T.P. Gabb

Glenn Research Center, Cleveland, Ohio

A. Garg

University of Toledo, Toledo, Ohio 

D.R. Miller

The Ohio State University, Columbus, Ohio

C.K. Sudbrack, D.R. Hull, D. Johnson, R.B. Rogers, and J. Gayda

Glenn Research Center, Cleveland, Ohio

S.L. Semiatin

U.S. Air Force Research Laboratory, Wright-Patterson Air Force Base, Dayton, Ohio

Formation of Minor Phases in a 

Nickel-Based Disk Superalloy

NASA/TM—2012-217604

July 2012
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T.P. Gabb

Glenn Research Center, Cleveland, Ohio

A. Garg

University of Toledo, Toledo, Ohio 

D.R. Miller

The Ohio State University, Columbus, Ohio

C.K. Sudbrack, D.R. Hull, D. Johnson, R.B. Rogers, and J. Gayda

Glenn Research Center, Cleveland, Ohio

S.L. Semiatin

U.S. Air Force Research Laboratory, Wright-Patterson Air Force Base, Dayton, Ohio

Formation of Minor Phases in a 

Nickel-Based Disk Superalloy

NASA/TM—2012-217604

July 2012

National Aeronautics and

Space Administration

Glenn Research Center

Cleveland, Ohio 44135
Acknowledgments

This work was supported by the NASA Aircraft Aging and Durability task within the Aviation Safety Program, and the NASA 

Undergraduate Student Research Program. The authors would like acknowledge Terry McCue for his help in the operation and 

maintenance of the FESEM. The authors would like to acknowledge Anna Palzer for acquiring the DTA data. The authors would 

also like to acknowledge helpful discussions with Greg Olsen and Herng Jou of Questek Innovations, LLC, Jeff Simmons of Air 

Force Research Laboratory, Wright-Patterson Air Force Base, Mark Hardy and Randy Helmink of Rolls-Royce, and Rebecca 

MacKay of NASA GRC. 

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This work was sponsored by the Fundamental Aeronautics Program 

at the NASA Glenn Research Center.

Level of Review: This material has been technically reviewed by technical management. 
Formation of Minor Phases in a 

Nickel-Based Disk Superalloy 

T.P. Gabb 

National Aeronautics and Space Administration

Glenn Research Center

Cleveland, Ohio 44135 

A. Garg

University of Toledo

Toledo, Ohio 43606 

D.R. Miller

The Ohio State University

Columbus, Ohio 43210 

C.K. Sudbrack, D.R. Hull, D. Johnson, R.B. Rogers, and J. Gayda 

National Aeronautics and Space Administration

Glenn Research Center

Cleveland, Ohio 44135 

S.L. Semiatin

U.S. Air Force Research Laboratory

Wright-Patterson Air Force Base

Dayton, Ohio 45433 

Abstract

The minor phases of powder metallurgy disk superalloy LSHR were studied. Samples were 

consistently heat treated at three different temperatures for long times to approximate equilibrium.

Additional heat treatments were also performed for shorter times, to then assess non-equilibrium 

conditions. Minor phases including MC carbides, M23C6 carbides, M3B2 borides, and sigma were 

identified. Their transformation temperatures, lattice parameters, compositions, average sizes and total 

area fractions were determined, and compared to estimates of an existing phase prediction software 

package. Parameters measured at equilibrium sometimes agreed reasonably well with software model 

estimates, with potential for further improvements. Results for shorter times representing non-equilibrium 

indicated significant potential for further extension of the software to such conditions, which are more 

commonly observed during heat treatments and service at high temperatures for disk applications. 

1.0 Introduction

Structure, composition, contents and size distributions of carbide and boride phases can significantly 

influence the processing and mechanical properties of nickel-base superalloys (Refs. 1 to 6), especially at 

elevated temperatures where grain boundaries influence strength as well as time-dependent deformation 

and failure modes. Cubic MC carbides are stable up to very high temperatures near superalloy 

solidification, and therefore can help constrain grain growth during solution heat treatments above the 

gamma prime solvus (Ref. 7). The elements Ta, Ti, Nb, and W often combine with carbon to form this 

carbide. M23C6 carbides are stable at lower temperatures up to near 900 °C, and can improve the 

resistance to grain boundary sliding, rupture and dwell crack growth in service temperatures of 600 to 

815 °C. Cr, Mo, and W can react with carbon to form this carbide. Boron in solid solution within the γ

NASA/TM—2012-217604 1
matrix also improves the resistance to grain boundary sliding, rupture, and dwell crack growth (Ref. 5).

Boride phase M3B2 has not been proven to specifically improve properties, but is often observed in disk 

superalloys containing boron. This phase can have a melting point only slightly higher than solution heat 

treatment temperatures (Ref. 6), to limit solution heat treatment upper tolerance temperatures. Mo, Cr, 

and W often combine with B to form this phase. 

Harmful topological close packed (TCP) phases can also form in disk superalloys (Ref. 1 and 2).

These phases such as σ, μ, Laves, and P have low inherent ductility due to limited slip systems, and can 

have needle or lathe morphologies, resulting in lower tensile and rupture ductilities at application 

temperatures (Refs. 8 and 9). Therefore, several models have been developed over the years to predict the 

formation of such phases (Refs. 10 and 11). 

Recent advances in calculating thermodynamic properties of multi-component systems have enabled 

microstructural modeling software packages such as Thermo-Calc (Thermo-Calc Software, Inc.) 

(Ref. 12), JMatPro (Sente Software Ltd.) (Ref. 13), and Pandat (CompuTherm LLC) (Ref. 14) to estimate 

relative percentages and compositions of phases at equilibrium in superalloys at various temperatures. 

However, for use with a specific alloy of interest, this work has shown such estimates need to be 

compared with experimentally-measured phase identities and contents after known heat treatments 

approximating equilibrium, for model verification and potential refinements (Ref. 15). Experimental 

measurements at non-equilibrium conditions could also be used to help understand and model formation 

kinetics of minor phases. The primary objective of this work was to provide experimental composition, 

content, and size measurements of the predominant minor phases of the disk superalloy LSHR after 

varied heat treatments, for comparison with typical estimates of such software packages. 

2.0 Materials and Procedures

Disk superalloy LSHR having a composition in weight percent of 3.46Al, 0.028B, 0.029C, 20.7Co, 

12.52Cr, 0.07Fe, 2.73Mo, 1.45Nb, 1.6Ta, 3.50Ti, 4.33W, 0.049Zr, bal. Ni and trace impurities was 

produced through powder processing. It was atomized in argon, hot compacted, extruded and then 

isothermally forged into a flat, uniform disk 5.08 cm thick and 30.48 cm in diameter. 

Phase extraction samples were removed by electro-discharge machining pins of about 10 g weight 

and 10 mm diameter in the circumferential direction from the rim of the disk. These samples were 

solution heat treated at 1199 °C for 1 h and furnace cooled at an average rate of 21 °C/min. They were 

then aged at 843 °C for 1,000 h and water quenched. Electrochemical extractions of minor phase were 

performed using a 10 ml HCl-1 g tartaric acid-90 ml methanol electrolyte. Triplicate extractions were 

performed in each case. Extracted phases and lattice parameters were identified using x-ray diffraction

(XRD). XRD data was gathered with a Bruker D8 Advance (Bruker Corporation) diffractometer using Cu 

Kα radiation in a Bragg-Brentano geometry with a solid state linear position sensitive detector.

