T.P. Gabb, C.K. Sudbrack, S.L. Draper, R.A. MacKay, and J. Telesman

Glenn Research Center, Cleveland, Ohio

Effects of Long Term Exposures

on PM Disk Superalloys

NASA/TM—2013-216614

December 2013
http://www.sti.nasa.gov

help@sti.nasa.gov
T.P. Gabb, C.K. Sudbrack, S.L. Draper, R.A. MacKay, and J. Telesman

Glenn Research Center, Cleveland, Ohio

Effects of Long Term Exposures

on PM Disk Superalloys

NASA/TM—2013-216614

December 2013
Acknowledgments

Level of Review
Effects of Long Term Exposures on PM Disk Superalloys

T.P. Gabb, C.K. Sudbrack, S.L. Draper, R.A. MacKay, and J. Telesman

National Aeronautics and Space Administration

Glenn Research Center

Cleveland, Ohio 44135

Abstract

Turbine disks in some advanced engine applications may be exposed to temperatures above 700 C

for extended periods of time, approaching 1,000 h. These exposures could affect the near-surface

composition and microstructure through formation of damaged and often embrittled layers. The creation

of such damaged layers could significantly affect local mechanical properties. Powder metal disk

superalloys LSHR and ME3 were exposed at temperatures of 704, 760, and 815 C for times up to

2,020 h, and the types and depths of environmental attacked were measured. Fatigue tests were performed

for selected cases at 704 and 760 C, to determine the impact of these exposures on fatigue life. Fatigue

resistance was reduced up to 98% in both superalloys for some exposure conditions. Tensile tests were

also performed to help understand fatigue responses, and showed corresponding reductions in ductility.

The changes in surface composition and phases, depths of these changed layers, failure responses, and

failure initiation modes were compared.

Introduction

To screen the fatigue resistance of nickel-based disk superalloys, conventional low cycle fatigue tests

are often performed with cycles having periods of 0.1 to 120 s, in the interests of combining cyclic lives

of up to 10,000 cycles with affordable test durations of near 100 h. However, environmental exposures at

high temperatures of 650 to 815 C in actual service can approach 1000 h in aerospace gas turbine

engines, and 10,000 h in land-based gas turbine engines. The disk temperatures in both classes of engines

are anticipated to rise, as higher compressor discharge air temperature is known to improve fuel efficiency

and performance (Ref. 1). Turbine disk rims are cooled by compressor discharge air, and their

temperature tracks closely with compressor discharge air temperature.

Exposures at a temperature of 870 °C have been shown to produce significant reductions in fatigue

resistance of blade superalloy Rene 80 (Ref. 2). Oxide layers were identified as the cause, along with an

 ́ precipitates had been dissolved. Fatigue lives could be largely recovered by

removing the oxidized layers. Exposures at disk application temperatures of 650 to 704 °C have also been

shown to reduce fatigue lives of powder metal disk superalloys Udimet 720 and ME3 (Ref. 3). It was

shown that the exposures could cause a change in the locations of failure initiation sites from internal

flaws to surface oxidation damage, which in turn reduced the fatigue lives. However, this response was

intermittent depending on exposure conditions, and often confounded the effects of exposures with that of

crack location on fatigue life. It has often been observed that disk fatigue specimens failing from cracks

initiating at defects on the surface have lower corresponding fatigue lives than those failing from internal

cracks initiating at defects far from the surface (Ref. 4). This has been associated with differences in

stress intensity due to longer surface crack initiation sites, and environmental interactions for accelerated

crack initiation and growth at surface versus internal cracks (Ref. 5). More recent evaluations of exposure

effects on disk superalloys have been performed on notched fatigue specimens, where the stress

concentration of the notch encourages cracking at the notch root. Here, the superalloy RR1000 (Ref. 6)

 ́ precipitates

had been dissolved, while ME3 (Ref. 7) instead had fingers of Al2O3 extending into a recrystallized zone.

Both alloys had significant reductions in fatigue life after exposures at 700 to 815 C. However, it is not

clear if these exposure effects were enhanced by the concentrated stresses at the exposed notch surfaces.
The objective of this study was to screen the effects of high temperature exposures in air on the

microstructure and fatigue properties of two powder metal disk superalloys. Fully machined mechanical

test specimens of LSHR and ME3 were exposed at 704, 760, and 815 C, for times up to 2,020 h.

Uniform gage and notch gage fatigue tests were performed at 704 or 760 C, allowing the effect of

concentrated notch stresses to be ascertained. Tensile tests were subsequently performed to further

investigate cracking of the identified environment-affected surface layers. The effects of exposures on

fatigue life, tensile strength, and tensile elongation were compared. The associated failure modes and their

relationships to exposure-induced changes in compositions and phases near the surface were determined.

Materials and Test Procedures

Materials

The compositions in weight percent of the tested materials are listed in Table 1. The compositions of

LSHR and ME3 are similar, with the biggest differences in Mo, Nb, Ta, and W contents. LSHR contains

less Mo and Ta, but more W and Nb than ME3. LSHR (Ref. 8) superalloy powder was obtained from

Special Metals Corp. This powder was atomized in argon, screened to –270 mesh, sealed in a stainless

steel container, and then consolidated by hot isostatic pressing. The consolidated material was

subsequently extruded and isothermally forged into several flat disks. Stacked rows of rectangular blanks

each about 13 mm square and 66 mm long were extracted throughout the forged disks, with their lengths

oriented parallel to the disk plane. The blanks were placed vertically within an enclosing fixture and

supersolvus solution heat treated at 1171 C for 2 h in a resistance heating furnace. The assembly of

blanks and enclosing fixture was then removed to cool in static air. This gave a near-linear cooling rate of

the blank cores, averaging 72 C per minute cooling rate down to 870 C. The blanks were subsequently

given an aging heat treatment of 855 C for 4 h followed by 775 C for 8 h. Blanks of similar dimensions

were also extracted from the web and rim of several fully heat treated disks of ME3 (Ref. 7), which had

been subjected to similar thermo-mechanical processing conditions. The ME3 blanks were selected from

the disks at locations having comparable cooling rates after the solution heat treatment to that of the

LSHR blanks. But unlike LSHR disks, the ME3 disks were fully heat treated before blank extraction, so

ME3 blanks were directly machined into notch fatigue and tensile specimens.

Fatigue and tensile specimens were machined using low stress grinding procedures, with the gage

sections finally polished parallel to the loadi

roughness. LSHR uniform gage fatigue specimens had a gage diameter of 4.8 mm across a gage length of

13 mm. ME3 cylindrical notched specimens (Fig. 1) had minimum diameter of 5.1 mm, with a geometric

elastic stress concentration factor Kt = 2. Several LSHR specimens were also prepared with the

cylindrical notch configuration of Figure 1, for consistent fatigue test comparisons with ME3.

Accompanying LSHR and ME3 tensile specimens having a nominal gage diameter of 4.1 mm across a

gage length of 21 mm were also machined.

Test Procedures

Most exposures were conducted in conventional resistance heating box furnaces in lab air, with all

specimens and blanks air cooled after exposures. Fully machined specimens were exposed after being

wiped with cotton soaked in acetone, and then ethanol. Several additional exposure conditions were

investigated, including: 1) oversized specimen blanks were sometimes exposed in air for selected long

exposures at 704, 760, and 815 C before machining into specimens, in order to assess the effects of the

exposures on superalloy microstructure and properties remote from the surface; 2) fully machined ME3

notched gage fatigue specimens were wrapped with Ta foil to getter remnant oxygen and then were

exposed in a vacuum pressure not exceeding 9×10-7 torr; 3) cyclic exposures in air were conducted on

uniform-gage LSHR fatigue specimens to assess cyclic effects. All cyclic specimens were heated to

815 C and cooled to approximately 32 C. Here, the machined specimens were suspended from a
horizontal alumina tube while a standard resistance-heating horizontal tube furnace automatically

translated over the specimens during the 60 min heating cycle. The furnace automatically translated away

from the specimens during the 20 min cooling cycle, to cool in static air. A typical heating and cooling

cycle is shown in Figure 2, showing outputs from attached platinum–rhodium “Type R” thermocouples.

These specimens were exposed to 440 cycles.

LSHR and ME3 were tested in different, yet complementary manners to screen the effects of the

exposures on fatigue life for two different fatigue conditions. Fatigue tests of LSHR specimens with

uniform gages were performed at 760 C, and tests of ME3 specimens with notched gages were

performed at 704 °C. Temperature was measured using chromel-alumel “Type K” thermocouples

contacting the specimens. Fatigue tests were performed in accordance with fatigue test specification

ASTM E466-07. All fatigue tests used a sinusoidal waveform cycling stress at a constant frequency of

0.33 Hz. LSHR specimens were tested using a uniaxial electro-mechanical testing machine having a

resistance heating furnace. A maximum stress of 841 MPa and minimum stress of –428 MPa was applied

in all tests. Preliminary testing of LSHR showed these cyclic stresses are typical stabilized values

generated in strain-controlled tests at a total strain range of 0.76 percent and minimum/maximum strain

ratio of 0, which typically resulted in surface-initiated fatigue failures and limited the fatigue life for

specimens without long prior exposures in air (Ref. 8). However, no extensometer for strain

measurements was attached to the specimens in the present fatigue tests, in order to avoid contacting the

exposed surfaces.

