ADVANCING DEVELOPMENT AND APPLICATION OF SUPERALLOYS

Planar Front Growth of Single Crystal Ni-Based SuperalloyRene N515

SAE MATSUNAGA ,1,3 DUO HUANG,1 SAMUEL B. INMAN,1

JACK C. MASON,1 DOUG KONITZER,2 DAVID R. JOHNSON,1

and MICHAEL S. TITUS1

1.—School of Materials Engineering, Purdue University, West Lafayette, IN 47907, USA. 2.—GEAviation, Evendale, OH 45215, USA. 3.—e-mail: smatsun@purdue.edu

A Ni-based superalloy, Rene N515, was fabricated by the levitation zonemelting process to produce planar front single crystal growth. Microstructure,alloy composition, and microhardness properties of the as-zone melted andsolution heat treated alloy were investigated and compared with the direc-tionally solidified state to study the effect of microsegregation on these alloycharacteristics. In the planar front region, the dendritic structure andmicrosegregation were eliminated, and the density of casting porosity de-creased as well. We modeled the zone melting/planar growth process bysolving the zone melting mass balance equation numerically, and the mea-sured and predicted composition profile of the planar front were in goodagreement. The Vickers microhardness did not vary throughout the planarfront region. These results indicate that a nearly homogeneous planar front ofa 2nd-generation single Ni-based superalloy can be fabricated via levitationzone melting, and common casting defects can be eliminated entirely.

INTRODUCTION

Ni-based superalloys have been widely used forhigh-temperature applications such as turbineblades for jet propulsion and power generationplants due to their superior balance of propertiesat elevated temperature that include resistance tocreep, fatigue, and oxidation.1 These properties arein part obtained by the unique microstructure of Ni-based superalloys which is composed of a solid-solution disordered face-centered cubic c matrixwith coherently embedded c¢ precipitates with aL12-Pm�3m crystal structure.1 The high pressureturbine blades in these turbine engines are exposedto corrosive environments at elevated temperatureand are subjected to severe stresses resulting fromcentrifugal and aerodynamic forces.2 These extremeoperating conditions motivated the development ofnew materials that significantly improve the creepresistance of these alloys.

The advent of new materials and manufacturingmethods culminated in the development ofadvanced directional solidification processes in the1970 s, which included the Bridgman process that isable to produce ‘‘single crystal’’ turbine blades

containing virtually no grain boundaries.2 In thismethod, a Ni-based superalloy charge is melted (viainduction or convection heating) and then pouredinto an investment mold held at a temperatureabove the liquidus temperature of the alloy. Themold and superalloy are then withdrawn downwardinto vacuum or other coolant as shown in Fig. 1a. Tomaintain dendritic solidification, the mold andsuperalloy are withdrawn at a constant withdrawalrate and thermal gradient, as shown in the v-Gprocessing map in Fig. 1b. Maintaining a steadythermal gradient enables the growth of highlyoriented [001] dendrites, which nearly eliminatesgrain boundaries in the single crystal. The absenceof grain boundaries prevents the formation of voidsand cavitation which are primary crack initiationsites during creep in equiaxed microstructures.3

Additionally, the absence of grain boundaries allowsfor the removal of grain boundary strengtheningelements in the alloy composition that form carbidesand borides, which can serve as initiation sites infatigue in Ni-based superalloys.4

In addition to new solidification processes, enginemanufacturers have improved creep resistance byalloying with specific elements, particularly

JOM

https://doi.org/10.1007/s11837-020-04091-x� 2020 The Minerals, Metals & Materials Society

refractory elements such as Ti, Nb, Ta, W, Hf, andRe.2 Rhenium, in particular, is a potent alloyingelement that significantly improves the creep resis-tance of these alloys, though the exact mechanism isstill not well understood.5 Alloying of heavy ele-ments can negatively influence the castability ofsuperalloys due to a tendency for these alloyingelements to create convective instabilities in themushy zone (liquid + solid) and significant residualmicrosegregation. During directional solidification,microsegregation occurs due to non-equilibriumsolidification and is observed as compositional inho-mogeneity throughout the dendritic microstruc-ture.2,6 Microsegregation can cause the formationof low-melting temperature or brittle phases and anon-uniform distribution of c¢ precipitates,7 whichare detrimental to mechanical properties.8 Further-more, because the dendrite growth direction isantiparallel to gravity in the Bridgman process, adensity inversion can form between the liquid metalin the mushy zone and the liquid metal outside ofthe mushy zone. This inversion can initiate ther-mosolutal convection and cause the formation of thecasting defects known as freckles.9 Freckle chainsare small chains of equiaxed grains oriented alongthe solidification direction and form due to highlyenriched solute content in the liquid that suddenlyfreezes.10 The compositions of these grains differfrom the bulk composition; they are typically

enriched in the elements concentrated at the inter-dendritic region and may contain high angle grainboundaries that can act as crack initiation sites.1,9

