Homogenization and Dissolution Kinetics of Fusion Welds in INCONEL Ò Alloy 740H Ò

Homogenization and Dissolution Kinetics of Fusion Welds in INCONEL Ò Alloy 740H Ò DANIEL H. BECHETTI, JOHN N. DUPONT, JOHN J. debarbadillo, and BRIAN A. BAKER Thermodynamic and kinetic modeling were used to determine appropriate heat treatment schedules for homogenization and second phase dissolution in INCONEL Ò alloy 740H Ò (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 a single pass gas tungsten arc weld (GTAW). Scanning electron microscopy (SEM) has been used to assess the changes in the distribution of alloying elements as well as changes in the fraction of second phase particles within the fusion zone. Experimental results demonstrate that adequate homogenization 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 good agreement with thermodynamic and kinetic predictions. DOI: 10.1007/s11661-014-2243-z Ó The Minerals, Metals & Materials Society and ASM International 2014 I. INTRODUCTION THE International Energy Agency (IEA) has predicted that the amount of electrical power generated by the burning of coal will increase [1] from 17,500 tera-watt hours (TWh) in 2005 to over 30,000 TWh by the year 2030. This, in combination with growing environmental concerns about the impact of CO 2 emissions from the burning of fossil fuels, has spurred the enactment of initiatives in Europe (Thermie AD700, MARKO) and the United States (DOE Vision 21) aimed at developing a new generation of coal fired power plants. [1,2] These plants, 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 steam pressures (35 to 45 MPa) than current generation Ultra Supercritical (USC) plants, which operate around [1 3] 873 K (600 C) and 24 MPa. Given that the efficiency of power generation by burning coal is proportional to the hottest temperature within the process cycle, the operating temperature/pressure increases are aimed at raising the overall efficiency of this new generation of plants. It is projected that A-USC boilers will achieve 46 to 50 pct process efficiency, which represents a 10 to 15 pct relative improvement over current generation plants. [2,4,5] Increased process efficiency directly reduces the amount of coal that needs to be burned to generate a given amount of power and therefore reduces the quantity of carbon dioxide released in generating that DANIEL H. BECHETTI, Research Assistant, and JOHN N. DUPONT, Professor, are with Lehigh University, Bethlehem, PA 18015. Contact e-mail: dhb210@lehigh.edu JOHN J. debarbadillo, Manager, Product & Process Development, and BRIAN A. BAKER, Product and Application Development Engineer, are with Special Metals Corporation, Huntington, WV 25705. Manuscript submitted February 13, 2013. Article published online March 18, 2014 power. Thus, A-USC plants are expected to release 40 to 50 pct less CO 2 than current generation boilers. [2] The proposed operating conditions within A-USC plants will place high demands on the materials used within these boilers. In the US, 100,000 hours creep strength greater than 100 MPa at 1023 K (750 C) and 200,000 hours coal-ash corrosion resistance of less than 2 mm metal loss have been set forth as property requirements of materials for use in A-USC boiler tubing. [2,4] These conditions disqualify ferritic and austenitic steels, as well as most of the solid solution strengthened nickel-based alloys, from use as boiler tubing materials. Thus, attention has turned to the c precipitation strengthened Ni-based alloys for use in the most severe regions of A-USC plants. [2,4,6] INCONEL alloy 740 was developed by Special Metals Corporation for use in the hottest regions of A-USC plants. Nimonic 263, a c strengthened nickel alloy used in aircraft engines, was used as a starting point for the development of alloy 740 because it possessed the required creep rupture strength. However, Nimonic 263 is not capable of achieving the necessary level of corrosion resistance, so compositional changes were explored in an effort to satisfy both A-USC property requirements. Significant additions of chromium and niobium were added to alloy 263 to increase corrosion resistance, and molybdenum content was reduced, as it is known to reduce coal-ash corrosion resistance. [7] Silicon, boron, and niobium contents were also adjusted to maximize alloy weldability. In addition, concerns about the stability of the deleterious g phase prompted adjustments to the Al/Ti ratio. The modified alloy, with increased corrosion resistance, weldability, and liquation cracking resistance along with decreased g phase stability, has been termed INCONEL alloy 740H. The as-processed alloy has been extensively characterized by corrosion studies and longterm creep rupture testing. [2,4,6 10] METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 45A, JUNE 2014 3051

