertimes and therefore larger creep strains (Figure 5(d))indicates that the bulk plastic strain does affect theextent to which the DC reaction occurs. The bulk strainis also relevant in that creep deformation by GBSprovides moving grain boundaries which are necessaryfor and promote DC over traditional coarsening.As discussed above and demonstrated in Figure 8,

significant compositional gradients are induced during

Fig. 15—DICTRA simulation of homogenization of a discontinuous concentration profile at a CZ reaction front: (a) initial austenite matrixcomposition profile generated by simulating an 1073 K/4 h (800 �C/4 h) precipitation heat treatment on the calculated dendrite core and inter-dendritic composition sets shown in Fig. 8, and (b) calculated evolution of the discontinuous concentration profile shown in (a) after 2434 h at1023 K (750 �C), the length of the creep test associated with the TEM specimen shown in Fig. 9(b).


weld metal solidification, which causes the austenitematrix to be supersaturated with solute and couldsignificantly affect the driving force for DC given inEq. [1]. Ju and Fornelle[73] described this idea anddeveloped an expression for the DC driving force thatincludes a term that is dependent on the matrixsupersaturation, DCo (wt pct), which is defined asfollows:

DCo ¼ xav � xe; ½2�

where xav is the actual composition of the matrix(wt pct), and xe is the equilibrium composition of thematrix (wt pct). Their derivation, however, was per-formed under the assumption that DC was occurringafter DP, which does not appear to be the case in alloy740H weldments. Still, it is intuitive that any matrixsupersaturation, say, that which remains after continu-ous precipitation during the direct aging treatment,would promote DC by providing some of the solutenecessary for coarsening. To address this, the degree ofsupersaturation in alloy 740H fusion welds was quan-tified by applying Eq. [2] to the data shown in Figure 8.Values for xav were taken to be the as-solidifiedcompositions across the simulated half dendrite inFigure 8, and values for xe were determined by calcu-lating the equilibrium matrix solid solubility across thehalf dendrite. The equilibrium calculations were per-formed assuming a temperature of 1073 K (800 �C),which is within the creep test temperature range of thespecimens investigated in this study. The resulting plotof DCo vs fraction solid for the c¢-formers immediatelyafter solidification is given in Figure 16. Next, the effectof c¢ precipitation on matrix supersaturation during the1073 K/4 hours (800 �C/4 hours) direct age treatmentwas assessed by using DICTRA to simulate the agingtreatment on the as-solidified concentration profileshown in Figure 8. DCo in this condition was thencalculated in the same manner as described above forthe as-solidified case, and the values for the c¢-formersare given in the inset to Figure 16. As shown, there isaround a two order of magnitude reduction in the

matrix supersaturation after aging. This, combined withthe lack of discontinuous reaction products in the as-aged weld metal shown in Figure 3(b), indicates that thechemical free energy contribution to the driving forcefor DC is likely low. This is in agreement with thoseauthors who have argued in favor of interfacial energyas the main driving force for DC[52,74,75] as well as thecalculations of Aaronson and Clark,[76] which predict a46- to 335-fold reduction in chemical driving force(albeit for DP) due to continuous precipitation. Despitethese indications that the chemical contribution toDGDC is low, the influence of the chemical gradientsintroduced during weld metal solidification must not beoverlooked. First, as demonstrated by the large c¢particles along the grain boundary in Figure 3(b) anddescribed by DuPont, et al., microsegregation results inan inhomogeneous precipitate size distribution, withlarger precipitates forming in the regions of increasedmatrix solute content.[18] This directly sets up a situationfavoring coarsening of grain boundary precipitates. Inaddition, the compositional gradients lead to a gradientin c¢ volume fraction (Figures 4(b) and (c)) which peaksin the interdendritic regions (and grain boundaries).This provides an increased source of local solute whichcan be readily absorbed by the large grain boundaryprecipitates during coarsening. Both of these points willincrease the propensity for DC in the weld metal ascompared to the base metal. Thus, even though com-positional gradients imposed by solidification duringwelding may not be a large part of the thermodynamicdriving force for DC in alloy 740H, they play anextremely important role in promoting the DC reaction.The interfacial energy component of the DC driving

