T.P. Gabb, J. Gayda, D.F. Johnson, R.A. MacKay, R.B. Rogers, and C.K. Sudbrack

Glenn Research Center, Cleveland, Ohio

A. Garg and I.E. Locci

University of Toledo, Toledo, Ohio

S.L. Semiatin

Air Force Research Laboratory, Wright-Patterson Air Force Base, Ohio

E. Kang

Northwestern University, Evanston, Illinois

Comparison of γ-γ′ Phase Coarsening Responses

of Three Powder Metal Disk Superalloys

NASA/TM—2016-218936

February 2016
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T.P. Gabb, J. Gayda, D.F. Johnson, R.A. MacKay, R.B. Rogers, and C.K. Sudbrack

Glenn Research Center, Cleveland, Ohio

A. Garg and I.E. Locci

University of Toledo, Toledo, Ohio

S.L. Semiatin

Air Force Research Laboratory, Wright-Patterson Air Force Base, Ohio

E. Kang

Northwestern University, Evanston, Illinois

Comparison of γ-γ′ Phase Coarsening Responses

of Three Powder Metal Disk Superalloys

NASA/TM—2016-218936

February 2016

National Aeronautics and

Space Administration

Glenn Research Center

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NASA/TM—2016-218936 1 

Comparison of γ-γ′ Phase Coarsening Responses 

of Three Powder Metal Disk Superalloys

T.P. Gabb, J. Gayda,* D.F. Johnson, R.A. MacKay, R.B. Rogers, and C.K. Sudbrack 

National Aeronautics and Space Administration

Glenn Research Center

Cleveland, Ohio 44135 

A. Garg and I.E. Locci

University of Toledo

Toledo, Ohio 43606 

S.L. Semiatin

Air Force Research Laboratory

Wright-Patterson Air Force Base, Ohio 45433 

E. Kang 

Northwestern University

Evanston, Illinois 60208 

Abstract

The phase microstructures of several powder metal (PM) disk superalloys were quantitatively 

evaluated. Contents, chemistries, and lattice parameters of γ and γ′ strengthening phase were determined

for conventionally heat treated Alloy 10, LSHR, and ME3 superalloys, after electrolytic phase 

extractions. Several of long term heat treatments were then performed, to allow quantification of the 

precipitation, content, and size distribution of γ′ at a long time interval to approximate equilibrium 

conditions. Additional coarsening heat treatments were performed at multiple temperatures and shorter 

time intervals, to allow quantification of the precipitation, contents and size distributions of γ′ at 

conditions diverging from equilibrium. Modest differences in γ and γ′ lattice parameters and their 

mismatch were observed among the alloys, which varied with heat treatment. Yet, γ′ coarsening rates 

were very similar for all three alloys in the heat treatment conditions examined. Alloy 10 had higher γ′

dissolution and formation temperatures than LSHR and ME3, but a lower lattice mismatch, which was 

slightly positive for all three alloys at room temperature. The γ′ precipitates of Alloy 10 appeared to 

remain coherent at higher temperatures than for LSHR and ME3. Higher coarsening rates were observed 

for γ′ precipitates residing along grain boundaries than for those within grains in all three alloys, during 

slow-moderate quenching from supersolvus solution heat treatments, and during aging at temperatures of 

843 °C and higher. 

Introduction

Nickel-based turbine disk superalloys are fundamentally made up of a Ni face centered cubic (γ) 

phase matrix reinforced by Ni3Al ordered L12 (γ′) phase precipitates (Ref. 1). The composition of both γ′

and γ phases can influence the mechanical properties and environmental resistance of each phase, and the 

superalloy as well (Ref. 2). The contents and size distributions of γ′ strengthening precipitates have been 

 

* Currently retired.
NASA/TM—2016-218936 2 

shown to strongly influence mechanical properties of disk superalloys (Refs. 1, 2, 3, and 4). When 

powder metallurgy alloys are processed for uniform composition, these morphology parameters are 

dictated by alloy composition and thermo-mechanical processing. More recent powder metallurgy disk 

superalloys for high temperature applications have been designed with 50 to 55 percent of γ′ precipitates, 

usually less than 0.3 μm in diameter, having refractory element additions of Ti, Ta, and Nb for additional 

inherent phase strength. The γ matrix phase also has element additions for processing (Co, Mo), strength 

(W, Mo), and environmental resistance (Cr). 

Microstructure modeling software such as ThermoCalc (Ref. 5), JMatPro (Ref. 6), and Pandat 

(Ref. 7) have been developed over recent years, which can estimate equilibrium phase compositions, 

quantities, and sizes as functions of temperature for input alloy compositions. Other software packages 

such as PrecipiCalc (Ref. 8) have been developed to specifically predict γ′ content and size distribution 

for input alloy composition and temperature-time paths. In order to assess and calibrate such tools for 

specific needs, it is necessary to consistently generate experimental data on representative alloys of 

interest, quantifying phase compositions, contents, and sizes after varying thermal histories. Furthermore, 

it would be useful to consistently apply the same isothermal coarsening experiments to disk superalloys 

having varied compositions, and then to consistently quantify the microstructures and compare coarsening 

responses. The objective of this study was to compare the γ′ phase contents and sizes present in three 

powder metallurgy disk superalloys consistently heat treated near and far from equilibrium conditions, 

specifically after varied quenching and isothermal coarsening heat treatments.

Materials and Procedures

Compositions in weight percent of the three disk superalloys examined are listed in Table 1 in weight 

percent. These superalloys were all produced using powder metallurgy processing, with powder 

atomization in argon atmosphere. Alloy 10 (Refs. 9 and 10) powder was passed through –270 mesh 

screens, then sealed in a container, hot isostatically pressed, and extruded at a reduction ratio of 6:1. A 

segment of the extrusion billet was machined, then isothermally forged into a flat disk about 35.6 cm in 

diameter and 5 cm thick. LSHR (Ref. 11) powder was passed through –270 mesh screens, then sealed in a 

container, hot compacted, and extruded at a reduction ratio of 6:1. A segment of the extrusion billet was

machined to 15 cm diameter and 20 cm long, then isothermally forged into a flat disk about 30 cm in 

diameter and 5 cm thick. ME3 (Refs. 12 and 13), was processed in a comparable manner, but the larger 

diameter extrusion segment was machined to 23 cm diameter and 20 cm long, then isothermally forged 

into a flat disk about 46 cm in diameter and 5 cm thick. Samples were obtained from each disk in the asforged condition for a variety of heat treatments. 

Phase Extractions

Samples of the three alloys weighing about 10 g each were solution heat treated at 1199 °C for 1.17 h, 

then cooled in the furnace with the power off to below 870 °C at an average cooling rate of 25 °C/min, as 

shown in the temperature-time cooling curves of Figure 1. They were then aged at 855 °C for 4 h and 

775 °C for 8 h, again with furnace off cools from each temperature. Compositions and contents of γ and γ′

phases were determined after separation by electrolytic phase extractions. Triplicate γ′ phase extractions 

were performed for each alloy using a 1 g ammonium sulfate-1 g citric acid-100 mL H2O electrolyte. The 

extractions were passed through a filter having mean mesh size of 0.1 μm. Extracted phases and lattice

parameters were identified using X-ray diffraction (XRD). The compositions of samples prior to 

extractions and of the phase extractions were determined by inductively coupled plasma (ICP) chemical
NASA/TM—2016-218936 3 

analyses. The composition of the γ phase was determined by chemical analysis of the ammonium 

sulfate/citric acid electrolyte immediately after the γ′ extractions were completed. The composition of the 

extracted γ′ phase was determined after analyses by X-ray diffraction. 

X-Ray Diffraction 

Lattice parameters in γ′ phase extractions and metallographically prepared samples were determined 

using x-ray diffraction (XRD). All XRD data was gathered using Cu Kα radiation and Bragg-Brentano 

(focusing) geometry. Data for the phase extraction samples were gathered using a Philips APD 

diffractometer equipped with a post-beam monochrometer. Six peaks in the 70° to 150° 2θ range were 

used to calculate γ′ phase extraction lattice parameters with the Le Bail full profile fitting method 

(Ref. 14) employing fundamental parameters peak fitting (Ref. 15) as implemented in TOPAS v.4 

software (Bruker AXS). Results are shown in Table 2(a). This technique was first qualified through lattice 

parameter measurements of a certified NIST sample of known lattice parameter (NIST SRM660a, LaB6). 

