GE Wukusick
Property-balanced nickel-base superalloys for producing single crystal articles
(R N5) 2000 US6074602A
Publication
Number: US6074602A
Publication
Date: 2000-06-13
Priority
Number: US1985790439A |
US1988253097A | US1991668816A | US199356597A | US1993152077A | US1994270528A
Application
Date: 1994-07-05
Title:
Property-balanced nickel-base superalloys
for producing single crystal articles
Inventor
- w/address: Wukusick Carl
Stephen,Cincinnati,OH,US | Buchakjian Jr. Leo,Loveland,OH,US
Assignee/Applicant:
General Electric Company,Cincinnati,OH,US
Front Page Drawing:
Abstract:
The present invention is directed to the achievement of
increased gas turbine engine efficiencies through further improvements in
nickel-base superalloys used to make parts and components for gas turbine
engines. The present invention comprises nickel-base superalloys for producing
single crystal articles having a significant increase in temperature
capability, based on stress rupture strength and low and high cycle fatigue
properties, over single crystal articles made from current production
nickel-base superalloys. Further, because of their superior resistance to
degradation by cyclic oxidation, and their resistance to hot corrosion, the
superalloys of this invention possess a balance in mechanical and environmental
properties which is unique and has not heretofore been obtained.
First Claim:
1. A nickel-base
single-crystal superalloy article consisting essentially of, in percentages by
weight, 5-10 Cr, 5-10 Co, 0-2 Mo, 3-8 W, 3-8 Ta, 0-2 Ti, 5-7 Al, Re in an
amount of up to 6, 0.08 to 0.2 Hf, 0.03-0.07 C, 0.003-0.006 B, and 0.0-0.04 Y,
the balance being nickel and incidental impurities.
Description w/Pub Language: This invention pertains generally to nickel-base superalloys
castable as single crystal articles of manufacture, which articles are
especially useful as hot-section components of aircraft gas turbine engines,
particularly rotating blades.
The efficiency of gas
turbine engines depends significantly on the operating temperature of the
various engine components with increased operating temperatures resulting in
increased efficiencies. One means by which the operating temperature capability
can be increased is by casting the components which operate at the highest
temperatures, e.g., turbine blades and vanes, with complex hollow passageways
therein so that cooling air can be forced through the component and out through
holes in the leading and trailing edges. Thus, internal cooling is achieved by
conduction and external cooling is achieved by film or boundary layer cooling.
The search for
increased efficiencies has also led to the development of heat-resistant
superalloys which can withstand increasingly high temperatures yet maintain
their basic material properties. Oftentimes, the development of such
superalloys has been done in conjunction with the design, development and
manufacture of the aforementioned cast components having intricate air cooling
passageways therein.
The present invention
is directed to the achievement of increased efficiencies through further
improvements in nickel-base superalloys. According, there is provided by the
present invention nickel-base superalloys for producing single crystal articles
having a significant increase in temperature capability, based on stress
rupture strength and low and high cycle fatigue properties, over single crystal
articles made from current production nickel-base superalloys. Further, because
of their superior resistance to degradation by cyclic oxidation, and their
resistance to hot corrosion, the superalloys of this invention possess a
balance in mechanical and environmental properties which is unique and has not
heretofore been obtained.
According to the
present invention, superalloys suitable for making single-crystal castings
comprise the elements shown in Table I below, by weight percent (weight %),
with the balance being nickel (Ni) plus incidental impurities:
TABLE I (See PDF)
The invention also includes cast single-crystal articles,
such as gas turbine engine turbine blades and vanes, made of an alloy having a
composition falling within the foregoing range of compositions.
There are two basic
directional solidification (DS) methods now well-known in the art by which
single crystal castings may be made. They generally comprise either the use of
a seed crystal or the use of a labyrinthine passage which serves to select a
single crystal of the alloy which grows to form the single crystal article
("choke" process).
In order to develop
and test alloys of the invention, three series of 3000 gram heats of the alloys
listed in Table II were vacuum induction melted and cast into 11/2" dia.
copper molds to form ingots. The ingots were subsequently remelted and cast
into 1/2"×2"× 4" single crystal slabs using the choke process,
although the other previously mentioned process could have been used.
In a series of
separate experiments, it was determined that yttrium retention in the single
crystal slabs was about 30% of that present in the initial ingots. Hence, in
preparing the series I, II, and III alloys shown in Table II, sufficient excess
yttrium was added to the initially cast alloys so as to achieve the yttrium
levels shown in Table II taking into account the 30% retention factor.