Use of alloy 740H in A-USC boilers will require joining by fusion welding processes, and post weld heat treatments (PWHT) may be required to eliminate microsegregation and dissolve undesirable phases that form during solidification. However, the PWHT response of alloy 740H fusion welds has not been investigated in detail. Thus, the objectives of this research are to develop a basic understanding of the homogenization and dissolution kinetics of alloy 740H fusion welds and to use this information to design an effective PWHT schedule for the welded alloy. II. EXPERIMENTAL A single pass gas tungsten arc weld was prepared on a 7.3 mm thick plate of alloy 740H using 1.14 mm diameter alloy 740H filler metal wire. The weld was made at 198 A and 11 V under 75/25 Ar/He shielding gas at 51 cubic meters per hour using a 3.18 mm diameter W-2 pct Th electrode. The filler wire was fed into the weld pool at a rate of 19.05 mm per second. The compositions of the base metal, filler metal, and as-deposited weld metal were determined using wet chemical and OES techniques and are given in Table I. Also included in Table I are experimentally measured compositions (using the experimental conditions described below) of the base metal and as-deposited weld metal, for use in comparing the analytical performance of the microscopy equipment used in this study to the known material compositions. Minor differences in the compositions of base metal and filler wire can be attributed to heat-to-heat variations. Microstructural imaging and X-ray energy dispersive spectroscopy (XEDS) were performed in a Hitachi 4300SE/N Schottky field emission scanning electron microscope (SEM) equipped with a silicon drift detector. The microscope was operated at an accelerating voltage of 15 kev for imaging and 20 kev for XEDS. Given the composition of the alloy and the operating conditions, Monte Carlo simulations using the CASINO program [11] predict an electron beam interaction depth of approximately 650 nm with lateral spatial resolution of approximately 850 nm, as well as an X-ray generation depth of approximately 1.2 lm with lateral spatial resolution of approximately 1 lm. A baseline microsegregation profile was acquired via an XEDS line scan across a series of parallel dendrites in the as-welded microstructure on an as-polished sample. The agreement of this baseline segregation profile with Scheil solidification calculations was assessed using the ThermoCalc software package [12] and the Thermo Tech TTNi7 thermodynamic database. [13] Appropriate homogenization and dissolution treatments were then developed using the DICTRA software package [12] in conjunction with the TTNi7 thermodynamic database and the MOB2 mobility database. [14] During solidification simulations, all phases included in the relevant databases were allowed to be active. Only phases that were predicted to form in the alloy (c, c, g, MC, M 23 C 6, Laves) were active during equilibrium and kinetic Table I. INCONEL Ò Alloy 740H Ò Compositions (Wt Pct) Technique Material Ni Cr Co Nb Ti Al Mo Fe Wet chemical/oes analysis 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-deposited weld metal 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 as-deposited weld metal 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 C Mn V W Zr Ta P Cu S Wet chemical/oes analysis 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 as-deposited weld metal 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-deposited weld metal 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 simulations, P and S, which are not included in the TTNi7 database, were excluded. Carbon was entered as a fastdiffusing component. It should be noted that the Scheil solidification simulation assumes a planar solid/liquid interface, no diffusion of substitutional elements in the solid, complete mixing in the liquid, and equilibrium at the solid/liquid interface. It, therefore, represents the worst case of microsegregation for substitutional elements, which could result in overestimation of the actual degree of microsegregation in a real fusion weld. In a real fusion weld, there may be reductions in the extent of microsegregation due to several factors, including solute enrichment at the dendrite tip, tip undercooling, and back diffusion toward the dendrite core. However, as shown later by comparison to experimental measurements, these effects are believed to be very small. A 1373 K/4 h (1100 C/4 h) homogenization heat treatment was applied to a 100 mm long segment of