Analysis of repeated measurements of the NIST standard yielded a 95 percent confidence interval of 

0.0001 Å for a future singleton measurement. 

Lattice mismatch measurements were performed on multiple samples heat treated at very fast and 

very slow cooling rates (Table 2(b)). A Bruker D8 Advance diffractometer equipped with a linear position 

sensitive detector and a Ni β-filter was used to gather data over the 40° to 150° 2θ range of 

metallographically polished samples. Lattice parameters were obtained with TOPAS via both the 

Fundamental Parameters and Split Pearson VII approaches using the (331) peak, which provided the best 

combination of intensity and resolution between γ and γ′ peaks. For single peak fitting, the procedure is 

interactive where the initial peak positions are set by the user. The output of the data fitting step consisted 

of peak positions for the γ and γ′ phases, which were used to calculate the lattice parameter of each phase 

and the lattice mismatch as shown in Table 2(b). This single-peak fitting approach for lattice mismatch 

measurements was chosen over multiple-peak fitting after examining the results of the two methods.

Comparison of the fitting results of these two models led to the conclusion that single peak fitting 

provided the best lattice mismatch values. With the multiple peak fitting approach, the significantly 

higher intensities of the lower angle peaks where γ and γ’ are unresolved influenced the least-squares 

fitting such that noticeably poorer fitting occurred for high angle peaks where γ and γ′ are resolved. 

Although absolute lattice parameter accuracy might be lower by fitting only a single peak, better relative 

lattice parameter values were obtained, which is more important for lattice mismatch calculations. 

Development of this approach for nickel-based superalloy single crystal samples indicated lattice 

mismatch error bars (1σ) in the range of ± 0.002 to ± 0.07 percent for repeat measurements where the 

error increased with decreasing difference between γ and γ′ lattice parameters. For some samples, the 

lattice mismatch was so small that the separate lattice parameters could not be measured with confidence.

This was indicated by the results being highly dependent on the user-defined starting lattice parameters.

For these samples, a composite, averaged lattice parameter is reported in Table 2(b). It was not possible 

from XRD data alone to determine which peak belonged to which phase in a given sample. However, 

comparison of results from Table 2(a) and (b) indicates that the a1 phase in Table 2(b) corresponds to γ′

for samples where both lattice parameters could be measured. Lattice mismatch was calculated using the 

following relationship:

( )

γγ

γγ

+

− =δ

aa

aa

'

2 '
NASA/TM—2016-218936 4 

Here, a is the lattice parameter and the subscripts identify different phases (Ref. 16). 

Differential Thermal Analyses

Differential thermal analysis (DTA) experiments were performed using a Netzsch STA 409C. Some 

of the LSHR results on the same tool were previously reported elsewhere (Ref. 17). Triple cycle 

experiments were used to help estimate the transformation temperatures that the γ′ phase precipitates and 

borides dissolve on heating and form on cooling, while single cycle experiments were used to estimate the 

solidus and liquidus temperatures on heating. Cylindrical samples were extracted from as-forged disks. 

The majority of these samples were 5 mm in diameter, ~8 mm in length, and weighed approximately 

1400 to 1600 mg (see Table 3). The only exception was the LSHR sample for the single cycle solidus / 

liquidus experiment, which was 5 mm in diameter, ~4 mm in length and weighed 683.1 mg. The DTA 

chamber was backfilled twice with inert He to remove oxygen before starting to heat for all experiments. 

Samples for triple-cycle experiments under a 40 to 60 mL/min He flow were heated quickly, either at 

20 or at 10 °C/min, to a starting temperature of 1000 °C for LSHR and ME3 and 1030 °C for Alloy 10 to 

minimize weight gains. Then, using slower rates selected among 10, 5, and 2 °C/min (Table 3), they were 

heated and cooled three times between the starting tempeature and maximum temperature of 1200 to

1220 °C for LSHR and ME3, and 1250 °C for Alloy 10. The samples were quickly cooled, at either 

20 °C/min or 10 °C/min, to room temperature. The weight gains after the triple cycle DTA experiments 

ranged between 0.1 to 0.3 mg (see Table 3), which gives confidence that any surface oxidation due to 

residual oxygen in the inert environment did not influence DTA measurements. The transformation 

temperatures of each cycle were determined from the intersections of tangential lines at the onset and at 

the tail of the transformation in the thermograms. The peaks were distinct for the boride transformation on 

cooling, while the γ′ phase transformations exhibited more gradual slopes with temperature. A mean and 

standard deviation value were determined from the three thermal cycles during the DTA experiment on 

one disk sample. Mean values as a function of temperature rate were extrapolated linearly to 0 °C/min as 

described in more detail in the Results section below. 

The maximum temperature for the single-cycle experiments was 1370 °C, where samples were run 

continuously at 5 °C/min in the temperature range of interest (Table 3) under a 40 to 60 mL/min He flow. 

The liquidus on heating was determined from the intersections of tangential lines at the onset and at the 

tail of each transformation in the thermograms. Since the change in heat capacity at the solidus 

temperature is gradual, the solidus is difficult to determine from DTA measurements. Solidus was thus 

determined by inspecting in an enlarged view the intersection of the base line and the extrapolated tangent 

approximation of the thermogram. The weight gains after the single cycle DTA experiments were also 

reasonalbly small and ranged between 0.4 to 0.7 mg (Table 3). 

Heat Treatments to Quantify γ′ Coarsening as Functions of Temperature and Time

Numerous cylindrical samples 3.2 mm diameter and 22 mm long of Alloy 10, LSHR, and ME3 were 

extracted by electro-discharge machining from the as-forged disks for heat treatments, listed in Table 4. 

Several pins were solution heat treated in argon atmosphere at 1199 °C for 1 h, and water quenched or 

furnace cooled for examination, as shown in the temperature-time cooling curves of Figure 1. The 

average cooling rate to 650 °C of water quenched specimens was 5,000 °C/min. Several other specimens 

were cooled in the furnace by simply turning the furnace off. The average cooling rate of specimens 

cooled with the furnace power off was 25 °C/min to 870 °C. Finally, several specimens were given a very 

slow controlled furnace cooling ramp at a constant cooling rate 2 °C/h (0.033 °C/min) to 870 °C followed 

by a constant cooling rate of 0.19 °C/h (0.0032 °C/min) to 650 °C, in an attempt to allow cooling 
NASA/TM—2016-218936 5 

conditions to approach equilibrium. Additional solution heat treated and water quenched pins were given 

aging heat treatments in argon at 760, 843, 927, and 1093 °C for 1,000 h and water quenched, to 

determine γ′ phase content and size at long term aging conditions approaching equilibrium. Aging heat 

treatments were also performed at these temperatures for shorter times, in conditions diverging from 

equilibrium. Sections of each sample were metallographically prepared, then etched with a solution of 

33 mL HCl - 33 mL HNO3 - 33 mL acetic acid - 33 mL H2O - 1 mL HF for examining γ′. These sections 

were examined in a Hitachi S-4700 Field Emission scanning electron microscope (SEM). For conditions 

producing very fine γ′ sizes, thin disks of 3 mm diameter and ~ 0.5 mm thick were sliced and 

mechanically ground and polished. These disks were then electrochemically thinned to electron 

transparency using a solution of 10 percent perchloric acid and 90 percent methanol cooled to –30 °C. A 

FEI CM200 Transmission Electron Microscope (TEM) operating at 200 kV was used for imaging γ′

precipitates in these cases. 

SEM and TEM evaluations were performed in a complementary manner. For γ′ phase quantification, 

TEM was used to image γ′ precipitates when they were less than 0.5 μm in diameter. TEM dark field 

imaging conditions were used on grains oriented with the [001] crystallographic direction near normal to 

the electron beam, to illuminate the γ′ phase at very thin locations of the foil. SEM was used to image γ′

precipitates when they were larger. Precipitates were measured within the grains having more cuboidal 

precipitate morphologies, signifying grains oriented with the [001] crystallographic direction near normal 

to the electron beam. Precipitates along those grain boundaries were also measured and tabulated 

separately. While their limited numbers prevented construction of size-frequency histograms, mean sizes 

could still be determined for initial comparisons.

Particle sizes were measured using SigmaScan Pro image analysis software. The difference in 

brightness between γ′ and γ phases enabled brightness-based local thresholding. γ′ precipitates residing at 

grain boundaries were thresholded separately from those within grains, to enable separate evaluations of 

precipitate sizes. Major axis, minor axis, and area were determined for each particle and tabulated. Phase 

area fraction was determined by combining the areas of all particles within an image.