The series I alloys
were designed to evaluate the interactions between tungsten, molybdenum and
rhenium as gamma (γ) matrix alloying elements. The series II alloys were
designed somewhat independently from series I in order to accommodate
additional variables. Aluminum was maintained at a high level and titanium and
tantalum were varied to accomplish a range of gamma prime (γ') levels up
to about 63 volume percent and chromium was reduced in order to permit the
increased γ ' contents. Since it was determined that the 8% Cr series I
alloys as a group were less stable than the series II alloys, the base Cr level
was reduced from the 8% in series I to 7% to achieve better stability. Co was
varied in alloys 812-814 to evaluate the effect of Co on stability.
The series III alloys
were based on evaluations of the series I and II alloys. From series II, the
upper limit in γ' content, based on . gamma.' solutioning, was about 60
volume percent. Alloys 824-826 were based on alloy 820 which had 5.5% Re and
high strength, but was unstable. Thus, the Re content was reduced to achieve
stability. Alloys 827-829 were based on alloy 821 (0% Ti), but in which W and
Re were varied. Alloys 830-833 were based on alloy 800 (1.5% Ti), but in which
Re, W and Mo were varied. Alloys 834 and 835 contained increased Al at the
expense of Ta and Ti. In all the series III alloys, the Co content was
maintained at 10%, based on the evaluation of alloys 812-814 in series II.
The series I, II and
III alloys were evaluated for stress rupture strength and the results of the
tests are set forth in Table III. Prior to testing, the alloys, except for the
"R" series noted in Table III, were heat treated as 1/2" thick
single crystal slabs according to the following schedule: solutionizing at
2350-2400° F. for two hours to achieve solutioning of at least 95% of the
γ' phase followed by an intermediate age at 1975° F. for 4 hrs. and a
final age at 1650. degree. F. for 16 hrs.
TABLE II (See PDF)
TABLE III (See PDF)
The series III alloys
were initially tested at 1600° F./80 ksi and 1800° F./40 ksi. Based on other
tests, such as those reported in Table VII, additional test specimens were
resolutioned at 2390° F. for two hours, given a more rapid cool and aged at
2050. degree. F./4 hours+1650° F./4 hours, the "R" treatment listed
in Table III, and stress rupture tested at 1800° F./40 ksi and 2000. degree. F./20
ksi. The reheat treatment resulted in an average increase in rupture life at
1800° F./40 ksi of about 30%. At the critical parameter of 1800. degree. F./40
ksi for gas turbine engine applications, it is expected that the series I and
II alloys would also exhibit a 30% increase in life when given the
"R" treatment.
Other experiments have
shown that cooling rates from the solutionizing temperature to 2000° F. in the
range of 100-600. degree. F./min have only a slight effect on the stress
rupture properties of the alloys of the invention with higher rates tending to
improve the life at 1800. degree. F./40 ksi slightly. The data are shown in
Table IV.
TABLE IV (See PDF)
Thus, for the
superalloys of the invention, the presently preferred heat treatment is as
follows: solutionize in a temperature range sufficient to achieve solutioning
of at least 95% of the γ' phase, preferably 2385-2395° F., for 2 hrs.,
cool to 2000° F. at 100° F./min. minimum, furnace cool to 1200° F. in 60 min.
or less and thereafter cool to room temperature; heat to 2050.+-.25. degree. F.
for 4 hrs., furnace cool to below 1200° F. in 6 min. or less and thereafter to
room temperature; and heat to 1650.+-.25. degree. F. for 4 hrs. and thereafter
furnace cool to room temperature. All heat treatment steps are performed in
vacuum or an inert atmosphere, and in lieu of the steps calling for cooling to
room temperature the treatment may proceed directly to the next heating step.
The stress rupture
data from the series I, II, and III alloys indicates that about 5% Re provides
the highest rupture strength at 1800. degree. F./40 ksi. The data also show,
when rupture life is graphed as a function of rhenium content at constant
tungsten contents, that high rupture life at 1800° F./40 ksi can be obtained
with rhenium plus tungsten levels in the (3Re+7W) to (5Re+3W) ranges. In the
most preferred embodiment, Alloy 821, the presently preferred (Re+W)
combination is (3Re+ 5W) due to the present relative costs of rhenium and
tungsten.