the weld while sealed in a tube furnace under flowing Ar gas. After the homogenization step, the specimen was sectioned into smaller segments which were subsequently solution heat treated using one of the following conditions: 1473 K, 1513 K, 1553 K, or 1578 K/1 h (1200 C, 1240 C, 1280 C, or 1305 C/1 h). The solution heat treatment was performed using the same furnace setup as the homogenization treatment. A type- K thermocouple was mechanically attached to the center of the weld to monitor temperature during all heat treatments. All specimens were cooled via agitated water quench. After dissolution heat treatment, the volume fraction of second phase particles remaining in the microstructure was determined via area fraction analysis using the Image J software package on 10 fields captured in the SEM. All metallographic specimens were prepared using standard metallographic techniques and electrolytically etched at 6 V in a solution comprised of 20 ml H 3 PO 4 and 150 ml H 2 SO 4 saturated with CrO 3. Light optical microscopy (LOM) was performed using and Olympus BH-2 microscope in conjunction with Pax-It micrograph acquisition software. III. RESULTS AND DISCUSSION A. Microsegregation and Homogenization Figure 1 shows a low magnification stereomicrograph of the weld analyzed in this study. The fusion zone, heat affected zone, and base metal are highlighted in the image. Figure 2 demonstrates the as-solidified concentration profile across a series of parallel dendrites within this weld. The field of interest is shown in the light optical and SEM micrographs of Figures 2(a) and (b), respectively. The variation in major alloying element content across the region of interest is shown in Figure 2(c), and the variation in c former content across the field is shown in Figure 2(d). Examples of dendrite core and interdendritic locations are labeled in Figures 2(c) and (d). Other minor alloying additions have been omitted for clarity and because of minimal Fig. 1 Stereomicrograph of a single pass alloy 740H GTA weld analyzed in this study. demonstrated partitioning. As indicated, the interdendritic regions are enriched in Ti and Nb and slightly depleted in Cr and Co. Ni and Al do not show significant partitioning behavior. ThermoCalc s Scheil solidification module was also used to assess the expected composition profiles for non-equilibrium solidification of alloy 740H. As mentioned earlier, this model assumed infinitely fast C diffusion and negligible diffusion of the substitutional alloying elements in the primary c phase during solidification. This behavior is justified based the established diffusion behavior of these elements [15] in Ni. The austenite composition during solidification is predicted to vary as shown in Figure 3. The leftmost coordinate on the x-axis (0 fraction solid) represents the predicted composition at a dendrite core, and the interdendritic composition is given by the point where the fraction solid is nearly equal to unity. It is expected that Nb, Ti, and Cr, will be the strongest segregants to the liquid during primary solidification, while Ni and Co are predicted to partition moderately to the solid. The large predicted changes in austenite composition that occur near the end of solidification are associated with solute redistribution as secondary phases form beyond 0.83 fraction solid. At this eutectic point, the calculated composition of the remaining liquid is given in Table II. The full solidification path of alloy 740H is predicted to be L! L þ c! L þ c þ MC! L þ c þ MC þ Laves! L þ c þ MC þ Laves þ g! c þ MC þ Laves þ g ½1Š Figure 4 compares the experimentally measured compositions across several dendrites, all normalized to the width 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 from around 0.5 wt pct at the dendrite core to a maximum measured value of 5.4 wt pct at the interdendritic region. METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 45A, JUNE 2014 3053

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 optical micrograph of region of interest, (b) SEM micrograph of region of interest, (c) concentration profile for major alloying elements, and (d) concentration 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 forming elements within the austenite phase in alloy 740H. 