force has been derived by Livingston and Cahn,[74] andrelated to the c/c¢ system by Funkenbusch.[52] Thesetreatments are based off of the classical treatment ofinterphase interface curvature driving coarseningthrough the Gibbs–Thomson effect with the assumptionthat interface (i.e., grain boundary) movement andparticle size increases are the result of diffusive fluxalong the grain boundary rather than through the bulk.The result of this is directional coarsening of theparticles that are either dragging the boundary withthem or being dragged by the moving boundary. In theevent of a small influence of matrix supersaturation onDGDC, which appears to be the case in this system due tothe points discussed above, interfacial energy is the mostwidely accepted contributor to the driving force fordiscontinuous reactions.[52,74,77–79] Thus, based on re-ports from the literature and the dismissal of some of theother driving forces given in Eq. [1], it is most likely thatthe largest contributor to the driving force for the DCreaction in fusion welds of alloy 740H is DGc

DC.Finally, it should be noted that the nature of weld

metal grain boundaries serves a source of increased CZsusceptibility in alloy 740H fusion welds. As-solidifiedboundaries in the fusion zone arise from the intersectionof dendrites growing in different orientations. Themorphology of these dendrites causes the grain bound-aries to be very tortuous and have a large degree ofcurvature, as demonstrated in Figures 3, 4, and 5. Atelevated temperatures, the boundaries will act to reduce

Fig. 16—Calculated degree of austenite supersaturation before and(inset) after c¢ precipitation. Ce for austenite calculated at 1073 K(800 �C).


their area by straightening, which provides an additionalsource of moving boundaries at which a discontinuousreaction can begin, beyond the boundary motion due toGBS that is experienced in the sample as a whole.

IV. CONCLUSIONS

1. Creep damage accumulation at grain boundariescontaining c¢ coarsened zones indicates that theseregions are related to the reduced creep rupture lifeof INCONEL alloy 740H fusion welds.

2. Microstructural observation of creep specimensshowed that microstructure, applied stress, time attemperature, and exposure temperature affect theevolution of these coarsened zones. This indicatesthat the mechanism of their formation is likely dif-fusional in nature and exacerbated by deformationduring creep.

3. Elongated particles within these grain boundary fea-tures have been identified as coarse c¢.

4. Microstructural evidence indicates that Nabarro–Herring creep, Nb/Ti carbide formation, and Crcarbide formation are unlikely to be the primarycauses of CZ formation in alloy 740H.

5. The primary formation mechanism of the grainboundary features observed in alloy 740H creepspecimens is discontinuous coarsening of the c¢phase. This is supported by the identity of the elon-gated particles within the CZs, their asymmetricnature, and their lack of an orientation relationshiprelative to the axis of applied stress.

6. GBS and migration during creep likely play a rolein catalyzing CZ formation by discontinuous coars-ening because they induce the grain boundary mo-tion that is necessary for the discontinuous reactionto occur.

7. Discontinuous concentration profiles induced bydiscontinuous reactions in alloy 740H may homoge-nize during long-term thermal exposure in creep.

8. The largest driving force for CZ formation in alloy740H weld metal is most likely the interfacial en-ergy considerations which drive all coarsening reac-tions. However, CZ susceptibility in alloy 740Hweld metal is significantly higher than that of thebase metal because of the inhomogeneous weld me-tal microstructure that results from solute partition-ing during solidification. In addition, the grainboundary motion which is necessary for discontinu-ous reactions to occur is more significant in theweld metal due to the combined effects of tortuousboundary straightening and creep deformation.

ACKNOWLEDGMENTS

The authors gratefully acknowledge the financialsupport of the NSF I/UCRC Center for IntegrativeMaterials Joining Science for Energy Applications(CIMJSEA) under contract #IIP-1034703. They would

also like to acknowledge the financial support pro-vided by Special Metals Corporation, Huntington,WV. Additional thanks are given to Binay Patel at Le-high University for assistance with the STEM in SEMcharacterization, as well as to Ronnie Gollihue at Spe-cial Metals, Jim Tanzosh at Babcock and WilcoxCompany, and Paul Mason at ThermoCalc USA forthe technical discussion and assistance.

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