Results

Phase Compositions 

The microstructures of samples used in phase extractions are shown in Figure 2. Both cuboidal 

secondary γ′ phase precipitates of 0.1 to 0.25 μm width and spherical tertiary γ′ phase precipitates near 

0.01 to 0.03 μm diameter were observed in all three alloys. Analyses by x-ray diffraction indicated the γ′

phase extractions were very clean, and did not contain other phases in significant quantities. Due to the 

pore size of the filter paper, this is expected to be mostly made up of secondary γ′ precipitates. Therefore, 

compositions of the γ′ phase extractions were measured as a representation of average chemistry for 

secondary γ′, Table 5. The solvent of the remaining phase extraction solution was evaporated away, and 

the solid remains were analyzed to determine average γ phase compositions, also compared in Table 5. 

Resulting partitioning ratios are also included. The measured γ and γ′ compositions closely agree with 

prior measurements on the three alloys using an atom probe (Ref. 8), and the associated partitioning 

ratios. Al, Ti, Ta, and Nb consistently partitioned preferentially to the γ′, while Cr, Co, Mo partitioned 

more to the γ phase, consistent with results reported on other superalloys (Refs. 1 and 18). As also 

reported before, W was near evenly partitioned between the γ and γ′ phases.
NASA/TM—2016-218936 6 

γ- γ′ Lattice Parameters and Mismatch

The clean γ′ phase extractions for each alloy allowed refined measurement of their lattice parameters 

using multiple peaks. These should be considered unconstrained lattice parameters, as the γ′ precipitates 

were extracted, and not constrained by the surrounding γ matrix phase. Unconstrained lattice parameters 

of the γ phase are compared for the alloys in Table 2. LSHR had the largest γ′ phase lattice parameter, 

followed in decreasing order by ME3 and Alloy 10. 

It would be expected that such constraint is a function of heat treatment (Ref. 19).Therefore, samples of 

all three alloys were solution heat treated at 1199 °C and then given highly divergent cooling paths as shown 

in the temperature-time cooling curves of Figure 1, and measured for lattice mismatch (Table 2(a)).The 

lattice mismatch values in the samples quenched at 5,000 °C/min (AW, LW, MW in Table 2(b)) were too 

small for our program to distinguish between peaks for the γ and γ′ phases. These results indicate the lattice 

mismatch magnitudes were less than 0.02 percent, the estimated limit of peak resolution at room 

temperature. Companion samples were given a very slow cooling rate of 2 °C/h (0.033 °C/min) to 870 °C 

followed by 0.19 °C/h (0.003 °C/min) from 870 to 650 °C. Microstructural examinations of the latter very 

slowly cooled specimens indicated the γ′ phase precipitates no longer consistently had a regular cuboidal 

morphology, and showed coalescence of the γ′ phase precipitates along grain boundaries, as seen in Figure 

3. This indicated a loss of coherency between the phases, with reduced constraint between phases. The 

lattice parameters of phase 1 in the mismatch measurements of these specimens approached that of the 

extracted γ′ phase for each alloy. This suggested that phase 1 having the larger lattice parameter was the γ′

phase in each alloy, and that all three alloys had small positive values of lattice mismatch at room 

temperature that ranged from +0.19 to +0.27 percent. Alloy 10 had the lowest mismatch, with LSHR and 

ME3 having similar, slightly higher values of mismatch, Table 2(b). 

γ′ Dissolution and Formation Temperatures

Heating and cooling curves from the cycling DTA experiments are shown in Figure 4 for as-forged 

LSHR, ME3 and Alloy 10 samples. The measured γ′ and boride transformation temperatures from these 

curves are listed in Table 3, where both the temperatures for the γ′ dissolution on heating and γ′ formation 

on cooling temperatures are reported. Note, the dissolution temperatures are most commonly reported as 

the γ′-solvus temperature and boride liquation temperature.

On the first pass of cycling DTA (Figure 4), broad peaks at approximately 1070, 1080, and 1090 °C 

are observed for the as-forged LSHR, ME3 and Alloy 10, respectively. These peaks are associated with a 

smaller population of secondary γ′ precipitates from the native bimodal distribution in the as-forged 

alloys. Then, with cooling from the first pass and continued cycling with the second and third passes, no 

such peaks are observed indicating that the γ′ precipitate distribution is or is nearly unimodal. The 

maximum temperature that was reached for the LSHR sample cycled at 10 °C/min was only 1200 °C, 

which meant the boride dissolution temperature was out of range. It followed that the borides were still 

intact on cooling. The boron content is higher in LSHR than in the Alloy 10 and ME3 alloys that were 

examined here (Table 1). Accordingly, LSHR was the only alloy where the boride formation temperatures 

were observable on cooling, as the volume fraction of borides that form on cooling is smaller. 

In all three alloys, the fastest heating rate of 10 °C/min for the cycling DTA consistently gave the 

highest dissolution temperatures for both the γ′ precipitates and borides, where these temperatures 

decreased as the heating rate decreased. Accordingly, the faster the cooling rates, the lower the formation 

temperatures during cooling were. LSHR and ME3 had very similar responses for the γ′ phase, with γ′

dissolution temperatures ranging from 1151 to 1162 °C, and γ′ formation temperatures ranging from 1129 
NASA/TM—2016-218936 7 

to 1123 °C. For the γ′ phase, Alloy 10 had consistently higher temperatures, with a γ′ dissolution 

temperature ranging from 1190 to 1193 °C, and a γ′ formation temperatures ranging from 1168 to 

1166 °C. The γ′ dissolution temperature for Alloy 10 is close to the boride dissolution temperature, where 

the former is only 10 to 13 °C lower than the latter, while for LSHR and ME3 they are 48 to 54 °C lower, 

Table 3. For the Alloy 10, the proximity of the boride dissolution peak led to overlap with the γ′

dissolution peak. The two measured γ′ dissolution values for Alloy 10 instead of decreasing with the 

slower rate like the other alloys are about same, ~ 1192 °C. This is shown clearly in Figure 4, which plots 

the mean transformation temperatures for all the γ′ dissolution and γ′ formation temperature measured by 

DTA as a function of temperature rate. 

The as-forged LSHR was the only condition where three different temperture rates (x) were anlayzed 

with the cycling DTA, which allowed the mean transformation temperatures (y) to be extrapolated 

linearly, y = mx + b, to 0 °C/min (b) in Figure 4(b). The slope (m) of LSHR γ′-dissolution temperature 

and LSHR γ′-formation temperature were determined to be 0.94 ± 0.09 min and 0.57 ± 0.09 min, 

respectively. Since the as-forged ME3 and Alloy 10 were only analyzed at two temperature rates, the

slopes determined for LSHR were employed for linear extrapolation of data from ME3. For Alloy 10, a γ′

dissolution temperature where a slope equal to zero was used. Extrapolated values are listed in Table 3a, 

and shown in Figure 4(b) with a fitting error.

The solidus and liquidus temperatures were determined for all three alloys using DTA, Figure 4 and 

Table 3(b). The solidus values were 1240, 1264 and 1257 °C for LSHR, ME3 and Alloy 10, respectively. 

The liquidus temperatures were extremely close for ME3 and Alloy 10, 1339.0 and 1339.6 °C, and 

slightly lower for LSHR at 1334.5 °C. The chemistry of ME3 more closely resembles LSHR than Alloy 

10, so closer agreement in liquidus measurements might be expected. The observation that LSHR’s

liquidus is 5 °C lower may be due to the fact that the sample size is nearly three times smaller than the 

other two samples. Such an effect was observed in a prior report with LSHR DTA, where some of the 

results presented here were previously reported (Ref. 17). 

γ′ Phase Morphology Versus Cooling Rate 

Typical images of the γ′ microstructures are compared for water quenched, furnace cooled, and very 

slowly cooled specimens in Figure 5, Figure 6, and Figure 2. Water quenched specimens cooled at 

5,000 °C/min had uniform, very small spherical precipitates. Specimens cooled in the furnace at 

25 °C/min had larger precipitates that were more cuboidal, yet the cubes were sometimes segmented into 

octets. Precipitates coarsened faster at grain boundaries than within grains for the slower cooling rate of 

the latter furnace cooled specimens. Very slowly cooled specimens had a complex morphology of γ′ sizes 

within grains, with large precipitates of several microns diameter and smaller precipitates at less than 

1 μm observed in the grains of each alloy. Excessive precipitate coarsening along as well as across the 

grain boundaries was also observed here. Water quenched specimens had the smallest, most uniform γ′

size for each alloy, and had equivalent radius versus frequency that could be approximated as normal 

distributions. Therefore, this solution quenching heat treatment was selected for subsequent coarsening 

experiments. 