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 3054 VOLUME 45A, JUNE 2014 METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 4 Composition profiles from Fig. 2 normalized to the average experimentally observed half-dendrite width and overlaid with the calculated concentration profiles from Fig. 3: (a) major alloying elements and (b) c forming elements. Data points are averages of 4 readings, and error bars indicate 95 pct confidence interval. Black curves are overlaid Scheil solidification predictions. Table III. Partition Coefficients in Alloy 740H at the Start of Solidification Element C core, Calculated C avg core, Experimental C o k, Calculated k, Experimental Ni 51.45 48.76 50.20 1.02 0.97 Cr 23.51 21.45 23.90 0.98 0.90 Co 21.06 18.76 19.40 1.09 0.97 Nb 0.33 0.37 1.52 0.22 0.24 Ti 0.66 0.89 1.28 0.52 0.70 Al 1.34 1.38 1.31 1.02 1.05 All concentration values given in wt pct. In comparison, the Scheil simulation predicts enrichment from 0.4 to 8.4 wt pct. Similarly, experimental measurements indicate enrichment of Ti from 1.0 wt pct at the dendrite core to a maximum recorded value of 3.3 wt pct, while the Scheil calculations predict enrichment from 0.8 to 4.7 wt pct. The minor disagreement for Nb, Ti, and Cr across the majority of the dendrite could be due to experimental error in the EDS measurements and/or inaccuracy in the Scheil simulations. It is likely that the larger discrepancies in the measured and calculated compositions at the interdendritic region result from the sharp change in composition in this region that cannot be accurately detected within the spatial resolution of the SEM. As described above, the lateral spatial resolution of the electron beam under the given SEM operating conditions is on the order of 1 lm, while the distance over which the sharp composition change occurs is significantly smaller. Therefore, it is to be expected that accurate detection of compositional changes in these areas is difficult. Note that this effect does not occur in the dendrite core, because the composition change with distance is small in these regions. The measured and calculated compositions are therefore in much better agreement. Overall, when considering the entire half-dendrite given in Figure 4, the solidification simulations and experimental data are in reasonably good agreement. As a second comparison between the calculated and experimental solute segregation behavior, the equilibrium partition coefficients at the start of solidification were determined for each element from both the calculated and experimental datasets. The equilibrium partition coefficient for an element, k, is defined as: k ¼ C S ; ½2Š C L where C S and C L are the composition of the solid and liquid, respectively. At the start of solidification, C S is the composition of the first solid to form (i.e., the dendrite core), and C L is the initial liquid composition (i.e., the nominal composition). Using the experimental data given in Figure 2, the values of C S were taken as the average of the minimum compositions of six dendrites. The values of C S for the calculated data were taken as the compositions at 0 fraction solid. Table III shows the calculated and experimentally determined k-values for the alloy 740H system at the start of solidification as well as the nominal and dendrite core compositions used to calculate the partition coefficients. The experimental and predicted k-values are in relatively good agreement and are also in agreement with the work of Tung and Lippold. [16] The discrepancies between the experimental and calculated k-values in Table III, especially those of Co METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 45A, JUNE 2014 3055

Fig. 5 Calculated phase stabilities in alloy 740H: (a) using nominal alloy composition, and (b) same as (a) but magnified to show phase stabilities 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. and Ni where the experimental results indicate partitioning to the liquid, while the calculated values predict partitioning to the solid, can be rationalized by considering the error associated with the XEDS measurements. As noted above, the k-values determined in this experiment are for the beginning of solidification only. As such, Ni, Cr, and Co, which typically exhibit partition coefficients very close to unity during the initial stages of solidification, are particularly vulnerable to slight variations in the experimentally measured compositions. These slight variations may push the experimentally observed partition coefficients to values marginally higher or lower than those predicted by the solidification model. In general, however, it can be seen by observation of the composition profile shown in Figure 4 that the calculated segregation trends for Co and Ni are in agreement with the experimentally observed one. Following the assessment of the segregation behavior of alloy 740H, the single phase homogenization model [12] within the DICTRA kinetic modeling package was used to develop an effective homogenization heat treatment. This model simulated