Stray Grain Formation in Welds of Single-Crystal Ni-Base Superalloy CMSX-4

Transcription

1 Stray Grain Formation in Welds of Single-Crystal Ni-Base Superalloy CMSX-4 T.D. ANDERSON, J.N. DUPONT, and T. DEBROY Autogenous welds on the single-crystal (SX) alloy CMSX-4 were prepared over a wide range of welding parameters and processes to investigate the formation and behavior of stray grains (SGs). The quantity and location of SGs in the welds were analyzed by orientation imaging microscopy (OIM). Heat- and fluid-flow modeling was conducted to understand the influence of welding parameters on the local solidification conditions and resultant SG formation tendency. The results indicate that constitutional supercooling and SG formation are generally reduced in low-power, high-travel-speed welds. Because of the complex effect of travel speed on temperature gradient and solidification velocity, the worst conditions for SG formation in alloy CMSX-4 for the conditions examined here occur at intermediate travel speeds of ~6 mm/s. These findings were corroborated with heat-transfer/fluid-flow modeling simulations that were coupled with SG predictions. These calculations also indicate that SG formation will be greatest where different regions of dendrite growth intersect, due to the so-called off-axis heat flow. For a given set of welding conditions, the amount of SGs will also vary with substrate orientation. This effect is attributed to differences in the number and location of dendrite growth intersection regions within the melt pool that occur with changes in substrate orientation. DOI: /s Ó The Minerals, Metals & Materials Society and ASM International 2009 I. INTRODUCTION SUPERALLOYS with an Ni base have long been used in the manufacture of gas turbine engine components because of their high melting temperatures and excellent high-temperature mechanical properties. Advanced single-crystal (SX) casting techniques have been developed that extend the maximum service temperatures within the engine for improved efficiency, and the resulting improvements in creep resistance have also extended their service lifetimes. [1] However, replacement of SX components is still necessary due to the effects of thermal fatigue and erosion. Owing to a high manufacturing cost (~$30,000 for a single turbine blade [2] ), the development of an effective weld repair strategy for failed or miscast SX components has been widely pursued due to the potential cost savings in the engine industry. The primary goal has been to produce weld fusion zones that are free of equiaxed grains that compromise the integrity of the SX component by introducing grain boundaries. Because these equiaxed grains have crystal orientations that differ from that of the base metal, they are often termed stray grains (SGs). Early SX welding research performed by Rappaz et al. [3 5] described the major features of an SX weld T.D. ANDERSON, Research Engineer, is with the Upstream Research Division, ExxonMobil, Houston, TX J.N. DUPONT, R.D. Stout Distinguished Professor, is with the Materials Science and Engineering Department, Lehigh University, Bethlehem, PA Contact jnd1@lehigh.edu T. DEBROY, Professor, is with the Materials Science and Engineering Department, Pennsylvania State University, College Park, PA Manuscript submitted March 3, Article published online November 6, 2009 zone using electron-beam (EB) welds conducted on a high-purity Fe-Ni-Cr stainless-steel-type alloy. Solidification of the weld metal was observed to proceed via epitaxial growth from the base metal along the h100i (i.e., easy ) growth directions. They developed geometric models that demonstrate that the active h100i growth direction will be that which is most closely aligned to the maximum heat-flow direction. Knowledge of the growth direction permits calculation of the dendrite tip velocity, which is a function of the heatsource travel speed and the weld-pool geometry. Experimental weld trials also showed that the weld-pool shapes were independent of orientation with respect to the SX substrate, indicating that heat flow was isotropic in the SX base metal. The root cause for equiaxed grain formation was not discussed, because these formations were rarely observed in these high-purity alloys. Subsequent studies indicated that the simplified stainless steel composition did not induce the necessary extent of constitutional supercooling required for stray grain (SG) formation. Vitek et al. [6,7] expanded upon the work of Rappaz, and showed that all of the central features of SX solidification were present in SX Ni-base superalloy welds. The theory of constitutional supercooling was used to describe the formation of equiaxed grains as a function of solidification conditions. The relevant parameters for determining the degree of constitutional supercooling are the dendrite tip velocity, V hkl, and the temperature gradient parallel to the dendrite, G hkl, which are both functions of welding parameters. Experimental welds were conducted over a range of welding parameters in order to affect the range of G and V within the melt pool. The results confirmed that SG METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 41A, JANUARY

2 formation is a function of G/V and the resultant degree of constitutional supercooling. [8] The alloy composition influences SG formation by controlling the nucleation parameters and solidification temperature range, the latter of which was shown to be significantly higher in Ni superalloys than in the high-purity stainless steels. [9] Significantly more SGs were observed in the multicomponent superalloy weld zones than in the relatively pure Fe-Ni-Cr stainless steel welds studied by Rappaz, due to this effect. The dendrite fragmentation theory for SG formation was not supported by the results, which showed a significantly asymmetric SG distribution in welds conducted along asymmetric crystal orientations. [10] Dendrite fragmentation would not occur asymmetrically, but the limited growth directions would cause significantly different degrees of constitutional supercooling on each side of the weld. However, the effect of the substrate