γ′ Phase Morphology Versus Coarsening Temperature and Time

Typical images of γ′ precipitates after water quenching and then aging for 1,000 h at temperatures of 

760, 843, 927, and 1093 °C are shown in Figure 7 to Figure 10. Precipitates remained spherical during 

coarsening at 760, 843, and 927 °C. However, lateral joining or coalescence of precipitates became 
NASA/TM—2016-218936 8 

evident at 1093 °C, especially for LSHR and ME3. Brief evaluations using transmission electron 

microscopy of the later specimen aged 1,000 h at 1093 °C indicated misfit dislocations were arranged 

along the precipitate-matrix interfaces, Figure 11. 

Equivalent radius, defined as Req=(A/π)0.5, versus frequency could be approximated as normal 

distributions (Figure 12), and mean sizes are compared after aging at 1,000 h in Figure 13, for precipitates 

measured within grains and for precipitates measured along grain boundaries. Measurements of the 

equilibrium volume fraction of γ′ as a function of temperature (f (T)) for LSHR and Alloy 10 were 

compared to the analytical expression suggested by Payton (Ref. 20), i.e., 





















×

− −





















×

− −





















 





















γ

γ

γ

γ

−=

'

'

'

'

*41 exp

*4 1 exp

TT

TT

R

Q C

TT

TT

R

Q C

f . (1) 

Here, C* denotes the atomic fraction of γ′ formers in the alloy (i.e., Al+Ti+Nb+Ta). This value was taken 

to be the total of the four elements, disregarding the fact that small levels of each element are also found 

in the γ matrix. However, even with this assumption, it has been found that the quantity 4C* typically 

provides relatively-accurate predictions of the total volume fraction of γ′ in many PM superalloys. 

Regarding the other terms in Equation (1), Q is a fitting parameter, T is absolute temperature (in K), R is 

the gas constant, and Tγ′ is the γ′ solvus temperature. 

For LSHR with the composition in Table 1, C* = 0.528, and Tγ′ = 1157 °C = 1430 K (Table 3). 

Previous work (Ref. 21) has shown that Q = 60 kJ/mol provides an excellent fit for the equilibrium 

solvus-approach curve for this material. Figure 14(a) compares the analytical fit of Equation (1) using 

these values of C*, Tγ′ , and Q with the measurements from Figure 13(b). The fit is very good except for 

the lowest temperature (760 °C).

For Alloy 10 (Table 1), C* = 0.572, and Tγ′ = 1157 °C = 1466 K (Table 3). In Figure 14(b), analytical 

fits based on Q = 60 kJ/mol, 67.5 kJ/mol, and 75 kJ/mol (the highest of these values having been found to 

be appropriate for IN-100 which has a relatively high level of γ′) are compared to the measurements from 

Figure 13(b). The fit showed a relatively-limited sensitivity to the specific value of Q, and mirrored the 

measured trend except for the datum at 927 °C.

The measured γ′ solvus temperatures for LSHR and IN-100 as a function of the quantity 4C* were 

also compared to the corresponding values for Waspaloy (with either a lean or rich content of Al+Ti 

(Ref. 22), Rene 88DT, and IN-100 (Figure 15). Figure 15 shows the expected general trend of increasing 

solvus temperature with increasing hardener content (= 4C*). However, the solvus of Alloy 10 appeared 

somewhat high based on its value of 4C*. In this regard, the solvus for a more recent, modified version of 

this alloy composition has been reported in the literature (Ref. 23) to be 1166 °C, not 1193 °C as found in 

the present work and material. 

As aging temperature increased, it is apparent that γ′ precipitates coarsened during aging more rapidly 

at grain boundaries than within grains. This could be attributed to enhanced diffusion near grain 

boundaries of the elements that stabilize γ′: Al, Ti, Ta, Nb. Area fractions are also compared after long 

term aging for 1,000 h in Figure 13. The higher area fraction of γ′ precipitates measured for Alloy 10 after 
NASA/TM—2016-218936 9 

1,000 h at 1093 °C is consistent with this alloy’s higher γ′ solvus temperature measured in DTA, in 

comparison to LSHR and ME3.

Mean γ′ precipitate sizes separately measured within grains and along grain boundaries are compared

for all coarsening test temperatures and times in Figure 16. Coarsening rates were approximated using the 

LSW theory (Ref. 24) on the mean equivalent radius measurements, where rt

3

 – ro

3

= k t, as shown in 

Figure 17. The resulting coarsening constant k is shown for each alloy and temperature, in the form log k 

versus 1/T, in Figure 18. The slope of this relationship represented the activation energy for coarsening of 

coherent precipitates. This relationship diverged from linearity for each alloy at the highest temperature of 

1093 °C, where precipitates were observed to coalescence and lose coherency (Figure 10 and Figure 11). 

Such divergence has been observed in other superalloys when precipitates coalescence and lose their

coherency (Ref. 25). This divergence in linearity is reflected in the lower coefficients of determination 

calculated when including the results at 1093 °C. Activation energies were determined both with and 

without inclusion of the results at 1093 °C, Table 6. 

The resulting activation energies are compared in Table 6, along with the slopes and associated 

coefficients of determination (R2

). The activation energies calculated with the results from 760 to 1093 °C 

were slightly higher than published values of 127 to 280 J/mole-K reported in References 26 and 27. The 

activation energies calculated from only the results from 760 to 927 °C were consistently higher than 

these values. The values calculated for grain boundaries were higher than those calculated for within 

grains in both cases, for all three alloys.

The challenge of applying a first-order activation (Arrhenius) analysis such as that utilized in Table 6 

and Figure 18 may lie with the fact that the classical LSW analysis is strictly applicable to a distribution 

of precipitates whose composition is a terminal solid solution and whose volume fraction is infinitesimal 

(Refs. 28 and 29). For PM γ′-strengthened superalloys, such assumptions are not met. In these instances, 

an alternate analysis similar to the LSW approach shows that coarsening kinetics still follow an r3

-versust dependence (assuming bulk-diffusion control), but the rate constant (kMLSW) takes the following modified 

form (Refs. 30 to 38): 

( ) ( )

( ) [ γ ] γγ

γγ

∂∂+

−φ = −CCRT v/ CVCDw

k M

MLSW

9 1 lnln

8 σ C1

2 '

(2) 

In this expression, w(φ) is a factor used to correct for the finite volume fraction of precipitates φ (Refs. 28 

to 31), D is the effective diffusivity for the rate-limiting solute, σ is the γ - γ′ interface energy, Cγ and Cγ′

are the equilibrium concentrations of the rate-limiting solute in the matrix and precipitate, respectively, 

while R and T have the same definitions as above. The bracketed term in the denominator of Equation (2) 

is the thermodynamic factor in which v denotes the activity coefficient for the rate-limiting solute in the γ

matrix. In addition to the diffusivity D (and T in the denominator), various other terms in Equation (2) 

depend in a complex fashion on temperature. These include w(φ) (through the dependence of γ′ volume 

fraction on T), Cγ, and the thermodynamic factor. Each of these terms can confound an Arrhenius-type 

activation analysis.

Kuehmann and Voorhees (Ref. 37) demonstrated that rapidly-diffusing solutes can also affect the 

coarsening rate and the related diffusional processes for ternary alloys. For multi-component alloys, their 

work suggests that an effective rate constant (keff) equal to the inverse of the sum of the inverse rate 

constants for the individual solutes can be defined as the following:

... 321

1111

+++=

keff k kk (3) 
NASA/TM—2016-218936 10

Input Data

Theoretical estimates of the rate of coarsening of intragranular precipitates were obtained for LSHR 

and Alloy 10. Because of its similarity to LSHR in terms of composition, alloying-element partition 

ratios, solvus temperature, and observed coarsening rates, such calculations were not done for ME3. The 

values of the various material coefficients needed to apply Equation (2) were obtained from the present 

measurements (i.e., the overall alloy and precipitate compositions) and the data in References 21 and 33. 