Qualitative phase identification was performed using commercial software. Because of significant peaks 

overlapping, Rietveld refinement (Ref. 16) was then employed to determine lattice parameters using a 

fundamental parameters approach as implemented in the Bruker TOPAS (Bruker Corporation) software 

program. The Rietveld method employs (Ref. 17) a least-squares fitting approach using a mathematical 

model based on the scattered intensities of constituent crystalline phases. In the complex mixtures of 

phases examined in this work, the accuracy of the lattice parameters of any individual phase was limited 

by the amount of the phase present. The accuracy of major phases was estimated to be as high as 5 

significant digits and that of trace phases was as low as 2 significant digits. 

Small pins of 3 mm diameter and 16 mm long were also extracted in the circumferential direction 

near the rim of the disk, for various heat treatments. The conditions of these samples are summarized in 

Table 1. Samples were extracted from the extrusion before forging, and from the forging, for examination 

without further material processing. Single samples were only solution heat treated at 1199 °C for 1 h, 

then water quenched at an average cooling rate of 5,000 °C/min. or furnace cooled at a controlled average 

cooling rate of 21 °C/min. Additional samples were solution heat treated at 1199 °C for 1 h, water 

NASA/TM—2012-217604 2
quenched at an average cooling rate of about 5,000 °C/min to minimize phase nucleation and growth 

during the cooling process, and subsequently given aging heat treatments. They were aged at temperatures 

of 760, 843, and 927 °C for times of 10, 100, and 1,000 h and again water quenched, yielding nine 

different aged cases. 

Differential thermal analysis (DTA) experiments were performed using a Netzsch STA 409C to help 

estimate the formation temperatures of minor phases, the solidus and the liquidus. Cylindral samples with

5 mm diameter and lengths ranging from 4 to 12 mm and weights from 683 to 1587 mg were extracted 

from the as-forged disk. The STA chamber was backfilled twice with Ar to remove oxygen before 

starting to heat. Under a 40 to 60 mL/min He flow, samples were heated quickly, either at 20 °C/min or at 

10 °C/min, to a starting temperature, Tstart, of 700 to 1000 °C. Then, using two rates, 5 and 2 °C/min, the 

samples were either i) heated and cooled three times between Tstart and a maximum temperature (Tmax) of 

1220 °C or ii) heated through melting to Tmax of 1342 to 1480 °C then cooled to Tstart. Formation 

temperatures with distinct peaks, i.e. maximum M3B2 boride dissolution on heating, maximum M3B2

boride formation on cooling and the liquidus on heating, were 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 temperature at which the MC 

carbides liquidate fully was estimated by the tangential intersections of the approximately flat slopes on 

either side of transition. To confirm this estimate, cylindrical samples with a 3 mm diameter and a 16 mm 

length were heat treated in air for 1 h at 1250, 1300, and 1325 °C, water quenched, metallographic 

prepared and examined using back-scattered SEM at 15 kV and energy dispersive X-ray spectrometry. 

Selected heat treated and quenched samples were also sectioned and conventionally 

metallographically prepared. The sections were etched using a waterless Kallings etch. Linear intercept 

grain size was measured from optical images according to ASTM E112. Scanning electron microscope 

(SEM) evaluations were used to select samples for determination of minor phase compositions using 

TEM and microprobe. Thin disks of 3mm diameter and ~ 0.5mm thick were sliced and mechanically 

ground and polished from selected heat-treated samples. These disks were then electrochemically thinned 

to electron transparency using a solution of 10 percent perchloric and 90 percent methanol cooled to –30 

°C. A FEI CM200 Transmission Electron Microscope (TEM) operating at 200 kV was used for imaging 

and to identify the crystal structure of predominant minor phases using selected area electron diffraction

patterns (SAEDP) and energy dispersive X-ray spectrometry (EDS). Associated bright field images were 

used to estimate the typical size range and morphology of each phase. Sections having larger size minor 

phase particles were usually selected for estimation of minor phase compositions using a JEOL JXA-8200 

scanning electron microprobe. At least 5 particles of each minor phase were analyzed. 

Sizes and area fractions of various second phases were quantified via SEM using metallographically 

prepared sections, in order to sample sufficient areas. Samples were imaged in a Hitachi S-4700 Field 

Emission SEM at 20 kV using a working distance of 12 mm. Secondary electron and backscattered 

electron images were obtained at 2,000 x magnification, so that 4-8 grains and their associated boundaries 

would be captured within each image. This enabled statistically valid area fraction estimates to be 

determined for each phase in each image, even for phases such as M23C6 which was observed only at 

grain boundaries. Five adjacent images were acquired for each sample, with no bias of imaging towards a 

certain region of particles or grain boundaries. An EDAX CDU LEAP (EDAX, Inc.) detector attached to

the SEM was used to obtain semi- quantitative chemical analysis of each particle visible in these images.

Compositions determined from EDS spectra were then compared to results previously obtained in thin 

foil TEM evaluations. While the EDS spectra acquired using TEM and SEM were not identical, relative 

major elemental compositions for a given phase were consistent. Based on these analyses, the phase 

identity of each particle was recorded, and later used for size determinations of various phases. The 

combination of these techniques enabled discrimination and measurement of minor phase contents and 

particle sizes over a much larger imaged area than what would have been feasible using TEM alone. 

NASA/TM—2012-217604 3
Particle sizes were measured using Sigma Scan Pro 5.0 image analysis software. The difference in 

brightness between minor and γ-γ ′ phases in secondary electron SEM images enabled brightness-based 

local thresholding. Each minor phase was measured separately, and also classified by location, either 

within grains or at grain boundaries. Area, equivalent radius, major axis, minor axis, and perimeter were 

determined for each particle. The equivalent radius calculation assumed a spherical shape and used the 

formula: 

Equivalent Radius= Area / π

Phase area fraction was determined by combining the areas of all particles of an identity and class within 

an image. Size-frequency histograms were generated for each phase with more than 10 particles acquired. 

JMatPro-Ni6.0 and Pandat 8.1 software packages were used to predict minor phase characteristics, 

assuming thermodynamic equilibrium. However, only phases identified experimentally, using both 

powder X-ray diffraction of minor phase extractions and selected area electron diffraction patterns of 

particles within thin foils in the TEM, were selected in the phase catalogues for each package. The 

packages were otherwise used with no modifications or custom enhancements to the models or input 

databases, for these initial estimates. 

3.0 Results and Discussion

3.1 Heat Treatments

Typical SEM images of samples after extrusion and after subsequent forging are compared in Figure

1. Linear intercept grain sizes of these samples were 4.1 and 5.5 μm, respectively. The large dark phase 

particles are primary γ ′, with mean radius of 0.4 and 1.1 μm, respectively. These particles acted to 

constrain grain growth during extrusion and forging. Samples subsequently solution heat treated at 

1199 °C for 1 h and either furnace cooled or water quenched are compared in Figure 2. The primary γ ′

particles have dissolved during this solution treatment. Supersolvus solution heat treatments thereby 

allowed grain growth to a larger linear intercept grain size, near 25 μm. Only MC carbides and M3B2

borides were identified at significant quantities in these samples. Images of samples aged 1,000 h to 

approach equilibrium at 760, 843, and 927 °C are compared in Figure 3. MC and M23C6 carbides, M3B2

borides, and σ phase were identified in these conditions. Images of samples aged for shorter times, 10 and 

near 100 h, are also compared for 760, 843, and 927 °C in Figure 4, 5, and 6, respectively. The formation 

and growth of some of these phases with increasing aging time are evident.