orientation on the overall SG fraction was not studied experimentally. Other SX welding experiments were performed by Gaumann et al. [11 13] Autogenous welds were used to validate constitutional supercooling-based models produced previously by Hunt. [13] Powder injection was also included in other weld trials to show the efficacy of cladding techniques in SX weld repair. [12] Elsewhere, a similar process known as laser-engineered net shaping (LENS) was used to form multilayer SX claddings with only a shallow layer of SGs on the top surface. [14] Yoshihiro et al. [15] recently investigated the microstructure of welds on alloy CMSX-4 prepared over a wide range of powers and travel speeds using both the laser and gas tungsten arc (GTA) heat sources. They identified three types of morphologies: single crystals with directional dendrites that only grow in the [001] direction from the bottom of the weld, SX with disoriented dendrites, and welds with SGs. The SX welds with disoriented dendrites simply indicate the presence of dendrites that grew in directions orthogonal to the [001] direction. Their results generally demonstrated that a reduction in power and an increase in travel speed are beneficial for preserving the SX structure. The successful processing window for the GTA welds was slightly smaller than for the laser welds. This is probably associated with the higher-intensity heat source of the laser that produces a higher temperature gradient. The substrate orientation relative to the direction of the heat-source travel speed can also influence the formation of SGs. For relatively simple conditions in which the direction of the heat-source travel is coincident with one of the crystallographic h100i easy-growth directions, the dendrite growth velocities across the line of symmetry in the fusion zone are equivalent. As a result, the tendency for SG formation is symmetrical about the weld centerline. There are likely to be conditions in which welding is required in an asymmetrical direction relative to the easy-growth directions. Park et al. investigated this effect on welds of the SX alloy Rene N5, [8] and demonstrated that the SG formation tendency will be different on each half of the weld for this condition. This effect is caused by differences in the growth angle and associated growth velocity across the weld centerline that occur due to the asymmetric welding condition. In addition, the temperature gradient in the direction of dendrite growth will decrease as the growth angle increases. Therefore, the G/V ratio will be different on each side of the weld. Park et al. showed that conditions can exist in which the critical G/V ratio required to reach the columnar to equiaxed transition (CET) is obtained on one side of the weld, while the G/V ratio on the other side is high enough to generally avoid the CET. Additional research by the same group [10] has demonstrated that the cooling rate (product of GV) is, indeed, symmetrical across the weld centerline. This carries important implications for identifying the mechanism of the CET. It is thought that the CET could also be induced by fragmented dendrites that are pushed into the solidification front by convection. Such dendrites would induce the CET by forming heterogeneous nucleation sites. In this case, convection would become more prevalent as GV decreases, and a decrease in the cooling rate would therefore be expected to lead to more SGs. The value of GV was shown to be constant across the fusion line for the experimental welds examined, but the extent of SG formation is not. Therefore, the dendrite fragmentation mechanism is not likely to be operable in fusion welding. Liu and DuPont [16,17] recently extended the analysis of Rappaz by combining the dendrite growth model with a mathematical model of the melt pool. The threedimensional (3-D) shape of the melt pool was modeled as the segment of an ellipsoid and was represented by four geometrical parameters. The melt-pool geometrical parameters are controlled by the heat- and fluid-flow conditions during processing and can be determined either computationally from a heat- and fluid-flow simulation or experimentally by directly measuring these parameters in situ or after processing. Coupling of this 3-D melt-pool model to the dendrite growth analysis permitted detailed investigations of the effects of both the melt-pool shape and the substrate orientation on the dendrite growth directions and velocities. A comparison of experimental and calculated dendrite growth directions were made and good agreement was observed. The results showed that the pool shape has a significant effect on the operable dendrite growth direction and resultant velocity. When the weld is relatively deep, growth is generally activated in four h100i-type directions. Dendrites grow from the bottom of the pool in the [001] direction, from the sides of the pool in the [010] and 0 10 directions, and from the back of the pool (along the heat-source travel direction) in the [100] direction. The favored growth along the [100] direction that is coincident with the heat-source direction causes the maximum growth velocity to equal that of the heat-source velocity in this location. When the weld is shallow, growth generally occurs from the bottom of the pool, and the maximum growth velocity is generally less than the heat-source velocity. Under asymmetrical welding conditions, the dendrite growth velocity can reach values that are up to 1.4 times that of the heat-source velocity near the top of the pool. 182 VOLUME 41A, JANUARY 2010 METALLURGICAL AND MATERIALS TRANSACTIONS A