Fatigue tests of notched gage ME3 specimens were performed using uniaxial servo-hydraulic testing

machines with resistance heating furnaces. A maximum stress of 855 MPa and minimum stress of

43 MPa were applied in all tests, which were performed at 704 C. Previous testing of ME3 using this

notched gage specimen with these test conditions produced mean fatigue lives of 24,000-78,000 cycles

and encouraged transgranular surface-initiated fatigue cracks to limit fatigue. (Ref. 9). For all fatigue

tests, one-way statistical analyses of variance in fatigue lives were performed using JMP Version 10 (SAS

Institute Inc., Cary, NC, 1989-2013) software, to test for significant differences in mean life response.

Stepwise multiple linear regressions were also performed using this software, with a 90% probability of

significance required for inclusion of any variable. This software was used to “code” all variables (V), by

normalizing them using Equation (1):

V mid)/(Vrange/2) (1)

This normalized each variable V to values of –1 to 1, so comparisons could be made of the relative

influence of each significant variable, by directly comparing the magnitudes of their coefficients in the

regression equation (Ref. 10).

Tensile tests were performed in general accordance with ASTM E21-05 and ASTM G129-00. Tests

on LSHR were performed at 704 and 760 C in air on a servo-hydraulic testing machine, using a

conventional resistance heating furnace. Tensile tests of LSHR performed at 760 C used a displacement

rate of 0.1041 mm/s, which produced an average strain rate 5.0×10-3 s-1, to simulate conditions

experienced in the uniform gage fatigue tests. Companion tensile tests of LSHR were also performed at

704 C with a much slower displacement rate of 0.0017 mm/s and average strain rate of 8.3×10-5 s-1 per

ASTM G129-00, to screen strain rate dependence and temperature dependence of exposure effects on

tensile response. Tests on ME3 were performed at 704 C in a vacuum chamber held at a vacuum

pressure not exceeding 8×10-6 torr, integrated in an electro-mechanical universal testing machine.

Resistance-heating elements were used to heat the specimen. These tensile tests in vacuum were

performed at displacement rates of 0.0017 mm/s and 0.017 mm/s, which produced average strain rates of

8.3 10-5 s-1and 8.3 10-4 s-1, respectively. However, no extensometer for strain measurements was attached

to the specimens in the tensile tests, to avoid contacting the exposed surfaces.

ASTM grain sizes were determined from etched metallographic sections using linear intercepts of

circular grid overlaps on optical images in accordance with ASTM E112-10 linear intercept procedures

using circular grid overlays on optical images. Precipitate microstructures were characterized using field
emission scanning electron microscopy on metallographically prepared and etched sections.

Metallographic sections were swab etched with 25% acetic acid, 25% nitric acid, 25% HCl, 25% H2O,

1% HF by volume. While it is recognized that imaging of the finest tertiary  ́ precipitates is

compromised using this approach, it was still considered sufficient for this study to make relative

comparisons among exposure conditions. The area of each precipitate was measured using SigmaScan

Pro (Systat Software Inc., San Jose, CA) image analysis software by thresholding based on image

brightness. The equivalent radius of a circular precipitate was also calculated from each precipitate’s area.

Fracture surfaces were examined to determine failure initiation sites using scanning electron microscopy.

Results and Discussion

Microstructures

Internal Microstructures

Internal microstructures for unexposed test materials are shown in Figure 3, including magnified

views of the secondary and tertiary ’-precipitates. Mean linear intercept ASTM grain size and the

measured dimensions of the secondary and tertiary  ́ precipitate are listed in Table 2. LSHR had a finer

average grain size of 15 μm, compared to ME3 at 29 μm. Secondary  ́ precipitates in LSHR were a

mixture of rounded cubes and moderately extended cubes, having lobes growing at the corners. The area

of each precipitate was measured, and the equivalent radius of a spherical particle is also compared in

Table 2. ME3 predominantly had larger precipitates on either size basis, with much more extended

growth of lobes at the cube corners. This was due to the different and slower initial cooling path after

solution heat treatment for the ME3 disk material, heat treated as large disks (Ref. 11). Tertiary  ́

precipitates were consistently near spherical, and slightly larger for ME3 than for LSHR. MC carbides

were often observed along with fewer M3B2 borides within grains of both superalloys, while MC and

M23C6 carbides and M3B2 borides were observed along the grain boundaries of LSHR (Ref. 8 to 12).

Internal microstructures are also shown after an exposure of 815 °C at 2,020 h in Figure 3e and 3f.

Secondary and tertiary  ́ precipitate dimensions are also listed for this material condition in Table 2.

Secondary  ́ precipitate size increased after aging LSHR, while precipitate size actually decreased for

ME3. This was apparently related to the stability of the highly extended, lobed precipitates in ME3. Work

on other disk superalloys has shown these particle shapes can be unstable in certain exposure conditions,

and that they evolved into smaller, more equiaxed shapes (Ref. 13). Tertiary  ́ precipitates were no

longer observed in both alloys, and were apparently dissolved during coarsening of secondary  ́

precipitates. This was observed both for specimens and blanks subjected to this exposure.

Surface Microstructures

To assess influence of aging exposures on mechanical properties, surface microstructure were

examined in detail. With increasing exposure time and temperature, oxides grew on the surfaces of LSHR

and ME3 in a very similar manner, Figure 4 and Figure 5. Cross section evaluations of accompanying

exposed coupons indicated that complex changes occurred near the surface during exposures for both

ME3 and LSHR. Continuous Cr2O3 was observed on the surface, along with faceted grains of TiO2. The

TiO2 grains became noticeably larger and more abundant with exposure time. Internal branched Al2O3

fingers were evident beneath the outer Cr2O3-TiO2 scale, even for the shortest exposures. These Al2O3

fingers extended into a zone of the alloy that was recrystallized, and where precipitates were also

dissolved. The depth of ution extended slightly further than the Al2O3 fingers.

Microprobe evaluations revealed that the exposures led to chemical and phase alterations near the

surface that extended beyond the oxide layers. This is illustrated in Figure 6, for an ME3 specimen

exposed at 815 °C at 2,020 h. -dissolution layer was depleted in Cr, Ti, Al, and Ta in

comparison to the deep interior, due to the formation of the major oxides. Within 30 μm of the oxidized
surface, the grain boundaries are depleted in Cr, Mo, Co and enriched in Ni relative to remote interior

grain boundaries. Due to their high Cr, Mo, and Co content, M23C6 carbides image brightly in backscattered SEM (BSE) due to their high refractory content and primarily reside at the grain boundaries.

More infrequent M3B2 borides, enriched in Mo, Cr, and W, and MC carbides, enriched in Ti, Ta, and Nb,

reside mostly within the grains (Ref. 7 and12). The BSE image in Figure 6 shows a region extending in

further than the precipitate-free zone is depleted in the bright M23C6 carbides at the grain boundaries. For

this exposure condition, at roughly 50 μm depth there is a clear transition from negligible minor phase

particle number density at the grain boundaries to full density. Near the surface, Cr from M23C6 carbides

and M3B2 borides apparently diffused out to help sustain the Cr2O3 scale growth, resulting in the local

dissolution of these minor phases. Mo and Co also accumulated underneath the external scale in the -

dissolution layer. This resulted in adjacent superalloy grain boundaries that were depleted in Cr, Mo, and

Co. The MC carbides appear to be stable during exposure and were present within the minor phase

dissolution zone.

Tables 3 and 4 compare the thicknesses of the outer oxide scale, alumina finger penetration, -

dissolution layer, and the minor phase dissolution zone, for various exposure conditions of each alloy.

-dissolution layer for both alloys

as functions of exposure time are shown in Figure 7. For ME3 (Ref. 7), invariant of time and temperature,

the minor phase dissolution zone thickness was three times greater than that of the -dissolution layer

(where Al2O3 fingers reside), which were roughly three times greater than the external oxide scale

thicknesses. As these layer thicknesses are additive, for the most aggressive 815 °C at 2,020 h exposure

imaged in Figure 6, removal of material to a depth of ~20 μm was necessary to eliminate both the

external Cr2O3-TiO2 scale and internal branched Al2O3 fingers, while material removal to a depth of

~50 was necessary to eliminate the minor phase dissolution zone. Measurements for the mean oxide

scale thickness and alumina finger length did not vary statistically between the two alloys for the majority

of the exposure conditions, with the exception of 704 °C 100 h and 704 °C 440 h conditions. The mean

’-dissolution layer thicknesses for LSHR showed agreement with ME3 for the 815 °C isotherm,

however, showed slight differences for the 704 and 760 °C exposure conditions, possibly an indication of

coupon-to-coupon variability.

Effect of Exposures on Fatigue Life

Uniform-Gage Fatigue of LSHR

Fatigue tests of uniform-gage specimens of LSHR were performed at 760 °C, with a maximum stress

of 841 MPa and minimum stress of –428 MPa applied in all tests. Preliminary testing of LSHR had

shown these cyclic stresses are typically stabilized values generated in strain-controlled tests at a total

strain range of 0.76 percent and minimum/maximum strain ratio of 0. These preliminary tests were

performed at a frequency of 0.5 Hz for 24 h, then continued at a higher frequency of 10 Hz. The short

total test times and associated exposures encouraged internal- initiated failures at grain facets or nonmetallic inclusions in tests at 704 and 760 °C, but surface initiated failures at 815 °C (Ref. 14). Such

response has also been observed in ME3 (Ref. 15). The subsequent tests performed here on LSHR at

760 °C with the same applied stresses but a slower, constant frequency of 0.33 Hz, however, encouraged

surface-initiated cracks to limit fatigue life, even for specimens without prior exposures. This testing

temperature of 760 °C and frequency of 0.33 Hz allowed only minor surface oxidation during fatigue test

durations of up to about 50 h, but this likely combined with cyclic relaxation of compressive residual

stresses near the surface (Ref. 16) and was sufficient to typically result in surface-initiated failures.