Figure 7 shows thermograms from the DTA experiments. SEM imaging of forged samples that were 

heat treated in air for 1 h at 1250, 1300, and 1325 °C and then water quenched are compared in Figure 8. 

These heat treatment temperatures are all above the average DTA solidus value of 1239.5 °C, which was 

near the respective JMatPro and Pandat predictions of 1234.8 and 1260.5 °C, Table 2. This SEM 

evaluation established that the DTA solidus is reasonable. It is clear that some of grain boundaries have 

melted during the 1250 °C heat treat and then resolidified upon quenching. With increasing temperature, 

more of this melting and resolidification was observed. Mean liquidus measured with an average of four 

DTA runs was 1334.1 ± 2.3 °C, near the JMatPro prediction of 1340 °C and significantly below that of 

Pandat, 1402.8 °C. 

3.2 Minor Phases

Predominant phases detected by x-ray diffaction and TEM observations are identified with 

measured lattice parameters in Table 3. MC and M23C6 carbides, M3B2 boride, and σ phases were 

identified by these approaches, and are shown in Figure 9 to 12. Compositions measured in atomic 

percent are compared to those predicted by JMatPro and Pandat in Table 4. 

NASA/TM—2012-217604 4
Subsequent measurements of minor phases over multiple SEM images indicated the equivalent radius 

of each minor phase could be assumed normally distributed (Figs. 13 to 16), so that the size distribution 

could be approximated using mean and standard deviation for each phase. Resulting SEM measurements 

of equivalent radius and area fraction for MC and M23C6 carbides are compared in Tables 5 and 6.

Equivalent radius and area fraction results for M3B2 boride and σ phase particles are compared in Tables 7 

and 8. Total area fractions of these minor phases measured after aging for 1,000 h are compared to those 

predicted by JMatPro and Pandat in Table 9. 

3.3 MC Carbides

Figure 9 shows a typical TEM image and diffraction pattern of MC carbides. These carbides have an 

FCC Fm3m crystal structure, with a lattice parameter of a = 4.3820 angstroms (Å). This was comparable 

to that reported for several other disk superalloys (Ref. 18), but slightly higher than that reported for Rene 

88DT, an alloy somewhat similar in composition to LSHR (Ref. 19]. Typical size-frequency histograms 

acquired for MC carbides are shown in Figure 13. Equivalent radius and area fraction are tabulated versus 

sample condition in Table 5. MC carbide average equivalent radius did not excede 0.27 μm, making 

quantitative microprobe measurements of the composition difficult. EDS measurements of these carbides 

in TEM thin foils indicated this phase had an approximate composition of 12Nb-12Ta-26Ti-50C in 

atomic percent, Table 4. Quantitative compositions are not often reported for minor phases such as 

carbides and borides existing at such small sizes, due to difficulties of in-situ measurements using 

scanning electron microscopy. However, surveys indicate many other disk superalloys also have MC 

carbides which contain Ti, Ta, and Nb (Refs. 1, 2, 18, and 19). The slightly higher lattice parameter 

compared to Rene 88DT could be related to the presence of Ta in MC carbides for LSHR. The measured 

composition was near that predicted at equilibrium by Pandat, but markedly differed in Ti and Zr contents 

from JMatPro estimates. Overall MC carbide size, morphology, and area fraction appeared comparable in 

most samples, indicating high overall phase stability for the various conditions examined. While most 

appeared spherical or slightly elliptical in the SEM images presented, TEM examinations indicated the 

MC carbides usually had short cylindrical or blocky shapes, each made up of a single crystal.

Yet, some changes in MC carbides were observed when comparing those within grains versus at grain 

boundaries. Over 80 percent of the carbides were observed at grain boundaries for the as-extruded 

condition, Figure 1(a). About 35 percent of the carbides were observed at grain boundaries after forging, 

Figure 1(b). However, no more than 23 percent were observed at grain boundaries for samples 

supersolvus heat treated and subsequently furnace cooled or water quenched and aged, Figure 2 to 6. Here 

they more often appeared within the grains, sometimes loosely distributed in clusters. Therefore, a 

majority of these carbides did not appear to pin grain boundaries, to ultimately constrain grain growth

during these solution heat treatments. The MC carbides at grain boundaries were often slightly larger than

those within grains. This suggested that larger MC carbides may be more effective at pinning grain 

boundaries. MC carbide content at grain boundaries decreased with increasing aging time at 843 °C, 

which will be further considered below. 

Overall MC carbide content was estimated at equilibrium conditions using JMatPro and Pandat in 

Table 9. The estimates agree reasonably well with experimental measurements at 927 and 843 °C, Figure

17, and suggest that MC carbides remained near equilibrium for samples aged at these temperatures for 

1,000 h up to 1200 °C. A much lower content was estimated by JMatPro and Pandat at 760 °C, not 

observed in the experimental measurements. However, MC carbide content at grain boundaries was 

observed to decrease with increasing aging time at 843 °C (Table 5), which will be further considered 

below. 

The average measured MC phase transformation temperature of 1310.2 °C agreed reasonably well 

with that predicted by JMatPro and Pandat, 1303.1 and 1316.7 °C respectively, as indicated in Table 2. 

The 1300 and 1325 °C heat treatments are lower and upper bounds for the DTA extracted temperature for 

full liquation of the MC carbides at 1310 °C; however locating MC carbides to image and analyze with 

NASA/TM—2012-217604 5
EDS was challenging for these conditions. The MC carbides that were located in the three samples 

quenched from the high temperature are compared in Figure 8. 

3.4 M23C6 Carbides

Figure 10 shows a typical TEM image and diffraction pattern of M23C6 carbides. These carbides have 

an FCC Fm3m crystal structure, with lattice parameter estimated using TEM of 10.6 Å. This is very close 

to that reported for Rene 88DT (Ref. 19). The average phase composition was measured by microprobe 

on a sample aged 843 °C/1,000 h, having particles sometimes as large as 0.5μm radius, as shown in 

Figure 5. Microprobe quantitative measurements indicated this phase had an approximate composition of 

20.6Co-33.0Cr-5.4Mo-14.3Ni-5-2.5W-24.2C in atomic percent, Table 4. Surveys indicate other disk 

superalloys also have M23C6 carbides which are principally high in Cr and Mo (Refs. 1, 2, 18, and 19). 

The high Co and Ni content measured here implies these elements can substitute for Cr in LSHR.

Measured composition significantly differed from JMatPro and Pandat equilibrium predictions in Cr 

content. Co can apparently substitute for Cr in M23C6 far more than predicted at equilibrium by JMatPro 

and Pandat. The M23C6 carbides populated this alloy as particles often elongated along the grain 

boundaries. However, larger particles sometimes resided at triple points. This gave a skewed sizefrequency distribution, as shown by the size-frequency histograms in Figure 14. TEM evaluations 

indicated the elongated particles were often made up of multiple grains. These carbides were not found 

away from grain boundaries. While some of these particles were highly elongated, equivalent radius was 

calculated to enable consistent comparisons among all particles and minor phases in this study. 