3 The successful repair of SX alloys requires both a general understanding of the influence of possessing conditions on SG formation as well as a predictive tool to aid in identifying process parameters that will avoid/ minimize the occurrence of SGs. The formation of solidification cracks within the repair area also needs to be avoided. The development and application of heat-/ fluid-flow and solidification modeling techniques for predicting SG formation in welds have recently been described in a companion article. [18] In that work, a detailed heat-/fluid-flow model was first validated for prediction of the melt-pool shape and the variation in the temperature gradient around the melt pool. The heat-/fluid-flow results were then integrated into a solidification model for determining the active growth directions as well as the temperature gradient and solidification velocity along the dendrite growth direction as an aid to predicting conditions that lead to the formation of SGs. Details of the modeling approach and validation are explained in a separate article. [18] The influence of welding parameters on solidification cracking susceptibility has also been investigated. [19] In that work, it was observed that solidification cracking susceptibility can be minimized by utilizing low heat input welding conditions. For the conditions considered, process maps were developed for producing crack-free welds with both the GTA and laser welding processes. The successful processing window was larger for the laser process. This was attributed to the higher temperature gradient associated with the laser process, which helps minimize SGs and the associated formation of solidification cracks that form along the grain boundaries. The objective of this article is to describe the general effects of welding parameters, substrate orientation, and welding process type on the development of SGs over a wide range of conditions and also to describe the formation of different types of SGs that were observed experimentally. Selected results from the heat-/ fluid-flow and solidification model are used to help develop a detailed understanding of these effects. II. EXPERIMENTAL PROCEDURE A series of 105 welds was prepared with a variety of welding processes and welding parameters. The amount of power actually absorbed by the substrate will depend on the heat-source output power (P output ) and the transfer efficiency (g) of the particular process: P absorbed ¼ g P output ½1Š The absorbed power is the term that affects heat and fluid flow within the melt pool. Thus, all weld powers discussed in this article are the absorbed powers. The output welding power was selected such that the absorbed power for each process would be similar. The value of g used in calculating P absorbed for each process was taken from measurements previously available in the literature. Table I summarizes the welding power and travel speed of the autogenous welds prepared. Table I. Selected Welding Parameters for EB (E), Laser (L), and GTA (G) Welds Travel Absorbed Power (W) Speed (mm/s) ELG ELG EL E E 6 ELG ELG ELG E E E 25 ELG ELG ELG E E E 50 EL ELG ELG ELG E E 75 E E E E E 100 EG EG E E Table II. Composition of Base Metal CMSX-4 Used in This Study (Values in Weight Percent) Element CMSX-4 Ni bal C Cr 6.36 Co 9.68 Mo 0.63 W 6.34 Ta 6.52 Ti 1 Al 5.62 B 0 Zr 0 Hf 0.1 Re 2.87 The Ni-base superalloy CMSX-4 was selected as a representative alloy due to its use in industrial SX applications and the availability of material properties. The composition of the CMSX-4 substrate was determined through wet-chemical analysis and is provided in Table II. Substrates with dimensions of mm were cast such that the (001) crystal planes were parallel to the surface. The substrates were solution heat treated with a schedule used in industrial practice for alloy CMSX-4 (heated at 1310 C for 7 hours under vacuum and gas fan cooled to 982 C). Laser weld trials were conducted using a 750W Nd:YAG laser housed within a LENS unit manufactured by Optomec, Inc., Albuquerque, NM. The transfer efficiency of the LENS unit was previously measured to be g ~ 0.5. [20] Weld passes were conducted with absorbed powers (P) of 120, 180, and 250 W at travel speeds (S) of 1, 6, 24, and 47 mm/s. Welds were conducted along the [100], [110], and [120] crystal directions along the (001) crystal plane. Electron-beam (EB) welds were prepared with absorbed powers of P = 120, 180, 250, 500, 1000, and 1500 W and travel speeds of S = 1, 6, 24, 47, 75, and 95 mm/s. The transfer efficiency for EB welds is known to be very high, [21] so it was taken as ~1.0. Several of the lowest heat input welds were not studied due to lack of melting; the highest heat input weld (1500 W, 1.1 mm/s) was also omitted due to complete penetration and melt through of the substrate. The EB was underfocused in METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 41A, JANUARY

4 order to minimize the keyhole mode of energy transfer. The resulting beam produced elliptical weld pools similar to those made by the laser beam. Finally, autogenous welds were also conducted with the GTA welding weld technique. The transfer efficiency was taken as g ~ 0.7. [22] Welds were performed using absorbed powers ranging from 120 to 500 W, and the travel speeds were set to 1, 6, 24, 47, and 100 mm/s. All welds were prepared using standard metallographic techniques to a polish of 0.05 lm using colloidal SiO 2. The weld structures were revealed by immersing the specimens for 5 seconds in a reagent consisting of 50 ml HCl, 50 ml H 2 O, and 2.5 g CuCl 2. Light optical microscopy (LOM) photomicrographs were taken from each weld cross section in order to measure the weldpool dimensions and observe the particular arrangement of the h100i growth variants. The samples were then repolished to their prior finish in order to conduct orientation imaging microscopy (OIM). This technique was used to map the crystallographic orientation of the entire weld cross section; SGs were demarcated by a deviation of 5 deg from the base metal orientation. The OIM was conducted on a Hitachi 4300 Schottky (Pleasanton, CA) field emission gun scanning electron microscope operating at an accelerating voltage of 15 kv. The microscope was equipped with a Hikari (Pleasanton, CA) electron backscattered diffraction (EBSD) detector and OIM Data Collection software (Edax, Mahwah, NJ), version 5.2. The OIM analysis was used to map the location of SGs and quantify the overall SG content. The SG area fraction was determined by dividing the area of SGs (as determined by OIM) by the total fusion zone area (as measured from LOM photomicrographs), and the average of three different cross sections was acquired