All the alloys were
evaluated for microstructural stability. Specimens were heat treated by
solutionizing at 2375-2400° F./2 hrs. and aging at 1975° F./4 hrs. and at 1650°
F./16 hrs. Thereafter, different sets of specimens were heated for 1000 hrs. at
1800° F. and for 1000 hrs. at 2000° F. After preparation, including etching
with diluted Murakami's electrolyte, the specimens were examined
metallographically and the relative amount of topologically close packed phase
(TCP) was determined visually. The series II alloys, except alloys 818 and 819,
showed either no TCP precipitation or only traces of precipitation (821) and,
as a group, were less prone to microstructural instability than the series III
alloys and much less prone than the series I alloys at both 1800° F. and 2000°
F.
Table V presents the
results of cyclic oxidation tests on uncoated 1/4" dia.×3" long pin
specimens conducted at 2150° F. using a natural gas flame at Mach 1 gas
velocity. The specimens were rotated for uniform exposure and cycled out of the
flame once per hour to cool the specimens to room temperature. External metal
loss was measured on a section cut transverse to the length dimension of the
specimen. Metal loss per side was found by dividing the difference between the
pin diameter before and after test by two. The data in the table are the
average of two such measurements at 90° to each other across the diameter of
the specimen.
The two series I
alloys that contained yttrium (802 and 808) had exceptional oxidation
resistance. The series II alloys, all of which were yttrium-bearing, exhibited
no metal loss after 200 hours of high velocity oxidation (Mach I) at 2150° F.
and only 2-3 mils γ ' depletion, demonstrating that a synergistic Y+Hf
effect was operating. These data also demonstrate that Re improves the
oxidation resistance or at least is less detrimental than W which it has
replaced in the alloys and, from metallographic studies, also results in
improved γ' stability.
TABLE V (See PDF)
The hot corrosion
resistance of the alloys of the invention was evaluated alongside three alloys
used to produce production turbine blades, Rene' 125, B1900, and MM200(Hf), in
tests wherein specimens of the alloys were exposed to a JP-5 fuel-fired flame
at 1600° F. with a nominal 1 ppm salt added to the combustion products. The
test was first run at ˜1 ppm for 1040 hrs., and then at ˜2 ppm, for 578 hrs.
The chemical determination of NaCl on calibration pins at every 200 hours
indicated that the salt level was between 0 and 1 ppm during the first 1000
hours, between 1 and 2 ppm during the next 300-400 hours and about 2 ppm during
the remaining 300 hours. The following conclusions were drawn from these hot
corrosion tests: 1) B1900 was least resistant to hot corrosion at all salt
levels, 2) MM200(Hf) was the next least resistant alloy at all salt levels, 3)
the alloys of the invention, especially alloy 821, and Rene'125 exhibit similar
hot corrosion behavior, with the alloys of the invention being slightly less
resistant than Rene' 125, and 4) as is the case for Rene' 125 and other alloys,
the alloys of the invention appear to be sensitive to salt level in the
corrosion test with increased salt level resulting in poorer corrosion
resistance. Thus, the difference between B1900, MM200(Hf), Rene' 125, and the
alloys of the invention narrows at high salt levels. These results are
consistent with prior experience and indicate that the hot corrosion resistance
of the alloys of the invention will be adequate for applications where Rene'
125 equivalency is required.
Alloy 821 was scaled
up as a 300 lb master heat having the composition given in Table VI. No yttrium
was added to the master heat; rather, yttrium was added when the master heat
material was remelted and molten prior to DS'ing to produce single crystal
slabs and turbine blades. For the test specimens used to obtain the data of
Tables VII, VIII, IX, and X, yttrium in the amount of 400 ppm was added. Stress
rupture strength data for alloy 821 from the 300 lb master heat and the 12 lb.
laboratory heat are presented in Table IX.
TABLE VI (See PDF)
TABLE VII (See PDF)
Tensile strength, low
cycle fatigue and high cycle fatigue tests were performed on single crystal
material from the 300 lb heat of alloy 821 solutioned at 2390° F./2 hrs. and aged
at 1975° F./4 hrs. and 1650° F./16 hrs., with the results shown in Tables VIII,
IX, and X, respectively, where UTS is ultimate tensile strength; YS is yield
strength at 0.2% strain offset; El is elongation; and RA is reduction in area.
