U.S. patent application number 12/494896 was filed with the patent office on 2010-12-30 for method of controlling and refining final grain size in supersolvus heat treated nickel-base superalloys.
This patent application is currently assigned to GENERAL ELECTRIC COMPANY. Invention is credited to Kenneth Rees Bain, David Paul Mourer.
Application Number | 20100329883 12/494896 |
Document ID | / |
Family ID | 42797247 |
Filed Date | 2010-12-30 |
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United States Patent
Application |
20100329883 |
Kind Code |
A1 |
Mourer; David Paul ; et
al. |
December 30, 2010 |
METHOD OF CONTROLLING AND REFINING FINAL GRAIN SIZE IN SUPERSOLVUS
HEAT TREATED NICKEL-BASE SUPERALLOYS
Abstract
A gamma prime precipitation-strengthened nickel-base superalloy
and method of forging an article from the superalloy to promote a
low cycle fatigue resistance and high temperature dwell behavior of
the article. The superalloy has a composition of, by weight,
16.0-22.4% cobalt, 6.6-14.3% chromium, 2.6-4.8% aluminum, 2.4-4.6%
titanium, 1.4-3.5% tantalum, 0.9-3.0% niobium, 1.9-4.0% tungsten,
1.9-3.9% molybdenum, 0.0-2.5% rhenium, greater than 0.05% carbon,
at least 0.1% hafnium, 0.02-0.10% boron, 0.03-0.10% zirconium, the
balance nickel and incidental impurities. A billet is formed of the
superalloy and worked at a temperature below the gamma prime solvus
temperature of the superalloy so as to form a worked article, which
is then heat treated above the gamma prime solvus temperature of
the superalloy to uniformly coarsen the grains of the article,
after which the article is cooled to reprecipitate gamma prime. The
article has an average grain size of not coarser than ASTM 7 and is
substantially free of critical grain growth.
Inventors: |
Mourer; David Paul;
(Beverly, MA) ; Bain; Kenneth Rees; (Loveland,
OH) |
Correspondence
Address: |
HARTMAN AND HARTMAN, P.C.
552 EAST 700 NORTH
VALPARAISO
IN
46383
US
|
Assignee: |
GENERAL ELECTRIC COMPANY
Schenectady
NY
|
Family ID: |
42797247 |
Appl. No.: |
12/494896 |
Filed: |
June 30, 2009 |
Current U.S.
Class: |
416/241R ;
148/538; 148/556; 148/677; 148/707; 419/28 |
Current CPC
Class: |
C22C 19/056 20130101;
C22C 19/057 20130101; F05B 2230/21 20130101; C22F 1/10 20130101;
C22C 19/03 20130101 |
Class at
Publication: |
416/241.R ;
148/677; 148/707; 148/538; 148/556; 419/28 |
International
Class: |
F01D 5/28 20060101
F01D005/28; C22F 1/10 20060101 C22F001/10; C22F 1/16 20060101
C22F001/16; B22D 25/06 20060101 B22D025/06; B22F 3/24 20060101
B22F003/24 |
Claims
1. A method of forming an article from a gamma prime
precipitation-strengthened nickel-base superalloy having a gamma
prime solvus temperature, the method comprising the steps of:
formulating the gamma prime precipitation-strengthened nickel-base
superalloy to have a composition of, by weight, about 16.0-22.4%
cobalt, about 6.6-14.3% chromium, about 2.6-4.8% aluminum, about
2.4-4.6% titanium, about 1.4-3.5% tantalum, about 0.9-3.0% niobium,
about 1.9-4.0% tungsten, about 1.9-3.9% molybdenum, about 0.0-2.5%
rhenium, greater than 0.05% carbon, at least 0.1% hafnium, about
0.02-0.10% boron, about 0.03-0.10% zirconium, the balance nickel
and incidental impurities; forming a billet of the superalloy;
working the billet at a temperature below the gamma prime solvus
temperature of the superalloy so as to form a worked article,
wherein the billet is worked to undergo deformation and to achieve
a maximum strain rate that is below an upper strain rate limit to
avoid critical grain growth yet sufficiently high to control
average grain size; heat treating the worked article at a
temperature above the gamma prime solvus temperature of the
superalloy for a duration sufficient to uniformly coarsen the
grains of the worked article; and cooling the worked article at a
rate sufficient to reprecipitate gamma prime within the worked
article, wherein the worked article has an average grain size of
not coarser than ASTM 7 and is substantially free of grains in
excess of three ASTM units coarser than the average grain size.