exposure of the composition profile shown in Figure 2 to 1373 K (1100 C) for times from 0.25 to 4 hours. This temperature was chosen based on industry experience and the thermodynamic calculations shown in Figure 5. These results show the phase fraction as a function of temperature for the nominal alloy composition (Figures 5(a) and (b)) and for the calculated composition at the interdendritic region where the fraction solid is 0.99 (Figures 5(c) and (d)). This interdendritic composition is given in Table IV, and a brief description of the relevant phases shown in Figure 5 is given in Table V. These results indicate that the microsegregation at the interdendritic regions significantly depresses the solidus temperature of the alloy from 1582 K (1309 C) for the nominal composition to 1425 K (1152 C) within the interdendritic region. For this reason, it is necessary to employ a two-step 3056 VOLUME 45A, JUNE 2014 METALLURGICAL AND MATERIALS TRANSACTIONS A

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 Phase Crystal Structure (Space Group) Approximate Lattice Parameter [17,18] (A ) Generalized Composition c (austenite) FCC (Fm 3m) 3.60 Ni solid solution c FCC (Pm 3m) 3.57 Ni 3 (Al,Ti,Nb) g Hexagonal (P6 3 =mmc) a = 5.1, c = 8.3 Ni 3 (Ti,Nb) MC FCC (Fm 3m) 4.5 (Nb,Ti)C M 23 C 6 FCC (Fm 3m) 10.6 Cr 23 C 6 homogenization/dissolution treatment consisting of a lower temperature homogenization followed by a higher temperature dissolution as described in the following section. If a single high temperature heat treatment step were used on the as-solidified microstructure, localized melting at the interdendritic regions could occur before homogenization or second phase dissolution. This effect is commonly observed in superalloys, and therefore, twostep homogenization treatments are typical. The effectiveness of a homogenization treatment is typically defined by the index of residual segregation, d, which is given by d ¼ C M C m C 0 M ; ½3Š C0 m where C M and C 0 M are the concentrations at the interdendritic region before and after homogenization, respectively, and C m and C 0 m are the concentrations at the dendrite core before and after homogenization, respectively. Preliminary calculations using the alloy 740H system indicated that the Nb concentration profile should be the slowest to homogenize (and would therefore be rate-limiting), so the variation in Nb concentration was used to calculate d. As shown in Figure 6, homogenization of the as-welded composition profile is predicted after 4 hours at 1373 K (1100 C). This treatment is predicted to reduce residual Nb segregation in the alloy to a d value less than 3.5 pct. Heat treatment of the alloy 740H weldment under these conditions yielded the measured concentration profile as shown in Figure 7. As indicated, the microsegregation within the dendrites has been sufficiently eliminated. Any remaining regions of local elemental enrichment are a result of electron beam interaction with second phase particles. It is recognized that homogenization may also be possible with a slightly shorter time than that considered here experimentally due to variations in the dendrite arm spacing within the weldment. Figure 8 demonstrates the predicted variation in homogenization time as defined above during a 1373 K (1100 C) treatment as a function of dendrite arm spacing. In general, the curve shown in Figure 8 can be used to Fig. 6 Simulated homogenization of Nb in an alloy 740H half-dendrite at 1373 K (1100 C) for t = 0 to 4 h. estimate the time required for homogenization of the alloy for a range of dendrite spacings. This has significant practical implications, as dendrite arm spacing will depend strongly on processing history. Thus, Figure 8 can be used for the design of alloy 740H heat treatment schedules when the dendrite arm spacing differs from the 7.5 lm investigated in this study. B. Dissolution Kinetics The homogenization treatment described in the previous section will eliminate local changes in the primary c-austenite composition and increase the local solidus temperature, thus reducing the likelihood of localized melting during subsequent higher temperature dissolution treatments. The calculation results shown in Figure 5 indicate that the MC-type carbide has the greatest stability near potential dissolution temperatures. By plotting the calculated variation in MC composition with temperature as shown in Figure 9, this phase is predicted to be of the (Nb,Ti)C type, with Nb:Ti 3.70 at the homogenization temperature of 1373 K (1100 C). EDS METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 45A, JUNE 2014 3057