TABLE VIII (See PDF)
TABLE X (See PDF)
As discussed at
greater length in co-pending co-assigned application Ser. No. 595,854, the
superalloys of this invention break with the long-standing wisdom of the single
crystal superalloy arts that grain boundary strengthening elements such as B,
Zr and C are to be avoided, i.e., kept to the lowest levels possible consistent
with commercial melting and alloying practice and technology. One general
reason given for restricting such elements is to increase the incipient melting
temperature in relation to the γ' solves temperature thus permitting
solutionizing heat treatments to be performed at temperatures where complete
solutionizing of the γ' phase is possible in reasonable times without
causing localized melting of solute-rich regions. Another is to minimize or
preclude the formation of deleterious TCP phases.
As noted in the Ser.
No. 595,854 application, single crystal articles are not necessarily wholly of
a single crystal as there may be present therein grain boundaries referred to
as low angle grain boundaries wherein the crystallographic mismatch across the
boundary is generally accepted to be less than about 5 to 6 degrees. Low angle
grain boundaries are to be distinguished from high angle grain boundaries which
are generally regarded as boundaries between adjacent grains whose
crystallographic orientation differs by more than about 5-6 degrees. High angle
grain boundaries are regions of high surface energy, i.e., on the order of
several hundreds of ergs/cm.sup.2, and of such high random misfit that the
structure cannot easily be described or modeled.
As also noted therein,
the discovery that small, but controlled, amounts of such previously prohibited
elements can be tolerated resulted in the single crystal superalloys of the
Ser. No. 595,854 application which have improved tolerance to low angle grain
boundaries, i.e., have greater grain boundary strength than the
state-of-the-art single crystal superalloys. As one result of this increased
grain boundary strength, grain boundary mismatches far greater than the 6°
limit for prior art single crystal superalloy articles can be tolerated in
single crystal articles made from the nickel-base superalloys of that
invention. This translates, for example, into better in-service reliability,
lower inspection costs and higher yields as grain boundaries over a broader
range can be accepted by the usual inspection techniques. The novel features of
that invention have been embodied in the novel superalloys of the present
invention; thus, the superalloys of the present invention also exhibit improved
tolerance to low angle grain boundaries and also have the above-described
benefits.
The superalloys of
this invention are also alloyed with yttrium which renders them more highly
reactive with respect to ceramic molds and cores used in the investment casting
process than nickel-base superalloys not alloyed with yttrium. Ceramic/metal
instability is controlled by the bulk thermodynamic condition of the system.
The more negative the free energy of formation, ΔG°.sub.f, the greater the
affinity for oxygen. It has been found that the free energy of formation for
oxides becomes more negative as more reactive elements, such as yttrium, are
added resulting in a greater potential for metal/ceramic reaction than when
typical SiO.sub.2 and ZrO.sub.2 ceramic mold and core systems are used. Based
on thermodynamic considerations and the work reported in U.S. Department of the
Air Force publication AFML-TR-77-211, "Development of Advanced Core and
Mold Materials for Directional Solidification of Eutectics" (1977),
alumina is less reactive and is, therefore, a preferred material for molds,
cores and face coats when casting superalloys containing reactive elements.
It has also been found
that melt/mold and core interactions are decreased, the retention of yttrium
increased and the uniformity of yttrium distribution improved by the use of low
investment casting parameters and temperatures. This translates to the use of
the lowest possible superheat and mold preheat and a high withdrawal rate in
the casting of the single crystal articles of this invention.
Several uncored small
turbine blades were investment cast using alloy 821 material from the
previously mentioned 300 lb scale-up master heat. Those blades measured about
1.5" from tip to root with a span of approximately 0.75". Blade tip to
platform distance was 1". As noted earlier, yttrium was added to the
master heat material while molten and prior to DS'ing--in this case the amount
was 2000 ppm. In general, most blades exhibited acceptable crystal structure
and, as shown in Table XI, those cast using low casting parameters had better
yttrium retention. Also, it appeared that surface to volume ratio influences
yttrium retention; as the ratio increases, the yttrium retention decreases.
This is illustrated by comparison of yttrium retention at the leading and
trailing edges; the surface to volume ratio is lower in the leading edge
compared to the trailing edge, and the yttrium retention in the leading edge is
consistently higher than at the trailing edge.
TABLE XI (See PDF)
Additional single
crystal investment castings of large solid turbine blades (43/4"
tip-to-root) and small and large turbine blades having cores therein to define
serpentine passageways for the provision of cooling air were also made. The
large solid turbine blades required late yttrium additions of up to 2400 ppm in
order to obtain yttrium distributions within the desired 50-300 ppm level.