Specifically, w(φ) was taken to be that derived by Voorhees and Glicksman (Ref. 33) as depicted in 

Figure 19, and σ was assumed to be 23 mJ/m2 (Ref. 21). The molar volume VM was estimated as 

7.22×10–6 and 7.04×10–6 m3

/mol for LSHR and Alloy 10, respectively. In addition, chromium was taken 

to be the rate-limiting solute for coarsening, and its “effective” diffusivity was assumed to be one-third 

that of Cr in the Ni-Cr binary, per Reference 21. 

Deff (m2

/s) = A × DB = 0.33 × 0.00036 exp (–34278/T(K)). (4) 

The activation energy for diffusion of Cr in Ni (285 kJ/mol) is similar to the values in Table 6(a) 

for intragranular coarsening. The diffusivities of the various other solutes in each alloy were taken to be 

in the ratios (relative to that of Cr) suggested by the work of Campbell, et al. (Ref. 39). The 

thermodynamic factors for Equation (2) were estimated using the commercial code Pandat (CompuTherm 

LLC, Madison, WI).

Coarsening-Rate Predictions

Calculated per Equation (2), the coarsening rate constant at 1093 °C for each substitutional solute 

in LSHR is summarized in Table 7. The results indicated that the rate constant for chromium 

(1.81×10–7 μm3

/s) was indeed smaller than those for each of the other elements by a factor of four or 

more. Hence, as has been found for both the nucleation and growth of γ′ precipitates in LSHR (Ref. 21), 

the diffusion of chromium was concluded to be the principal alloying element which limits the rate of 

coarsening of the intragranular precipitates in LSHR (and the other program materials). When the effect 

of all of the solutes on the coarsening rate was taken into account (Eq. (3)), an effective rate constant of 

1.14×10–7 μm3

/s was estimated. This value is approximately 37 percent lower than that associated with 

Cr alone. 

Based on Equation (2) for chromium, without or with the contribution of the other solutes taken into 

account via Equation (3), predicted rate constants for both LSHR and Alloy 10 for the three highest heattreatment temperatures (843, 927, and 1093 °C) were estimated and are summarized in Table 8. These 

values are compared in the table to the measured rate constants based on the measured average radii from 

the TEM and SEM observations in Table 4. In these instances, it was assumed that the measured radii 

were the true 3D values. This assumption is valid for observations taken from TEM foils. For the SEM 

observations on 2D sections, on the other hand, the actual 3D radius is more difficult to determine. For a 

distribution of mono-size spheres, the true 3D radii are larger (by a factor of approximately 22 percent) 

than the average deduced from a 2D section. However, recent geometric analysis by Payton (Ref. 40) for 

distributions that have finite breadth, as is pertinent for a distribution of precipitates undergoing 

coarsening, suggests that such a correction provides a large overestimation of the actual average 3D size. 

Hence, as a first approximation, the radii from SEM observations were not corrected for stereological 

effects in obtaining the measured coarsening rates listed in Table 8. Furthermore, the measurements in 

Table 8 were based on data for the 0 to 1000 h or 0 to 100 h time intervals. In all cases, except the 

1093 °C heat treatment of Alloy 10, the measurements for the two different intervals were very similar to 
NASA/TM—2016-218936 11

each other. The divergence for Alloy 10 at the highest test temperature merits further investigation; the 

datum for the longest time (1000 h) appeared to be the source of the discrepancy. 

The results in Table 8 exhibited good to excellent agreement between the measured and predicted rate 

constants, especially for comparisons using predictions incorporating all of the solutes. For LSHR, the 

agreement was excellent at 843 and 1093 °C. For 927 °C, the prediction also showed excellent agreement 

with the measurement based on the 0 to 1000 h interval, but was high by a factor of ~2, when compared 

to the measurement for the 0 to 100 h interval; such behavior can still be considered reasonable based on 

typical differences cited in the coarsening literature. For Alloy 10, the agreement was very good at 843 

and 927 °C. At 1093 °C, the predicted value showed excellent agreement with the measurement based on 

the 0 h to 100 h interval, but a noticeable difference from the measurement based on the 0 to 1000 h 

interval.

The predicted rate constants were also compared to measurements for which the radii determined 

specifically from SEM 2D sections were corrected to 3D using a factor of 1.225, which would provide an

upper bound for the stereological adjustment. Such comparisons Table 9 provided agreement between 

measurements and predictions which is reasonable but not as good as that in Table 8. 

At 760 °C, the predicted coarsening rate constant for LSHR (from Eq. (2)) overestimated the 

measurement by approximately a factor of 5. This difference was likely due to a change in the coarsening 

mechanism from bulk-diffusion-control (at the higher temperatures) to more sluggish trans-interfacediffusion-control (Ref. 41) which typifies behavior at temperatures comparable those seen during jetengine service. 

Summary

The phase microstructures of powder metal (PM) disk superalloys Alloy 10, LSHR, and ME3 were 

quantitatively evaluated. The γ and γ′ phase transformation temperatures, compositions, lattice 

parameters, and morphologies were first assessed. Alloy 10 had higher γ′ dissolution and formation 

temperatures than LSHR and ME3, but a lower lattice mismatch at room temperature. Yet, all three alloys 

had slightly positive lattice mismatch at room temperature. Solution heat treatments with varied cooling 

rates and coarsening heat treatments were performed at multiple temperatures and time intervals, to 

enable assessments of γ′ morphology and size. γ′ coarsening rates were then compared. Coarsening rates 

were very similar for all three alloys in the heat treatment conditions examined. Higher coarsening rates 

were observed for γ′ precipitates residing along grain boundaries than for those within grains in all three 

alloys, during slow-moderate quenching from supersolvus solution heat treatments, and during aging at 

temperatures of 843 °C and higher. The γ′ precipitates of Alloy 10 appeared to remain coherent to higher 

temperatures than for LSHR and ME3.

Conclusions

1. For powder metal superalloys having similar γ′ contents, alloys having higher γ′ solvus temperature 

can have a higher content of γ′ precipitates at 900 to 1000 °C, which remain coherent up to these high 

temperatures. 

2. For such alloys having similar γ′ contents and γ- γ′ lattice mismatches, in spite of the different γ′

solvus temperatures, coarsening rates of γ′ precipitates within grains can be quite comparable. The 

coarsening rates can be estimated well using static coarsening experiments, and agree well with 

theoretical predictions.
NASA/TM—2016-218936 12

3. The coarsening rate of γ′ precipitates along the grain boundaries in these alloys can be significantly 

higher than for that within grains, with the enhancement of coarsening rate at grain boundaries increasing 

with temperature. 

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NASA/TM—2016-218936 13

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NASA/TM—2016-218936 14

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NASA/TM—2016-218936 15

TABLE 1.—COMPOSITIONS OF ALLOY 10, LSHR, AND ME3 FORGING SAMPLES, 

USING INDUCTIVELY COUPLED PLASMA MASS SPECTROMETRY MEASUREMENTS

wt% ME3 LSHR Alloy 10

Al 3.36 3.4 3.73

B 0.02 0.03 0.022

C 0.053 0.024 0.033

Co 21.04 21.3 15.23

Cr 12.95 12.9 10.66

Fe 0.036 0.056 0.027

Mo 3.71 2.66 2.64

N 0.0021 NR 0.0015

Nb 0.8 1.35 1.68

Ni 49.6 48.5 54.9

O 0.02155 0.0215 0.0125

S 0.0003 0.002 <0.0001

Si 0.033 0.035 0.023

Ta 2.4 1.66 0.88

Ti 3.83 3.67 4 

W 2.07 4.3 6.06

Zr 0.045 0.048 0.093

TABLE 2.—LATTICE PARAMETERS OF AND ASSOCIATED MISMATCH VALUES OF THE ALLOYS

(a) γ′ extracted from Alloy 10, LSHR, and ME3 specimens solution heat treated 1199 °C for 1.17 h, furnace off cooled

at an average cooling rate of 25 °C/min, then aging heat treated at 855 °C for 4 h followed by 775 °C for 8 h

Sample Alloy GP-1x 

(Å)

GP-2x 

(Å)

GP-3x 

(Å)

Mean 

(Å)

95% conf. interval on 

mean 

(Å)