2. The method according to claim 1, wherein the forming step
comprises a process chosen from the group consisting of powder
metallurgy, cast and wrought, and spraycast forming techniques.
3. The method according to claim 1, wherein the forming step
comprises hot isostatic pressing or extrusion consolidation of a
powder of the superalloy to form the billet.
4. The method according to claim 1, wherein the superalloy contains
greater than 0.1 weight percent carbon.
5. The method according to claim 1, wherein the superalloy contains
greater than 0.1 weight percent up to about 0.125 weight percent
carbon.
6. The method according to claim 1, wherein the superalloy contains
0.1 to 0.6 weight percent hafnium.
7. The worked article formed by the method of claim 1, wherein the
worked article is a component chosen from the group consisting of
turbine disks and compressor disks and blisks of gas turbine
engines.
8. The method according to claim 1, wherein the maximum strain rate
is at least 0.003 per second.
9. The method according to claim 8, wherein the worked article has
an average grain size of not coarser than ASTM 8.
10. The worked article formed by the method of claim 9, wherein the
worked article is a component chosen from the group consisting of
turbine disks and compressor disks and blisks of gas turbine
engines.
11. The method according to claim 1, wherein the maximum strain
rate is at least 0.03 per second.
12. The method according to claim 11, wherein the worked article
has an average grain size of not coarser than ASTM 8.
13. The worked article formed by the method of claim 12, wherein
the worked article is a component chosen from the group consisting
of turbine disks and compressor disks and blisks of gas turbine
engines.
14. A method of forming an article from a gamma prime
precipitation-strengthened nickel-base superalloy having a gamma
prime solvus temperature, the method comprising the steps of:
formulating the gamma prime precipitation-strengthened nickel-base
superalloy to have a composition of, by weight, about 16.0-22.4%
cobalt, about 6.6-14.3% chromium, about 2.6-4.8% aluminum, about
2.4-4.6% titanium, about 1.4-3.5% tantalum, about 0.9-3.0% niobium,
about 1.9-4.0% tungsten, about 1.9-3.9% molybdenum, about 0.0-2.5%
rhenium, greater than 0.05% to about 0.125% carbon, about 0.1-0.6%
hafnium, about 0.02-0.10% boron, about 0.03-0.10% zirconium, the
balance nickel and incidental impurities; forming a billet of the
superalloy to have a fine grain size; working the billet at a
temperature below the gamma prime solvus temperature of the
superalloy so as to form a worked article, the working step being
performed so that the billet undergoes non-superplastic deformation
and achieves a maximum strain rate that is below an upper strain
rate limit to avoid critical grain growth yet sufficiently high to
control average grain size, wherein the maximum strain rate is at
least 0.03 per second; heat treating the worked article at a
temperature above the gamma prime solvus temperature of the
superalloy for a duration sufficient to uniformly coarsen the
grains of the worked article; and cooling the worked article at a
rate sufficient to reprecipitate gamma prime within the worked
article, wherein the worked article has an average grain size of
not coarser than ASTM 7 and is substantially free of grains in
excess of two ASTM units coarser than the average grain size.
15. The method according to claim 14, wherein the superalloy
contains greater than 0.10 weight percent carbon.
16. The method according to claim 14, wherein the worked article
has an average grain size of not coarser than ASTM 8.
17. The worked article formed by the method of claim 16, wherein
the worked article is a component chosen from the group consisting
of turbine disks and compressor disks and blisks of gas turbine
engines.
18. The method according to claim 14, wherein the maximum strain
rate is at least 0.03 to about 0.3 per second.
19. The method according to claim 18, wherein the worked article
has an average grain size of not coarser than ASTM 8.
20. The worked article formed by the method of claim 19, wherein
the worked article is a component chosen from the group consisting
of turbine disks and compressor disks and blisks of gas turbine
engines.
Description
BACKGROUND OF THE INVENTION
[0001] The present invention generally relates to nickel-base
superalloys and methods for processing such superalloys. More
particularly, this invention relates to a nickel-base superalloy
and a method of forging an article from the nickel-base superalloy
to promote a more controlled grain growth during supersolvus heat
treatment, such that the article is characterized by a
microstructure with a finer uniform grain size and exhibits
improved low cycle fatigue behavior.
[0002] The turbine section of a gas turbine engine is located
downstream of a combustor section and contains a rotor shaft and
one or more turbine stages, each having a turbine disk (rotor)
mounted or otherwise carried by the shaft and turbine blades
mounted to and radially extending from the periphery of the disk.