The lives for fatigue tests of LSHR are compared in Table 5 and Figure 8. Stepwise linear regressions

were performed on log(life), using variables of temperature, log(time), and their product. The resulting

linear regression equation is given in Figure 9a, along with comparisons of estimated and actual lives.

This equation indicated that increasing exposure temperature and increasing time reduced fatigue life.

Their coefficients indicated they had comparable effects, when increased from minimum to maximum
settings employed here. The product of temperature and time also had a significant effect, indication their

effects were enhanced for combinations of high temperature and long time. This regression equation gave

a fairly high coefficient of determination (R2) of about 0.87, but had a lack of fit that was significant at a

probability of over 95%. As shown in Figure 9a, this lack of fit was evident at intermediate exposure

conditions, where only small reductions in mean fatigue life were observed. Regressions of fatigue live

versus depths of oxide scales, fingers, and dissolution zones each also had this problem of lack of fit, with

poor agreement at intermediate exposure conditions. The coefficient of determination was not improved

using any of these depths, or when using their sum. This is exemplified in Figure 9b, using the sum of

oxide scales, fingers, and dissolution zones depths for a total depth of damage. Fatigue life decreased with

increasing total depth of damage.

Fatigue lives for each condition were compared by one-way analyses of variance, and then grouped

where no significant difference in life was observed. Exposures at 704 °C for 100 h gave no significant

reduction in mean life for these limited tests, and could be grouped with unexposed specimens (Group N).

However, longer exposures at 704 °C and exposures at higher temperatures significantly reduced mean

fatigue life, at a statistical confidence of at least 95%. Specifically, exposures at 704 °C for 440 h and

2,020 h, 760 °C for 440 h, and 815 °C for 100 h, had comparable mean fatigue lives representing

relatively low effects on life (Group L), about 50% of unexposed life. The mean lives for these different

exposure conditions were statistically equivalent at a confidence of over 90% and significantly less than

Group N fatigue life. But isothermal or cyclic exposures at 815 °C for 440 h produced significantly lower

fatigue lives than Groups N or L. These conditions gave equivalent mean fatigue lives at a confidence of

over 95% (Group M), and more severely reduced fatigue mean life to 20% of unexposed life. The highest

reduction in mean fatigue life was observed for an exposure of 815 °C for 2,020 h (Group H), which

reduced mean fatigue life by 98% or about 2% of the unexposed mean fatigue life.

Material blanks were also aged at 815 °C for 2,020 h before machining into specimens, to assess

microstructure aging effects. These LSHR specimens had about 25% lower mean fatigue lives than

unexposed specimens, at a statistical significance of over 95%. The data is plotted in Figure 8 for

comparison. This reduction is likely due to reductions in strength, often observed after overaging (Fig. 3)

of the strengthening precipitates (Ref. 17). However, the fatigue lives of samples extracted from these

aged blanks were significantly higher than for machined specimens that were tested after an exposure for

the same duration and temperature.

Several specimens were subjected to cyclic exposures giving an accumulated time of 440 h at 815 °C,

for comparison to specimens given static exposures of 440 h at 815 °C. The surfaces of fatigue specimens

typically looked very similar after these two exposure conditions, although the outer oxide layers

appeared to spall in several locations during cyclic exposure, Figure 10a and b. Energy-dispersive X-ray

spectroscopy (EDS) indicated that the Cr-rich oxide layer spalled off in some locations, leaving the

underlying Al2O3-rich layer uncovered. Nevertheless, corresponding mean fatigue lives of the isothermal

and cyclic exposures were equivalent at a confidence of 95% (Group M, Fig. 8).

Notch Fatigue of ME3

Fatigue tests of ME3 were performed on notch specimens at 704 °C. The stress concentration of the

notch forced the failure initiations to remain at the surface, and did not allow internal crack initiation sites

such as inclusions or large grains to limit life. The testing temperature was not expected to cause rapid

surface changes during the fatigue test durations of up to about 50 h, based on the examinations of

coupons exposed at 704 °C. This fatigue test condition could represent a disk feature such as a bolt hole

or fillet for a service application near highest current disk temperatures.

The ME3 fatigue lives are compared in Table 6 and Figure 11. For unexposed and exposed

conditions, fatigue life varied more for ME3 notched fatigue tests than for LSHR uniform gage fatigue

specimens. This gave increased root mean square error in regression equations, but still allowed

inspection of trends. Stepwise linear regression was again performed on log (life), using variables of

temperature, log(time), and their product. The resulting linear regression equation is given in Figure 12a,
along with comparisons of lives estimated by the equation and actual lives. This equation again indicated

that increasing temperature and time reduced fatigue life. Temperature had the strongest influence on

fatigue life here, based on the higher magnitude of the coefficient for temperature than for time. The

product of temperature and log(time) did not have a significant effect on life, indicating no significant

enhancement for combinations of high temperature and high time as observed for LSHR uniform gage

fatigue specimens. Regressions of fatigue life versus depths of oxide scales, fingers, dissolution zones,

and their sum each did not give improved coefficients of determination or root mean square errors. This is

illustrated in Figure 12b, again showing fatigue life decreased with total depth of damage.

As for uniform gage fatigue of LSHR, ME3 notch fatigue lives for each condition were compared by

one-way analyses of variance, and then grouped where no significant difference in life was observed,

Figure 11. After exposures at 704 °C for 100 h and 440 h, no significant reductions in mean life were

identified with these limited notched fatigue tests. However, exposures of 704 °C for 2,020 h and those at

760 and 815 C gave significantly lower mean life than unexposed specimens, at a statistical confidence

of 95%. Yet, the exposure conditions of 704 °C for 2,020 h and 760 °C for 440 h appeared to be

transitional, and gave more widely varied lives that spanned between multiple groups. These required

comparisons of failure modes, to confirm the group classifications. The largest reduction in mean fatigue

life was again observed for the exposure condition producing the largest surface layer thicknesses, 815 °C

for 2,020 h, which reduced mean fatigue life by 99%. Also, specimens exposed at conditions giving

comparable intermediate values of surface layer thicknesses, 704 °C for 2,020 h, 760 °C for 440 h, and

815 °C for 100 h, had statistically equivalent mean fatigue lives at a confidence of 95%. This indicated

that fatigue life was related to the depth of environmental attack. However, as will be shown in failure

mode evaluations, both crack initiation and crack growth modes were affected by the exposures.

Material blanks aged at exposure conditions of 815 °C for 440 h and 2,020 h and then machined into

specimens (Table 6), as well as specimens exposed in vacuum at 815 °C for 440 h (Table 7), all had

fatigue lives not significantly reduced from that of unexposed specimens. These results are included for

comparison in Figure 11. Therefore, the debits in fatigue life associated with exposures for notched

specimens were due to environmental attack, and not due to aging of the remote interior microstructure

during pre-exposures. Several specimens were exposed at 815 °C for 2,020 h and then polished to remove

only the outer oxide scales (20 μm), or all oxide scales and the minor phase dissolution zones (50 μm).

For these samples, fatigue lives are also included in Figure 11. Polishing to a depth of 20 μm only slightly

improved fatigue life for this exposure condition, while removing 50 μm restored life to unexposed levels.

The effects of exposure on fatigue lives of LSHR and ME3 could be compared by normalizing fatigue

lives of exposed specimens by mean fatigue life of unexposed specimens, Figure 13. Notch fatigue life of

ME3 appeared to be influenced more by exposures at intermediate temperatures and times than for

uniform gage life of LSHR. This could be due to differences in how the fatigue lives of the varied

materials, specimen configuration, and test conditions were influenced by intermediate exposures.

Therefore, several notch fatigue configuration specimens were prepared from LSHR blanks, some

exposed at an intermediate condition of 815 °C for 440 h, then all were fatigue tested at 704 °C using the

same conditions as for the ME3 notched specimens. Unexposed LSHR specimens had about 35% higher

notched fatigue life than for ME3 in these conditions, attributable to the finer LSHR grain size (Fig. 3).

Nevertheless, the normalized mean fatigue life of exposed notch specimens of LSHR were reduced by

this exposure just as for ME3, and did not appear to be alloy dependent, Figure 13b. This may be related

to the concentration of stresses produced by the notch at the exposed surface. However, dissimilar failure

modes could also be operative for the two alloys, and should to be compared.

Effect of Exposures on Fatigue Failure Modes

Uniform-Gage Fatigue of LSHR

Failure initiation modes of the LSHR tested specimens are compared in Figures 14 to 18. To gain

further information of failure initiation modes, typical secondary cracks are also presented in the gage
sections of variously exposed specimens in Figure 19. Figures 14 to 17 show the failure modes of

specimen failures were grouped according to failure mode and associated fatigue life, indicated as Groups

N, L, M, and H in the fatigue life plot of Figure 8. For each unexposed specimen, usually one

transgranular crack initiated at the specimen surface, normal to the loading axis, to cause failure, Figure

14. In some cases, these cracks initiated at angled faces of grains adjacent to the surface, but these shifted

to transgranular crack growth within 10 m of the surface. These cracks then grew in a predominantly

transgranular mode normal to the loading axis. This failure mode was also usually observed after

exposures of 100 h at 704 °C, and could be grouped with unexposed specimens in showing no consistent

change in failure mode or life (Group N).