8045-1 Alloy 10 3.59400 3.59400 3.59396 3.59399 0.00006

8045-2 LSHR 3.59559 3.59561 3.59559 3.59560 0.00003

8045-3 ME3 3.59452 3.59452 3.59456 3.59453 0.00005

(b) Mismatch of γ and γ′ lattice parameters within Alloy 10, LSHR, and ME3 after solution heat treatment and water

quenching or very slow furnace ramp cool. Based on comparison of the results for AEQ, LEQ, and MEQ and 

the results for extracted γ′ in Table 2(a), a1 corresponds to the γ′ phase, and a2 corresponds to the γ phase

Alloy Heat Treatment Sample Fundamental parameters Split-Pearson VII

a1 

(Å)

a2 

(Å)

δ 

(%)

a1

(Å)

a2 

(Å)

δ 

(%) 

Alloy 10 1199 °C/1 h/quenched 5,000 °C/min to 650 °C AW (M294) 3.59946 3.59904 

Alloy 10

1199 °C /1 h/quenched 0.033 °C/min to 

870 °C + 0.003 °C/min to 650 °C

AEQ (M291) 3.59592 3.5899 0.16739 3.59592 3.58899 0.19307

LSHR 1199 °C/1 h/quenched 5,000 °C/min to 650 °C LW (M295) 3.59912 3.59871 

LSHR

1199 °C /1 h/quenched 0.033 °C/min to 

870 °C + 0.003 °C/min to 650 °C

LEQ (M292) 3.59882 3.5904 0.23448 3.59918 3.58959 0.26662

ME3 1199 °C/1 h/quenched 5,000 °C/min to 650 °C MW (M296) 3.59789 3.59764 

ME3

1199 °C/1 h/quenched 0.033 °C/min to 870 °C 

+ 0.003 °C/min to 650 °C

MEQ (M293) 3.59709 3.5884 0.24178 3.59736 3.58911 0.22949
NASA/TM—2016-218936 16

TABLE 3.—RESULTS OF DIFFERENTIAL THERMAL ANALYSES (DTA) 

(a) The γ′ dissolution, γ′ formation, boride dissolution and boride formation temperatures measured from triple-cycled DTA 

of LSHR, ME3, and Alloy 10. The γ′ dissolution temperature is commonly reported as the γ′ solvus temperature, 

while the boride dissolution is reported as the boride liquation temperature

Alloy

temperature rate

initial weight 

(final weight gain)

Cycle Secondary γ′

dissolution 

(heating) 

°C

Valley between 

secondary and

primary γ′

dissolution 

(heating) 

°C

γ′ dissolution 

(heating)

°C

γ′ formation

(cooling)

°C

Boride 

dissolution 

(heating)

°C

Boride 

formation

(cooling)

°C

LSHR First 1068.0 1084.1 1162.4 1123.6 Out of range ND

10 °C/min Second -------- -------- 1158.7 1124.7 Out of range ND

1516.5 (0.1) mg Third -------- -------- 1158.7 1125.3 Out of range ND

Mean ± σ -------- -------- 1159.9 ± 2.1 1124.5 ± 0.9 --- ---

LSHR First 1067.6 1090.7 1154.4 1126.9 1207.6 1175.2

5 °C/min Second -------- -------- 1156.8 1126.8 1201.3 1173.1

1586.5 (0.2) mg Third -------- -------- 1155.8 1126.8 1202.4 1174.8

Mean ± σ -------- -------- 1155.8 ± 1.2 1126.8 ± 0.06 1203.8 ± 3.4 1174.4 ± 1.1

LSHR First 1069.6 1078.8 1151.2 1129.3 1204 1179.6

2 °C/min Second -------- -------- 1154.5 1129.3 1200.5 1177.6

1501.2 (0.7) mg Third -------- -------- 1151.2 1129.1 1201.0 1177.6

Mean ± σ -------- -------- 1152.3 ± 1.3 1129.2 ± 0.12 1201.8 ± 1.9 1178.3 ± 1.1

LSHR, 0 °C/min Extrapolated 1150.7 ± 0.9 1130.1 ± 0.6

ME3 First 1077.4 1100.8 1165.2 1125.9 1212 ND

10 °C/min Second -------- -------- 1162.9 1126.2 1207.2 ND

1530.9 (0.3) mg Third -------- -------- 1163.1 1126.2 1206.8 ND

Mean ± σ -------- -------- 1163.7 ± 1.3 1126.1 ± 0.17 1208.7 ± 2.4 ---

ME3 First 1080.3 1096.4 1158.0 1128.3 1208.7 ND

5 °C/min Second -------- -------- 1157.5 1128.3 1204.4 ND

1554.6 (0.2) mg Third -------- -------- 1156.4 1128.4 1204.8 ND

Mean ± σ -------- -------- 1157.3 ± 0.8 1128.3 ± 0.06 1206.0 ± 2.4 ---

ME3, 0 °C/min Extrapolated 1153.5 ± 0.9 1131.5 ± 0.3

Alloy 10 First 1089.3 1104.2 1193.8 1166.6 1205.5 ND

10 °C/min Second -------- -------- 1190.8 1166.2 1204.7 ND

1525.1 (0.1) mg Third -------- -------- 1190.6 1166.6 1204.7 ND

Mean ± σ -------- -------- 1191.7 ± 1.8 1166.5 ± 0.23 1205.0 ± 0.5 ---

Alloy 10 First 1089.5 1099.2 1193.4 1168.2 1203.9 ND

5 °C/min Second -------- -------- 1192.7 1168.2 1203.2 ND

1552.7 (0.3) mg Third -------- -------- 1191.6 1168.3 1202.5 ND

Mean ± σ -------- -------- 1192.6 ± 0.9 1168.2 ± 0.06 1203.2 ± 0.7 ---

Alloy 10, 0 °C/min Extrapolated ~1192 1171.7 ± 0.6

(b) The solidus and liquidus temperatures measured from the heating thermograms

at 5 °C/min of single-cycled DTA of as-forged LSHR, ME3, and Alloy 10

Alloy Thermal path on heating during DTA analysis Initial 

weight, 

mg

Weight 

gain, 

mg

Solidus 

(heating)

°C

Liquidus 

(heating)

°C

LSHR 25 to 700 °C at 20 °C/min, 700 to 1370 °C at 5 °C/min 683.1 0.4 1240 1334.5

ME3 25 to 1370 °C at 5 °C/min 1390.8 0.5 1264 1339.0

Alloy 10 25 to 1370 °C at 5 °C/min 1396.1 0.7 1257 1339.6
NASA/TM—2016-218936 17

TABLE 4.—SIZES AND AREA FRACTIONS OF Γ′ IN ALLOY 10, LSHR, 

AND ME3 AFTER DIFFERENT HEAT TREATMENTS

Prepared 

sample 

ID

Alloy Solution 

quench

Aging 

temperature,

C 

Aging 

Time,

h 

Aging 

quench

Met. mount 

ID

Inside 

grain sec. 

γ′ radius, 

avg., 

μm

Inside 

grain sec. 

γ′ radius, 

StDev, 

mm 

Inside 

grain sec. 

γ′ radius, 

count, 

n 

Grain 

bound. 

sec. γ′

radius, 

avg., 

mm 

Grain 

bound. 

sec. γ′

radius, 

StDev, 

mm 

Grain 

bound. 

sec. γ′

radius, 

count, 

n 

Overall 

sec. γ′

area 

fraction

Overall 

tert. γ′

area 

fraction

MEQ ME3 Furnace ramp ---- ---- ------- I379,M293 1.916 0.71 49 ------- ------- ----- ------- ~0

LEQ LSHR Furnace ramp ---- ---- ------- I380,M292 2.054 1.086 58 ------- ------- ----- ------- ~0

AEQ Alloy 10 Furnace ramp ---- ---- ------- J023,M291 3.41 0.918 138 ------- ------- ----- ------- ~0

MF ME3 Furnace ---- ---- ------- I877 0.152 0.049 342 0.361 0.147 719 0.536 0.038

LF LSHR Furnace ---- ---- ------- I876 0.162 0.040 306 0.371 0.127 77 0.542 0.054

AF Alloy 10 Furnace ---- ---- ------- I875 0.188 0.073 397 0.427 0.171 724 0.539 ~0

MW ME3 Water ---- ---- ------- I767 0.022 0.004 163 ------- ------- ----- ~0.57 ~0

LW LSHR Water ---- ---- ------- I766 0.014 0.003 153 ------- ------- ----- ~0.57 ~0

AW Alloy 10 Water ---- ---- ------- I765 0.022 0.004 175 ------- ------- ----- ~0.57 ~0