Components within the combustor and turbine sections are often
formed of superalloy materials in order to achieve acceptable
mechanical properties while at elevated temperatures resulting from
the hot combustion gases. Higher compressor exit temperatures in
modern high pressure ratio gas turbine engines can also necessitate
the use of high performance nickel superalloys for compressor
disks, blisks, and other components. Suitable alloy compositions
and microstructures for a given component are dependent on the
particular temperatures, stresses, and other conditions to which
the component is subjected. For example, airfoil components such as
blades and vanes are often formed of equiaxed, directionally
solidified (DS), or single crystal (SX) superalloys, whereas
turbine disks are typically formed of superalloys that must undergo
carefully controlled forging, heat treatments, and surface
treatments such as peening to produce a polycrystalline
microstructure having a controlled grain structure and desirable
mechanical properties.
[0003] Turbine disks are often formed of gamma prime (.gamma.')
precipitation-strengthened nickel-base superalloys (hereinafter,
gamma prime nickel-base superalloys) containing chromium, tungsten,
molybdenum, rhenium and/or cobalt as principal elements that
combine with nickel to form the gamma (.gamma.) matrix, and contain
aluminum, titanium, tantalum, niobium, and/or vanadium as principal
elements that combine with nickel to form the desirable gamma prime
precipitate strengthening phase, principally Ni.sub.3(Al,Ti).
Particularly notable gamma prime nickel-base superalloys include
Rene 88DT (R88DT; U.S. Pat. No. 4,957,567 to Krueger et al.) and
Rene 104 (R104; U.S. Pat. No. 6,521,175 to Mourer et al.), as well
as certain nickel-base superalloys commercially available under the
trademarks Inconel.RTM., Nimonic.RTM., and Udimet.RTM.. R88DT has a
composition of, by weight, about 15.0-17.0% chromium, about
12.0-14.0% cobalt, about 3.5-4.5% molybdenum, about 3.5-4.5%
tungsten, about 1.5-2.5% aluminum, about 3.2-4.2% titanium, about
0.5.0-1.0% niobium, about 0.010-0.060% carbon, about 0.010-0.060%
zirconium, about 0.010-0.040% boron, about 0.0-0.3% hafnium, about
0.0-0.01 vanadium, and about 0.0-0.01 yttrium, the balance nickel
and incidental impurities. R104 has a nominal composition of, by
weight, about 16.0-22.4% cobalt, about 6.6-14.3% chromium, about
2.6-4.8% aluminum, about 2.4-4.6% titanium, about 1.4-3.5%
tantalum, about 0.9-3.0% niobium, about 1.9-4.0% tungsten, about
1.9-3.9% molybdenum, about 0.0-2.5% rhenium, about 0.02-0.10%
carbon, about 0.02-0.10% boron, about 0.03-0.10% zirconium, the
balance nickel and incidental impurities. Another notable gamma
prime nickel-base superalloy is disclosed in European Patent
Application EP1195446, and has a composition of, by weight, about
14-23% cobalt, about 11-15% chromium, about 0.5-4% tantalum, about
0.5-3% tungsten, about 2.7-5% molybdenum, about 0.25-3% niobium,
about 3-6% titanium, about 2-5% aluminum, up to about 2.5% rhenium,
up to about 2% vanadium, up to about 2% iron, up to about 2%
hafnium, up to about 0.1% magnesium, about 0.015-0.1% carbon, about
0.015-0.045% boron, about 0.015-0.15% zirconium, the balance nickel
and incidental impurities.
[0004] Disks and other critical gas turbine engine components are
often forged from billets produced by powder metallurgy (P/M),
conventional cast and wrought processing, and spraycast or
nucleated casting forming techniques. Gamma prime nickel-base
superalloys formed by powder metallurgy are particularly capable of
providing a good balance of creep, tensile, and fatigue crack
growth properties to meet the performance requirements of turbine
disks and certain other gas turbine engine components. In a typical
powder metallurgy process, a powder of the desired superalloy
undergoes consolidation, such as by hot isostatic pressing (HIP)
and/or extrusion consolidation. The resulting billet is then
isothermally forged at temperatures slightly below the gamma prime
solvus temperature of the alloy to approach superplastic forming
conditions, which allows the filling of the die cavity through the
accumulation of high geometric strains without the accumulation of
significant metallurgical strains. These processing steps are
designed to retain the fine grain size originally within the billet
(for example, ASTM 10 to 13 or finer), achieve high plasticity to
fill near-net-shape forging dies, avoid fracture during forging,
and maintain relatively low forging and die stresses. (Reference
throughout to ASTM grain sizes is in accordance with the scale
established in ASTM Standard E 112.) In order to improve fatigue
crack growth resistance and mechanical properties at elevated
temperatures, these alloys are then heat treated above their gamma
prime solvus temperature (generally referred to as supersolvus heat
treatment), to cause significant, uniform coarsening of the
grains.