However, for Group L specimens exposed at 704 °C for 440 h and 2,020 h, 760 °C for 440 h, and

815 °C for 100 h, consistent changes in failure mode were observed, Figure 15. For this group, more

cracks were present in the oxide scales coating the sides of these specimens, Figure 19b to 19e. While

shorter cracks were roughly normal to the loading axis, they sometimes joined to follow the boundaries of

underlying grains. Usually, a singular failure initiation point was observed on each fracture surface as in

Figure 15a, but occasionally multiple crack initiation points were observed, Figure 15b. The main crack

became more torturous and mixed in mode, with several grain boundary surfaces evident adjacent to the

cracked oxide scales of the specimen surface as shown in Figure 15, as indicated by rounded contours of

grain boundaries and the rougher texture produced by coarser secondary  ́ particles precipitating and

growing along grain boundaries. However, these initial cracks at surface grain boundaries only grew to a

depth of about 30 m, then transitioned to transgranular crack propagation through adjoining interior

grains.

Group M specimens, subjected to a constant exposure for 440 h at 815 °C and cyclic exposure for a

cumulative hot time of 440 h at 815 °C, had a more extended intergranular failure initiation mode, Figure

16. More frequent cracking of the outer oxide scales was observed, with a mixture of shorter cracks

roughly normal to the loading axis, and longer cracks linked to follow the boundaries of underlying,

unrecrystallized grains, Figure 19j and 19k. The main failure initiation point had grain boundaries

exposed by intergranular crack growth to a depth of about 60 μm, Figure 16. This was again followed by

a transition to predominantly transgranular crack propagation further into the specimen interior. Their

mean fatigue lives represented about 20% of unexposed mean fatigue life.

Finally, for the most severe exposure at 815 °C for 2,020 h of Group H specimens, extensive cracking

of the oxide layers and adjoining recrystallized grains was evident on the specimens’ sides (Fig. 19g) and

along the fracture surfaces, Figure 17. Crack initiation and propagation of the adjoining unrecrystallized

grains remained predominantly intergranular until final overload failure. This was associated with the

highest effect on mean fatigue life of all tested conditions, with the 98% reduction giving about 2% of

unexposed mean fatigue life. This indicated that both crack initiation and crack growth modes could be

affected by the exposures, and explained why estimations of fatigue life assuming only enhanced crack

initiation at oxide layers of varied depth were insufficient.

There were several exceptions to the above trends. One unexposed specimen, which failed from an

aluminum oxide granulated inclusion at the specimen surface shown in Figure 18a, had a lower life than

Group N specimens at a statistical confidence of over 95%, yet it exhibited transgranular crack initiation

and growth. Previous work has shown that inclusions intersecting a disk superalloy specimen surface can

inordinately reduce fatigue life, and this should be considered a separate failure mode whose probability

of occurrence can depend on the inclusion number density, inclusion size, and surface area of the

specimen gage section (Ref. 4), as well as the specimen surface and fatigue test conditions. Specimens

aged as blanks at 815 °C for 2,020 h had lower mean life than Group N specimens, at a statistical

confidence of over 95%, yet transgranular crack initiation and growth. One of these specimens failed

from a near-surface pore (Fig. 18b), while another failed from an internal inclusion. However, three other

specimens tested after short exposures at low temperatures had cracks initiating from internal inclusions,

but no significant divergence in life from others at the same exposure condition. The lower life after aging

the blanks could also be related to a reduction in strength brought on by extended coarsening of the

precipitates producing the changes evident in Figure 3.
Longitudinal metallographic sections were also prepared for selected failed specimens, to examine

gage cross sections containing secondary surface cracks adjacent to the dominant failure initiation site,

and are shown in Figure 20. These locations would be unloaded as the dominant crack grew, and not

subjected to higher stresses often generated near final specimen failure. These locations were also

observed before sectioning for comparisons, shown previously in Figure 19. As shown in Figure 19a, 19h

and 20a, group N specimens had very few secondary cracks, which grew from the surface across grains

roughly normal to the loading axis. Group L specimens had more frequent secondary cracks, Figure 19c

through Figure 19e and Figure 20b. These were often still relatively flat and normal to the loading axis,

initiating in the outer oxide layers. Secondary cracks for Group M specimens sometimes linked together

to outline underlying grain boundaries, and then grew along grain boundaries for a short distance, as

shown in Figure 19j and 19k. Finally Group H specimens had very many cracks in the recrystallized zone

along the alumina fingers, Figure 20d. Some of these cracks then propagated as intergranular cracks

extending into the specimen interior, evident on the section’s adjacent fracture surface.

Notch Fatigue of ME3

Failure initiation modes of ME3 in the notch fatigue tests at 704 °C are compared in Figure 21-24. All

failures initiated from the notch surface. The dominant failure initiation modes could be divided into three

groups of N, M, and H (Fig. 11) according to failure modes and fatigue lives, but their associations with

exposure conditions were less rigid than for the uniform gage fatigue specimens. This could be due in part

to the fact that a far smaller volume was highly stressed near the notch root for ME3, in comparison to

that for the uniform gage section for LSHR. Therefore, representative images are presented to describe

each group for ME3, rather than exposure condition.

Group N specimens failed from primary cracks initiating and propagating in a predominantly

transgranular fashion through grains, with crack growth roughly normal to the loading axis, Figure 21.

This most often occurred for unexposed specimens, and sometimes after exposures of 704 °C for 100 h,

440 h, and 2020 h. In some cases, cracks initiated at angled faces of grains adjacent to the surface, but

these shifted to transgranular crack growth within 10 m of the surface. Specimens failed from either one

or two primary crack initiation points, indicated by the arrows in Figure 21. Fatigue life for these

specimens remained within 28,326 to 100,256 cycles.

Group M specimens failed from primary cracks that started more torturous and mixed in mode, with

several grain boundary surfaces evident adjacent to the cracked oxide scales of the specimen surface as

shown in Figure 22. These initial cracks at surface grain boundaries usually only grew up to one grain

deep, typically about 30 m, then transitioned to transgranular crack propagation through grains further

interior. This most often occurred for specimens exposed at 704 °C for 440 h and 2020 h, 704 °C for

440 h, and 815 °C for 100 h. With the thicker oxides generated by some exposures, fatigue encouraged a

higher number of cracks along the notch surface than for Group N specimens, which produced more

numerous crack initiation sites on the fracture surface. These surface cracks were similar to those of the

uniform gage specimens, although restricted to the highly stressed notch root section. But this failure

mode was also observed for one unexposed specimen and for one specimen exposed at 815 °C for 440 h,

showing the mode was not dependent only on exposure condition and associated outer surface layers for

notched specimens. Fatigue life for these specimens ranged from 2,430 to 30,472 cycles.

Group H specimens had extensive cracking across more expansive oxide layers and recrystallized

grain zones along the notch surface and at many failure initiation sites on the fracture surfaces, as shown

in Figure 23. Environmental exposures caused the adjacent unrecrystallized grains’ boundaries to be

degraded, leading to intergranular propagation that extended into the interior for several hundred microns.

This failure mode was observed for exposures at 815 C for 440 and 2,020 h. Fatigue life for these

specimens ranged from 397 to 874 cycles.

Additional tests were performed with altered exposure and testing conditions, to further examine the

relationships between fatigue life and failure modes, Table 7. ME3 notch specimens were exposed in

vacuum at 815 °C for 440 h, and others were aged as blanks at exposure conditions of 815 °C for 440 h
and 2,020 h and then machined into specimens for testing. Their fatigue lives were not significantly

reduced from that of unexposed specimens. This indicated the debits associated with exposures were due

to environmental attack and that aging of the remote interior microstructure during pre-exposures did not

significantly affect notch fatigue life in these test conditions. Yet, they had a different failure mode. After

transgranular crack initiation as for unexposed specimens, crack propagation transitioned to become

predominantly intergranular, as shown in Figure 24. An intergranular crack propagation mode was also

observed for specimens exposed at 815 C for 2,020 h and then re-polished to selectively remove surface

layers, which will be discussed below. Long term aging of the internal microstructure at these conditions

had apparently increased the potential for intergranular cracking during notch fatigue tests. It is not clear

why this shift in crack propagation mode was not observed in tests of uniform gage LSHR specimens

aged at 815 C for 2,020 h, but this apparently was related to the different testing configuration and

conditions.

Prior exposures at intermediate conditions of ME3 notch fatigue specimens appeared to have greater

effects on fatigue life at 704 C than for LSHR uniform gage fatigue specimens tested at 760 C, shown

in Figure 13a. However, several LSHR specimens machined into the notch fatigue specimen

configuration, exposed at 815 C for an intermediate time of 440 h, and tested at 704 oC had comparable

normalized fatigue lives to ME3, Figure 13b. As shown in Figure 25b, LSHR notch fatigue specimens

given this prior exposure had a very similar failure initiation mode as for ME3 specimens in Group H,

with extensive cracking across the oxide layers and recrystallized grain zones along the notch surface, at

many failure initiation sites. Intergranular propagation extended into the interior for 15 to 25 m, not

nearly as deep as for ME3 specimens. But overall, this showed that notched specimens of both alloys

fatigue tested at 704 C had lives affected more by prior exposures at these intermediate conditions.