for each weld condition. The OIM measurements were not performed on the to 1500-W welds with S 6 mm/s, due to the large cross-sectional area of these fusion zones. Heat-transfer and fluid-flow simulations were conducted in order to determine the melt-pool shape and local solidification conditions in the welds. The modeling approach is described in detail in a separate article [18] and is only briefly summarized here. Modeling of the melt-pool temperature fields and fluid velocities was conducted using the heat-transfer and fluid-flow code developed by DebRoy et al. [23,24] These calculations require the input of several material properties of the substrate, and these were found in the literature. [25,26] The calculated weld-pool shape was used to determine the local dendrite growth directions based on the minimum undercooling criterion, in which dendrites grow in the h100i direction that is nearest to the solidification interface normal. [3] Temperature data derived from the simulations were used to calculate the temperature gradient in the liquid parallel to the local dendrite growth direction (G hkl ). The local dendrite tip velocity (V hkl ) was determined from the weld-pool shape. The calculated solidification parameter data were used to predict the propensity to form SGs throughout the weld pool. The expression for the area fraction of equiaxed grains (/) was first implemented by Hunt [27] in the description of the columnar-to-equiaxed transition observed in castings. It was calculated in the same manner as Vitek [28] from equations derived by Gaumann [13] and is given by 8! 9 < / ¼ 1 exp 4pN = : 3 ðn þ 1ÞðG n =avþ 1=n ; ¼ 2: V 3=3:4 G 3:4 ½2Š where a = s K 3.4 /m, n = 3.4, and the nuclei density N 0 = /m 3 for alloy CMSX-4, as determined by Gaumann. The / value was calculated along each point of the solid/liquid interface using the calculated G hkl and V hkl as inputs in Eq. [2]. Following the method proposed by Vitek, [28] a single value, F, representing the SG area fraction for a particular weld condition, was determined by calculating the areaweighted average of / given by P i U ¼ A P i/ i i A ½3Š i where A and / are the area of solid/liquid interface and the SG area fraction for a particular location I, respectively. Because F takes into account the entire solidification front, it is more indicative of SG formation than G/V, which varies throughout the weld pool. A more detailed discussion of the modeling efforts is provided in a separate article. [18] III. RESULTS Fig. 1 contains typical examples of SX weld microstructures produced in this study. The OIM-generated SG maps are superimposed upon LOM photomicrographs of the laser weld microstructure from which they were gathered. Fig. 1(a) displays a microstructure generated with low heat input (120 W, 24 mm/s), while Fig. 1(b) was conducted with a higher heat input (250 W, 1 mm/s). Fig. 1(c) was also produced with a low heat input, but in this case, along an asymmetric welding orientation ([120], 250 W, 47 mm/s). Uncolored regions of the weld zone indicate regions of SX solidification, because the crystal orientation there matched that of the base metal. In order to facilitate later discussion on the different varieties of SGs, Fig. 1 has different regions of SGs labeled based on their relative location within the fusion zone. A solidification crack is also evident in Fig. 1(b). This type of cracking can generally be avoided by minimizing SG formation, because these cracks form predominately in the grain boundaries. The cracking behavior will be the subject of a separate article. [19] Fig. 2 describes the overall SG area fraction as a function of the travel speed for the laser-beam welds conducted in the following crystal directions: (1) [100], (2) [120], and (3) [110]. Each curve represents a particular welding power, and the error bars for each data point are the standard deviation from three separate measurements for each condition. The maximum 184 VOLUME 41A, JANUARY 2010 METALLURGICAL AND MATERIALS TRANSACTIONS A

5 Fig. 1 Three laser weld microstructures with superimposed OIM maps: (a) [110]-120 W-24 mm/s, (b) [110]-250 W-1 mm/s, and (c) [120]-250 W-47 mm/s. SG area fraction is generally observed at the lowest travel speeds for the [110] and [120] welds but is not reached until an intermediate travel speed of ~6 mm/s for the [100] welds. Fig. 3 shows a similar plot for the EB welds. All EB weld data are plotted in Fig. 3(a), while those welds with an SG content of less than 0.2 are shown in Fig. 3(b). A limited set of SG measurements were performed on the GTA weld structures. Those results are shown relative to equivalent sets of data from welds conducted using high-energy density (HED) processes in Fig. 4. IV. DISCUSSION A. Effect of Welding Parameters The majority of welds were conducted parallel to the [100] crystal direction, so these welds will be used to describe the effects of the welding parameters on SG formation. The influence of substrate orientation and process type are discussed in Sections IV B and IV C. For the conditions examined here, the maximum SG fraction for all beam powers occurred at a travel speed of ~6 mm/s. Beyond this travel speed, the SG fraction was observed to decrease. The transition in behavior can be explained based on the changes in the temperature gradient (G), growth rate (V), G/V ratio, and resultant degree of constitutional supercooling. At low travel speeds, the increase in travel speed causes an increase in V with only a moderate change in G. As a result, the G/V ratio is generally reduced, and more SGs form due to increased constitutional supercooling. Further increases in travel speed lead to a significant increase in G with a corresponding increase in the G/V ratio, which leads to a reduction in the SG content. This observed effect of the travel speed is generally consistent with the previous results reported in the literature. For example, Vitek performed a theoretical analysis [28] of the effects of welding parameters on SG content. His results demonstrated a similar transition in behavior at ~14 mm/s. The effect of beam power is not as apparent, which may be due to the small range of powers utilized for the laserbeam welds (120 to 250 W). The EB welds were produced over a wider range of welding parameters, up to a 1500-W beam power and 95-mm/s travel speed. The OIM data plotted in Fig. 3(a) reveal that the high-power welds (1000 to 1500 W) contained very high SG area fractions at each level of travel speed. This is a direct result of a reduction in G around the weld pool under a highpower beam. Welds conducted with lower beam power METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 41A, JANUARY