MW760-100 ME3 Water 760 100 Water TEM 0.021 0.006 152 0.022 0.005 13 ~0.57 ~0

LW760-100 LSHR Water 760 100 Water TEM 0.022 0.008 152 0.020 0.007 11 ~0.57 ~0

AW760-100 Alloy 10 Water 760 100 Water TEM 0.027 0.007 150 0.034 0.007 13 ~0.57 ~0

MW760-1k ME3 Water 760 1000 Water TEM 0.029 0.010 104 0.031 0.006 7 ~0.57 ~0

LW760-1k LSHR Water 760 1000 Water TEM 0.031 0.013 110 0.026 0.008 7 ~0.57 ~0

AW760-1k Alloy 10 Water 760 1000 Water TEM 0.031 0.012 102 0.025 0.005 7 ~0.57 ~0

LW843-1k LSHR Water 843 1000 Water TEM 0.115 0.039 153 0.130 0.066 31 ~0.57 ~0

MW843-1k ME3 Water 843 1000 Water TEM 0.116 0.040 167 0.106 0.037 48 ~0.57 ~0

AW843-1k Alloy 10 Water 843 1000 Water TEM 0.11 0.040 151 0.133 0.039 24 ~0.57 ~0

LW843-10 LSHR Water 843 10 Water TEM 0.02 0.007 153 0.025 0.007 10 ~0.57 ~0

MW843-10 ME3 Water 843 10 Water TEM 0.022 0.008 157 0.024 0.006 21 ~0.57 ~0

AW843-10 Alloy 10 Water 843 10 Water TEM 0.023 0.009 152 0.028 0.010 10 ~0.57 ~0

LW843-100 LSHR Water 843 100 Water TEM 0.051 0.018 158 0.064 0.023 14 ~0.57 ~0

MW843-100 ME3 Water 843 100 Water TEM 0.052 0.021 159 0.058 0.019 36 ~0.57 ~0

AW843-100 Alloy 10 Water 843 100 Water TEM 0.058 0.022 151 0.073 0.023 15 ~0.57 ~0

MW927-10 ME3 Water 927 10 Water TEM 0.063 0.017 156 0.074 0.023 13 ------- ~0

LW927-10 LSHR Water 927 10 Water TEM 0.060 0.019 155 0.079 0.019 17 ------- ~0

AW927-10 Alloy 10 Water 927 10 Water TEM 0.073 0.027 158 0.099 0.038 12 ------- ~0

MW927-99 ME3 Water 927 99 Water TEM 0.149 0.047 159 0.155 0.047 10 ------- ~0

LW927-99 LSHR Water 927 99 Water TEM 0.142 0.054 157 0.156 0.053 20 ------- ~0

AW927-99 Alloy 10 Water 927 99 Water TEM 0.144 0.046 160 0.141 0.045 16 ------- ~0

MW927-1k ME3 Water 927 1000 Water I850 0.238 0.072 365 0.324 0.071 100 0.398 ~0

LW927-1k LSHR Water 927 1000 Water I852 0.23 0.073 410 0.274 0.063 100 0.383 ~0

AW927-1k Alloy 10 Water 927 1000 Water I851 0.264 0.073 394 0.306 0.080 100 0.394 ~0

MW1093-5 ME3 Water 1093 5 Water K202 0.149 0.045 368 0.206 0.071 165 0.26 ~0

LW1093-5 LSHR Water 1093 5 Water K201 0.156 0.040 347 0.223 0.066 121 0.29 ~0

AW1093-5 Alloy 10 Water 1093 5 Water K200 0.153 0.040 336 0.243 0.094 213 0.41 ~0

MW1093-20 ME3 Water 1093 20.1 Water I847 0.255 0.074 514 0.405 0.133 615 0.184 ~0

LW1093-20 LSHR Water 1093 20.1 Water I849 0.252 0.064 1879 0.492 0.161 334 0.189 ~0

AW1093-20 Alloy 10 Water 1093 20.1 Water I848 0.273 0.086 723 0.542 0.244 659 0.322 ~0

MW1093-100 ME3 Water 1093 100 Water K205 0.369 0.115 307 0.675 0.278 353 0.27 ~0

LW1093-100 LSHR Water 1093 100 Water K204 0.360 0.106 325 0.662 0.293 299 0.28 ~0

AW1093-100 Alloy 10 Water 1093 100 Water K203 0.404 0.145 322 0.783 0.330 384 0.4 ~0

MW1093-1k ME3 Water 1093 1000 Water I764 0.69 0.284 285 1.42 0.543 192 0.224 ~0

LW1093-1k LSHR Water 1093 1000 Water I769 0.7 0.282 652 1.498 0.716 164 0.205 ~0

AW1093-1k Alloy 10 Water 1093 1000 Water I768 0.447 0.202 636 1.309 0.690 236 0.361 ~0
NASA/TM—2016-218936 18

TABLE 5.—COMPOSITIONS OF EXTRACTED γ′ AND γ, AND ASSOCIATED 

γ PARTITIONING RATIOS FOR ALLOY 10, LSHR, AND ME3

Alloy 10 LSHR ME3

γ γ ' γ 

partitioning 

ratio

γ γ ' γ 

partitioning 

ratio

γ γ ' γ 

partitioning 

ratio wt% at.% wt% at.% wt% at.% wt% at.% wt% at.% wt% at.%

Al 0.90 2.00 5.65 12.11 0.17 0.95 2.08 5.64 12.00 0.17 0.96 2.09 5.74 12.10 0.17

Co 25.09 25.43 8.20 8.05 3.16 31.99 32.16 11.44 11.15 2.88 31.08 31.00 10.43 10.07 3.08

Cr 23.50 26.99 1.89 2.10 12.84 24.90 28.37 1.91 2.12 13.41 24.16 27.31 1.84 2.01 13.60

Fe 0.05 0.05 0.02 0.02 2.35 0.10 0.10 0.02 0.02 4.19 0.07 0.07 0.01 0.02 4.90

Mo 4.43 2.76 1.53 0.92 2.99 4.36 2.69 1.24 0.74 3.63 6.31 3.86 1.42 0.84 4.58

Nb 0.19 0.12 2.95 1.84 0.07 0.24 0.15 2.60 1.61 0.10 0.18 0.11 1.58 0.97 0.12

Si 0.00 0.00 0.03 0.06 0.00 0.00 0.00 0.04 0.09 0.00 0.00 0.00 0.05 0.11 0.00

Ta 0.03 0.01 1.42 0.45 0.02 0.09 0.03 3.01 0.95 0.03 0.15 0.05 4.53 1.42 0.03

Ti 0.22 0.27 6.34 7.65 0.04 0.34 0.43 6.54 7.85 0.05 0.41 0.50 6.96 8.27 0.06

W 5.74 1.87 6.06 1.91 0.98 4.77 1.54 3.96 1.24 1.24 2.45 0.78 1.67 0.52 1.52

Zr 0.00 0.00 0.15 0.09 0.01 0.00 0.00 0.09 0.05 0.01 0.00 0.00 0.08 0.05 0.01

Ni 39.85 40.55 65.75 64.78 0.63 32.26 32.56 63.50 62.13 0.52 34.24 34.28 65.66 63.60 0.54

TABLE 6.—CALCULATED ACTIVATION ENERGIES FOR γ′ COARSENING IN ALLOY 10, LSHR, AND ME3

(a) Results for γ′ coarsening with data at 1093 °C included. 

Alloy ln k versus 1/T slopes for activation energy

Within grains Coefficient of 

determination

Grain boundaries Coefficient of 

determination

Alloy 10 –35132 0.898 –48940 0.983

LSHR –39557 0.979 –48922 0.992

ME3 –41144 0.964 –51403 0.993

Alloy Q (kJ/mole) = –(Slope*8.314)

Within grains -------- Grain boundaries

Alloy 10 292 -------- 407

LSHR 329 -------- 407

ME3 342 -------- 427

(b) Results for γ′ coarsening with data at 1093 °C omitted, due to loss of coherency at that temperature.