It is curious that aging at 815 C for 2,020 h in the same furnace reduced fatigue life for LSHR

uniform gage specimens, but did not consistently do so for ME3 notch specimens, Table 6. This could be

due in part to the different coa  ́ precipitates, Figure 3. After this

 ́ precipitates than LSHR, which could explain

the smaller reductions in tensile strength. This also could be associated with the differing effects of aginginduced reductions in strength on uniform gage versus notched specimens. In the uniform gage of LSHR

specimens, maximum cyclic stresses were maintained by the test control waveform, and reduced yield

and ultimate strength could lower the fatigue stress-life response of aged specimens. But in the notches of

ME3 specimens, maximum cyclic stresses of aged specimens could be reduced in comparison to those of

unexposed specimens, due to more plastic flow in the notch. For the same applied net section stress, this

would produce lower maximum cyclic stresses in the notch of aged specimens than for those without the

extended aging. This could compensate for the lower fatigue stress-life response of aged specimens.

Longitudinal metallographic sections were also prepared across a primary failure initiation location

for selected ME3 specimens, and are compared in Figure 26. As shown in Figure 26a, group N specimens

had few, transgranular secondary cracks. Group M specimens had more numerous secondary cracks,

which sometimes started at grain boundaries but became transgranular during crack growth, Figure 26b.

As observed for uniform gage specimens, Group H specimens had very many cracks of the outer oxide

scales and recrystallize zone, which usually grew along the alumina fingers, Figure 26c. Intergranular

cracks sometimes then extended into the specimen interior.

Effects of Individual Exposed Surface Layers on Fatigue Failures in Notch Specimens

The different surface layers had different compositions and phases, so could have different cracking

properties. Therefore, the effects on fatigue life of individual surface layers generated by these exposures

were examined, by polishing away selected surface layers before fatigue testing. These experiments were

performed on ME3 notch fatigue specimens, in order to ensure all fatigue failures occurred at the polished

notch surface. Samples given the extreme exposure condition of 815 °C for 2,020 h were studied here, as

this exposure gave the largest dimensions of layers and consistently gave the Group H failure mode.
Polishing away material to a total depth of 20 μm removed the outer Cr2O3-TiO2 scale plus inner

Al2O3 branched fingers. However, this gave only marginally improved fatigue life, to an average of 1,305

cycles, included for comparison in Figure 11. Fractographic evaluation revealed that the specimens failed

by multiple intergranular crack initiations, with the cracks continuing to propagate along grain boundaries

into the interior, Figure 27a. This life and failure mode remained consistent with that of Group H

specimens.

Polishing away more material to a total depth of about 50 μm removed the above layers and also the

minor phase dissolution zone. This gave a significant increase in mean fatigue life to 35,998 cycles,

which is nearly comparable to the mean fatigue life of the unexposed specimens, Figure 11. Furthermore,

a single, transgranular crack initiated failure, Fig 27b, which is typical of unexposed failures. These

aspects of fatigue life and crack initiation mode were consistent with Group N specimens. However, the

deeply polished specimens had intergranular propagation extending over 1 mm in distance before

overload. This crack propagation mode was consistent with Group H ME3 specimens. It would appear the

deeper intergranular crack propagation mode was clearly related to aging of the internal microstructure

during this exposure, as shown for specimens exposed in vacuum and blanks aged at 815 °C for 440 to

2,020 h before machining into specimens.

Effects of Exposures on Tensile Response

Slow strain rate tensile tests described in ASTM G129-00 have often been used to screen the effects

of dynamic environment-assisted cracking. But the standard indicates these tests have also been

successfully used to screen for embrittlement due to prior hydrogen charging and plating processes.

Therefore, tensile tests of varied displacement rates and corresponding strain rates were performed after

exposures of LSHR and ME3, to examine exposure effects on tensile strength and ductility.

LSHR tensile specimens were exposed in air at 704 °C for 2,020 h, 760 °C for 440 h, and 815 °C for

2,020 h, along with blanks that were later machined into specimens. Typical tensile stress-strain curves

are shown in Figure 28a for tests at 760 °C performed with a displacement rate of 0.1041 mm/s and

average strain rate of 5.0 10-3 s-1, approximating the rates present for the fatigue tests. Companion tensile

tests of LSHR performed at 704 °C with a much slower displacement rate of 0.0017 mm/s are compared

in Figure 28b. Measured ultimate tensile strength and inelastic strain at fracture are listed in Table 8.

Strengths were not strongly affected by the exposures at 704 and 760 °C. However, the exposure of

815 °C for 2,020 h reduced tensile strength by 150 to 190 MPa in the tensile tests, with greater effects

observed for the slow strain rate testing condition. Yet, blanks exposed 815 °C for 2,020 h and

subsequently machined into specimens produced comparable strengths for those given equivalent

exposures in air. This indicated reductions in strength were primarily due to aging of the overall

microstructure, and is consistent with the lower fatigue life observed for LSHR blanks aged at these

conditions.

Ductility, as indicated by inelastic fracture strain, was reduced by exposures, going from 14 to 20% in

unexposed specimens to as low as 1.9 to 4.2% for specimens exposed 815 °C for 2,020 h. Greater

reductions in ductility were observed in tests at the slower strain rate. This suggested that slow strain rate

tensile tests could be useful for screening exposure effects. On the other hand, blanks exposed at 704, 760

and 815 °C all had higher ductilities than unexposed specimens, at both fast and slow strain rate

conditions. This indicated the reductions in ductility were driven by prior exposures in air, not aging of

the overall microstructure.

Several tensile tests of ME3 were performed in air at 704 °C with a displacement rate of 0.0017

mm/s, just as for LSHR. The response and effects of prior exposure or blank aging of 815 °C for 2,020 h

were quite comparable to LSHR, Figure 28b. Therefore, most tensile tests of ME3 were performed at

704 °C in a vacuum pressure not exceeding 8 10-6 torr, to assess the effects of prior exposures in air

without the influence of an air environment during the tensile test, Table 9. Typical tensile stress-strain

curves are shown in Figure 29, for displacement rates of 0.017 mm/s and 0.0017 mm/s. Ultimate tensile

strength and inelastic strain at fracture are compared in Table 9 and Figure 29. Tensile strength was not
reduced in tensile tests by exposure of 2,020 h at 704 °C, but was reduced about 80 MPa by exposure of

2,020 h at 815 °C. Unlike tests in air, ductility was only moderately reduced by prior exposures for tensile

tests performed in vacuum, going from 22 to 25% in unexposed specimens to 18 to 19% in specimens

exposed 2,020 h at 815 °C. So exposure effects on ductility were clearly more acute in air. The slower

strain rate decreased ultimate strength by 150 to 200 MPa, for each specimen exposure condition.

Several ME3 specimens were again re-polished after exposure, to remove the external oxide and

oxide subscale from the gage surface (20 μm depth removed), and these showed a small improvement in

tensile ductility. However, removal of all discussed layers including the minor phase dissolution zone,

(50 m depth removed), gave a full recovery in tensile ductility for each strain rate. These re-polished

specimens still had tensile strengths reduced about 80 to 100 MPa (Table 9) from unexposed specimens,

so the strength reductions were confirmed to be due to extended aging of the interior microstructure.

Therefore, while ductility was reduced by the surface zones produced by exposures in air, strength was

reduced in both LSHR and ME3 by extended aging of the interior microstructure, Figure 3.

Typical tensile fracture surfaces are compared in Figure 30. LSHR specimens exposed at 815 °C for

2,020 h and tested in air at 704 and 760 °C had very many surface cracks initiated at the outer oxide

layers and then grown by intergranular cracking through the minor phase depleted zone and underlying

superalloy. ME3 specimens exposed at the same conditions in air and tested in vacuum at 704 °C also had

very many surface cracks initiated in the oxide layers, with subsequent intergranular cracking through the

minor phase depleted zone. However, the underlying ME3 superalloy failed by transgranular microvoid

coalescence. Additional exposed ME3 specimens which were polished to remove the outer oxide layers

and minor phase depleted zone also failed by transgranular microvoid coalescence. This indicated that the

intergranular cracking observed for exposed as well as aged specimens required an air environment, and

appeared related to an increased environment susceptibility of aged grain boundaries.

Summary and Conclusions

The effects of high temperature exposures in air on the microstructure and fatigue properties of two

powder metal disk superalloys were determined. Fully machined mechanical test specimens of LSHR and

ME3 were exposed at 704, 760, and 815 C, for times up to 2,020 h. Coupons and blanks of both alloys

were also exposed at the same conditions prior to machining to isolate the effect of the aging treatment on

microstructures and fatigue properties. The depths of environmental attack were measured for each case.

Uniform gage and notch gage fatigue tests were performed at 704 and 760 C, for selected exposure

conditions.

Fatigue resistance was reduced up to 98% in both superalloys by selected exposures. This was

associated with enhanced cracking of surface oxide layers, followed by cracking of underlying grain

boundaries depleted of minor phases, which then continued further into the interior. Tensile tests were

subsequently performed to further investigate effects of the exposures on tensile response in air and in

vacuum. The exposures reduced strength due to aging of the internal microstructure, and reduced ductility

due to the enhanced cracking of surface layers. But intergranular cracking of the underlying superalloy

was suppressed for tensile tests in vacuum. The effects of exposures on fatigue life, tensile strength, and

tensile ductility were compared. The associated failure modes and their relationships to exposure-induced

changes in compositions and phases near the surface were determined.