Mean equivalent radius and area fraction of M23C6 carbides are tabulated versus sample condition in 

Table 6. These carbides were not found in any samples that were not given aging heat treatments. This 

indicates that significant time is required at the aging temperatures for carbide formation to occur, 

exclusively at grain boundaries. Results in Table 6 show consistent trends in growth and dissolution for 

the different times and temperatures investigated. At both 760 and 843 °C, the average equivalent radius 

and area fraction increased with time. This trend was enhanced at 843 °C, where the longest aging time of 

1,000 h produced the highest phase area fraction of 0.0042 and equivalent mean radius near 0.29 μm for 

this phase. Significant growth occurred at this temperature, where number density decreased while 

particle radius increased. Recall from Table 5 that MC carbide content at grain boundaries decreased with 

increasing time at 843 °C. No evidence of MC carbides being partially transformed to M23C6 carbides was

directly observed for these samples, as previously reported for several other nickel-base superalloys 

(Refs. 1 and 2). Due to the difficulties in isolating such particles in SEM and TEM, this transformation 

could still exist. However, based on the much wider spacing of prior MC carbides compared to that of 

M23C6, it appears that nucleation and growth of new particles accounted for a large majority of the M23C6

phase particles observed in aged specimens. No M23C6 carbides were observed within grains in any of the 

aged conditions, however many MC carbides could be found. At 927 °C, M23C6 formed after 10 h aging, 

but not at longer times, indicating it is not stable at this temperature. Thus as the aging time increased and 

conditions approached equilibrium, the size and area fraction were driven to zero. 

Total M23C6 carbide content is compared to that estimated at equilibrium conditions using JMatPro

and Pandat in Table 9 and Figure 17. The estimates agree fairly well with experimental measurements, 

but suggest that M23C6 carbide contents were still approaching equilibrium even for samples aged 1,000 h.

The models predicted equilibrium of M23C6 carbides and elimination of MC carbides at temperatures of 

760 to near 900 °C. However, this transformation appears to be very slow at these temperatures, and there 

was no evidence observed supporting this transformation for MC carbide particles away from grain 

boundaries. 

M23C6 carbide formation temperature was predicted to be 851 and 987.9 °C using JMatPro and 

Pandat; however, no change in the thermograms was detected between 800 to 1000 °C during heating at 

5, 10, and 20 °C/min, due to the lack of detectable M23C6 carbides in the tested as-forged samples. 

NASA/TM—2012-217604 6
3.5 M3B2 Borides

Figure 11 shows the a typical TEM image and diffraction pattern of M3B2 borides. These borides 

have a tetragonal P4/mbm crystal structure, with lattice parameters estimated as a = 5.791 Å, b = 3.11 Å. 

These lattice parameters are comparable to those reported for other disk alloys (Refs. 18, and 19). EDS 

and microprobe measurements on a sample aged 927 °C for 1,000 h indicated this phase had an 

approximate composition of 3.2Co-18.8Cr-15.8Mo-2.4Nb-4.5Ni-0.7Ta-2.4Ti-8.7W-3.8C-39.7B in atomic 

percent, Table 4. Surveys in the literature have also reported M3B2 borides rich in Cr, Mo and W (Refs. 1, 

2, 18, and 19). The measurements indicate W can substitute for Mo significantly more than predicted at 

equilibrium by JMatPro and Pandat. Typical size-frequency histograms acquired for M3B2 borides are 

shown in Figure 15. Mean equivalent radius and area fraction are tabulated versus sample condition in 

Table 7 for M3B2 borides. These borides were most numerous in the as-extruded sample, as fine particles 

less than 0.4 μm radius chiefly along the grain boundaries. The quantity observed along grain boundaries 

was greatly reduced by forging. Most borides dissolved after the supersolvus solution heat treatment, but 

widely scattered borides less than 0.5 μm radius reappeared both within grains and at grain boundaries for 

long aging times at 760 and 843 °C. The highest area fraction for each temperature was found at the 

longest aging time of 1,000 h. The overall highest area fraction and average size of these borides was 

observed for 1,000 h aging at 927 °C, where they grew into the largest minor phase particles found in this 

study, of 1.19 μm radius. 

Overall M3B2 boride content is estimated at equilibrium conditions using JMatPro and Pandat in 

Table 9. This was estimated to be essentially constant from 760 to 927 °C. The estimates were nearest

experimental measurements at 927 °C, where measured content exceeded predictions. However at lower 

temperatures of 760 and 843 °C, M3B2 boride content was far below equilibrium estimates. This may be 

related in part to differences in equilibrium composition predicted versus that measured. 

The average dissolution temperature on initial heating was estimated by DTA to be 1206.7 °C for 

heating at 5 °C/min, and 1202.8 °C for heating at 2 °C/min. The first cycle in the heating thermogram 

shows a slightly higher temperature for the maximum rate of M3B2 liquation than subsequent cycles (Fig. 

7, Table 2). This is believed to be a result of a microstructural change in state associated with cycling. As 

is typical for DTA measurements, heavier samples and samples heated at faster rates show temperatures 

slightly above, within 3.6 °C here, than lighter, more slowly heated samples. Melting leads to an increase 

of 3 to 8 °C in the maximum rate of M3B2 formation observed during cooling. Overall mean M3B2

liquation measured with an average of three single cycle DTA runs was 1201.3 ± 0.4 °C. Mean M3B2

formation temperature during cooling was measured to be 1179.2 °C, which is significantly lower than 

respective JMatPro and Pandat equilibrium predictions of 1243.8 and 1269.0 °C. This also may be related 

in part to differences in equilibrium compositions predicted versus those measured. But boride formation 

was apparently very slow for these lower temperature conditions.

3.6 σ

Figure 12 shows a TEM image and diffraction pattern of σ phase particles. This phase had a 

tetragonal P42/mnm crystal structure, with lattice parameters a = 8.888 Å, c = 4.593 Å. These lattice 

parameter measurements are about 3.5 percent smaller than those made on σ phase particles in N18 

(Ref. 20). Observed σ precipitates were too small for composition measurements using the microprobe. 

Particles were separated by electrochemical phase extraction, and then analyzed by EDS in the SEM. This 

required careful backscatter electron imaging to locate small acicular σ particles that were isolated away 

from more cuboidal MC and spherical M3B2 particles. Their composition was found to be 27.6Co-26.1Cr14.2Mo-11.7Ni-4.1Ta-2.4Ti-14.0W in atomic percent, Table 4. The current measurements differ 

significantly from that measured in prior work for many superalloys having different overall compositions 

than LSHR (Refs. 1, 2, 8, and 9). Yet, roughly similar Co, Ni, and Mo contents were measured in the σ

phase particles analyzed in N18 (Ref. 20), although at a higher Cr level. The measured composition also 

NASA/TM—2012-217604 7
differed from that estimated at equilibrium using JMatPro and Pandat, which predicted higher Cr, lower 

W, and no Ta or Ti.

Typical size-frequency histograms acquired for these σ phase particles are shown in Figure 16. Mean 

equivalent radius and area fraction are tabulated versus sample condition in Table 8. σ phase particles 

were found in a minority of examined cases, attaining a maximum area fraction of 0.00327 at 843 °C for 

1,000 h. They appeared as irregular shapes at the grain boundaries, and seemed more elongated when 

nucleated in the grains. Some had distinct needle-like shapes. Several σ phase particles were also found to 

be attached to small MC carbides. These σ particles sometime appeared to have nucleated on the MC 

carbides, as this could require less energy. Figure 18 depicts two examples of this, observed in thin foils 

using TEM. During aging at 843 and 927 °C, the σ particles appeared to precipitate first at the triple 

points of grain boundaries, and later within the grains. However, they were not observed after 1,000 h at 

927 °C, indicating they were not stable at this temperature. 