6 Fig. 2 Effect of welding parameters on SG area fraction within the laser-beam weld structures for several substrate orientations: (a) [100] (001), (b) [120] (001), and (c) [110] (001). Fig. 3 Effect of welding parameters on SG area fraction within the EB weld structures for the [100] (001) substrate orientation. (Fig. 3(b)) showed less differentiation in SG content, exactly like in the laser welds. The effect of travel speed was also very similar. While the low-power welds showed a fairly constant SG fraction at travel speeds beyond 24 mm/s, they did contain a maximum SG fraction at ~6 mm/s. 186 VOLUME 41A, JANUARY 2010 METALLURGICAL AND MATERIALS TRANSACTIONS A

7 Fig. 4 Effect of welding technique on SG area fraction for an equivalent melting power over a range of travel speeds. Fig. 5 Effect of welding parameters on the predicted SG area fraction within simulated EB weld pools for the [100] (001) substrate orientation. The general trends in both sets of data show that the preferred welding parameter selection for promoting SX solidification is low power and high travel speed. Although this conclusion matches previous research, [29] the OIM data here also indicate that similar levels of SG may be possible with low powers and low travel speeds. For example, the EB weld with P = 120 W and S = 1 mm/s contained an SG fraction similar to welds produced at an equivalent power with higher travel speeds. These results may be important for selecting the process parameters needed to cover the wide range of repair conditions encountered in practice. These findings are corroborated by the results derived from heattransfer and fluid-flow modeling. Fig. 5 shows the calculated SG fraction F as a function of travel speed for various beam powers. Although the magnitude of the F values differ from the experimental values, the simulation results exhibit the major features of the Fig. 6 Effect of substrate orientation on the 12 sets of laser welding parameter combinations. experimental data. The SG fraction decreases at high travel speed, and there is an initial increase under lowpower/low-travel-speed conditions. The calculations also indicate that higher beam powers induce more SG formation. The 120-W, 1-mm/s calculated weld contained the lowest overall SG content, similar to that found experimentally. A more thorough discussion of the modeling work is provided in a separate article. [18] B. Effect of Substrate Orientation The SG content of the [100] laser welds was studied using a set of 12 combinations of welding parameters. The same set of parameter combinations were conducted along [110] and [120] crystal directions, allowing observations on the effects of substrate orientation. Because SX growth is limited to the h100i directions, changing the substrate orientation alters the distribution of G hkl and V hkl across the solidification front. It is important to note that while the [100] and [110] directions are symmetrical about the weld centerline, the [120] welding direction is asymmetrical. The SX growth in the [120] weld pool will thus not be equivalent on each side of the weld centerline. Fig. 2(b) and (c) show the effect of welding parameters on the [120] and [110] laser welds, respectively. Comparison is made of the SG fractions in all three substrates for the 12 combinations of welding parameters in Fig. 6. The data show that the overall SG content generally decreases as the welding direction shifts from [100] to [120] to [110]. Of the 12 sets of welding parameters considered, only one of the welding parameter combinations increased over this range, while two combinations produced a maximum SG fraction under the [120] orientation. The root cause for this effect lays in the distribution of dendrite growth regions. While G and V normal to the solidification interface (i.e., G sl and V sl ) do not change with substrate orientation, the distribution of the critical values G hkl and V hkl (i.e., G and V along the growth direction) will be strongly affected. Weld-pool simulations revealed that the regions in which G hkl /V hkl are lowest are located along the junctions between dendrite METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 41A, JANUARY

8 Fig. 7 (a) Distribution of the h100i dendrite growth variants along the solid/liquid interface of the simulated [100] 250-W 1.5-mm/s EB weld pool. (b) Predicted stray grain volume fraction parameter / along the solid/liquid interface for the identical conditions. growth regions. This is demonstrated in Fig. 7, which plots the active h100i growth regions (Fig. 7(a)) and SG formation tendency (Fig. 7(b)) across the entire solidification interface. Note that the greatest SG concentrations (in blue) are generally present where the dendrite growth regions converge. In these areas, the dendrite growth axes do not run parallel to the direction of greatest heat flow. Previous experimental research by Mokadem et al. [30] described the significance of this and referred to this phenomenon as off-axis heat flow. The off-axis heat flow will reduce G hkl and will induce a faster V hkl than if the dendrites grow parallel to the maximum heat-flow direction. The resulting low G/V at these locations will cause substantial SG formation. Mokadem reached those conclusions by performing [100] welds along a multitude of crystal planes. Here, a variety of directions were studied upon an SX plane. The experiments are complementary to one another, but the conclusions are the same. The amount and location of these junction points will control the degree of SG formation. In these welds, the amount of [001] dendrites growing