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) for 4 h:(a) light optical micrograph of region of interest, (b) SEM micrograph of region of interest, (c) concentration profile for major alloying elements, and (d) concentration profile for c forming elements. Fig. 8 Calculated master curve for homogenization of single pass GTA weld in alloy 740H at 1373 K (1100 C). measurements of the second phase particles remaining in the microstructure after the 1373 K (1100 C) homogenization and water quench were shown to be of the (Nb,Ti)C type with Nb:Ti = 3.87 ± 0.25 which bounds the calculated result and empirically demonstrates that the particles remaining in the microstructure after homogenization are MC carbides. It is worth noting that the only other second phase that is predicted to Fig. 9 Calculated variation in MC carbide composition with temperature for the nominal alloy 740H composition. be thermodynamically stable near the heat treatment temperatures investigated in this study is g. For the as-deposited weld metal chemistry (Figures 5(a) and (b)), there is no predicted stability of g across the chosen temperature range. Since g is enriched in Nb and Ti (Table V), this prediction is consistent with the reduction in Nb content in the 740H variant of this alloy. However, 3058 VOLUME 45A, JUNE 2014 METALLURGICAL AND MATERIALS TRANSACTIONS A

the significant enrichment of Nb and Ti during solidification is predicted to result in an as-solidified g content of 1.5 9 10 3 wt pct, and an equilibrium stability of g up to 20 wt pct at the interdendritic regions. Despite these predictions, g has not been experimentally observed during this study, even at the interdendritic regions. This is likely the consequence of the homogenization treatment and kinetic effects. First, as given in Figure 5, the stability of g decreases with decreasing enrichment of Ni, Nb, and Ti. Thus, when the as-welded microstructure is subjected to the 1373 K/4 h (1100 C/4 h) homogenization treatment, the solvus temperature of g will decrease from its initial value of 1423 K (1150 C), which is given in Figure 5(d). Eventually, the g solvus will drop below the homogenization temperature, resulting in the dissolution of any g in the microstructure. In addition, g is a kinetically slow-developing phase in alloy 740H at these temperatures. [17] As such, while up to 20 wt pct g at the interdendritic regions may be predicted from the equilibrium calculations given in Figure 5, it is highly unlikely that such a large g content would evolve before the g solvus drops below the homogenization temperature. Therefore, while g is a concern because it is an undesirable Fig. 10 Variation in MC mass fraction across a half-dendrite in alloy 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 between 4 and 5 h, and 5 h. phase whose presence is generally deleterious for mechanical properties, [19] and it was present in significant amounts in the original variant of alloy 740, it is unlikely to be present after the heat treatment described in this study. In contrast to the behavior of g, the MC solidus temperature is predicted to increase from 1505 K to 1587 K (1232 C to 1314 C) as the alloy is homogenized, so it is expected that the second phase particles remaining in the microstructure after homogenization would be MC, 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 both composition sets shown in Figure 5. This poses the same practical issue as encountered in the choice of homogenization temperature, namely that there is a possibility of localized melting in the weldment before complete dissolution of the carbides takes place. Thus, kinetic modeling of the carbide dissolution followed by experimental heat treatments is necessary to validate whether a full solution treatment can actually be applied without localized melting. First, kinetic simulations using the DICTRA software package were used to estimate the evolution of MC mass (weight) fraction as a function of time at various exposure temperatures. These calculations were conducted using the dispersed phase model described by Andersson et al. [12] and simulated the homogenization treatment described in the previous section, followed by an instantaneous ramp up to a specified solution temperature. In doing this, the starting composition profile and MC phase fraction for the dissolution step are identical to the final composition profile and MC phase fraction from the homogenization simulation. A typical result of these calculations is shown in Figure 10 for the maximum dissolution temperature considered. The plot demonstrates the variation in phase fraction of MC as a function of exposure time and distance across a half-dendrite during a 1373 K/4 h + 1578 K/1 