homogenization and dissolution kinetics of fusion … and dissolution kinetics of fusion welds in...

Homogenization and Dissolution Kinetics of Fusion Weldsin INCONEL� Alloy 740H�

DANIEL H. BECHETTI, JOHN N. DUPONT, JOHN J. deBARBADILLO,and BRIAN A. BAKER

Thermodynamic and kinetic modeling were used to determine appropriate heat treatmentschedules for homogenization and second phase dissolution in INCONEL� alloy 740H�

(INCONEL and 740H are registered trademarks of Special Metals Corporation) fusion welds.Following these simulations, a two-step heat treatment process was applied to specimens from asingle pass gas tungsten arc weld (GTAW). Scanning electron microscopy (SEM) has been usedto assess the changes in the distribution of alloying elements as well as changes in the fraction ofsecond phase particles within the fusion zone. Experimental results demonstrate that adequatehomogenization of alloy 740H weld metal can be achieved by a 1373 K/4 h (1100 �C/4 h)treatment. Complete dissolution of second phase particles could not be completely achieved,even at exposure to temperatures near the alloy’s solidus temperature. These results are in goodagreement with thermodynamic and kinetic predictions.

DOI: 10.1007/s11661-014-2243-z� The Minerals, Metals & Materials Society and ASM International 2014

I. INTRODUCTION

THE International Energy Agency (IEA) has pre-dicted that the amount of electrical power generated bythe burning of coal will increase[1] from 17,500 tera-watthours (TWh) in 2005 to over 30,000 TWh by the year2030. This, in combination with growing environmentalconcerns about the impact of CO2 emissions from theburning of fossil fuels, has spurred the enactment ofinitiatives in Europe (Thermie AD700, MARKO) andthe United States (DOE Vision 21) aimed at developinga new generation of coal fired power plants.[1,2] Theseplants, termed Advanced Ultra Supercritical (A-USC),are designed to operate at higher steam temperatures[973 K to 1033 K (700 �C to 760 �C)] and higher steampressures (35 to 45 MPa) than current generation UltraSupercritical (USC) plants, which operate around[1–3]

873 K (600 �C) and 24 MPa. Given that the efficiency ofpower generation by burning coal is proportional to thehottest temperature within the process cycle, the oper-ating temperature/pressure increases are aimed at rais-ing the overall efficiency of this new generation ofplants. It is projected that A-USC boilers will achieve 46to 50 pct process efficiency, which represents a 10 to 15pct relative improvement over current generationplants.[2,4,5] Increased process efficiency directly reducesthe amount of coal that needs to be burned to generate agiven amount of power and therefore reduces thequantity of carbon dioxide released in generating that

power. Thus, A-USC plants are expected to release 40 to50 pct less CO2 than current generation boilers.[2]

The proposed operating conditions within A-USCplants will place high demands on the materials usedwithin these boilers. In the US, 100,000 hours creepstrength greater than 100 MPa at 1023 K (750 �C) and200,000 hours coal-ash corrosion resistance of less than2 mm metal loss have been set forth as propertyrequirements of materials for use in A-USC boilertubing.[2,4] These conditions disqualify ferritic andaustenitic steels, as well as most of the solid solutionstrengthened nickel-based alloys, from use as boilertubing materials. Thus, attention has turned to the c¢precipitation strengthened Ni-based alloys for use in themost severe regions of A-USC plants.[2,4,6]

INCONEL alloy 740 was developed by Special MetalsCorporation for use in the hottest regions of A-USCplants. Nimonic 263, a c¢ strengthened nickel alloy used inaircraft engines, was used as a starting point for thedevelopment of alloy 740 because it possessed therequired creep rupture strength. However, Nimonic 263is not capable of achieving the necessary level of corrosionresistance, so compositional changes were explored in aneffort to satisfy both A-USC property requirements.Significant additions of chromium and niobium wereadded to alloy 263 to increase corrosion resistance, andmolybdenum content was reduced, as it is known toreduce coal-ash corrosion resistance.[7] Silicon, boron,and niobium contents were also adjusted to maximizealloyweldability. In addition, concerns about the stabilityof the deleterious g phase prompted adjustments to theAl/Ti ratio. The modified alloy, with increased corrosionresistance, weldability, and liquation cracking resistancealong with decreased g phase stability, has been termedINCONEL alloy 740H. The as-processed alloy has beenextensively characterized by corrosion studies and long-term creep rupture testing.[2,4,6–10]

DANIEL H. BECHETTI, Research Assistant, and JOHN N.DUPONT, Professor, are with Lehigh University, Bethlehem, PA18015. Contact e-mail: [email protected] JOHN J. deBARBADILLO,Manager, Product & Process Development, and BRIAN A. BAKER,Product and Application Development Engineer, are with SpecialMetals Corporation, Huntington, WV 25705.