[0005] Forged gas turbine engine components often contain grains
with sizes of about ASTM 9 and coarser, such as ASTM 2 to 9, though
a much tighter range is typically preferred, such as grain sizes
within a limited range of 2 to 3 ASTM units. Such a limited range
can be considered uniform, which as used herein refers to grain
size and growth characterized by the substantial absence of
non-uniform critical grain growth. As used herein, critical grain
growth (CGG) refers to localized excessive grain growth in an alloy
that results in the formation of grains outside typical uniform
grain size distributions whose size sufficiently exceeds the
average grain size in the alloy (such as regions as coarse as ASTM
00 in a field of ASTM 6-10) to negatively affect the low cycle
fatigue (LCF) properties of an article formed from the alloy,
manifested by early preferential crack nucleation in the CGG
regions. Critical grain growth can also have a negative impact on
other mechanical properties, such as tensile strength. Critical
grain growth occurs during supersolvus heat treatment following hot
forging operations in which a wide range of local strains and
strain rates are introduced into the material. Though not wishing
to be held to any particular theory, critical grain growth is
believed to be driven by excessive stored energy within the worked
article, and may involve individual grains, multiple individual
grains within a small region, or large areas of adjacent grains.
The grain diameters of the effected grains are often substantially
coarser than the desired grain size. Disks and other critical gas
turbine engine components forged from billets produced by powder
metallurgy and extrusion consolidation have appeared to exhibit a
lesser propensity for critical grain growth than if forged from
billets produced by conventional cast and wrought processing or
spraycast forming techniques, but in any event are susceptible to
critical grain growth during supersolvus heat treatment.
[0006] The above-noted U.S. Pat. No. 4,957,567 to Krueger et al.
teaches a process for eliminating critical (abnormal) grain growth
in fine grained component formed of R88DT by controlling the
localized strain rates experienced during the hot forging
operation. Strain rate is defined as the instantaneous rate of
change of geometric strain with time. Krueger et al. teach that
local strain rates must generally remain below a critical value,
{dot over (.epsilon.)}.sub.c, in order to avoid detrimental
critical grain growth during subsequent supersolvus heat treatment.
According to Krueger et al., the maximum strain rate is
composition, microstructure, and temperature dependent, and can be
determined for a given superalloy by deforming test samples under
various strain rate conditions, followed by a suitable supersolvus
heat treatment. The maximum (critical) strain rate is then defined
as the strain rate that, if exceeded during deformation and working
of a superalloy and accompanied by a sufficient amount of total
strain, will result in critical grain growth after supersolvus heat
treatment.
[0007] Another processing limitation identified by Krueger et al.
as avoiding critical grain growth in a nickel-base superalloy
having a gamma prime content of, for example, 30-46 volume percent
and higher, is to ensure superplastic deformation of the billet
during forging. For this purpose, the billet is processed to have a
fine grain microstructure that achieves a minimum strain rate
sensitivity (m) of about 0.3 or greater for the superalloy within
the forging temperature and strain rate ranges. As known in the
art, the ability of a fine grain billet to deform superplastically
is dependent on strain rate sensitivity, and superplastic materials
exhibit a low flow stress as represented by the following
equation:
.sigma.=K{dot over (.epsilon.)}.sup.m
where .sigma. is the flow stress, K is a constant, {dot over
(.epsilon.)} is the strain rate, and m is the strain rate
sensitivity, with higher values of m corresponding to greater
superplasticity.
[0008] Further improvements in the control of final grain size have
been achieved with the teachings of commonly-assigned U.S. Pat. No.