Fabrication of a truly homogeneous microstruc-ture free of secondary phases, grain boundaries,porosity, and microsegregation may then improvethe mechanical and environmental properties ofadvanced Ni-based superalloys. Some work hasbeen done to elucidate the creep and fatigue lifecredit afforded by eliminating casting defects.11–13

This study has been conducted to understand theplanar front solidification process of advanced Ni-based superalloys. In this study, a Ni-based super-alloy was fabricated via planar front growth in orderto explore the possibility of eliminating microsegre-gation and secondary phase formation. The compo-sition, microstructure, and microhardness profileswill be presented. A thermodynamic model wasdeveloped to predict the as-solidified concentrationand phase fraction profile. The results from exper-imental studies and modeling will be compared, andthe influence of planar front microstructure tocompositional and mechanical properties will bediscussed.

METHODOLOGY

A Rene N515 [Ni – 6.95 Cr – 7.87 Co – 1.29 Mo –2.10 W – 14.43 Al – 2.22 Ta – 0.498 Re – 0.208 Hf –0.206 C – 0.023 B (at.%)] directionally solidified bar,15 mm in diameter and approximately 150 mm

Fig. 1. Schematic images of (a) Bridgman casting process, and (b) the processing map of the directional solidification of single crystal Ni-basedsuperalloys. (b) Was adapted from references7 and,24 and the actual solidification rate for the levitation zone melting process was added usingthe dotted line and square symbol. The thermal gradient can be varied within the planar front region with a constant solidification rate.

Matsunaga, Huang, Inman, Mason, Konitzer, Johnson, and Titus

long, was provided in the as-cast state by GEAviation and used to fabricate a planar front grownbar. The as-cast bar was sectioned in half andfastened in a levitation zone melting furnace to topand bottom grips. Levitation zone melting14–16 wascarried out using a 25-mm single turn induction coiloperating at 17 kW under an Ar-5%H2 atmosphere;a schematic of the levitation zone melter is shown inFig. 2.15 The two ends of the bars were brought towithin 5 mm of each other until the molten zonewas formed. Once the top of the bottom bar startedto melt, the top bar was translated downward tocreate a molten zone between the two bars. Inlevitation zone melting, the liquid is partiallylevitated between the upper and lower bars andprovides for the processing of crystal diameters forwhich surface tension alone would be insufficient tocontain the molten zone in a containerless environ-ment.14 The bottom bar, which is used to seed thecrystal growth, consisted of a Rene N515 singlecrystal with a h100i orientation. During processing,the bottom bar was rotated between 27 rpm and124 rpm to average the power input from theinduction coil and promote a flat solid–liquid inter-face.16 Both bars were translated through theinduction coil at a rate of 3 mm/h (8.3 9 10�5 cm/s) to fabricate planar front grown single crystals(PG). The size of the molten zone was maintained ata diameter equal to that of the as-cast bars bymoving the top bar relative to the bottom throughstretch/squeeze control.17 While the upper solid/liquid interface was melted, the bottom interfacewas frozen and a water-cooled radiation jacketsurrounding the solidified bar promoted a highthermal gradient.

The planar growth bar was sectioned longitudi-nally using electrical discharge machining for bothmicrostructural characterization and compositionalanalysis. One half of the sectioned material wasadditionally subjected to a solution anneal at1313�C for 1 h. The microstructure characterization

for both the as-received, as-grown, and solution heattreated (SHT) samples was carried out using an FEIQuanta 650 and FEI Quanta 3D FEG scanningelectron microscopes (SEM) with a backscatterelectron detector. The specimens were mechanicallyground using SiC paper to 2000 grit, polished from6 lm to 1 lm with diamond paste, and vibratorypolished for 2 h using 0.06 lm colloidal silica.Electron probe microanalysis (EPMA) was per-formed using a JEOL JXA8200 WD/ED CombinedMicroanalyzer Superprobe to investigate the segre-gation characteristics of the zone-melted samplebefore and after solution heat treatment. EPMAmeasurements were made over 16.75 mm at a stepsize of 50 lm. EPMA measurements shown hereonly include those from the SHT specimen. The as-grown and SHT specimens exhibited comparablecomposition profiles; however, the profile was lessnoisy in the SHT specimen due to the absence ofcoarsened c¢ precipitates. Microhardness of the as-grown material was measured by Vickers microin-dentation using a Wilson� TukonTM Vickers hard-ness tester. The indent was spaced every 200 lmwith a load of 500 g. The measurements wererepeated once.