Manuscript submitted February 13, 2013.Article published online March 18, 2014

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 45A, JUNE 2014—3051

Use of alloy 740H inA-USC boilers will require joiningby fusion welding processes, and post weld heat treat-ments (PWHT) may be required to eliminate microseg-regation and dissolve undesirable phases that formduring solidification. However, the PWHT response ofalloy 740H fusion welds has not been investigated indetail. Thus, the objectives of this research are to developa basic understanding of the homogenization and disso-lution kinetics of alloy 740H fusion welds and to use thisinformation to design an effective PWHT schedule for thewelded alloy.

II. EXPERIMENTAL

A single pass gas tungsten arc weld was prepared on a7.3 mm thick plate of alloy 740H using 1.14 mmdiameter alloy 740H filler metal wire. The weld wasmade at 198 A and 11 V under 75/25 Ar/He shieldinggas at 51 cubic meters per hour using a 3.18 mmdiameter W-2 pct Th electrode. The filler wire was fedinto the weld pool at a rate of 19.05 mm per second. Thecompositions of the base metal, filler metal, andas-deposited weld metal were determined using wet chem-ical and OES techniques and are given in Table I. Alsoincluded in Table I are experimentally measured com-positions (using the experimental conditions describedbelow) of the base metal and as-deposited weld metal,for use in comparing the analytical performance of themicroscopy equipment used in this study to the knownmaterial compositions. Minor differences in the compo-sitions of base metal and filler wire can be attributed toheat-to-heat variations.

Microstructural imaging and X-ray energy dispersivespectroscopy (XEDS) were performed in a Hitachi4300SE/N Schottky field emission scanning electronmicroscope (SEM) equipped with a silicon drift detec-tor. The microscope was operated at an acceleratingvoltage of 15 keV for imaging and 20 keV for XEDS.Given the composition of the alloy and the operatingconditions, Monte Carlo simulations using the CASINOprogram[11] predict an electron beam interaction depthof approximately 650 nm with lateral spatial resolutionof approximately 850 nm, as well as an X-ray generationdepth of approximately 1.2 lm with lateral spatialresolution of approximately 1 lm. A baseline microseg-regation profile was acquired via an XEDS line scanacross a series of parallel dendrites in the as-weldedmicrostructure on an as-polished sample. The agreementof this baseline segregation profile with Scheil solidifi-cation calculations was assessed using the ThermoCalcsoftware package[12] and the Thermo Tech TTNi7thermodynamic database.[13] Appropriate homogeniza-tion and dissolution treatments were then developedusing the DICTRA software package[12] in conjunctionwith the TTNi7 thermodynamic database and theMOB2 mobility database.[14] During solidification sim-ulations, all phases included in the relevant databaseswere allowed to be active. Only phases that werepredicted to form in the alloy (c, c¢, g, MC, M23C6,Laves) were active during equilibrium and kinetic

Table

I.IN

CONEL

�Alloy740H

�Compositions(W

tPct)

Technique

Material

Ni

Cr

Co

Nb

Ti

Al

Mo

Fe

Wet

chem

ical/OESanalysis

base

metal

49.17

24.35

20.08

1.53

1.45

1.28

0.53

1.07

filler

metal

50.23

23.88

19.52

1.50

1.31

1.36

0.52

1.04

as-depositedweldmetal

50.20

23.90

19.40

1.52

1.28

1.31

0.54

1.10

XEDS(5

readings)

base

metal

48.90

±0.21

24.41

±0.17

19.85

±0.20

1.61

±0.07

1.47

±0.06

1.21

±0.09

0.70

±0.07

1.16

±0.06


48.83

±0.21

24.45

±0.17

19.91

±0.21

1.60

±0.13

1.46

±0.05

1.21

±0.07

0.69

±0.05

1.12

±0.05

Technique

Material

Si

CMn

VW

Zr

Ta

PCu

S

Wet

chem

ical/OESanalysis

base

metal

0.20

0.05

0.30

0.007

0.008

0.02

<0.001

0.002

<0.001

—filler

metal

0.21

0.05

0.30

——

0.029

<0.0001

0.005

—<0.01


0.22

0.05

0.29

0.006

0.013

0.018

<0.001

0.0022

0.105

0.0013

XEDS(5

readings)