from the weld root was generally similar regardless of welding orientation, as shown in Fig. 8. Because the extent of the transition in dendrite growth direction caused by the [001] dendrites is similar, discussion here will concentrate on the junctions between [100] and [010] dendrites along the weld surface. The schematic in Fig. 9 illustrates the general number and location of the dendrite growth regions along the weld surface for each substrate orientation studied. The photomicrograph presented for each orientation in Fig. 9 was taken from the 180-W, 6-mm/s laser weld condition. (As discussed later in this article, the centerline SGs seen in the [120] and [110] orientations do not represent epitaxial SX growth. The different regions arise because epitaxial growth is limited to the h100i easy-growth directions.) The presence of two transitions in the dendrite growth direction in the [100] welds produced more SGs than the single transition in the [110] welds (Figure 6). The [120] welds represent the intermediate case. Because one of the junctions is positioned near the fusion line at which G sl /V sl is relatively high, it will have less of an overall impact on SG formation. These trends agree very well with the trends in SG contents displayed in Fig. 6. The shapes of the curves in Fig. 2 also imply that the [120] direction is an intermediate condition: the low-power curves appear similar to the [110] data, while the 250-W curve is very similar to the [100] data. The transition in behavior is likely a function of the weld-pool shape produced by each welding power and the relative location of the transition in dendrite growth direction on that shape. The high laser power would produce a more circular weld pool, so the second transition in dendrite growth direction would be closer to the back of the weld pool, producing a state similar to the [100] substrate orientation. It appears that reducing the amount of transitions in dendrite growth directions causes a decrease in the SG fraction as the travel speed increases. With the exception of the [120], 250-W weld, the OIM data indicate a maximum SG area fraction at the minimum travel speed in both the [110] and [120] data. Increasing travel speed always had a positive effect on SX weld development by increasing the G/V ratio. The effect of beam power is also more pronounced in these welds, especially at low travel speeds. Increased power induced more SG nucleation by reducing G. C. Effect of Welding Process Type The measured SG area fractions for the three welding processes at travel speeds from 1 to 24 mm/s are compared in Fig. 4. The data for the GTA welds showed a maximum SG content at the intermediate travel speed, just as was observed in the HED weld grain structures. The SG area fraction of the GTA welds was greater than that found in either of the HED weld structures for all travel speeds, and the SG amount in the laser-beam welds is always greater than that observed in the EB welds. The trend in the SG content 188 VOLUME 41A, JANUARY 2010 METALLURGICAL AND MATERIALS TRANSACTIONS A

9 Fig. 8 LOM micrographs from three 180-W, 6-mm/s EB welds: (a) [100], (b) [120], and (c) [110]. The [001] dendrite region growing vertically from the weld root is roughly identical in all substrate orientations. Fig. 9 Distribution of dendrite growth regions along the weld surface for three substrate orientations: (a) [100] (001), (b) [120] (001), and (c) [110] (001). METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 41A, JANUARY

10 between the three processes coincides with the effect of the energy density on the temperature gradient. The energy density of the heat source influences the temperature gradient in the weld pool, where welds produced with higher-energy density processes will experience higher temperature gradients. Thus, welds produced with higher-energy density processes are expected to exhibit lower SG contents than lower-energy density processes at equivalent levels of input power and travel speed. D. Strain Grain Characterization The three types of SGs defined in Fig. 1 provide information about the formation of SGs and their pertinent location within the weld. Each type of SG will be discussed separately, in order to effectively highlight the appearance and properties characteristic to each variety. 1. Type 1: Crown SGs Type 1 SGs (SG1) are confined to the crown of the weld. The SG1s were observed in the fusion zone of welds in which dendrites grew parallel to the substrate surface (e.g., [100], [010], etc.). The most significant exception was low heat input welds in which [001] dendrites dominated the fusion zone from the root to the crown. The crown of the weld is the region of the fusion zone in which SGs are most commonly anticipated [7] due to the relatively low value of G/V in this region. While the magnitude of G/V in this region indicates a high degree of SG formation, its position as the final location of weld metal solidification also affects the character of the SGs observed in this area of the weldment. The SGs form in the liquid ahead of the columnar front and thus may be susceptible to being swept about the melt pool via convection currents. The SGs observed at the weld crown were large, indicating that they must have had significant time to grow before being locked into place by the oncoming dendrite front. Although nuclei trajectories were not calculated in this research, previous work has shown that particles often undergo considerable recirculatory motion before being trapped by the solidifying track. [31,32] The extra time may have been provided by the movement of the recently formed SGs within the weld pool until they reached the surface and were met by the SX front. Figure 10 displays typical fluid velocity vectors along the solidifying side of a simulated weld pool along the substrate surface and along the depth of the pool at the weld centerline. Fluid currents move from the heat-source location outward to the fusion line, driven by Marangoni convection. The SGs that reach the surface would be transported to the columnar dendritic front, but SGs below the surface would be carried away by