It can be concluded from this work that:

1) Extended exposures on machined specimens in air at high temperatures can produce surface oxide

layers and oxidation-effected zones with enhanced susceptibility to cracking.

2) This can result in significantly lower fatigue lives, strength, and ductility for disk applications

exposed at these conditions, in uniform sections and more acutely at notches.

3) Exposures can also produce aged microstructures having reduced strength and higher susceptibility

to environment-assisted intergranular cracking in air.
4) Such aged microstructures can still have fatigue lives and ductilities near that of unaged specimens.

5) Cyclic exposures for equivalent cumulative times at temperatures up to 815 °C do not accelerate

this damage.

6) Service exposure limits of temperatures below 704 °C and times below 440 h could be needed in

some disk applications to prevent these effects, as longer exposures near 704 °C or higher temperature

exposures can substantially degrade fatigue lives and ductility without effective environmental barrier

coatings.

References

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Table 1. Measured compositions in weight percent of tested materials.

Alloy -wt.% Al B C Co Cr Fe Mo Ni Nb Ta Ti V W ZrLSHR 3.54 0.027 0.045 20.4 12.3 0.1 2.71 Bal. 1.49 1.52 3.45 0.0055 4.28 0.049ME3 3.40 0.025 0.050 20.7 12.8 

± 3.80 Bal. 0.90 2.30 3.70 

± 2.00 0.05
Table 2. Summary of linear intercept grain size and precipitate dimensions for test materials.

 Secondary Tertiary

Alloy

Mean

Linear

Intercept

Grain Size

- m

Count

-n

Median

Area -

m2

Mean

Area ±

95% Conf.

Interval -

m2

Median

Equiv.

Radius -

m

Mean

Equiv.

Radius ±

95% Conf.

Interval -

m

Count

-n

Median

Area -

nm2

Mean

Area ±

95%

Conf.

Interval -

nm2

Median

Equiv.

Radius

-nm

Mean

Equiv.

Radius ±

95% Conf.

Interval -nm

LSHR 15 204 0.051

0.051 ±

0.003 0.128

0.124 ±

0.004 119 678

761.2 ±

77.0 14.7 15.1 ± 0.7

LSHR Aged815 °C-2,020 h 15 200 0.077

0.079 ±

0.005 0.157

0.154 ±

0.006 0

ME3 29 168 0.082

0.095 ±

0.011 0.162

0.161 ±

0.010 118 941.8

1283.0 ±

184.1 17.3 18.8 ± 1.3

ME3 Aged815 °C-2,020 h 29 191 0.044

0.049 ±

0.004 0.118

0.120 ±

0.005 0
Table 3. Thickness of layers developed near the surface after exposures of ME3 in air. Oxide scale measurements were lognormally distributed, while the finger depths and ͛-dissolution layer thicknesses were normally distributed. The standard

errors were determined by the variation in average values of the three areas.

Temperature

- °C

Time

(hour)

Mean Oxide Scale

Thickness - m (Log

normal)

Estimated

Standard

Error -

m

Mean Finger Depth -

m (Normal)

Estimated

Standard Error

- m

Mean ¶-Dissolution

Layer Thickness -

m

(Normal)

Estimated

Standard Error

- m

Mean Carbide

Dissolution Layer

Thicknessm

Standard

Deviation -

m

704 10 0.185 0.035 0.164 0.060 ND ± ± ±704 100 0.372 0.040 0.421 0.034 0.571 0.461 ± ±

704 440 0.556 0.146 1.450 0.434 2.074 0.715 ± ±704 1010 0.927 0.046 2.343 0.872 2.732 0.256 ± ±704 2020 1.340 0.108 2.971 0.382 3.109 0.764 9.6 2.3760 10 0.288 0.069 0.492 0.074 0.287 0.233 ± ±760 100 0.670 0.027 1.230 0.035 1.819 0.194 ± ±760 440 1.093 0.178 2.653 0.605 3.274 0.183 ± ±760 1010 1.457 0.140 3.830 0.437 4.232 0.616 ± ±760 2020 1.985 0.210 5.173 0.152 5.804 0.907 16.4 2.4

815 10 0.571 0.041 0.953 0.143 1.355 0.294 ± ±

815 100 1.192 0.045 2.448 0.474 3.441 0.467 ± ±815 440 1.886 0.243 5.715 0.352 5.649 0.805 ± ±815 1010 2.662 0.125 7.235 0.245 7.845 0.291 ± ±815 2020 3.770 0.161 9.965 0.251 10.584 0.412 32.6 4.2
Table 4. Thickness of layers developed near the surface after exposures of LSHR in air. Oxide scale measurements were

log-normally distributed, while the finger depths and ͛-dissolution layer thicknesses were normally distributed. The standarderrors were determined by the variation in average values of the three areas.

Temperature

- °C

Time

(hour)

Mean Oxide Scale

Thickness -

m

(Log normal)

Estimated

Standard

Error -

m

Mean Finger

Depth -

m

(Normal)

Estimated

Standard

Error -

m

Mean ¶-Dissolution

Layer Thickness -

m

(Normal)

Estimated

Standard

Error -

m

704 100 0.126 0.069 0.150 0.032 0.554 0.089704 440 0.428 0.116 0.613 0.242 0.979 0.230704 2020 0.929 0.106 2.680 0.502 5.003 0.180760 440 1.261 0.070 2.649 0.408 4.828 0.167815 100 1.428 0.055 3.666 0.345 4.561 0.177815 440 1.850 0.175 5.441 0.884 6.370 0.712815 2020 3.883 0.531 9.816 0.385 13.480 1.185
Table 5. Summary of uniform gage fatigue test results at 760 oC for LSHR.

Specimen Identification

Prior Exposure

Temperature -C Prior Exposure Time -h Life -cycles

T3-L12 22,257*

U3-L1 46,025

Z6-L14 49,879

T3-L17 55,123T3-L15 Blank Aged 815C 2020 36,481

Z6-L17 Blank Aged 815C 2020 36,482Z6-L18 Blank Aged 815C 2020 30,914

T3-L1 704 100 41,624

T3-L16 704 100 43,577

Z6-L11 704 100 45,108

T3-L2 704 440 25,734

Z6-L5 704 440 40,474

Z6-L16 704 440 33,584

T3-L4 704 2020 39,050

T3-L13 704 2020 39,831

Z6-L9 704 2020 16,896

T3-L11 760 440 43,119

Z6-L6 760 440 36,835

Z6-L15 760 440 33,222

T3-L5 815 100 25,894T3-L8 815 100 23,036

Z6-L1 815 100 31,138

T3-L10 815 440 5,417

Z6-L4 815 440 11,921Z6-L7 815 440 9,593

T3-L7 815 2020 669

T3-L14 815 2020 635

Z6-L2 815 2020 1,695T3-L3 815 440 cyclic 9,186T3-L9 815 440 cyclic 16,172Z6-L8 815 440 cyclic 13,762

Z6-L10 815 440 cyclic 6,973

* Premature failure at surface inclusion
Table 6. Summary of notch gage fatigue test results at 704 oC for ME3.

Specimen

Identification

Prior Exposure

Temperature -°C

Prior

Exposure

Time -h

Prior

Exposure

Environment

Fatigue Life -

cycles

H111-NNR4 84,588H101-NNR4 47,394

H111-NKR2 19,516

S101C-NER3 28,326

S101A -NER1 Blank Aged 815 440 Air 65,613S101A-NEW1 Blank Aged 815 2020 Air 71,480

S101A-NER4 704 100 Air 100,256

S101C-NER8 704 100 Air 47,108

S101C-NEW5 704 100 Air 23,562

S101A-NER9 704 440 Air 4,531

S101A-NEW3 704 440 Air 33,730

S101B-NER3 704 440 Air 41,064

S101B-NEW6 704 440 Air 39,712

S101A-NER7 704 2020 Air 6,028

S101A-NEW2 704 2020 Air 15,265

S101B-NER5 704 2020 Air 31,150S101B-NER9 704 2020 Air 4,810S101B-NER6 760 440 Air 4,355

S101C-NER5 760 440 Air 30,472

S101C-NEW3 760 440 Air 8,422

S101B-NEW2 815 100 Air 10,284

S101B-NEW5 815 100 Air 6,002

S101C-NER9 815 100 Air 3,160

H101-NKR5 815 440 Air 874

H101-NNR5 815 440 Air 796

S101A-NER5 815 440 Air 2,430

H101-NKR2 815 2020 Air 397

H101-NLR1 815 2020 Air 482

S101A-NER3 815 2020 Air 588
Table 7. Summary of additional notch gage fatigue test results at 704 oC for ME3.

Specimen

Identification

Prior Exposure

Temperature -°C

Prior

Exposure

Time -h

Prior

Exposure

Environment

Notch Depth

Removed -

m Fatigue Life -cycles

S101A-NEW5 815 440 Vac 53,309

S101B-NER7 815 440 Vac 62,520S101A-NER6 815 2020 Air 23 927S101B-NER2 815 2020 Air 19 1,943S101B-NER8 815 2020 Air 19 1,046S101A-NER8 815 2020 Air 48 32,980S101B-NER4 815 2020 Air 52 42,140

S101B-NEW3 815 2020 Air 52 32,873
Table 8. Summary of tensile test results at 704 oC and 760 oC in air for LSHR and ME3.