base

metal

0.47

±0.02

—0.22

±0.06

——

——

——

—as-depositedweldmetal

0.54

±0.02

—0.20

±0.10

——

——

——

—

3052—VOLUME 45A, JUNE 2014 METALLURGICAL AND MATERIALS TRANSACTIONS A

calculations to speed simulation time. During all simu-lations, P and S, which are not included in the TTNi7database, were excluded. Carbon was entered as a fast-diffusing component. It should be noted that the Scheilsolidification simulation assumes a planar solid/liquidinterface, no diffusion of substitutional elements in thesolid, complete mixing in the liquid, and equilibrium atthe solid/liquid interface. It, therefore, represents theworst case of microsegregation for substitutional ele-ments, which could result in overestimation of the actualdegree of microsegregation in a real fusion weld. In areal fusion weld, there may be reductions in the extent ofmicrosegregation due to several factors, including soluteenrichment at the dendrite tip, tip undercooling, andback diffusion toward the dendrite core. However, asshown later by comparison to experimental measure-ments, these effects are believed to be very small.

A 1373 K/4 h (1100 �C/4 h) homogenization heattreatment was applied to a 100 mm long segment ofthe weld while sealed in a tube furnace under flowing Argas. After the homogenization step, the specimen wassectioned into smaller segments which were subse-quently solution heat treated using one of the followingconditions: 1473 K, 1513 K, 1553 K, or 1578 K/1 h(1200 �C, 1240 �C, 1280 �C, or 1305 �C/1 h). The solu-tion heat treatment was performed using the samefurnace setup as the homogenization treatment. A type-K thermocouple was mechanically attached to the centerof the weld to monitor temperature during all heattreatments. All specimens were cooled via agitated waterquench. After dissolution heat treatment, the volumefraction of second phase particles remaining in themicrostructure was determined via area fraction analysisusing the Image J software package on 10 fields capturedin the SEM.

All metallographic specimens were prepared usingstandard metallographic techniques and electrolyticallyetched at 6 V in a solution comprised of 20 mL H3PO4

and 150 mL H2SO4 saturated with CrO3. Light opticalmicroscopy (LOM) was performed using and OlympusBH-2 microscope in conjunction with Pax-It micrographacquisition software.

III. RESULTS AND DISCUSSION

A. Microsegregation and Homogenization

Figure 1 shows a low magnification stereomicrographof the weld analyzed in this study. The fusion zone, heataffected zone, and base metal are highlighted in theimage. Figure 2 demonstrates the as-solidified concen-tration profile across a series of parallel dendrites withinthis weld. The field of interest is shown in the lightoptical and SEM micrographs of Figures 2(a) and (b),respectively. The variation in major alloying elementcontent across the region of interest is shown inFigure 2(c), and the variation in c¢ former contentacross the field is shown in Figure 2(d). Examples ofdendrite core and interdendritic locations are labeled inFigures 2(c) and (d). Other minor alloying additionshave been omitted for clarity and because of minimal

demonstrated partitioning. As indicated, the interden-dritic regions are enriched in Ti and Nb and slightlydepleted in Cr and Co. Ni and Al do not showsignificant partitioning behavior.ThermoCalc’s Scheil solidification module was also

used to assess the expected composition profiles fornon-equilibrium solidification of alloy 740H. As men-tioned earlier, this model assumed infinitely fast Cdiffusion and negligible diffusion of the substitutionalalloying elements in the primary c phase during solid-ification. This behavior is justified based the establisheddiffusion behavior of these elements[15] in Ni. Theaustenite composition during solidification is predictedto vary as shown in Figure 3. The leftmost coordinateon the x-axis (0 fraction solid) represents the predictedcomposition at a dendrite core, and the interdendriticcomposition is given by the point where the fractionsolid is nearly equal to unity. It is expected that Nb, Ti,and Cr, will be the strongest segregants to the liquidduring primary solidification, while Ni and Co arepredicted to partition moderately to the solid. The largepredicted changes in austenite composition that occurnear the end of solidification are associated with soluteredistribution as secondary phases form beyond 0.83fraction solid. At this eutectic point, the calculatedcomposition of the remaining liquid is given in Table II.The full solidification path of alloy 740H is predicted tobe

L! Lþ c! Lþ cþMC! Lþ cþMCþ Laves

! Lþ cþMCþ Lavesþ g! cþMCþ Lavesþ g

½1�

Figure 4 compares the experimentally measured com-positions across several dendrites, all normalized to thewidth of the average half-dendrite in the field of interest(~7.5 lm), with the calculated compositions. As shown,the experimental data indicate enrichment of Nb fromaround 0.5 wt pct at the dendrite core to a maximummeasured value of 5.4 wt pct at the interdendritic region.