The general solution to the zone melting equationfor a binary system can be expressed as:

CS zð Þ ¼ C0 1 � 1 � k0ð Þe�k0z

z0

� �ð1Þ

where z is the position along the bar, CS is thecomposition of the solid, C0 is the initial composi-tion, k0 is the partition coefficient, and z0 is the zonelength. The differential form of this zone meltingequation was coupled to a thermodynamic database(TCNI8) and TQ-Interface and numerically solvedto predict solute concentration along the length ofthe bar.

To model the concentration profile of the planarfront, the following assumptions were considered forsimplicity: (1) the model is one-dimensional anddoes not consider composition variation in the radialdirection, (2) the travel rate and liquid zone dimen-sion is constant, (3) diffusion is infinitely fast in theliquid but negligible in solid. Algorithm 1 outlinesthe computational process of the model. The initialcomposition of the alloy C0, zone length z0, length ofplanar growth region L, system pressure P, and thestep size dz were input as initial conditions. Aninitial phase equilibrium calculation was performedat the freezing liquid/solid interface while the liquidzone moved towards the bar end. As the liquid zonemoved towards the end, a phase equilibrium calcu-lation was performed for each step. The compositionof the solid at distance z; CSðzÞ, was determined bythe calculation and the corresponding liquid com-position CL zþ dzð Þ for the next step was calculatedby the following equation based on mass balance:

CL zð Þ � z0 � CS zð Þ � dzþ C0 � dz ¼ CL zþ dzð Þ � z0 ð2ÞFig. 2. Schematic image of levitation zone melting. Adapted fromRef.15.

Planar Front Growth of Single Crystal Ni-Based Superalloy Rene N515

The phase fractions were calculated at 1340�C,1300�C, and 1100�C using the predicted solid com-position profiles obtained from the zone meltingmodel with the Thermo-Calc TCNI8 database.

RESULTS

The as-received directionally solidified (DS)microstructure of Rene N515 is shown in Fig. 3. Atypical dendritic structure with interdendriticregions and carbides was observed. The secondarydendrite arm spacing was in the range of 70–85 lm.Carbides formed in the interdendritic regions, asseen by the high contrast phases, and the averagelength of the carbides was 10.6 lm.

The as-levitation zone melted alloy structure isdistinguished by four different regions as shown inFig. 4a: the directionally solidified (DS), meltback,planar front growth (PG), and power down (PD)regions. The as-cast dendritic structure began tofade when the zone melting process started (DS), asevidenced by the gradual loss of the dendriticmicrostructure and loss of contrast variation corre-sponding to residual microsegregation between theDS and meltback regions. The dendritic structurewas completely eliminated when a planar frontstarted growing (PG), and this single crystal,microsegregation-free region was observed to be3.6 mm long. During power down (PD), the den-dritic structure began to reappear first in the formof a cellular microstructure (Fig. 4b). In this region,carbides precipitated at cell boundaries in thecellular structure and eventually in interdendriticregions. The density of carbides increased signifi-cantly as the microstructure transitioned fromcellular to fine dendritic structure, and the densityof solidification-induced porosity increased as themicrostructure transitioned as well.

Cuboidal c¢ precipitates were observed in the PGregion as shown in Fig. 4c. The average edge lengthof c¢ precipitates was 1.0 lm and the c’ area fractionwas 62.7%. This revealed c¢ coarsening occurred dueto a low cooling rate during the zone meltingprocess. Grain boundaries, dendrites, and

microsegregation were not observed in the PGregion. Porosity was also observed in PG region(0.072%); however, the area fraction was signifi-cantly lower compared to the DS region (0.24%),meltback (0.33%), and PD regions (2.96%). Since theexisting porosity was not expected at the planarfront, it is assumed it was created by the sectioningprocess by the electrical discharge machining.