h (1100 C/4 h + 1305 C/1 h) treatment. During the simulations, the interdiffusion coefficients for Nb and Ti were predicted to vary as given in Table VI. As shown in Figure 10, the maximum predicted mass fraction of MC (before homogenization or dissolution) is 10.1 9 10 3. By integrating the t =0 curve given in Figure 10, the total predicted mass fraction of MC in the system after homogenization is 2.6 9 10 3, which was converted to a volume fraction of 2.9 9 10 3 using the predicted densities for each phase. A negligible decrease in the MC fraction is predicted during the 4 hour homogenization at 1373 K (1100 C), Table VI. Calculated Interdiffusion Coefficients During Homogenization and Dissolution Treatments (m 2 /s) Austenite MC Nb Ti Nb Ti 1373 K (1100 C) homogenization 5.89 9 10 15 to 9.77 9 10 15 4.98 9 10 15 to 8.71 9 10 15 N/A N/A (composition dependent) (composition dependent) 1473 K (1200 C) low 3.22 9 10 14 2.40 9 10 14 8.71 9 10 18 5.62 9 10 18 temperature dissolution 1578 K (1305 C) high temperature dissolution 1.47 9 10 13 9.20 9 10 14 6.76 9 10 17 4.68 9 10 17 METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 45A, JUNE 2014 3059

Fig. 11 Second phase evolution in alloy 740H: (left column) base metal and (right column) fusion zone of single pass GTA weld. Applied heat 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), and 1373 K/4 h + 1513 K/1 h (1100 C/4 h + 1240 C/1 h). while a moderate decrease in MC content upon exposure to a higher temperature solution treatment is predicted, as shown in the inset to Figure 10. This decrease happens quickly, such that it reaches its minimum value after about 30 minutes of exposure to the dissolution temperature. The simulation predicts virtually no change in MC content beyond 1 hour of exposure to the dissolution temperature (i.e., 5 hours total heat treatment). The predicted MC dissolution behavior was similar across all modeled homogenization/dissolution treatments. After the kinetic simulations, specimens taken from the single pass GTA weld homogenized as described above were heat treated for 1 hour at 1473 K, 1513 K, 1553 K, and 1578 K (1200 C, 1240 C, 1280 C and 1305 C). The evolution of the second phase particles in both the base metal and weld metal is shown in Figures 11 and 12. As shown in Figure 11(a), the base 3060 VOLUME 45A, JUNE 2014 METALLURGICAL AND MATERIALS TRANSACTIONS A

metal contains many large, blocky particles, while the homogenized weld metal in Figure 11(b) maintains a finer dispersion of smaller particles which still clearly outline the dendritic substructure. The measured volume fraction of second phase particles after the 1373 K (1100 C) treatment was 14 ± 2 9 10 3, which varies greatly from the predicted value given above. Following the 1473 K (1200 C) exposure, there is an observable decrease in the amount of second phase, but no significant change in the particles morphology is noted (Figures 11(c) and (d)). Kinetic calculations for dissolution at 1473 K (1200 C) predict a final MC volume fraction of 2.4 9 10 3, which also differs significantly from the measured value of 10 ± 1 9 10 3. After treatment at 1513 K (1240 C), there is a noticeable decrease in the size and amount of second phase particles in both the base metal and weld metal (Figures 11(e) and (f)). In addition, the dendritic substructure of the weld is no longer continuously outlined by the particles. At this temperature, the calculated MC volume fraction of 2.2 9 10 3 is again in disagreement with the experimental value of 6.7 ± 0.8 9 10 3. Exposure at 1553 K (1280 C) (Figures 12(a) and (b)) results in a further reduction in secondary phase fraction. At this temperature, the remaining particles in the base metal begin to break down from a large blocky morphology 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 is still outside the experimental error of the measured value of 4.6 ± 0.7 9 10 3. Heat treatment at 1578 K (1305 C) for 1 hour produces a further reduction in the fraction of second phase particles in the alloy 740H weldment (Figures 12(c) and (d)). Once again, the predicted MC volume fraction of 2.0 9 10 3 was not in agreement with the measured 3.6 ± 0.4 9 10 3 volume fraction of MC. The calculated and measured MC volume fractions are summarized in Figure 13. Although their magnitudes differ greatly at each temperature, the experimental inability to take the carbides into solution is consistent with the thermodynamic and kinetic predictions of their stability above the alloy s solidus temperature. As discussed, a large degree of mismatch between the calculated and experimental carbide volume fractions after a 1 hour dissolution treatment has been observed. However, because the experimental values were all larger 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. Applied heat 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). METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 45A, JUNE 2014 3061