Fig. 1—Stereomicrograph of a single pass alloy 740H GTA weldanalyzed in this study.


Table II. Calculated Composition of Liquid in Alloy 740H at the End of Primary (L fi L+ c) Solidification (Wt Pct)

Fraction Solid Ni Cr Co Nb Ti Al Mo Fe Si C

0.83 50.47 22.58 15.56 5.04 2.60 1.16 1.04 0.92 0.45 0.18

Fig. 2—XEDS line scan across dendrites in the fusion zone of a single pass alloy 740H GTA weld in the as-welded condition: (a) light opticalmicrograph of region of interest, (b) SEM micrograph of region of interest, (c) concentration profile for major alloying elements, and (d) concen-tration profile for c¢ forming elements.

Fig. 3—Scheil solidification calculations showing the predicted segregation behavior of: (a) the major alloying elements and (b) the c¢ formingelements within the austenite phase in alloy 740H.


In comparison, the Scheil simulation predicts enrich-ment from 0.4 to 8.4 wt pct. Similarly, experimentalmeasurements indicate enrichment of Ti from 1.0 wt pctat the dendrite core to a maximum recorded value of3.3 wt pct, while the Scheil calculations predict enrich-ment from 0.8 to 4.7 wt pct. The minor disagreement forNb, Ti, and Cr across the majority of the dendrite couldbe due to experimental error in the EDS measurementsand/or inaccuracy in the Scheil simulations. It is likelythat the larger discrepancies in the measured andcalculated compositions at the interdendritic regionresult from the sharp change in composition in thisregion that cannot be accurately detected within thespatial resolution of the SEM. As described above, thelateral spatial resolution of the electron beam under thegiven SEM operating conditions is on the order of 1 lm,while the distance over which the sharp compositionchange occurs is significantly smaller. Therefore, it is tobe expected that accurate detection of compositionalchanges in these areas is difficult. Note that this effectdoes not occur in the dendrite core, because thecomposition change with distance is small in theseregions. The measured and calculated compositions aretherefore in much better agreement. Overall, whenconsidering the entire half-dendrite given in Figure 4,the solidification simulations and experimental data arein reasonably good agreement.

As a second comparison between the calculated andexperimental solute segregation behavior, the equilib-rium partition coefficients at the start of solidificationwere determined for each element from both thecalculated and experimental datasets. The equilibriumpartition coefficient for an element, k, is defined as:

k ¼ CS

CL; ½2�

where CS and CL are the composition of the solid andliquid, respectively. At the start of solidification,CS is thecomposition of the first solid to form (i.e., the dendritecore), and CL is the initial liquid composition (i.e., thenominal composition). Using the experimental data givenin Figure 2, the values of CS were taken as the average oftheminimum compositions of six dendrites. The values ofCS for the calculated data were taken as the compositionsat 0 fraction solid. Table III shows the calculated andexperimentally determined k-values for the alloy 740Hsystem at the start of solidification as well as the nominaland dendrite core compositions used to calculate thepartition coefficients. The experimental and predictedk-values are in relatively good agreement and are also inagreement with the work of Tung and Lippold.[16]

The discrepancies between the experimental andcalculated k-values in Table III, especially those of Co

Table III. Partition Coefficients in Alloy 740H at the Start of Solidification

Element Ccore, Calculated Cavgcore, Experimental Co k, Calculated k, Experimental

Ni 51.45 48.76 50.20 1.02 0.97Cr 23.51 21.45 23.90 0.98 0.90Co 21.06 18.76 19.40 1.09 0.97Nb 0.33 0.37 1.52 0.22 0.24Ti 0.66 0.89 1.28 0.52 0.70Al 1.34 1.38 1.31 1.02 1.05

All concentration values given in wt pct.

Fig. 4—Composition profiles from Fig. 2 normalized to the average experimentally observed half-dendrite width and overlaid with the calculatedconcentration profiles from Fig. 3: (a) major alloying elements and (b) c¢ forming elements. Data points are averages of 4 readings, and errorbars indicate 95 pct confidence interval. Black curves are overlaid Scheil solidification predictions.


and Ni where the experimental results indicate parti-tioning to the liquid, while the calculated values predictpartitioning to the solid, can be rationalized by consid-ering the error associated with the XEDS measurements.As noted above, the k-values determined in this exper-iment are for the beginning of solidification only. Assuch, Ni, Cr, and Co, which typically exhibit partitioncoefficients very close to unity during the initial stages ofsolidification, are particularly vulnerable to slight vari-ations in the experimentally measured compositions.These slight variations may push the experimentallyobserved partition coefficients to values marginallyhigher or lower than those predicted by the solidificationmodel. In general, however, it can be seen by observa-tion of the composition profile shown in Figure 4 thatthe calculated segregation trends for Co and Ni are inagreement with the experimentally observed one.