Similar such levels, coupled with the use of low investment casting parameters,
were required to obtain acceptable yttrium levels in the cored blades. As was
the case with the uncored small turbine blades, the effect of surface to volume
ratio was evident; the leading edge retained higher yttrium levels compared to
the trailing edge.
Although the present
invention has been described in connection with specific examples, it will be
understood by those skilled in the art that the present invention is capable of
variations and modifications within the scope of the invention as represented
by the appended claims.
Claims: What is
claimed is:
1. A nickel-base
single-crystal superalloy article consisting essentially of, in percentages by
weight, 5-10 Cr, 5-10 Co, 0-2 Mo, 3-8 W, 3-8 Ta, 0-2 Ti, 5-7 Al, Re in an
amount of up to 6, 0.08 to 0.2 Hf, 0.03-0.07 C, 0.003-0.006 B, and 0.0-0.04 Y,
the balance being nickel and incidental impurities.
2. The superalloy
article of claim 1 consisting essentially of, in percentages by weight,
6.75-7.25 Cr, 7.0-8.0 Co, 1.3-1.7 Mo, 4.75-5.25 W, 6.3-6.7 Ta, 0.02 max. Ti,
6.0-6.4 Al, 2.75-3.25 Re, 0.12-0.18 Hf, 0.04-0.06 C, 0.003-0.005 B, and
0.005-0.02 Y, the balance being nickel and incidental impurities.
3. The superalloy
article of claim 2 consisting essentially of, in percentages by weight, 7 Cr,
7.5 Co, 1.5 Mo, 5 W, 6.5 Ta, 0 Ti, 6.2 Al, 3 Re, 0.15 Hf, 0.05 C, 0.004 B, and
0.01 Y, the balance being nickel and incidental impurities.
4. The superalloy
article of claim 1, wherein the Co and Re contents are, in percentages by
weight, 5-8 and up to 3.25, respectively.
5. The superalloy
article of claim 1, wherein the Cr and W contents are, in percentages by
weight, 5-9.75 and 3-7, respectively.
6. The superalloy
article of claim 1, wherein the article is an airfoil member for a gas turbine
engine.
7. The superalloy
article of claim 2, wherein the article is an airfoil member of a gas turbine
engine.
8. The superalloy
article of claim 3, wherein the article is an airfoil member of a gas turbine
engine.
9. The superalloy
article of claim 4, wherein the article is an airfoil member of a gas turbine
engine.
10. The superalloy
article of claim 5, wherein the article is an airfoil member of a gas turbine
engine.
11. The superalloy
article of claim 1, wherein the superalloy has a gamma prime content of up to
60 volume percent.
12. The superalloy
article of claim 1, wherein the superalloy is substantially free of a
topologically close-packed phase that would cause microstructural instability.
13. The superalloy
article of claim 1, wherein the superalloy exhibits no metal loss after 200
hours of high-velocity oxidation testing at about 2150° F. with a gas velocity
of Mach 1 and cooling to room temperature once each hour.
14. The superalloy
article of claim 1, wherein the superalloy has a grain boundary mismatch of
greater than 6 degrees.
15. The superalloy
article of claim 1, wherein the Y content is, in percentage by weight,
0.005-0.03.
16. The superalloy
article of claim 1, wherein the Y content is about 0 weight percent.
17. A gas turbine
blade case from a nickel-base single-crystal superalloy consisting essentially
of, in percentages by weight, 5-10 Cr, 5-10 Co, 0-2 Mo, 3-8 W, 3-8 Ta, 0-2 Ti,
5-7 Al, Re in an amount of up to 6, 0.08 to 0.2 Hf, 0.03-0.07 C, 0.003-0.006 B,
and 0.0-0.04 Y, the balance being nickel and incidental impurities.
18. The gas turbine
blade of claim 17, wherein the Co and Re contents are, in percentages by
weight, 5-8 and up to 3.25, respectively.
19. The gas turbine
engine component of claim 17, wherein the Cr and W contents are, in percentages
by weight, 5-9.75 and 3-7, respectively.
20. The gas turbine
engine component of claim 17, wherein the superalloy has a gamma prime content
of up to 60 volume percent.
21. The gas turbine
engine component of claim 18, wherein the superalloy has a gamma prime content
of up to 60 volume percent.