The composition profile of the SHT specimenmeasured from EPMA is in good agreement withtheoretical and model predictions obtained from thezone melting mass balance model integrated withThermo-Calc, as shown in Fig. 5. A significantchange of measured composition was observed inthe transient meltback and PD regions, as exempli-fied by the decrease of Al content from 14.0 at.% to12.0 at.% in the meltback region and increase of Alcontent from 12.1 at.% to 13.8 at.% in the PDregion. Note that the PD region corresponds to thelast liquid zone, and because the concentrationprofile was not measured for the complete PDregion, the mass balance did not fully recoverbetween DS and PD regions, as shown in Fig. 5.

In the PG region, the concentration of Al, Cr, andTa increased steadily in the direction parallel toplanar growth whereas Ni, Co, W, and Re concen-trations decreased. These results indicate that Al,Cr, and Ta partition to the liquid and exhibitpartition coefficients, ki ¼ Ci

S=CiL where Ci

S and CiL

are compositions of the solid and liquid of the ithelement, respectively, of less than unity, whereasNi, Co, W and Re partitioned to solid and exhibitpartition coefficients greater than unity. Theseresults are mostly in agreement with the partition-ing trends in CMSX-4 and PWA1484.18 Partitioncoefficients were directly calculated via Thermo-Calc using the TCNI8 database for predicting thecomposition during the levitation zone meltingprocess, as they can change with temperature andcomposition. The predicted partition coefficients arein good agreement with experimental results.

These experimentally measured compositionswithin the PG region are consistent with thethermodynamic modeling, which also predicted amonotonic change of composition of each alloyingelement as the planar front grew, as shown inFig. 5. The measured composition and predictedcomposition of the PG region agree well; however, anoticeable difference in Cr and Ni content betweenthe predicted and experimentally measured valueswas observed. The measured Cr and Ni composi-tions differed by 0.96% (Ni) and 0.60% (Cr) com-pared to the nominal alloy composition. The verygradual and monotonic change in the PG regionsolid composition shows that steady-state conditionswere not achieved, such that the liquid wouldbecome saturated in solute and the alloy wouldsolidify with a composition identical to the nominalalloy composition. For the given conditions, ourmodel revealed that steady-state planar growth,Fig. 3. SEM image of as-cast directionally solidified (DS)

microstructure.

Matsunaga, Huang, Inman, Mason, Konitzer, Johnson, and Titus

whereby the solid composition is within 5% of thenominal composition, would be achieved at a dis-tance of 12.7 mm.

Hardness profiles across the as-levitation zonemelted material were measured and the microstruc-ture corresponding to the solidification region areshown in Fig. 6. The corresponding yield stress wasestimated by:19

ry � HV=3 ð3Þ

where HV is the Vickers microhardness. The hard-ness varied minimally across the DS, meltback, andPG regions, but significantly decreased in PDregion. Although the variance in the first threeregions was relatively small in the DS, meltback,and PG regions, the overall measured hardnessvalues were comparable. In these three regions, adecent area fraction of c¢ precipitates in excess of60% was observed. c¢ coarsening occurred through-out DS, meltback, and PG regions, indicating diffu-sion-controlled c¢ precipitation, and growth occurredwith the transition of solid/liquid interface from DS,meltback, to PG region due to high temperature.20

The average edge length of c¢ precipitates was1.02 lm in the DS region, 1.99 lm in the meltback

region, and 1.21 lm in the PG region. Fine, irreg-ular-shaped c¢ precipitates were observed in the DSregion. The c¢ precipitate size increased, and themorphology changed from cuboidal to spherical inthe meltback region. Finally, the precipitates exhib-ited a cube morphology in the PG region.21 Large,eutectic-like c¢ phase was observed in the DS andmeltback regions whereas they were not observed inPG and PD regions, as shown in Fig. 6. Althoughcarbides were observed near indentations in the DSand meltback regions, there was no correlationbetween the carbide formation and hardness in bothtrials. The shape of the indentation significantlychanged in the PD region, and slip bands wereobserved emanating outwards from the edge of theindentation.

DISCUSSION

The presence of a dendritic structure, microseg-regation, and carbides in the directionally solidifiedmaterial indicate inhomogeneity in composition andproperties; the absence of these features in the PGmaterial confirms that the PG region is homoge-neous. This is also supported by a very low densityof solidification porosity observed in this region. The

Fig. 4. (a) Overview of as-levitation zone melted of Rene N515, and (b) planar front—cellular structure transition during power-down, and (c) as-zone melted microstructure of PG region of Rene N515.