Following the assessment of the segregation behaviorof alloy 740H, the single phase homogenization model[12]

within the DICTRA kinetic modeling package was usedto develop an effective homogenization heat treatment.This model simulated exposure of the compositionprofile shown in Figure 2 to 1373 K (1100 �C) for timesfrom 0.25 to 4 hours. This temperature was chosen basedon industry experience and the thermodynamic calcula-tions shown in Figure 5. These results show the phasefraction as a function of temperature for the nominalalloy composition (Figures 5(a) and (b)) and for thecalculated composition at the interdendritic region wherethe fraction solid is 0.99 (Figures 5(c) and (d)). Thisinterdendritic composition is given in Table IV, anda brief description of the relevant phases shown inFigure 5 is given in Table V. These results indicate thatthe microsegregation at the interdendritic regions signif-icantly depresses the solidus temperature of the alloyfrom 1582 K (1309 �C) for the nominal compositionto 1425 K (1152 �C) within the interdendritic region.For this reason, it is necessary to employ a two-step

Fig. 5—Calculated phase stabilities in alloy 740H: (a) using nominal alloy composition, and (b) same as (a) but magnified to show phase stabili-ties near solidus temperature, (c) using the Scheil-predicted austenite composition at an interdendritic region (0.99 fraction solid), (d) same as (c)but magnified to show phase stabilities near solidus temperature.


homogenization/dissolution treatment consisting of alower temperature homogenization followed by a highertemperature dissolution as described in the followingsection. If a single high temperature heat treatment stepwere used on the as-solidified microstructure, localizedmelting at the interdendritic regions could occur beforehomogenization or second phase dissolution. This effectis commonly observed in superalloys, and therefore, two-step homogenization treatments are typical.

The effectiveness of a homogenization treatment istypically defined by the index of residual segregation, d,which is given by

d ¼ CM � Cm

C0M � C0

m

; ½3�

where CM and C0M are the concentrations at the

interdendritic region before and after homogenization,respectively, and Cm and C0

m are the concentrations atthe dendrite core before and after homogenization,respectively. Preliminary calculations using the alloy740H system indicated that the Nb concentration profileshould be the slowest to homogenize (and wouldtherefore be rate-limiting), so the variation in Nbconcentration was used to calculate d. As shown inFigure 6, homogenization of the as-welded compositionprofile is predicted after 4 hours at 1373 K (1100 �C).This treatment is predicted to reduce residual Nbsegregation in the alloy to a d value less than 3.5 pct.Heat treatment of the alloy 740H weldment under theseconditions yielded the measured concentration profile asshown in Figure 7. As indicated, the microsegregationwithin the dendrites has been sufficiently eliminated.Any remaining regions of local elemental enrichment area result of electron beam interaction with second phaseparticles. It is recognized that homogenization may alsobe possible with a slightly shorter time than thatconsidered here experimentally due to variations in thedendrite arm spacing within the weldment. Figure 8demonstrates the predicted variation in homogenizationtime as defined above during a 1373 K (1100 �C)treatment as a function of dendrite arm spacing. Ingeneral, the curve shown in Figure 8 can be used to

estimate the time required for homogenization of thealloy for a range of dendrite spacings. This hassignificant practical implications, as dendrite arm spac-ing will depend strongly on processing history. Thus,Figure 8 can be used for the design of alloy 740H heattreatment schedules when the dendrite arm spacingdiffers from the 7.5 lm investigated in this study.

B. Dissolution Kinetics

The homogenization treatment described in the previ-ous section will eliminate local changes in the primaryc-austenite composition and increase the local solidustemperature, thus reducing the likelihood of localizedmelting during subsequent higher temperature dissolutiontreatments. The calculation results shown in Figure 5indicate that the MC-type carbide has the greateststability near potential dissolution temperatures. Byplotting the calculated variation inMC composition withtemperature as shown in Figure 9, this phase is predictedto be of the (Nb,Ti)C type, with Nb:Ti � 3.70 at thehomogenization temperature of 1373 K (1100 �C). EDS

Table IV. Calculated Composition of c in Alloy 740H at the End of Solidification (Wt Pct)