22. The gas turbine
engine component of claim 19, wherein the superalloy has a gamma prime content
of up to 60 volume percent.
23. The gas turbine
engine component of claim 17, wherein the superalloy is substantially free of a
topologically close-packed phase that would cause microstructural instability.
24. The gas turbine
engine component of claim 18, wherein the superalloy is substantially free of a
topologically close-packed phase that would cause microstructural instability.
25. The gas turbine
engine component of claim 19, wherein the superalloy is substantially free of a
topologically close-packed phase that would cause microstructural instability.
26. The gas turbine
engine component of claim 17, wherein the superalloy exhibits no metal loss
after 200 hours of high-velocity oxidation testing at about 2150° F. with a gas
velocity of Mach 1 and cooling to room temperature once each hour.
27. The gas turbine
engine component of claim 18, wherein the superalloy exhibits no metal loss
after 200 hours of high-velocity oxidation testing at about 2150° F. with a gas
velocity of Mach 1 and cooling to room temperature once each hour.
28. The gas turbine
engine component of claim 19, wherein the superalloy exhibits no metal loss
after 200 hours of high-velocity oxidation testing at about 2150° F. with a gas
velocity of Mach 1 and cooling to room temperature once each hour.
29. The gas turbine
engine component of claim 17, wherein the superalloy has a grain boundary mismatch
of greater than 6 degrees.
30. The gas turbine
engine component of claim 18, wherein the superalloy has a grain boundary
mismatch of greater than 6 degrees.
31. The gas turbine
engine component of claim 19, wherein the superalloy has a grain boundary
mismatch of greater than 6 degrees.
32. The gas turbine
engine component of claim 17, wherein the Y content is, in percentage by
weight, 0.005-0.03.
33. The gas turbine
engine component of claim 18, wherein the Y content is, in percentage by
weight, 0.005-0.03.
34. The gas turbine
engine component of claim 19, wherein the Y content is, in percentage by
weight, 0.005-0.03.
35. The gas turbine
engine component of claim 17, wherein the Y content is about 0 weight percent.
36. The gas turbine
engine component of claim 18, wherein the Y content is about 0 weight percent.
37. The gas turbine
engine component of claim 19, wherein the Y content is about 0 weight percent.
38. A gas turbine
engine component cast from a nickel-base single-crystal superalloy consisting
essentially of, in percentages by weight, 6.75-7.25 Cr, 7.0-8.0 Co, 1.3-1.7 Mo,
4.75-5.25 W, 6.3-6.7 Ta, 0.02 max. Ti, 6.0-6.4 Al, 2.75-3.25 Re, 0.12-0.18 Hf,
0.04-0.06 C, 0.003-0.005 B, and 0.005-0.02 Y, the balance being nickel and
incidental impurities.
39. The gas turbine
engine component of claim 38, wherein the superalloy has a gamma prime content
of up to 60 volume percent.
40. The gas turbine
engine component of claim 38, wherein the superalloy is substantially free of a
topologically close-packed phase that would cause microstructural instability.
41. The gas turbine
engine component of claim 38, wherein the superalloy exhibits no metal loss
after 200 hours of high-velocity oxidation testing at about 2150° F. with a gas
velocity of Mach 1 and cooling to room temperature once each hour.
42. The gas turbine
engine component of claim 38, wherein the superalloy has a grain boundary
mismatch of greater than 6 degrees.
43. The gas turbine
engine component cast from a nickel-base single-crystal superalloy consisting
essentially of, in percentages by weight, 7 Cr, 7.5 Co, 1.5 Mo, 5 W, 6.5 Ta, 0
Ti, 6.2 Al, 3 Re, 0.15 Hf, 0.05 C, 0.004 B, and 0.01 Y, the balance being
nickel and incidental impurities.
44. The gas turbine
engine component of claim 43, wherein the superalloy has a gamma prime content
of up to 60 volume percent.
45. The gas turbine
engine component of claim 43, wherein the superalloy is substantially free of a
topologically close-packed phase that would cause microstructural instability.
46. The gas turbine
engine component of claim 43, wherein the superalloy exhibits no metal loss
after 200 hours of high-velocity oxidation testing at about 2150° F. with a gas
velocity of Mach 1 and cooling to room temperature once each hour.
47. The gas turbine
engine component of claim 43, wherein the superalloy has a grain boundary
mismatch of greater than 6 degrees.