Planar Front Growth of Single Crystal Ni-Based Superalloy Rene N515

Fig. 5. Composition profile of the solution-annealed levitation zone melted Rene N515 as a function of position along the bar. The solid linerepresents EPMA measurement results, the dotted line represents thermodynamic concentration profile calculation for planar front growth region,and square symbols represent the nominal composition of Rene N515. One point in the DS regime was omitted because it measured the carbidecomposition.

Matsunaga, Huang, Inman, Mason, Konitzer, Johnson, and Titus

successful construction of the homogeneous struc-ture at the planar front is confirmed by the fact thatthere is no large, eutectic-like c¢ phase in the PGregion, as shown in Fig. 6, which is typicallyobserved in dendritic structure as well.11

The composition gradients of all elements arenearly constant throughout the PG region, as shownin Fig. 5. However, a significant change in compo-sition was observed in the DS, meltback, and PDregions due to transients in the melting and solid-ification behavior near these regions. Upon initialmelting, the measured solid composition altereduntil a kink in the composition profile was observedat 0 mm. This kink should correspond with the start

of solidification at z = 0 such that the composition ofthe solid is Ci

S ¼ kiCiL. The measured solid compo-

sitions at z = 0 agree well with calculated values ofthe solid composition using measured alloy compo-sitions from the DS region with measured partitioncoefficients in similar alloys from Ref.10 For exam-ple, for measured Al and W compositions of 14.1at.% and 2.2 at.%, the predicted initial solid compo-sitions using partition coefficients of kAl ¼ 0:87 andkW ¼ 1:42 are 12.3 at.% and 3.1 at.% for Al and W,respectively. The predicted composition was calcu-lated based on the nominal composition, whereasthe measured composition at the PG region repre-sents both directionally solidified and levitation

Fig. 6. Vickers hardness and estimated yield strength profile in the as-levitation zone melted specimen as a function of position along the bar.The yield stress was estimated by ry � HV=3, and the microstructure near the indentation at each region.

Planar Front Growth of Single Crystal Ni-Based Superalloy Rene N515

zone melted compositions. This can cause a notice-able difference between the predicted and themeasured composition. For instance, for Cr, thepredicted value was lower than the nominal value,and the measured value at the PG region was lowerthan the DS region. This trend is consistent withthe trend of Al and Ta which has a partitioncoefficient of k< 1 in this alloy system. It isconsidered that Cr was almost uniformly dis-tributed in the DS region with a higher contentthan nominal because of the directional solidifica-tion method, and was not significantly affected bylevitation zone melting, which can transport soluteto the liquid. A similar trend can be seen in Mo,which started with less content in the DS region dueto the partition coefficient k> 1. Conversely, thesubsequent excess or depletion of solute in the liquidshould be observed near the end of the PD region.However, because solidification in the PD regionoccurred over a large length and the excess ordepleted solute was distributed over this length,excess or depleted solute content was not observedin the EPMA measurements.

Significant changes to c¢ morphology and c¢ coars-ening occurred throughout the DS, meltback, andPG regions due to the low cooling rate during thezone melting process. The c¢ size generally increasedfrom the DS to meltback region and again decreasedin the PG region, as seen in Fig. 6. Furthermore, thearea fraction of the large eutectic-like c¢ phase washighest in the DS and meltback regions, likely dueto larger cooling rates.10,22 The c¢ morphologytransitioned from irregular cuboidal in the DSregion, to irregular in the meltback, to irregularsplit in the PG region. The average edge length of c¢precipitates was 1.02 lm in the DS region, 1.99 lmin the meltback region, and 1.21 lm in the PGregion. The largest c¢ precipitates were observed inthe meltback region because pre-existing

precipitates experienced the longest time for coars-ening throughout the levitation zone melting pro-cess (pre-heating + planar front fabrication). Fromthe DS to meltback region, the eutectic c¢ areafraction decreased from 18.5% to 10.2%. The eutec-tic c¢ fraction decreased, whereas the as-cast c¢precipitates fraction increased as the dendriticstructure disappeared and started to transition tothe planar front structure. Since the dendritic armspacing (DAS) gradually increased from DS tomeltback, microsegregation decreased from the DStowards the meltback region with the decrease ofeutectic c¢ area fraction.20