Fraction Solid Ni Cr Co Nb Ti Al Mo Fe Si C

0.99 54.93 18.82 8.43 8.51 4.70 1.63 1.83 0.71 0.44 0.001

Table V. Description of Common Phases in Alloy 740H

PhaseCrystal Structure(Space Group)

Approximate LatticeParameter[17,18] (A)

GeneralizedComposition

c (austenite) FCC (Fm�3m) 3.60 Ni solid solutionc¢ FCC (Pm�3m) 3.57 Ni3(Al,Ti,Nb)g Hexagonal (P63=mmc) a = 5.1, c = 8.3 Ni3(Ti,Nb)MC FCC (Fm�3m) 4.5 (Nb,Ti)CM23C6 FCC (Fm�3m) 10.6 Cr23C6

Fig. 6—Simulated homogenization of Nb in an alloy 740H half-den-drite at 1373 K (1100 �C) for t = 0 to 4 h.


measurements of the second phase particles remaining inthe microstructure after the 1373 K (1100 �C) homoge-nization and water quench were shown to be of the(Nb,Ti)C type with Nb:Ti = 3.87 ± 0.25 which boundsthe calculated result and empirically demonstrates thatthe particles remaining in the microstructure afterhomogenization are MC carbides. It is worth notingthat the only other second phase that is predicted to

be thermodynamically stable near the heat treatmenttemperatures investigated in this study is g. For theas-deposited weld metal chemistry (Figures 5(a) and (b)),there is no predicted stability of g across the chosentemperature range. Since g is enriched in Nb and Ti(Table V), this prediction is consistent with the reductionin Nb content in the 740H variant of this alloy. However,

Fig. 7—XEDS line scan across dendrites in the fusion zone of a single pass alloy 740H GTA weld after homogenization at 1373 K (1100 �C) for4 h: (a) light optical micrograph of region of interest, (b) SEM micrograph of region of interest, (c) concentration profile for major alloying ele-ments, and (d) concentration profile for c¢ forming elements.

Fig. 8—Calculated master curve for homogenization of single passGTA weld in alloy 740H at 1373 K (1100 �C). Fig. 9—Calculated variation in MC carbide composition with tem-

perature for the nominal alloy 740H composition.


the significant enrichment of Nb and Ti during solidifi-cation is predicted to result in an as-solidified g content of1.5 9 10�3 wt pct, and an equilibrium stability of g up to20 wt pct at the interdendritic regions. Despite thesepredictions, g has not been experimentally observedduring this study, even at the interdendritic regions. Thisis likely the consequence of the homogenization treat-ment and kinetic effects. First, as given in Figure 5, thestability of g decreases with decreasing enrichment of Ni,Nb, and Ti. Thus, when the as-welded microstructure issubjected to the 1373 K/4 h (1100 �C/4 h) homogeniza-tion treatment, the solvus temperature of g will decreasefrom its initial value of 1423 K (1150 �C), which is givenin Figure 5(d). Eventually, the g solvus will drop belowthe homogenization temperature, resulting in the disso-lution of any g in the microstructure. In addition, g is akinetically slow-developing phase in alloy 740H at thesetemperatures.[17] As such, while up to 20 wt pct g at theinterdendritic regions may be predicted from the equi-librium calculations given in Figure 5, it is highly unlikelythat such a large g content would evolve before the gsolvus drops below the homogenization temperature.Therefore, while g is a concern because it is an undesirable

phase whose presence is generally deleterious formechan-ical properties,[19] and it was present in significantamounts in the original variant of alloy 740, it is unlikelyto be present after the heat treatment described in thisstudy. In contrast to the behavior of g, the MC solidustemperature is predicted to increase from 1505 K to1587 K (1232 �C to 1314 �C) as the alloy is homogenized,so it is expected that the second phase particles remainingin the microstructure after homogenization would beMC, as has been compositionally confirmed above.Note also that the MC phase is predicted to be stable