Total σ content is estimated at equilibrium conditions using JMatPro and Pandat in Table 9. The 

estimates are significantly greater than experimental measurements. This may be largely related to the 

differences in composition predicted at equilibrium versus that measured. But the test data indicates that σ 

phase content was only slowly increasing and at very low levels, and did not reach equilibrium for these 

heat treatment conditions. Yet, formation of the very small quantities of σ as observed here appears 

possible for extended service times near 843 °C. 

σ phase formation temperature was predicted to be 865 and 1016 °C using JMatPro and Pandat ; 

however, no change in the thermograms was detected between 800 to 1000 °C during heating at 5, 10, 

and 20 °C/min, due to the lack of detectable σ phase in the tested as-forged samples.

3.6.1 Potential Future Work

Additional techniques to further verify the compositions of these minor phases as functions of heat 

treatment are necessary. Significant differences in phase compositions exist between current experimental 

measurements, findings for prior generation superalloys, and equilibrium model predictions. Errors in 

assumed phase compositions could significantly influence many model predictions based on mass 

conservation, thermodynamics, and kinetic properties. Additional aging heat treatments after slower, 

more varied cooling paths other than the water quench treatment employed here would also be useful, to 

understand cooling path dependencies on composition, precipitation, and coarsening of minor phases 

during quench and subsequent aging heat treatments. Similar aging heat treatments on other disk alloys 

could then allow understanding of alloy composition effects on minor phase precipitation. 

4.0 Conclusions

Carbide and boride phase contents and sizes can vary significantly with heat treatment temperature 

and time, and at short times significantly diverge from equilibrium conditions. Commercial software such 

as JMatPro and Pandat can provide estimates of equilibrium contents of MC, M23C6 carbides, M3B2 and σ

phases, representing upper limits for samples aged for a long time to simulate equilibrium conditions. The 

predictions were closest at 927 °C, where equilibrium conditions were most nearly attained. Longer times 

than 1,000 h appear necessary to approach equilibrium at lower temperatures of 843 and 760 °C, where 

predicted M3B2 and σ contents greatly exceeded measurements. Yet, the applied aging times are relevant, 

as they do encompass disk heat treatment times and accumulated service time at high temperatures. The 

predictions of phase formation temperatures reported here appear to be useful for initial estimates, 

although M3B2 predictions can be improved. The errors in predictions may be primarily related to 

divergences in predicted phase compositions from experimental measurements. In addition, the partially 

predicted total replacement of MC with M23C6 carbides at these lower temperatures appears to only be 

occurring at grain boundaries, which would ultimately exert lower limits to MC carbide content and upper 

limits to M23C6 carbide contents at equilibrium. Enhanced prediction software would be necessary to fully 

simulate these factors, and to predict phase contents and sizes at shorter aging times of 1 to 10 h, as 

NASA/TM—2012-217604 8
typically used in aging heat treatments. As maximum disk application temperatures rise, accurate 

simulation of these factors could become important for both processing and service life design. 

References

1. C. T. Sims, N. S. Stoloff, W. C. Hagel, Superalloys II, John Wiley & Sons, Inc., New York, 1987, 

pp. 91-133. 

2. R. C. Reed, The Superalloys, Cambridge University Press, 2006, pp. 35-54. 

3. R. F. Decker, “Strengthening Mechanismss in Nickel-Base Superalloys,” Steel Strengthening 

Mechanisms Symposium, Climax Molybdenum Company, 1969. 

4. S. Floreen, J. M. Davidson, “The Effects of B and Zr on the Creep and Fatigue Crack Growth 

Behavior of a Ni-base Superalloy,” Met Trans. A, V. 14A, 1983, pp. 895-901. 

5. T. J. Garosshen, T. D. Tillman, G. P. McCarthy, “Effects of B, C, and Zr on the Structure and 

Properties of a P/M Nickel Base Superalloy,” Met. Trans. A, V. 18A, 1987, pp. 69-77. 

6. E. S. Huron, K. R. Bain, D. P. Mourer, J. J. Schirra, P. L. Reynolds, E. E. Montero, “The Influence 

of Grain Boundary Elements on Properties and Microstructures of P/M Nickel Base Superalloys,” 

Superalloys 2004, ed. K. A. Green, T. M. Polock, H. Harada, T. E. Howson, R. C. Reed, J. J. 

Schirra, S. Walston, TMS, Warrendale, PA, 2004, pp. 73-81. 

7. E. L. Raymond, R. D. Kissinger, A. J. Paxson, E. S. Huron, “Method for Forming a Nickel-base 

Superalloy Having Improved Resistance to Abnormal Grain Growth,” U.S. Patent 5,584,947, U.S. 

Patent Office, 1996. 

8. R. L. Dreshfield, R. Ashbrook, “Sigma Phase Formation and Its Effect on Stress Rupture Properties 

of IN-100,” NASA TND-5185, Washington, D.C., April 1965. 

9. R. C. Reed, M. P. Jackson, Y. S. Na, “Characterization and Modeling of the Precipitation of the 

Sigma Phase in Udimet 720 and Udimet 720LI,” Met. Trans. A, V. 30A, 1999, pp. 521-533. 

10. H. Murphy, C. T. Sims, “Phacomp Revisited,” Symp. On Structural Stability in Superalloys, ASTMASME, 1968, pp. 47-66. 

11. M. Moringa, N. Yukawa, H. Ezaki, H. Adachi, “New PHACOMP and Its Applications to Alloy 

Design (Phase Computation),” Superalloys 1984, ed. M. Gell, C. S. Kortovich, R. H. Bricknell, W. 

B. Kent, J. F. Radavich, The Metallurgical Society of AIME, 1984, pp. 523-532. 

12. N. Saunders, “Phase Diagram Calculations for Ni-Based Superalloys,” Superalloys 1996, ed. R. D. 

Kissinger, D. J. Deye, D. L. Anton, A. D. Cetel, M. V. Nathal, T. M. Pollock, D. A. Woodford, The 

Minerals, Metals, and Materials Society, Warrendale, PA, 1996, pp. 101-110. 

13. N. Saunders, Z. Guo, X. Li, A. P. Miodownik, J.-P. Schille, “Modelling The Material Properties and 

Behavior of Ni-Based Superalloys,” Superalloys 2004, ed. K.A. Green, T.M. Pollock, H. Harada, 

T.E. Howson, R.C. Reed, J.J. Schirra, and S. Walston. The Minerals, Metals, and Materials Society, 

Warrendale, PA, 2004, pp. 849-858. 

14. W. Cao, S.-L. Chen, F. Zhang, K. Wu, Y. Yang, Y. A. Chang, R. Schmid-Fetzer, W. A. Oates, 

“PANDAT Software with PanEngine, PanOptimizer and PanPrecipitation for Multi-Component 

Phase Diagram Calculation and Materials Property Simulation,” Calphad, V. 33(2), 2009, pp. 328-

342. 

15. J.-C. Zhao, M. F. Henry, “The Thermodynamic Prediction of Phase Stability in Multicomponent 

Superalloys,” Journal of Metals, V. 54 (1), 2002, pp. 37-41. 

16. H.M. Rietveld, “A Profile Refinement Method for Nuclear and Magnetic Structures,” J. Appl. 

Crystallography, V. 2, 1969, pp. 65-71. 

17. 17. R.W. Cheary and A.A. Coleho, “A Fundamental Parameters Approach of X-ray Line-profile 

Fitting,” J. Appl. Crystallography, V. 25, 1992, pp. 109-121. 