The calculated phase fraction profiles at 1100�C,1300�C, 1340�C is shown in Fig. 7a. The formationof c¢ phase was only predicted at 1100�C, whereasthe formation of the c phase was predicted at alltemperatures. This indicated that the c¢ solvustemperature of this alloy was between 1100�C and1300�C. The c¢ solvus temperature was calculatedbased on the predicted composition along the bar. Itwas 1214�C at the beginning of the PG region, thenrapidly increased to 1278�C as the planar front grewto 30 mm in the prediction, as shown in Fig. 7b. Thepredicted c¢ phase fraction in the PG region was 30–33%, which is far less than the actual c¢ areafraction. This indicates that significant c¢ coarseningoccurred below the c¢ solvus temperature so that thec¢ phase fraction increased significantly.20 The c¢phase fraction slightly increased in the PG regionand eventually became constant as zone meltingcontinued, despite small changes to the alloy com-position in the PG region. This is consistent with theexperimentally obtained microstructure, where car-bides did not exist in PG region but began toprecipitate in the PD region between cells anddendrites, where a high concentration of solutecould be found.

Fig. 7. (a) Predicted phase fraction at 1100�C and 1300�C of c and c¢ phase, and (b) c¢ solves temperature as a function of position along the bar(as levitation zone-melted). The phase fraction of planar front fabrication of this study was predicted from 0 mm to 3.6 mm, and the predictedphase fraction was kept calculating until the composition become the steady-state condition. c¢ phase fraction is predicted to be 45% after thecomposition becomes steady-state.

Matsunaga, Huang, Inman, Mason, Konitzer, Johnson, and Titus

To achieve steady-state planar growth wherebythe composition of the solidifying solid matches thatof the nominal alloy composition, fabricating alonger planar front region is proposed. A longerplanar front region can be achieved by increasingthe solidification rate; however, the thermal gradi-ent also needs to be significantly increased as shownin Fig. 1b. Experimentally, the solidification ratecan be increased by decreasing the diameter of theinitial directionally solidified bar in order toincrease the thermal gradient significantly. In thereal manufacturing process, maintaining a highthermal gradient uniformly throughout the processwould be very challenging because of the relativelylarge size of components to fabricate.

The variance of the measured hardness is mini-mal across the DS, meltback, and PG regions. Therelatively constant hardness values are likely due tothe uniform distribution of c¢ precipitates acrossthese regions despite their morphological variances.The significant decrease in hardness observed in thePD region likely occurred due to the lack of coarse c¢precipitates in this region. The precipitates in thelast parts of the PD region cannot be expected toexhibit comparable strength compared to the DS,meltback, and PG regions due to a lack of time togrow the c¢ precipitates because the PD regioncooled much more quickly than other regions.

Although there is minimum change in the hard-ness, the PG region would be expected to possesssuperior fatigue resistance and creep propertiescompared with a conventional directionally solidi-fied material. Tien et al.23 argued that a uniformdistribution of c¢ precipitates at the planar frontwould contribute to improving creep resistancesignificantly by improved phase stability. Also, paststudies confirmed that crack initiation sites areusually due to casting porosity and carbides whenfatigue tests were performed.6,7,10,11 Since the den-sity of porosity is relatively low and carbides areentirely eliminated in PG region, it is expected thatthis region is most likely to show superior creep andfatigue resistance.

CONCLUSION

Analysis of microstructure, local composition, andmechanical properties of planar front Ni-basedsuperalloys have suggested the possibility of elim-inating microsegregation and casting defects whichare two significant problems of the conventionaldirectional solidification casting process. The fol-lowing conclusions are drawn from this work:

1. A 3.6-mm-long, dendrite-free, carbide-free, andmicrosegregation-free planar front Ni-basedsuperalloy was successfully fabricated by thelevitation zone melting process.

2. The density of casting porosity was significantlydecreased from 0.24% to 0.072% (area fraction)

in the planar front region.3. A monotonic change of the composition in the

planar front region was consistent with thermo-dynamic modeling predictions.

4. In the planar front region, Al, Cr, and Tapreferably partitioned to liquid, whereas Ni,Co, W, and Re preferably partitioned to solid.

5. Carbides only formed during the latter stages ofpower down, where rapid solidification, lowthermal gradients, and microsegregation be-tween cells and dendrites were observed.

ACKNOWLEDGEMENTS

S.M and M.S.T. acknowledge support from theNational Science Foundation (NSF CAREER DMR-1848128).

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Planar Front Growth of Single Crystal Ni-Based Superalloy Rene N515