at or above the matrix solidus temperature for bothcomposition sets shown in Figure 5. This poses the samepractical issue as encountered in the choice of homog-enization temperature, namely that there is a possibilityof localized melting in the weldment before completedissolution of the carbides takes place. Thus, kineticmodeling of the carbide dissolution followed by exper-imental heat treatments is necessary to validate whethera full solution treatment can actually be applied withoutlocalized melting. First, kinetic simulations using theDICTRA software package were used to estimate theevolution of MC mass (weight) fraction as a function oftime at various exposure temperatures. These calcula-tions were conducted using the dispersed phase modeldescribed by Andersson et al.[12] and simulated thehomogenization treatment described in the previoussection, followed by an instantaneous ramp up to aspecified solution temperature. In doing this, the startingcomposition profile and MC phase fraction for thedissolution step are identical to the final compositionprofile and MC phase fraction from the homogenizationsimulation. A typical result of these calculations isshown in Figure 10 for the maximum dissolutiontemperature considered. The plot demonstrates thevariation in phase fraction of MC as a function ofexposure time and distance across a half-dendrite duringa 1373 K/4 h+1578 K/1 h (1100 �C/4 h+1305 �C/1 h)treatment. During the simulations, the interdiffusioncoefficients for Nb and Ti were predicted to vary asgiven in Table VI. As shown in Figure 10, the maximumpredicted mass fraction of MC (before homogenizationor dissolution) is 10.1 9 10�3. By integrating the t = 0curve given in Figure 10, the total predicted massfraction of MC in the system after homogenization is2.6 9 10�3, which was converted to a volume fraction of2.9 9 10�3 using the predicted densities for each phase.A negligible decrease in the MC fraction is predictedduring the 4 hour homogenization at 1373 K (1100 �C),

Table VI. Calculated Interdiffusion Coefficients During Homogenization and Dissolution Treatments (m2/s)

Austenite MC

Nb Ti Nb Ti

1373 K (1100 �C)—homogenization 5.89 9 10�15 to 9.77 9 10�15

(composition dependent)4.98 9 10�15 to 8.71 9 10�15

(composition dependent)N/A N/A

1473 K (1200 �C)—lowtemperature dissolution

3.22 9 10�14 2.40 9 10�14 8.71 9 10�18 5.62 9 10�18

1578 K (1305 �C)—hightemperature dissolution

1.47 9 10�13 9.20 9 10�14 6.76 9 10�17 4.68 9 10�17

Fig. 10—Variation in MC mass fraction across a half-dendrite in al-loy 740H during exposure to a 1373 K/4 h+1578 K/1 h (1100 �C/4 h+1305 �C/1 h) heat treatment, for t = 0, 4 h, various times be-tween 4 and 5 h, and 5 h.


while a moderate decrease in MC content upon exposureto a higher temperature solution treatment is predicted,as shown in the inset to Figure 10. This decreasehappens quickly, such that it reaches its minimum valueafter about 30 minutes of exposure to the dissolutiontemperature. The simulation predicts virtually nochange in MC content beyond 1 hour of exposure tothe dissolution temperature (i.e., 5 hours total heattreatment). The predicted MC dissolution behavior was

similar across all modeled homogenization/dissolutiontreatments.After the kinetic simulations, specimens taken from

the single pass GTA weld homogenized as describedabove were heat treated for 1 hour at 1473 K, 1513 K,1553 K, and 1578 K (1200 �C, 1240 �C, 1280 �C and1305 �C). The evolution of the second phase particles inboth the base metal and weld metal is shown inFigures 11 and 12. As shown in Figure 11(a), the base

Fig. 11—Second phase evolution in alloy 740H: (left column) base metal and (right column) fusion zone of single pass GTA weld. Appliedheat treatments from top row to bottom row are 1373 K/4 h (1100 �C/4 h), 1373 K/4 h+1473 K/1 h (1100 �C/4 h+1200 �C/1 h), and1373 K/4 h+1513 K/1 h (1100 �C/4 h+1240 �C/1 h).


metal contains many large, blocky particles, while thehomogenized weld metal in Figure 11(b) maintains afiner dispersion of smaller particles which still clearlyoutline the dendritic substructure. The measured volumefraction of second phase particles after the 1373 K(1100 �C) treatment was 14 ± 2 9 10�3, which variesgreatly from the predicted value given above. Followingthe 1473 K (1200 �C) exposure, there is an observabledecrease in the amount of second phase, but nosignificant change in the particles’ morphology is noted(Figures 11(c) and (d)). Kinetic calculations for disso-lution at 1473 K (1200 �C) predict a final MC volumefraction of 2.4 9 10�3, which also differs significantlyfrom the measured value of 10 ± 1 9 10�3. Aftertreatment at 1513 K (1240 �C), there is a noticeabledecrease in the size and amount of second phaseparticles in both the base metal and weld metal(Figures 11(e) and (f)). In addition, the dendritic sub-structure of the weld is no longer continuously outlinedby the particles. At this temperature, the calculated MCvolume fraction of 2.2 9 10�3 is again in disagreementwith the experimental value of 6.7 ± 0.8 9 10�3. Expo-sure at 1553 K (1280 �C) (Figures 12(a) and (b)) resultsin a further reduction in secondary phase fraction. At