18. T. P. Gabb, A. Garg, D. L. Ellis, “Microstructural Evaluations of Baseline HSR/EPM Disk Alloys,” 

NASA/TM-2004-213123, Washington, D.C., 2004. 

NASA/TM—2012-217604 9
19. S. T. Wlodek, M. Kelly, D. A. Alden, “The Structure of Rene′ 88DT,” Superalloys 1996, ed. R. D. 

Kissinger, D. J. Deye, D. L. Anton, A. D. Cetel, M. V. Nathal, T. M. Pollock, D. A. Woodford, 

TMS, Warrendale, PA, 1996, pp. 129-136. 

20. S. T. Wlodek, M. Kelly, D. A. Alden, “The Structure of N18,” Superalloys 1992, ed. S. D. 

Antolovich, R. W. Stusrud, R. A. MacKay, D. L. Anton, T. Khan, D. L. Klarstrom, TMS, 

Warrendale, PA, 1992, pp. 467-476. 

TABLE 1.—CONDITIONS OF LSHR SAMPLES ANALYZED

Sample ID Solution 

temperature, 

°C

Solution 

temperature, 

°C

Solution 

Quench

Aging 

temperature, 

°C

Aging 

temperature, 

°C

Aging 

quench

LAE-As-extruded

LAF-Extruded+Forged

LFF 1199 1 Furnace

LFW 1199 1 Water

LFW760-10 1199 1 Water 760 10 Water

LFW760-100 1199 1 Water 760 100 Water

LFW760-1k 1199 1 Water 760 1000 Water

LFW843-10 1199 1 Water 843 10 Water

LFW843-100 1199 1 Water 843 100 Water

LFW843-1k 1199 1 Water 843 1000 Water

LFW927-10 1199 1 Water 927 10 Water

LFW927-99 1199 1 Water 927 99 Water

LFW927-1k 1199 1 Water 927 1000 Water

TABLE 2.—MEASURED TRANSFORMATION TEMPERATURES USING DIFFERENTIAL THERMAL ANALYSIS FOR 

AS-FORGED LSHR SAMPLES, COMPARED TO THOSE PREDICTED BY JMATPRO-NI 6.0 AND PANDAT 8.1 

Temperature rate,

initial weight

Tstart

°C

Tmax

°C

Cycle M3B2

liquation 

(heating)

M3B2

formation 

(cooling)

Solidus 

(heating)

MC 

liquation 

(heating)

Liquidus 

(heating)

5 °C/min

1586.5 mg 1000 1220

First 1207.6 1175.2 ------- ------- -------

Second 1201.3 1173.1 ------- ------- -------

Third 1202.4 1174.8 ------- ------- -------

683.1 mg 700 1370 Single 1205.8 1183.2 1240 1310.5 1334.5

2 °C/min

1.5012 mg 1000 1220

First 1204.0 1179.6 ------- ------- -------

Second 1200.5 1177.6 ------- ------- -------

Third 1201.0 1177.6 ------- ------- -------

915.6 mg 850 1342 Single 1202.8 1185.6 1239 1310 1330.8

816.2 mg 1000 1475 Single 1202.0 1181.9 1240 1310 1336.1

854.4 mg 1000 1480 Single 1202.1 1183.5 1239 ND 1335.0

JMatPro 1243.8 1234.8 1303.1 1340.0

Pandat 1269.0 1260.5 1316.7 1402.8

NASA/TM—2012-217604 10
TABLE 3.—PHASE LATTICE PARAMETERS ESTIMATED BY TEM SELECTED 

AREA ELECTRON DIFFRACTION PATTERNS (SAEDP) OF THIN FOILS, 

AND BY X-RAY DIFFRACTION OF MINOR PHASE EXTRACTIONS

Phase Structure TEM X-ray

MC FCC, Fm3m a (Å) 4.3 4.3820±0.0003

M23C6 FCC, Fm3m a (Å) 10.6 ------------------

σ Tetragonal, P4/mbm a (Å) 8.9 8.888±0.001

c (Å) 4.6 4.593±0.001

M3B2 Tetragonal, P42/mnm a (Å) 6.0 5.791±0.008

c (Å) 3.15 3.11±0.01

TABLE 4.—MEASURED COMPOSITIONS OF MINOR PHASES IN ATOMIC PERCENT, 

COMPARED TO THOSE PREDICTED BY JMATPRO-NI 6.0 AND PANDAT 8.1

Phase Condition Atomic % Al Co Cr Mo Nb Ni Ta Ti W C B Zr

MC LFW760-1k Mean 12 12 26 50

StanDev 1.4 1.5 1.6

 JMatPro (760 °C) 10 9.1 0.6 0.1 49.5 30.7

Pandat (978°C) 4.7 1.3 8.9 8 22.3 0.5 47.1 7.3

M23C6 LFW843-1k Mean 20.6 33.0 5.4 14.3 2.5 24.2 

StanDev 1.6 3.8 0.6 3.7 0.2 6.1

JMatPro (760 °C) 5.6 59.7 9.5 4.0 0.4 20.7

Pandat (760 °C) 5.6 70.3 1.9 1.1 0.4 20.7

M3B2 LFW927-1k Mean 3.2 18.8 15.8 2.4 4.5 0.7 2.4 8.7 3.8 39.7

StanDev 0.4 1.6 1.7 0.4 0.1 0.1 0.2 1.0 1.3 5.8

 JMatPro (927 °C) 0.4 20.7 37.6 0.2 1.1 40.0 

Pandat (927 °C) 1.7 19.9 36.6 0.7 1.1 40.0

Sigma LFW843C-1k Mean 27.6 26.1 14.2 11.7 4.1 2.4 14.0 

StanDev 1.7 1.6 1.5 0.6 1.4 0.7 1.5

JMatPro (843 °C) 28.8 47.7 10.2 11.5 1.7

Pandat (843 °C) 31.0 42.5 8.7 7.0 10.8

TABLE 5.—EQUIVALENT RADIUS AND AREA FRACTION OF MC CARBIDES WITHIN GRAINS (WG) 

AND AT GRAIN BOUNDARIES (GB), SAMPLE CONDITIONS DESCRIBED IN TABLE 1

Sample MC: 

WG 

average

radius,

μm 

WG 

StDev 

radius,

μm 

Count, 

n 

WG 

average

area 

fraction

WG 

StDev 

area 

fraction

GB 

average 

radius,

μm 

GB 

StDev 

radius,

μm 

Count, 

n 

GB 

average 

area 

fraction

GB 

StDev 

area 

fraction

Sum 

WG+GB 

area 

fraction

LAE-As extruded 0.139 0.057 24 0.00012 0.00009 0.137 0.048 145 0.00068 0.00013 0.00080

LAF-As forged 0.151 0.043 110 0.00060 0.00012 0.164 0.055 60 0.00040 0.00010 0.00100