Alloy ln k versus 1/T slopes for activation energy

Within grains Coefficient of 

determination

Grain boundaries Coefficient of 

determination

Alloy 10 –52443 0.989 –55938 0.978

LSHR –46024 0.977 –50011 0.977

ME3 –50696 0.97 –57588 0.999

Alloy Q (kJ/mole) = –(Slope*8.314)

Within grains -------- Grain boundaries

Alloy 10 436 -------- 465

LSHR 383 -------- 416

Alloy 10 421 -------- 479
NASA/TM—2016-218936 19

TABLE 7.—CALCULATIONS FOR COARSENING RATE CONSTANT k (μm3

/s) AT T = 1093 °C 

Element Cγ Cγ′ D/DCr Thermo 

factor

kMLSW×107 

(mm3/s)

Cr 0.1740 0.0212 1.00 1.26 1.81

Al 0.0624 0.1205 3.07 2.5 7.89

Ti 0.0366 0.0785 2.33 1.97 8.82

Co 0.2231 0.1115 4.90 0.93 27.2

Mo 0.0183 0.0074 1.88 0.96 110.2

Nb 0.0067 0.0161 2.14 1.25 47.7

Ta 0.0044 0.0095 1.00 1.2 53.1

W 0.0139 0.0124 1.57 1.14 3238

φ = 0.191, σ = 23 mJ/m2, f (φ) = 1.9 (Ref. 29), VM = 7.22×10–6 m3/mol; DCr = 1.502×10–15 m2/s 

TABLE 8.—COMPARISON OF MEASURED (WITHOUT STEREOLOGICAL CORRECTION)

AND PREDICTED COARSENING RATE CONSTANTS FOR CHROMIUM (μm3/s) 

Alloy Temperature

(°C)

Measurement 

(0 to 1000 h)

Measurement 

(0 to 100 h)

Prediction based on 

Cr

Prediction based on 

all elements

LSHR 843 4.22×10–10 3.61×10–10 6.91×10–10 3.51×10–10

LSHR 927 3.38×10–9 8.03×10–9 5.74×10–9 3.18×10–9

LSHR 1093 0.99×10–7 1.30×10–7 1.81×10–7 1.14×10–7

Alloy 10 843 3.67×10–10 5.12×10–10 8.95×10–10 4.58×10–10

Alloy 10 927 5.11×10–9 8.35×10–9 7.59×10–9 4.19×10–9

Alloy 10 1093 2.45×10–8 1.83×10–7 2.30×10–7 1.43×10–7

TABLE 9.—COMPARISON OF MEASURED (WITH STEREOLOGICAL CORRECTION) 

AND PREDICTEDCOARSENING RATE CONSTANTS FOR CHROMIUM (μM3/S)

Alloy Temperature 

(°C)

Measurement 

(0 to 1000 h)

Measurement 

(0 to 100 h)

Prediction based on 

Cr

Prediction based on 

all elements

LSHR 843 4.22×10–10 3.61×10–10 6.91×10–10 3.51×10–10

LSHR 927 6.21×10–9 8.02×10–9 5.74×10–9 3.18×10–9

LSHR 1093 1.83×10–7 2.38×10–7 1.81×10–7 1.14×10–7

Alloy 10 843 3.67×10–10 5.12×10–10 8.95×10–10 4.58×10–10

Alloy 10 927 9.39×10–9 8.35×10–9 7.59×10–9 4.19×10–9

Alloy 10 1093 4.50×10–8 3.37×10–7 2.30×10–7 1.43×10–7
Figure 1.—Typical cooling curves of temperature versus time measured after the various solution heat treatments: (a) Water quenched, (b) Furnace off cooled, 

(c) Very slow furnace ramp. 

NASA/TM—2016-218936 20
Figure 2.—Microstructures of specimens used in phase extractions: solution heat treated 1199 °C for 1.17 h, cooled in the furnace at an average cooling rate of 

25 °C/min, then aging heat treated at 855 °C for 4 h followed by 775 °C for 8 h: (a) Alloy 10, (b) LSHR, and (c) ME3.

NASA/TM—2016-218936 21
Figure 3.—Images of the complex γ′ microstructures produced within grains and at grain boundaries (GB) of samples solution heat treated at 1199 °C for 1 h 

followed by cooling at very slow cooling rates to approximate thermodynamic equilibrium, using 0.033 °C/min. to 870 °C, then 0.003 °C/min. to 650 °C: (a) Alloy 

10, (b) LSHR, and (c) ME3. 

NASA/TM—2016-218936 22
Figure 4.—DTA results. (a) Thermograms acquired from differential thermal analysis of as-forged samples of LSHR, ME3 and Alloy 10 that 

were triple cycled at either 10, 5, or 2 °C/min. The triple cycle allows for multiple measurements of the γ′ dissolution, γ′ formation, boride 

dissolution, and boride formation temperatures. 

NASA/TM—2016-218936 23
Figure 4.—(Concluded) (b) Mean γ′ dissolution and γ′ formation temperatures determined from differential thermal analyses (DTA) of 

(a) LSHR, ME3 and b) Alloy 10 as a function of heating or cooling rates. The values marked are the linear extrapolation to 0 °C/min 

and the standard error associated with this fit. (c) Heating thermograms acquired to obtain the solidus and liquidus temperatures 

from differential thermal analysis at 5 °C/min of as-forged samples of LSHR, ME3 and Alloy 10. 

NASA/TM—2016-218936 24
Figure 5.—Image and precipitate size distribution measured within grains for after solution heat treatment of 1199 °C for 1 h followed by water 

quench at an average cooling rate of 5,000 °C/min: (a) Alloy 10 (AW), (b) LSHR (LW), and (c) ME3 (MW).

NASA/TM—2016-218936 25
Figure 6.—Images at low and high magnifications of samples solution heat treated 1199 °C for 1 h followed by furnace off cooling at an average cooling rate of 

25 °C/min., showing segmented cuboidal morphology (S) often present within grains: (a) Alloy 10, (b) LSHR, and (c) ME3.

NASA/TM—2016-218936 26
Figure 7.—Solution heat treated and water quenched, then aged at 760 °C for 1,000 h: (a) Alloy 10, (b) LSHR, and (c) ME3. 

NASA/TM—2016-218936 27
Figure 8.—Solution heat treated and water quenched, then aged at 843 °C for 1,000 h: (a) Alloy 10, (b) LSHR, and (c) ME3. Coarser precipitates became evident 

along grain boundaries.

NASA/TM—2016-218936 28
Figure 9.—Solution heat treated and water quenched, then aged at 927 °C for 1,000 h: (a) Alloy 10, (b) LSHR, and (c) ME3. 

NASA/TM—2016-218936 29
Figure 10.—Solution heat treated and water quenched, then aged at 1093 °C for 1,000 h: (a) Alloy 10, (b) LSHR, (c) ME3, (d) Alloy 10 TEM, and (e) ME3 TEM. 

Coalescence between precipitates is apparent, and more advanced for LSHR and ME3. Very coarse precipitates are present at some locations along grain 

boundaries, indicating enhance coarsening there.

NASA/TM—2016-218936 30
Figure 11.—Misfit dislocations associated with loss of coherency 

for precipitates during aging at 1093 °C: ME3 water quenched, 

then aged 1093 °C for 1000 h.

NASA/TM—2016-218936 31
Figure 12.—Precipitate size-frequency distributions within grains: (a) 760 °C/1000 h, (b) 843 °C/1000 h, (c) 927 °C/1000 h, and (d) 1093 °C/1000 h. 

NASA/TM—2016-218936 32
Figure 13.—Mean sizes and area fractions of γ′ precipitates compared versus temperature after aging at 1,000 h: (a) mean equivalent radius of precipitates 

measured within grains (WG) and of precipitates measured along grain boundaries (GB);and (b) overall area fraction versus temperature aged for 1,000 h.

NASA/TM—2016-218936 33
Figure 14.—Comparison of analytical fits for the γ′ solvus-approach curve (smooth lines, Eq. (1)) and 

measurements (data points) from Figure 13(b) for: (a) LSHR and (b) Alloy 10.

Figure 15.—Plot of γ′ solvus temperature as a function of 4C* for 

various γ′-strengthened superalloys.

NASA/TM—2016-218936 34
Figure 16.—Comparisons of mean equivalent radius after coarsening treatments at all aging times and temperatures examined for each alloy. 

NASA/TM—2016-218936 35
NASA/TM—2016-218936 36
Figure 17.—Comparisons of mean equivalent radius cubed after coarsening treatments at all aging times and temperatures examined for each

alloy, as used for determining coarsening constant k. 

NASA/TM—2016-218936 37
Figure 18.—Comparisons of ln k versus reciprocal temperature all aging 

temperatures examined for each alloy, as used for determining activation 

energies Q.

NASA/TM—2016-218936 38