fluid currents in the opposite direction in a recirculatory motion toward the heat source. In this way, fluid flow in the weld pool can alter the locations at which the SGs appear and the SX weld microstructure from predictions based solely on local solidification parameters. The SG1s at the crown will grow only for a finite amount of time because of their Fig. 10 Fluid velocity vectors along surface and centerline of a simulated 180-W, 24-mm/s laser weld pool. unfavorable crystallographic orientations. Because the orientation of an SG1 at the crown is less favorable than the base metal orientation, it will eventually be overcome by dendrites growing in the preferred SX orientation. Comparing several cross sections showed that the SG1 grains in the crown were highly random in both distribution and quantity. The standard deviation error bars in Fig. 2 and Fig. 3 are a direct consequence of this random distribution. A special type of SG1 was observed in a subset of welds. The OIM maps of the 120- and 180-W laser welds conducted along the [110] and [120] directions at a travel speed of 24 mm/s showed an SG of similar dimensions in identical locations across multiple weld cross sections. Etching of the cross section showed it to be present directly at the junction of three dendrite growth regions (Fig. 1(a)). Furthermore, the grain was always found in the EBSD measurements to have a similar orientation, in which the [100] crystal direction was perpendicular to the metallographic plane. The presence of a single centerline grain over the entire length of the weld was confirmed by metallographic preparation of the weld surface in order to observe the SG distribution in the weld crown. The photomicrograph in Fig. 11 shows a single centerline SG growing parallel to the welding velocity along the weld centerline in the [120] 120-W 24-mm/s laser weld. Once such a grain is formed via typical equiaxed grain formation, it can proceed to grow along the entire length of a weld. This makes it different from the other SG1s, which can eventually be overcome by the SX columnar front of the preferred orientation. This type of centerline grain formation has previously been observed and explained by Li and Brooks. [33] As described in that work, a centerline grain can actually extend into the weld pool, making its removal by competitive grain growth exceedingly difficult. More recently, Dye et al. [21] proposed a geometric model for predicting the occurrence of centerline grains in 190 VOLUME 41A, JANUARY 2010 METALLURGICAL AND MATERIALS TRANSACTIONS A

11 Fig. 11 [120] 120-W 24-mm/s laser weld as viewed from the top surface. Centerline grain found in this weld was observed along the entire length of the metallographic mount. polycrystalline superalloys. In that work, the distance between the liquidus and the dendrite tip temperatures at the weld centerline (Dx und ) is compared to the radius of curvature (q) of the liquidus isotherm at the trailing edge. A centerline grain boundary is not expected to form when the ratio Dx und /q is small. An expression for Dx und /q was obtained from the two-dimensional conduction heat-flow solution and the dendrite growth law. From a practical perspective of weld repair applications, SG1s represent the least troublesome variety of SG. The weld repair of gas turbine blades will typically require deposits of multiple layers. Successive layers will thus remelt the crown of the previous passes. In this way, SG1s can be melted and replaced with SX solidification at the root of the subsequent pass. The final weld pass could then be machined off, leaving behind a fully SX build. This approach has recently been used to demonstrate direct deposition of a 12-layer SX deposit. [14] However, other types of SG were observed in the weld microstructures that do not occupy this preferred position. 2. Type 2: SGs at Growth Direction Transitions Type 2 SGs are found in regions representing the transition between different h100i dendrite zones. In Fig. 1(c), SG2s are observed between the [001] and [010] dendrite zones. This weld, having been conducted at a relatively high travel speed of 47 mm/s, solidified without a [100] dendrite zone that grows parallel to the heat-source velocity. This is likely a consequence of the weld-pool shape in which the weld pool becomes shallow and elongated and the [001] dendrites growing from the bottom of the weld pool dominate growth such that no [100] dendrites form. This appears to have inhibited SGs from occupying the crown of the weld upon final solidification (Fig. 1(c)). Instead, SGs are found along the transition between the vertical [001] dendrites and the horizontal [010] and 0 10 dendrites. This appeared in all of the high (>6-mm/s) travel speed welds. These results can be understood with the reference to Fig. 7 described earlier, which shows the distribution of / across the solidification interface for the simulated [100] 250-W 1-mm/s EB weld. The regions in which / is greatest coincide with the junctions of the dendrite growth regions. The off-axis heat flow described earlier is the primary cause for this phenomenon. These locations experience the lowest values of G hkl because the angle between the interface normal and the dendrite growth vectors is highest. [30] The SG2s defined in this study are located in precisely these locations. For this type of SG, it may be possible to presume that the observed location of the SG2s coincides with the site of their nucleation because of their small size. Based on this presumption, SG2s should also be nucleating along the weld centerline between the [001] and [100] growth regions. The distinction between these grains and the SG1s observed in the weld crown would be difficult, because the [001]/[100] intersection occurs near the crown of the weld. The SG2 label is therefore applied only to SGs observed between the [010] and [001] growth regions. 