LFF 0.179 0.065 199 0.00160 0.00021 0.248 0.087 23 0.00035 0.00027 0.00195

LFW 0.174 0.062 193 0.00150 0.00069 0.163 0.067 20 0.00014 0.00008 0.00164

LFW760-10 0.161 0.052 219 0.00139 0.00006 0.177 0.052 29 0.00022 0.00005 0.00161

LFW760-100 0.171 0.059 236 0.00172 0.00040 0.195 0.066 18 0.00017 0.00017 0.00189

LFW760-1k 0.175 0.058 247 0.00187 0.00039 0.223 0.079 24 0.00030 0.00016 0.00217

LFW843-10 0.189 0.061 237 0.00207 0.00018 0.222 0.062 44 0.00052 0.00007 0.00259

LFW843-100 0.195 0.052 211 0.00191 0.00039 0.223 0.078 23 0.00028 0.00006 0.00219

LFW843-1k 0.191 0.069 184 0.00169 0.00041 0.262 0.079 9 0.00015 0.00014 0.00184

LFW927-10 0.192 0.051 164 0.00143 0.00029 0.207 0.080 49 0.00054 0.00071 0.00197

LFW927-99 0.177 0.055 129 0.00099 0.00031 0.268 0.102 15 0.00027 0.00023 0.00126

LFW927-1k 0.192 0.072 265 0.00248 0.00111 0.239 0.094 36 0.00053 0.00030 0.00301

NASA/TM—2012-217604 11
TABLE 6.—EQUIVALENT RADIUS AND AREA FRACTION OF M23C6 CARBIDES WITHIN GRAINS (WG) 

AND AT GRAIN BOUNDARIES (GB), SAMPLE CONDITIONS DESCRIBED IN TABLE 1

Sample M23C6: 

WG 

average 

radius, 

μm 

WG 

StDev 

radius,

μm 

Count,

n 

WG 

average 

area 

fraction

WG 

StDev 

area 

fraction

GB 

average 

radius,

μm 

GB 

StDev 

radius,

μm 

Count, 

n 

GB 

average 

area 

fraction

GB 

StDev 

area 

fraction

Sum 

WG+GB 

area 

fraction

LAE-As extruded 0 0 0

LAF-As forged 0 0 0

LFF 0 0 0

LFW 0 0 0

LFW760-10 0 0.050 ------- 10 0.00001 ---------- 0.00001

LFW760-100 0 0.104 0.049 52 0.00015 0.00013 0.00015

LFW760-1k 0 0.171 0.115 157 0.00148 0.00100 0.00148

LFW843-10 0 0.077 0.038 17 0.00003 0.00002 0.00003

LFW843-100 0 0.144 0.039 193 0.00095 0.00043 0.00095

LFW843-1k 0 0.288 0.160 173 0.00418 0.00181 0.00418

LFW927-10 0 0.146 0.050 246 0.00131 0.00063 0.00131

LFW927-99 0 0 0

LFW927-1k 0 0 0

TABLE 7.—EQUIVALENT RADIUS AND AREA FRACTION OF M3B2 BORIDES WITHIN GRAINS (WG) 

AND AT GRAIN BOUNDARIES (GB), SAMPLE CONDITIONS DESCRIBED IN TABLE 1

Sample M3B2: 

WG 

average 

radius,

μm

WG 

StDev 

radius, 

μm 

Count,

n 

WG 

average 

area 

fraction

WG 

StDev 

area 

fraction

GB 

average 

radius,

μm 

GB 

StDev 

radius,

μm 

Count,

n 

GB 

average 

area 

fraction

GB 

StDev 

area 

fraction

Sum 

WG+GB 

area 

fraction

LAE-As extruded 0.200 0.084 17 0.00018 0.00015 0.262 0.113 176 0.00333 0.00029 0.00351

LAF-As forged 0.276 0.100 18 0.00034 0.00031 0.359 0.158 7 0.00023 0.00033 0.00057

LFF 0 0 0.00000

LFW 0.178 0.109 2 0.00002 0.00003 0 0.00002

LFW760-10 0 0 0.00000

LFW760-100 0.364 0.122 10 0.00033 0.00026 0 0.00033

LFW760-1k 0.103 0.040 11 0.00003 0.00003 0.215 0.226 24 0.00051 0.00075 0.00054

LFW843-10 0 0 0.00000

LFW843-100 0 0 0.00000

LFW843-1k 0.402 0.118 10 0.00039 0.00036 0.285 0.049 2 0.00004 0.00008 0.00043

LFW927-10 0 0 0.00000

LFW927-99 0 0 0.00000

LFW927-1k 0.140 1 0.00001 0.00001 1.186 0.618 21 0.00827 0.0055 0.00828

NASA/TM—2012-217604 12
TABLE 8.—EQUIVALENT RADIUS AND AREA FRACTION OF σ PHASE WITHIN GRAINS (WG) 

AND AT GRAIN BOUNDARIES (GB), SAMPLE CONDITIONS DESCRIBED IN TABLE 1

Sample σ: 

WG 

average 

radius,

μm

WG 

StDev 

radius, 

μm 

Count,

n 

WG 

average 

area 

fraction

WG 

StDev 

area 

fraction

GB 

average 

radius, 

μm 

GB 

StDev 

radius, 

μm 

Count,

n 

GB 

average 

area 

fraction

GB 

StDev 

area 

fraction

Sum 

WG+GB 

area 

fraction

LAE-As extruded 0 0 0

LAF-As forged 0 0 0

LFF 0 0 0

LFW 0 0 0

LFW760-10 0 0 0

LFW760-100 0 0 0 

LFW760-1k 0 0 0

LFW843-10 0 0.077 0.014 4 0.00001 0.00001 0.00001

LFW843-100 0.163 0.081 17 0.00012 0.00009 0.216 0.166 28 0.00045 0.00033 0.00057

LFW843-1k 0.237 0.088 129 0.00184 0.00039 0.333 0.122 51 0.00143 0.00093 0.00327

LFW927-10 0.161 0.083 8 0.00006 0.00006 0.175 0.062 69 0.00053 0.00031 0.00059

LFW927-99 0 0.191 0.095 24 0.00024 0.00045 0.00024

LFW927-1k 0 0 0

TABLE 9.—OVERALL AREA FRACTIONS OF EACH MINOR PHASE MEASURED IN SAMPLES AGED 1,000 h, 

COMPARED TO THOSE PREDICTED AT EQUILIBRIUM USING JMATPRO-NI 6.0 AND PANDAT 8.1 

Temperature,

°C 

Sample

Software

Total MC 

area fraction

Total M23C6

area fraction

Total M3B2

area fraction

Total σ

area fraction

760 LFW760-1k 0.00217 0.00148 0.00054 0.00000

760 JMatPro 0.00053 0.00546 0.00376 0.05249

760 Pandat 0.00000 0.00648 0.00376 0.08748

843 LFW843-1k 0.00184 0.00418 0.00042 0.00327

843 JMatPro 0.00206 0.00178 0.00376 0.01208

843 Pandat 0.00000 0.00615 0.00376 0.06263

927 LFW927-1k 0.00300 0.00000 0.00827 0.00000

927 JMatPro 0.00281 0.00000 0.00375 0.00000

927 Pandat 0.00000 0.00557 0.00375 0.03482

NASA/TM—2012-217604 13
Figure 1.Comparison of microstructures: a. LAE-as extruded, b. LAF-extruded and isothermally fo

b. a. 

M3B2 

MC 

γ′

M3B2 

MC 

NASA/TM—2012-217604 14
Figure 2.Comparison of microstructures subsequently heat treated: a. LFF, b. LFW.

b. a. 

MC 

MC 

NASA/TM—2012-217604 15
Figure 3.Images of 1199 °C /1 h/WQ samples additionally aged 1,000 h: a. LFW760-1k, b. LFW843-1k, and c

a. b.

M23C6 

M3B2 

MC 

M

M3B2 

MC 

σ

c.