this temperature, the remaining particles in the basemetal begin to break down from a large blockymorphology into agglomerations of smaller particles.Those that remain in the weld metal have spheroidized.The calculated carbide volume fraction of 2.2 9 10�3 isstill outside the experimental error of the measuredvalue of 4.6 ± 0.7 9 10�3. Heat treatment at 1578 K(1305 �C) for 1 hour produces a further reduction in thefraction of second phase particles in the alloy 740Hweldment (Figures 12(c) and (d)). Once again, thepredicted MC volume fraction of 2.0 9 10�3 was notin agreement with the measured 3.6 ± 0.4 9 10�3 vol-ume fraction of MC. The calculated and measured MCvolume fractions are summarized in Figure 13.Although their magnitudes differ greatly at each tem-perature, the experimental inability to take the carbidesinto solution is consistent with the thermodynamic andkinetic predictions of their stability above the alloy’ssolidus temperature.As discussed, a large degree of mismatch between the

calculated and experimental carbide volume fractionsafter a 1 hour dissolution treatment has been observed.However, because the experimental values were alllarger than the calculated values, and the magnitude of

Fig. 12—Second phase evolution in alloy 740H: (left column) base metal and (right column) fusion zone of single pass GTA weld. Appliedheat treatments from top row to bottom row are 1373 K/4 h+1553 K/1 h (1100 �C/4 h+1280 �C/1 h) and 1373 K/4 h+1578 K/1 h(1100 �C/4 h+1305 �C/1 h).


their mismatch decreased with increasing temperature, itwas surmised that the disagreement came simply as aresult of the kinetics in the alloy being slower thanpredicted by the model. To test this, an additional set ofhomogenized samples were exposed the same dissolu-tion temperatures for a longer duration (24 hours). Asshown in Figure 13, this longer exposure brought theexperimental results into better agreement with thecalculated ones. Although the measured volume frac-tions are still larger, the magnitude of their mismatchhas decreased significantly, such that the largest differ-ence between the two datasets (which is still at the lowestdissolution temperature, as expected) is less than 0.4 volpct. It is, therefore, concluded that the differencesbetween the observed and calculated volume fractionsof MC in alloy 740H during dissolution heat treatmentcan be attributed to slower than predicted kinetics in thealloy.

Based on these experimental observations and thethermodynamic and kinetic predictions, it is concludedthat while homogenization of alloy 740H weldments canbe achieved through a time/temperature treatment thatis practical during the manufacture of welded compo-nents, it is likely that complete dissolution cannot beachieved without exposure heat treatment above thealloy’s solidus temperature.

IV. CONCLUSIONS

The homogenization and dissolution behavior of asingle pass alloy 740H GTA weld were investigated.Thermodynamic and kinetic simulations were performedin combination with experimental heat treatments andcharacterization using electron microscopy techniques.The conclusions of this research are as follows:

1. In the as-welded condition, Nb and Ti display thehighest partitioning to the interdendritic regions.The measured Nb enrichment was approximately

5.4 wt pct, compared to its nominal value of 1.5 wtpct, while the measured Ti enrichment was approxi-mately 3.3 wt pct, compared to its nominal value of1.3 wt pct. These values are reasonably consistentwith the enrichment predicted by non-equilibriumsolidification calculations, which are 8.4 wt pct forNb and 4.7 wt pct for Ti. The actual interdendriticNb and Ti concentrations are likely to be higherthan those measured, because the interaction vol-ume of the electron probe was too large to capturethe quickly changing concentration profile withinthe interdendritic regions.

2. Homogenization of a single pass GTA weld onINCONEL alloy 740H with a 7.5 lm dendrite armspacing can be accomplished with a 1373 K/4 h(1100 �C/4 h) heat treatment. Elimination of micro-segregation with this treatment in is agreement withthermodynamic and kinetic predictions for alloyhomogenization.

3. Heat treatment of an alloy 740H GTA weld at1578 K (1305 �C) for 24 hours was insufficient tocompletely dissolve all second phase particles withinthe microstructure. Thus, while homogenization ofsuch weldments can be achieved, it is likely thatcomplete dissolution cannot be achieved unless thealloy is heated above its solidus temperature.

4. The kinetics of dissolution in the dispersed phasemodel of the DICTRA software package are fasterthan those observed experimentally in alloy 740H.

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 wouldalso like to acknowledge the financial support pro-vided by Special Metals Corporation, Huntington,WV. Additional thanks are given for the technical dis-cussion and assistance provided by Ronnie Gollihue atSpecial Metals, Jim Tanzosh at Babcock and WilcoxCompany, and Paul Mason at ThermoCalc USA.

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