5,529,643 to Yoon et al. and U.S. Pat. No. 5,584,947 to Raymond et
al. In addition to the requirement for superplasticity during
forging (in other words, maintaining a high m value), Raymond et
al. teach the importance of a maximum strain rate in combination
with chemistry control, particularly the carbon and/or yttrium
content of the alloy to achieve grain boundary pinning in alloys
having a gamma prime content of up to 65 volume percent. In a
particular example, Raymond et al. cites an upper limit strain rate
of below about 0.032 per second (s.sup.-1) for R88DT (identified by
Raymond et al. as Alloy D). In addition to maintaining a high m
value, Yoon et al. also identifies a maximum strain rate of not
more than about 0.032 s.sup.-1, particularly in reference to
forging R88DT (identified in Yoon et al. as Alloy A). Yoon et al.
further place an upper limit on the maximum strain rate gradient
during forging, and requires extended annealing of the forging at a
subsolvus temperature to remove stored strain energy prior to
performing a supersolvus heat treatment. Finally, Yoon et al.
achieve optimum superplasticity by forming the billet to have a
grain size of finer than about ASTM 12, and maintaining the billet
microstructure to achieve a minimum strain rate sensitivity of
about m=0.3 within the forging temperature range.
[0009] In addition to the absence of critical grain growth,
mechanical properties of components forged from fine grain
nickel-base superalloys further benefit from improved control of
the grain size distribution to achieve a distribution and average
grain size that are, respectively, as narrow and fine as possible.
Such a capability is particularly beneficial for high temperature,
high gamma prime content (e.g., about 30 volume percent and above)
superalloys, such as R88DT and R104, for which a desired uniform
grain size is generally not coarser than ASTM 6 for gas turbine
disks. Though prior forging practices of the type described above
have achieved grain sizes in a range of ASTM 5 to 8, less than
optimal mechanical properties can still result. For example, FIG. 1
is a graph evidencing that low cycle fatigue life tends to decrease
with coarser average grain sizes, even if uniform. The impact of
average grain size on low cycle fatigue properties of supersolvus
heat treated P/M superalloys is most apparent at low to
intermediate temperatures, such as in a range of about 400.degree.
F. to about 800.degree. F. (about 200.degree. C. to about
425.degree. C.) for R104. While the overall temperature capability
and balance of properties that R104 and other P/M alloys offer are
very attractive and relied on for the most advanced current engine
applications, even more benefit from these alloys could be obtained
if their low cycle fatigue properties and tensile behavior at low
to intermediate temperatures could be improved.
BRIEF DESCRIPTION OF THE INVENTION
[0010] The present invention provides a gamma prime
precipitation-strengthened nickel-base superalloy and a method of
forging an article from the superalloy to promote a more controlled
grain growth during supersolvus heat treatment, such that the
article is characterized by a microstructure with a finer uniform
grain size and exhibits improved low cycle fatigue behavior.
[0011] The method includes formulating the superalloy to have a
composition of, by weight, about 16.0-22.4% cobalt, about 6.6-14.3%
chromium, about 2.6-4.8% aluminum, about 2.4-4.6% titanium, about
1.4-3.5% tantalum, about 0.9-3.0% niobium, about 1.9-4.0% tungsten,
about 1.9-3.9% molybdenum, about 0.0-2.5% rhenium, greater than
0.05% and in certain embodiments greater than 0.1% carbon, at least
0.1% hafnium, about 0.02-0.10% boron, about 0.03-0.10% zirconium,
the balance nickel and incidental impurities. The superalloy is
similar in composition to R104, with the notable exceptions that
R104 does not contain hafnium and has a carbon content of 0.02-0.10
weight percent. A billet is formed of the superalloy and worked at
a temperature below the gamma prime solvus temperature of the
superalloy so as to form a worked article. In particular, the
billet is worked while maintaining strain rates as high as possible
to control average grain size, but below an upper strain rate limit
of greater than 0.03 per second to avoid critical grain growth. The
worked article is then heat treated at a temperature above the
gamma prime solvus temperature of the superalloy for a duration
sufficient to uniformly coarsen the grains of the worked article,
after which the worked article is cooled at a rate sufficient to
reprecipitate gamma prime within the worked article. The cooled
worked article has an average grain size of not coarser than ASTM 7
and preferably not coarser than ASTM 8, and is substantially free
of grains in excess of three ASTM units coarser than the average
grain size.