Alloy

Specimen

Identification

Prior

Exposure

Temperature

-C

Prior

Exposure

Time -h

Test

Temperature

-C

Test

Environment

Displace

ment

Rate -

mm/s

Average

Strain

Rate -s-1

Yield

Strength

- MPa

Ultimate

Strength

-MPa

Inelastic

Fracture

Strain -

%

LSHR T3-T1A None 704 Air 0.00173 8.3E-05 1057 1220 14.4LSHR T2-T5 None 704 Air 0.00173 8.3E-05 1028 1310 14.3LSHR T3-T2B Blank 704 2,020 704 Air 0.00173 8.3E-05 1052 1197 22.1LSHR T3-T1C 704 2,020 704 Air 0.00173 8.3E-05 1062 1212 11.1LSHR Z6-T4B Blank 760 440 704 Air 0.00173 8.3E-05 1009 1173 23.5LSHR T3-T4C 760 440 704 Air 0.00173 8.3E-05 1018 1182 12.2LSHR Z6-T1D Blank 815 2,020 704 Air 0.00173 8.3E-05 948 1125 22.9LSHR T3-T1B 815 2,020 704 Air 0.00173 8.3E-05 912 1077 1.9LSHR T3-T4A None 760 Air 0.104 5.0E-03 1038 1282 20.5LSHR T2-T3 None 760 Air 0.104 5.0E-03 1020 1284 17.7LSHR T3-T2A Blank 704 2,020 760 Air 0.104 5.0E-03 1029 1281 25.0LSHR Z6-T3B 704 2,020 760 Air 0.104 5.0E-03 1034 1262 15.2LSHR Z6-T4D Blank 760 440 760 Air 0.104 5.0E-03 993 1257 24.7LSHR Z6-T2B 760 440 760 Air 0.104 5.0E-03 1001 1248 16.1LSHR Z6-T1A Blank 815 2,020 760 Air 0.104 5.0E-03 902 1186 23.2LSHR Z6-T2A 815 2,020 760 Air 0.104 5.0E-03 903 1137 4.2

ME3 H111-SRR1 None 704 Air 0.00173 8.3E-05 1031 1188 15.4ME3 S101-SRW3 None 704 Air 0.00173 8.3E-05 1018 1180 18.0ME3 S101-R1ANB Blank 815 2,020 704 Air 0.00173 8.3E-05 870 1078 19.9ME3 S101-R1ANC Blank 815 2,020 704 Air 0.00173 8.3E-05 845 1043 25.3ME3 H111-SRR3 815 2,020 704 Air 0.00173 8.3E-05 880 1000 2.1ME3 H111-SRW5 815 2,020 704 Air 0.00173 8.3E-05 878 999 3.4
Table 9. Summary of tensile test results at 704 oC in vacuum for ME3.

Specimen

Identification

Prior

Exposure

Temperature

-C

Prior

Exposure

Time -h

Gage

Depth

Removed

- m

Test

Temperature

-C

Displacement

Rate -mm/s

Average

Strain

Rate -s-1

Yield

Strength

- MPa

Ultimate

Strength

-MPa

Inelastic

Fracture

Strain -

%

S101-SRR1 None 704 0.00173 8.3E-05 1048 1128 22.8S101-SRR3 None 704 0.00173 8.3E-05 1022 1145 21.9

S101-R94 704 2,020 704 0.00173 8.3E-05 1032 1118 26.3S101-R61 815 2,020 704 0.00173 8.3E-05 960 1054 18.1S101-R92 815 2,020 15 704 0.00173 8.3E-05 967 1058 19.6S101-R91 815 2,020 66 704 0.00173 8.3E-05 963 1060 21.7

S101-SRR2 None 704 0.0173 8.3E-04 1027 1307 24.9

H101-SRW1 None 704 0.0173 8.3E-04 1032 1281 22.8

S101-R64 704 2,020 704 0.0173 8.3E-04 1015 1282 27.5S101-R31 815 2,020 704 0.0173 8.3E-04 947 1209 19.5S101-R32 815 2,020 15 704 0.0173 8.3E-04 988 1246 20.5S101-R62 815 2,020 49 704 0.0173 8.3E-04 938 1225 30.3
Fig. 1. Notch configuration for fatigue specimens of ME3, dimensions in mm.
Fig. 2. Typical temperature versus time cycle for cyclic exposure of selected LSHR fatigue specimens, attaining

maximum temperature of 815 oC for cumulative time of 440 h.

0

900800700600500400300200100

0 20 40 60 80 100 120

Temperature (oC)

Time (min)
Fig. 3. Microstructures of the test materials: a) Optical microscope (OM) image of LSHR grain structure, b) OM

image of ME3 grain structure, c) Secondary electron (SE) image of LSHR secondary (S) and tertiary (T) precipitatesbefore aging or exposures, d) SE image of ME3 precipitates before aging or exposures, e) SE unage if LSHRprecipitates after aging 815 °C for 2,020 h, f) SE image of ME3 precipitates after aging 815 °C for 2,020 h.

TT

S

a) b)

c) d)

e) f)

S

T

S

S

50 m 50 m
a) b) c)

d)

e) f) g)

Fig. 4. Backscattered electron (BSE) images showing effects of prior exposures on LSHR surface microstructure (beforetesting): a) 704 °C-100 h., b) 704 °C-440 h, c) 704 °C-2,020 h, d) 760 °C-440 h, e) 815 °C-100 h, f) 815 °C-440 h, g) 815 °C2,020 h.

re-crystalized grains _oxide scale depth

alumina fingers

TiO2 grain

Cr O2 3

dissolved
a) b) c)

d)

e) f) g)

Fig. 5. BSE images showing effects of prior exposures on ME3 surface microstructure: a) 704 °C-100 h, b)704 °C-440 h, c) 704 °C-2,020 h, d) 760 °C-440 h, e) 815 °C-100 h, f) 815 °C-440 h, g) 815 °C-2,020 h.

re-crystalized grains_oxide scale depth

alumina fingers

TiO2 grain

Cr2O3
Fig. 6. Backscattered electron (BSE) image in a) shows area examined by electron microprobe in ME3

coupon exposed at 815 oC for 2,020 h. Composition dot maps indicate levels in weight % of O, Al, and Ti,which change near the surface.

O 20

m

Al 20

m

_minor phases depleted at grain boundaries

20

m

Ti 20

m

a)
Ni

Fig. 6 (cont.) Composition dot maps indicating levels in weight % of Ni, Cr, Co, and Mo, which change nearthe surface.

20

m Cr 20

m

<depleted grain boundaries

Co 20

m

<depleted grain boundaries

Mo 20

m

<depleted grain boundaries
Fig. 7. Comparison of thicknesses for different surface zones of exposed coupons for ME3 and LSHR, a) continuous oxide scale

thickness, b) alumina finger length, and c) ¶-dissolution layer thickness. Black symbols and lines demark the data and power-lawfits for ME3, brackets indicate standard errors.

b)

a)

c)

Dissolution

Layer

Thickness

Continuous

Oxide Scale

Alumina

Finger

Length

ME3
Fig. 8. Effect of prior exposures on fatigue lives of uniform-gage LSHR specimens, with failure groups N (no effects), L

(low effects), M (medium effects), and H (high effects) indicated. * Blanks aged 815 oC for 2,020 h before machiningspecimens. All other exposures were performed in air prior to testing on machined specimens.

N

L

M

H

*
Fig. 9. Linear regressions of fatigue lives of uniform-gage LSHR specimens: a) Actual versus

estimated log(life) using stepwise regression of exposure temperature, log(time), and theirinteraction; b) Actual versus estimated log(life) using regression of total damage depth = outeroxide scale thickness + inner finger depth + dissolution depth. 95% confidence intervals ofestimated values indicated in red, RMSE - root mean square error. Points of poor agreementindicated in orange circles.

Log(life)=4.2131 - 0.3875

T - 0.4100

log t - 0.3272 T log t

Log(life)=3.9614 - 0.8560

D

a) b)
Fig. 10. Comparison of LSHR specimen surfaces before fatigue testing: a) 815 °C-440 h isothermal,

and b) cyclic exposures, with a local discontinuity observed after cyclic exposure; c) chromium-rich oxidetop layer was generally present for both exposures, including near discontinuity, and d) underlyingaluminum-rich oxide layer was uncovered at discontinuity.

c)

b)

d)

a)

Cr

Al

Cr

Ti

Ni Co

c)

Cr

Cr

Ni

Ti

Al

Co

Mo

W
Fig. 11. Comparison of notch fatigue lives for ME3 in tests at 704 oC after varied exposure time and

temperature, with failure groups N (no effects), M (medium effects), and H (high effects) indicated. *Blanks aged 815 oC for 440 h before machining specimens, or exposed 815 oC for 440 h in vacuum.All other exposures were performed in air prior to testing on machined specimens.

100

1,000

10,000

100,000

Fatigue Life-cycles

 Not 704C 704C 704C 760C 815C 815C 815C 815C 815C 815Cexposed 100h 440h 2,020h 440h 100h 440h 2,020h 440h 2,020h 2,020h

 in

ME3 Fatigue at 704C, 855/-43MPa, 0.33 hertz

N

H

M

*
Fig. 12. Linear regressions of fatigue lives of ME3 notch fatigue specimens: a) Actual versus

estimated log(life) using stepwise regression of exposure temperature, log(time), and theirinteraction; b) Actual versus estimated log(life) using regression of total damage depth = outer oxidescale thickness + inner finger depth + dissolution depth. 95% confidence intervals of estimatedvalues indicated in red, RMSE - root mean square error.