3. Type 3: Root SGs The high-power, low-travel-speed weld microstructures exhibited an additional position for SGs to form. As shown in Fig. 1(b), welds made with high heat input were observed to contain SGs at the fusion line near the root of the weld. The high concentration of grain boundaries in these locations provided locations for solidification cracks, as highlighted in Fig. 1(b). Solidification cracks have long been associated with SG formation. [6,10] This type of weld defect can effectively be prevented in SX weld passes if SGs are avoided. A thorough discussion of solidification cracking is provided in a separate article. [19] The weld root is typically associated with relatively high temperature gradients and negligible growth rates. Because the magnitude of G/V would be high, SG nucleation would not normally be expected at this location. A representative higher-resolution OIM map of this type of SG is shown in Fig. 12, in which it is evident that epitaxial growth occurred at the fusion line. This indicates that SG3s are not the product of crystal imperfections or second phases in the base metal. Rather, the first appearance of SG3s appears to coincide with the breakdown of planar growth. The SGs then continue to grow perpendicular to the fusion line in a manner comparable to a polycrystalline material, in which competition between randomly oriented grains forces grains with the preferred h100i orientation to take over the solidifying front. [34] Analysis of the OIM data confirmed that the h100i crystal direction is coincident with the long axes of SG3s. Although SG3s were observed under a variety of welding conditions, the largest networks of this type of SG were observed in welds that also possessed an inflection point along the fusion line (labeled in Fig. 12). The fusion line discontinuity and SG3 networks both appear to be the result of fluid flow in the weld pool. As mentioned previously, it is important to note that the location of an SG in the solidified weld microstructure is not necessarily an indication of its METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 41A, JANUARY

12 Fig. 12 High-resolution map of SG3 network found in the [120] 180-W 1-mm/s weld pass. 1. For the conditions examined in this work, the SG formation tendency of alloy CMSX-4 is highest at a travel speed of ~6 mm/s. The amount of SGs decreases at higher travel speeds, due to an increase in the G/V ratio. The amount of SGs increases with increasing weld power. 2. The HED welding processes produced lower SG fractions than the GTA process at equivalent input powers. This is attributed to the higher temperature gradients associated with the HED processes. 3. Dendrite growth junctions near the back of the weld pool are the most likely location for SG formation due to off-axis heat flow in combination with relatively high G/V conditions. 4. For a given set of welding conditions, the [100]/ (001) substrate orientation generally produced the highest amount of SGs, while the [110]/(001) orientation produced the lowest amount. The [120]/(001) orientation produced intermediate amounts of SGs. This effect is attributed to differences in the number and location of dendrite growth intersection regions within the melt pool for each substrate orientation. Fig. 13 Fluid velocity vectors in cross section at the maximum width of a calculated 250-W, 1-mm/s laser weld pool. Convergence of two opposing currents may induce the high SG fractions observed at this location in the weld root. point of origin. Fluid flow in the weld pool has been shown to be a significant factor in the weld-pool shape and temperature fields. [23] Two opposing fluid currents along the weld root can produce this weld-pool shape. Heat-transfer and fluid-flow modeling of these weld pools indicated that this behavior is highly likely under these welding conditions. A cross section of a 250-W, 1-mm/s laser weld-pool simulation (i.e., same parameters as shown in Fig. 1(b)) displaying opposing fluid currents is shown in Fig. 13. The location of the fluid convergence coincides with the location of the inflection point observed in the experimental weld microstructures. These opposing fluid currents could also draw SGs from throughout the weld pool and deposit them where the opposing currents meet. The high concentration of SGs at the weld root is possibly due to the convection currents shown along the weld centerline in Fig. 10. V. CONCLUSIONS The influence of welding parameters, process type, and substrate orientation on SG formation in the SX alloy CMSX-4 has been investigated by LOM, OIM, and heat-/fluid-flow simulations. The following conclusions can be drawn from this work. ACKNOWLEDGMENTS The authors acknowledge the financial support of The National Science Foundation (Arlington, VA) through Grant No The SX substrates used for welding trials were cast by Howmet (Whitehall, MI), a division of Alcoa, Inc. The EB welding experiments were made possible by Thomas Lienert and Paul Burgardt of the Los Alamos National Laboratory (Los Alamos, NM). REFERENCES 1. R.C. Reed: The Superalloys: Fundamentals and Applications, Cambridge University Press, Cambridge, United Kingdom, 2006, pp A. Cullison: Weld. J., 2003, vol. 82, pp M. Rappaz, S.A. David, J.M. Vitek, and L.A. Boatner: Metall. Trans. A, 1989, vol. 20A, pp S.A. David, J.M. Vitek, M. Rappaz, and L.A. Boatner: Metall. Trans. A, 1990, vol. 21A, pp M. Rappaz, S.A. David, J.M. Vitek, and L.A. Boatner: Metall. Trans. A, 1990, vol. 21A, pp S.A. David, J.M. Vitek, S.S. Babu, L.A. Boatner, and R.W. Reed: Sci. Technol. Weld. Join., 1997, vol. 2, pp J.M. Vitek, S.A. David, and L.A. Boatner: Sci. Technol. Weld. Join., 1997, vol. 2, pp J.-W. Park, S.S. Babu, J.M. Vitek, E.A. Kenik, and S.A. David: J. Appl. Phys., 2003, vol. 94, pp J.M. Vitek, S.S. Babu, S.A. David, and J.-W. Park: Mater. Sci. Forum, 2003, vols , pp J.W. Park, J.M. Vitek, S.S. Babu, and S.A. David: Sci. Technol. Weld. Join., 2004, vol. 9, pp M. Gaumann, R. Trivedi, and W. Kurz: Mater. Sci. Eng., A, 1997, vols , pp M. Gaumann, S. Henry, F. Cleton, J.-D. Wagniere, and W. Kurz: Mater. Sci. Eng., 1999, vol. 271, pp M. Gaumann, C. Bezencon, P. Canalis, and W. Kurz: Acta Mater., 2001, vol. 49, pp W. Liu and J. N. DuPont: Metall. Mater. Trans. A, 2005, vol. 36A, pp VOLUME 41A, JANUARY 2010 METALLURGICAL AND MATERIALS TRANSACTIONS A

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