Fig. 13 DICTRA-predicted variation in MC content and experimentally measured variation in MC content with changing dissolution heat treatment temperature. Dashed lines are a linear fit to the experimental data points. Experimental error bars represent one standard deviation from the average MC content of 10 fields. their mismatch decreased with increasing temperature, it was surmised that the disagreement came simply as a result of the kinetics in the alloy being slower than predicted by the model. To test this, an additional set of homogenized samples were exposed the same dissolution temperatures for a longer duration (24 hours). As shown in Figure 13, this longer exposure brought the experimental results into better agreement with the calculated ones. Although the measured volume fractions are still larger, the magnitude of their mismatch has decreased significantly, such that the largest difference between the two datasets (which is still at the lowest dissolution temperature, as expected) is less than 0.4 vol pct. It is, therefore, concluded that the differences between the observed and calculated volume fractions of MC in alloy 740H during dissolution heat treatment can be attributed to slower than predicted kinetics in the alloy. Based on these experimental observations and the thermodynamic and kinetic predictions, it is concluded that while homogenization of alloy 740H weldments can be achieved through a time/temperature treatment that is practical during the manufacture of welded components, it is likely that complete dissolution cannot be achieved without exposure heat treatment above the alloy s solidus temperature. IV. CONCLUSIONS The homogenization and dissolution behavior of a single pass alloy 740H GTA weld were investigated. Thermodynamic and kinetic simulations were performed in combination with experimental heat treatments and characterization using electron microscopy techniques. The conclusions of this research are as follows: 1. In the as-welded condition, Nb and Ti display the highest partitioning to the interdendritic regions. The measured Nb enrichment was approximately 5.4 wt pct, compared to its nominal value of 1.5 wt pct, while the measured Ti enrichment was approximately 3.3 wt pct, compared to its nominal value of 1.3 wt pct. These values are reasonably consistent with the enrichment predicted by non-equilibrium solidification calculations, which are 8.4 wt pct for Nb and 4.7 wt pct for Ti. The actual interdendritic Nb and Ti concentrations are likely to be higher than those measured, because the interaction volume of the electron probe was too large to capture the quickly changing concentration profile within the interdendritic regions. 2. Homogenization of a single pass GTA weld on INCONEL alloy 740H with a 7.5 lm dendrite arm spacing can be accomplished with a 1373 K/4 h (1100 C/4 h) heat treatment. Elimination of microsegregation with this treatment in is agreement with thermodynamic and kinetic predictions for alloy homogenization. 3. Heat treatment of an alloy 740H GTA weld at 1578 K (1305 C) for 24 hours was insufficient to completely dissolve all second phase particles within the microstructure. Thus, while homogenization of such weldments can be achieved, it is likely that complete dissolution cannot be achieved unless the alloy is heated above its solidus temperature. 4. The kinetics of dissolution in the dispersed phase model of the DICTRA software package are faster than those observed experimentally in alloy 740H. ACKNOWLEDGMENTS The authors gratefully acknowledge the financial support of the NSF I/UCRC Center for Integrative Materials Joining Science for Energy Applications (CIMJSEA) under contract #IIP-1034703. They would also like to acknowledge the financial support provided by Special Metals Corporation, Huntington, WV. Additional thanks are given for the technical discussion and assistance provided by Ronnie Gollihue at Special Metals, Jim Tanzosh at Babcock and Wilcox Company, and Paul Mason at ThermoCalc USA. REFERENCES 1. R. Blum and J. Bugge: Proceedings of the 6th International Conference on Advances in Materials Technology for Fossil Power Plants, Santa Fe, NM, USA, August 31, 2010 September 3, 2010, ASM International, Materials Park, OH, 2010, pp. 1 10. 2. G. Smith and L. Shoemaker: Adv. Mater. Process., 2004, vol. 162, pp. 23 26. 3. 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