Log(life)=3.8000 - 0.5914

T - 0.4260 log t Log(life)=3.4875 

± 1.0157

D

a) b)
Fig. 13. Effects of prior exposures of machined specimens in air on fatigue lives, normalized by mean unexposed fatigue

lives: a) ME3 notched specimens tested at 704 °C (filled symbols) and LSHR uniform gage specimens tested at 760 °C(open symbols), with arrow indicating difference in response at intermediate exposure conditions, b) ME3 and LSHRnotched specimens tested at 704 °C, showing comparable response.

1.0000.1000.010

Normalized Fatigue Life

0.001

ME3 and LSHR Normalized Fatigue Lives, 704C

 Not 704C 704C 704C 760C 815C 815C 815CExposed 100h 440h 2,020h 440h 100h 440h 2,020h

ME3, LSHR 704C

1.0000.1000.010

Normalized Fatigue Life

0.001

 Not 704C 704C 704C 760C 815C 815C 815CExposed 100h 440h 2,020h 440h 100h 440h 2,020h

ME3 and LSHR Normalized Fatigue Lives

LSHR 760C

ME3 704C

b) a)
Fig. 14. Typical fracture surfaces for LSHR group N specimens: a) T3-L1, 704 °C-100 h, 41,624 cycles; b) T3-L17,

no exposure, 55,123 cycles.

a)

b)
Fig. 15. Typical fracture surfaces for LSHR group L specimens: a) T3-L2, 704 °C-440 h, 25,734 cycles; b) T3-L11,

760 °C-440 h, 43,119 cycles.

a)

b)

<Grain boundary

<Grain boundary

<Grain boundary
Fig. 16. Typical fracture surfaces for LSHR group M specimens: a) T3-L10, 815 °C-440 h constant exposure, 5,417

cycles; b) T3-L9, 815 °C-440 h cyclic exposure, 16,172 cycles.

a)

b)

<Grain boundaries

<Grain boundaries
Fig. 17. Typical fracture surfaces showing intergranular crack initiation and growth for LSHR group H specimens:

a) T3-L7, 815 °C-2,020 h, 669 cycles; b) Z6-L2, 815 °C-2,020 h, 1695 cycles.

a)b)
Fig. 18. Failure initiation sites for other LSHR specimens: a) T3-L12, no exposure, failed at

22,257 cycles from near-surface aluminum oxide inclusion; b) Z6-L17, Aged 815 °C-2,020 h,failed at 36,482 cycles from near-surface pore.

a) b)
Fig. 19. Typical surface cracks after fatigue tests of LSHR at 760 °C: a) 704 °C-100 h, b) 704 °C-440 h,

c) 704 °C-2,020 h, d) 760 °C-440 h, e) 815 °C-100 h, f) 815 °C-440 h, g) 815 °C-2,020 h.

a) b) c)

d)

e) f) g)
Fig. 19 (cont.). Typical surface cracks after fatigue tests of LSHR at 760 °C: h) No exposure,

i) Aged 815 °C-2,020 h, j) Static exposure 815 °C-440 h, k) Cyclic exposure 815 °C-440 h.

h)

j)

i)

k)
Fig. 20. Images showing surface cracks in longitudinal metallographic sections of LSHR specimens after fatigue

tests at 760 °C: a) Group N, T3-L17, no exposure, 55,123 cycles; b) Group L, T3-L11, 760 °C-440 h, 43,119 cycles;c) Group M, T3-L9, 815 °C-440 h cyclic, 16,172 cycles; d) Group H, T3-L14, 815 °C-2,020 h, 635 cycles. Exposedsurfaces were plated with Ni to preserve the oxide layers during metallographic preparation.

a) b)

c) d)

Ni plating

Ni plating

Ni plating Ni plating
Fig. 21. Typical ME3 notched specimen fracture surfaces for group N: a) S101C-NER3, no exposure, 28,326

cycles; b) S101A-NER4, 704 °C-100 h, 100,256 cycles.

a)b)
Fig. 22. Typical ME3 notched specimen fracture surfaces for group M, after static exposures: a) S101A-NER5,

815 °C-440 h, very many failure initiation points, 2,430 cycles; b) S101C-NER5, 760 °C-440 h, fewer failureinitiation points, 30,472 cycles.

a)b)
Fig. 23. Typical ME3 notched specimen fracture surfaces for group H, after static exposures: a) H101-NKR2,

815 °C-2,020 h, 397 cycles, b) H101-NKR5, 815 °C-440 h, 874 cycles.

a)

b)
Fig. 24. Typical failure initiation sites of ME3 notched gage specimens, after altered conditions: a) exposed 815 °C440 h in vacuum, S101A-NEW5, 53,309 cycles; b) Blank aged 815 °C-2,020 h before specimen machining, S101ANEW1, 71,480 cycles.

a)b)
Fig. 25. Typical failure initiation sites of LSHR notched gage specimens, fatigue tested at 704 °C as per ME3

notched specimens: a) no exposure, T1W-SPL2, 154,692 cycles; b) Exposed 815 °C-440 h, T1N-SPL2, 2,481cycles.

a)b)
Fig. 26. Optical images for longitudinal metallographic sections of surface cracks after fatigue tests of ME3 at

704 °C: a) Group N, S101A-NER4, 704 °C-100 h, very few cracks, 100,256 cycles; b) Group M, S101B-NER9,704 °C-2,020 h, showing more numerous cracks, 4,810 cycles; c) Group H, H101-NLR1, 815 °C-2,020 h, showingmany cracks in recrystallized grain zone, 482 cycles. Exposed surfaces were plated with Ni to preserve the oxidelayers during metallographic preparation.

a) b) c)

H101-NLR1 PE

815°C, 2020 h Nf=

482

M272

Ni plating

 S101B-NER9 PE

704°C, 2020 h

Nf= 4810

M510

Niplating

A-NER4 PE

704°C, 100 h

Nf= 100,256

M269

Ni plating
Fig. 27. Typical failure initiation sites of ME3 notched gage specimens, exposed at 815 °C-2,020 h in air, then

polished to remove surface layers before fatigue testing: a) 20 m depth removed, S101B-NER8, 1,046 cycles; b) 50

m depth removed, S101A-NER8, 32,980 cycles.

a)b)
Fig. 28. Effects of prior static exposures in air or aging on tensile response in air: a) LSHR at 760 °C and 0.104 mm/s, b)

LSHR and ME3 at 704 °C and 0.00173 mm/s. Conditions producing large reductions in strength ( uts) and elongation( f) indicated by arrows.

a) b)

0

800600400200

140012001000

0 5 10 15 20 25 30

Stress-MPa

Strain-%

T3-T4A NoneT2-T3 NoneT3-T2A Aged 704C/2,020hZ6-T3B Exposed 704C/2,020hZ6-T4D Aged 760C/440hZ6-T2B Exposed 760C/440hZ6-T1A Aged 815C/2,020hZ6-T2A Exposed 815C/2,020h

uts

f

0

800600400200

140012001000

0 5 10 15 20 25 30

Stress-MPa

Strain-%

T3-T1A NoneT2-T5 NoneT3-T2B Aged 704C/2,020hT3-T1C Exposed 704C/2,020hZ6-T4B Aged 760C/440hT3-T4C Exposed 760C/440hZ6-T1D Aged 815C/2,020hT3-T1B Exposed 815C/2,020hME3 H111-SRR1 NoneME3 S101-R1ANB Aged 815C-2020hME3 H111-SRR3 Exposed 815C/2,020h

uts

f
Fig. 29. Effect of prior static exposures in air on tensile response of ME3, tested in vacuum at 704 °C: a) 0.0173 mm/s,

b) 0.00173 mm/s. Maximum reductions in strength ( uts) and elongation ( f) indicated by arrows.

a) b)

0

800600400200

140012001000

0 5 10 15 20 25 30

Stress-MPa

Strain-%

S101-SRR1 No ExposureS101-SRR3 No ExposureS101-R94 704C/2,020hS101-R61 815C/2,020hS101-R92 815C/2,020h D-30umS101-R91 815C/2,020h D-112um

f

uts

0

800600400200

140012001000

0 5 10 15 20 25 30

Stress-MPa

Strain-%

S101-SRR2H101-SRW1S101-R64 704C/2,020hS101-R31 815C/2,020hS101-R32 815C/2,020h D-30umS101-R62 815C/2,020h D-97um

f

uts
Fig. 30. Effect of prior static exposures in air on tensile failure modes: a) LSHR, Z6-T2A, exposed 815 °C-2,020 h, tested in

air at 760 °C; b) LSHR, T3-T1B, exposed 815 °C-2,020 h, tested in air at 704 °C.

a)

b)
Fig. 30 (cont.) Effect of prior static exposures in air on tensile failure modes: c) ME3, S101-R31, exposed 815 °C -2,020 h,

tested in vacuum at 704 °C, showing surface cracking of the outer oxide layers (red arrow), intergranular cracking of theminor phase depleted zone (white) as for LSHR, but internal transgranular failure (green); d) ME3, S101-R62, exposed815 °C -2,020 h, 50 m removed, tested in vacuum at 704 °C showing transgranular failure and highest reduction in area.

c)

d)