M3B2 

MC 

NASA/TM—2012-217604 16
Figure 4.Images of 1199 °C/1 h/WQ samples additionally aged at 760 °C: a. LFW760-10, b. LFW760-100, and c. LF

a. b.

MC 

M23C6 

M3B2 

MC 

c.

•M23C

6 

M3B2 

MC 

NASA/TM—2012-217604 17
Figure 5.Images of 1199 °C/1 h/WQ samples additionally aged at 843 °C: a. LFW843-10 , b. LFW843-100, and c. 

a. b.

σ 

M23C6 

MC 

MC 

M23C6 

σ

c.

M23C6 

M3B2 

MC 

σ

NASA/TM—2012-217604 18
Figure 6.Images of 1199 °C/1 h/WQ samples additionally aged at 927 °C: a. LFW927-10, b. LFW927-99, c. LFW

a. b.

σ 

MC 

M23C6 

MC 

c.

M3B2 

MC 

NASA/TM—2012-217604 19
Figure 7.Thermograms acquired from differential thermal analysis of as-forged LSHR samples that were a) triple cycled 

at 5 °C/min, b) triple cycled at 2 °C/min and c) single cycled through melting at 5 and 2 °C/min. 

NASA/TM—2012-217604 20
a. 1250 °C 1300 °C 1325 °C 

b. 1250 °C 1300 °C 1325 °C 

Figure 8.Heated treated as-forged LSHR samples in air for 1 h at 1250 °C, 1300 °C, 1325 °C 

and water quenched showing: a. the degree of resolidification and b. MC carbides.

NASA/TM—2012-217604 21
Figure 9.MC carbide phase, FCC, Fm3m, a~4.3 Å: a. TEM image and b. SAEDP of LFW760-1k.

a.

MC

<110> FCC

b. 

NASA/TM—2012-217604 22
Figure 10.M23C6 carbide, FCC, Fm3m, a ~ 10.6 Å: a. TEM image and b. SAEDP of LFW760-1k.

[-122]FCC

M23C6 M23C6 

a.

b. 

NASA/TM—2012-217604 23
Figure 11.M3B2 boride phase, tetragonal, P4/mbm, a ~ 6.0 Å, c ~ 3.15 Å: a. TEM image and b. SAEDP of LFW

00

1 

1-

10

1-11

X 

M3B2 

MC

a.

11-1 

-120

03-

1 

X 

b.

NASA/TM—2012-217604 24
Figure 12.σ phase, tetragonal, P42/mnm, a ~ 8.9 Å, c ~ 4.6 Å : a. TEM image and b. SAEDP of LFW843

σ

MC 

[211]σ [100]σ 

a.

b. 

NASA/TM—2012-217604 25
Figure 13.Typical histograms of equivalent radius vs. frequency for MC carbides within grains (WG) and at grain b

a. LFW, b. LFW760-1k, c. LFW843-1k, d. LFW927-1k. 

a. b. c. d. 

0

0.5

1

1.5

2

2.5

3

3.5

0

0.05

0.1

0.15

0.2

0.25

0.3

Frequency

Equiv.Radius-microns

MC-GB

5

0

35

30

25

20

15

10

0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

Frequency

Equiv.Radius-microns

MC-WG

0

60

50

40

30

20

10

0.1

0.175

Frequency

Equiv.

MC -

5

0

50

45

40

35

30

25

20

15

10

0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

Frequency

Equiv.Radius-microns

MC -WG

9

8

7

6

5

4

3

2

1

0

0.1

0.15

0

2

Frequency

Equiv.

MC 

7

6

5

4

3

2

1

0

0.0750.1250.1750.2250.2750.3250.375

Frequency

Equiv.Radius-microns

MC -GB

0

70

60

50

40

30

20

10

0.050.150.250.350.45

Frequency

Equiv.Radius-microns

MC - WG

0

0.5

1

1.5

2

2.5

3

3.5

0.050.150.250.35

More

Frequency

Equiv. Radius-microns

MC-GB

NASA/TM—2012-217604 26
Figure 14.Typical histograms of equivalent radius vs. frequency for M23C6 carbides, only 

observed at grain boundaries (GB): a. LFW760-1k, b. LFW843-1k. 

a. b. 

5

0

40

35

30

25

20

15

10

0.050.150.250.350.450.550.650.75

Frequency

Equiv.Radius-microns

M23C6 GB 

5

0

30

25

20

15

10

0.0250.0750.1250.1750.2250.2750.325

More

Frequency

Equiv.Radius-microns

M23C6 -GB

NASA/TM—2012-217604 27
Figure 15.Typical histograms of equivalent radius vs. frequency for M3B2 borides within grains (W

and at grain boundaries (GB): a. LFW760-1k, b. LFW843-1k, c. LFW927-1k. 

a. b. c. 

0

0.5

1

1.5

2

2.5

3

3.5

4

4.5

0.30.71.11.51.9

More

Frequency

Equiv.Radius-microns

M3B2 -GB

8

6

4

2

0

18

16

14

12

10

0.050.150.250.350.450.550.65

Frequency

Equiv.Radius-microns

M3B2-GB

0

0.5

1

1.5

2

2.5

0.10.20.30.40.50.60.7

Frequency

Equiv.Radius-microns

M3B2-WG

NASA/TM—2012-217604 28
Figure 16.Histograms of equivalent radius vs. frequency for σ phase in LFW843-1k.

NASA/TM—2012-217604 29
Figure 17.Overall area fractions of each minor phase measured in samples aged 1,000 h, 

compared to those predicted at equilibrium using JMatPro-Ni 6.0 and Pandat 8.1. 

NASA/TM—2012-217604 30
Figure 18.TEM images from LFW843-1k of σ TCP phase attached to MC carbide, 

suggesting preferential σ phase nucleation there. 

MC

σ

σ

σ

σ

NASA/TM—2012-217604 31
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1. REPORT DATE (DD-MM-YYYY)

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4. TITLE AND SUBTITLE

Formation of Minor Phases in a Nickel-Based Disk Superalloy 

5a. CONTRACT NUMBER

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5c. PROGRAM ELEMENT NUMBER

6. AUTHOR(S)

Gabb, T., P.; Garg, A.; Miller, D., R.; Sudbrack, C., K.; Hull, D., R.; Johnson, D.; Rogers, R., 

B.; Gayda, J.; Semiatin, S., L. 

5d. PROJECT NUMBER

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WBS 561581.02.08.03.15.14 

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National Aeronautics and Space Administration 

John H. Glenn Research Center at Lewis Field 

Cleveland, Ohio 44135-3191 

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13. SUPPLEMENTARY NOTES

14. ABSTRACT

The minor phases of powder metallurgy disk superalloy LSHR were studied. Samples were consistently heat treated at three different 

temperatures for long times to approximate equilibrium. Additional heat treatments were also performed for shorter times, to then assess 

non-equilibrium conditions. Minor phases including MC carbides, M23C6 carbides, M3B2 borides, and sigma were identified. Their 

transformation temperatures, lattice parameters, compositions, average sizes and total area fractions were determined, and compared to 

estimates of an existing phase prediction software package. Parameters measured at equilibrium sometimes agreed reasonably well with 

software model estimates, with potential for further improvements. Results for shorter times representing non-equilibrium indicated 

significant potential for further extension of the software to such conditions, which are more commonly observed during heat treatments and 

service at high temperatures for disk applications. 

15. SUBJECT TERMS

Gas turbine engines; Rotating disks 

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