[0012] In view of the above, the superalloy has a sufficiently high
carbon content and is forged at sufficiently high local strain
rates so that, following a supersolvus heat treatment, the
resulting forged component is characterized by a fine and
substantially uniform grain size distribution. Also preferably
avoided is critical grain growth that would produce individual
grains or small regions of grains having grain sizes of more than
five and preferably three ASTM units coarser than the average grain
size in the component, or large regions that are uniform in grain
size but with a grain size coarser than a desired grain size range
of about two ASTM units. As a result, the forged component is
capable of exhibiting improved mechanical properties, particularly
low cycle fatigue behavior. Though not wishing to be held to any
particular theory, it is believed that formulating a superalloy to
have a chemistry similar to R104 but formulated to contain
relatively high carbon levels, especially carbon levels above the
upper limit of R104 (0.10 weight percent), allows the use of high
strain rates, resulting in a forged component capable of exhibiting
a more refined average grain size and substantially free of
critical grain growth, which together improve the low cycle fatigue
life of the component. Low cycle fatigue life can be particularly
improved within a temperature range of about 400.degree. F. to
about 800.degree. F. (about 200.degree. C. to about 425.degree. C.)
relative to R104 with a conventional carbon content of up to 0.10
weight percent. Other benefits of the finer average grain size
achieved with this invention include improved sonic inspection
capability due to lower sonic noise, and improved yield behavior in
service due to improved yield strength with finer grain size.
[0013] Other aspects and advantages of this invention will be
better appreciated from the following detailed description.
BRIEF DESCRIPTION OF THE DRAWINGS
[0014] FIG. 1 is a schematic graph representing low cycle fatigue
versus average grain size data for a variety of nickel-base
superalloys.
[0015] FIG. 2 is a perspective view of a turbine disk of a type
used in gas turbine engines.
[0016] FIG. 3 is a table listing a series of nickel-base superalloy
compositions initially identified to evaluate the effects of carbon
and hafnium contents on the low cycle fatigue behavior and hold
time fatigue crack growth rate behavior.
[0017] FIG. 4 is a table listing a series of nickel-base superalloy
compositions obtained and thermomechanically processed under
various conditions, including those in accordance with embodiments
of the present invention.
[0018] FIG. 5 is a table listing the compositions of FIG. 4 and
average grain size resulting from the use of different forging
conditions.
[0019] FIG. 6 shows four scanned images of two specimens from FIG.
4.
[0020] FIG. 7 is a graph plotting average grain size versus carbon
content, forging temperature, and forging rate for R104 and the
specimens of FIG. 4.
[0021] FIGS. 8 and 9 are graphs plotting the tensile strength
behavior of three specimens of FIG. 5 versus ASTM grain size whose
variation was achieved by the use of carbide enhanced grain size
control.
DETAILED DESCRIPTION OF THE INVENTION
[0022] The present invention is directed to gamma prime nickel-base
superalloys, and particular those suitable for components produced
by a hot working (e.g., forging) operation to have a
polycrystalline microstructure. A particular example represented in
FIG. 2 is a high pressure turbine disk 10 for a gas turbine engine.
The invention will be discussed in reference to processing of a
high-pressure turbine disk for a gas turbine engine, though those
skilled in the art will appreciate that the teachings and benefits
of this invention are also applicable to compressor disks and
blisks of gas turbine engines, as well as numerous other components
that are subjected to stresses at high temperatures and require low
cycle fatigue and high temperature dwell capabilities.
[0023] Disks of the type shown in FIG. 2 are typically produced by
isothermally forging a fine-grained billet formed by powder
metallurgy (PM), a cast and wrought processing, or a spraycast or
nucleated casting type technique. Such processes are carried out to
yield a billet with a fine grain size, typically about ASTM 10 or
finer, to achieve low flow stresses during forging. In a preferred
embodiment utilizing a powder metallurgy process, the billet can be
formed by consolidating a superalloy powder, such as by hot
isostatic pressing (HIP) or extrusion consolidation. The billet is
typically forged at a temperature at or near the recrystallization
temperature of the alloy but less than the gamma prime solvus
temperature of the alloy, and under conditions to enable filling of
the forging die cavity through the accumulation of high geometric
strains without the accumulation of significant metallurgical
strains. While superplastic forming conditions (corresponding to a
strain rate sensitivity (m) of 0.3 or higher at the forging
temperature) are often employed for this purpose, an aspect of the
invention is that the billet can be worked without the forging
process being fully superplastic, i.e., at strain rate sensitivity
values of less than about 0.3, for example, non-superplastically at
a strain rate sensitivity value of about 0.2 at the working (e.g.,
forging) temperature. After forging, a supersolvus (solution) heat
treatment is performed, during which grain growth occurs. The
supersolvus heat treatment is performed at a temperature above the
gamma prime solvus temperature (but below the incipient melting
temperature) of the superalloy to recrystallize the worked grain
structure and dissolve (solution) the gamma prime precipitates in
the superalloy. Following the supersolvus heat treatment, the
component is cooled at an appropriate rate to re-precipitate gamma
prime within the gamma matrix or at grain boundaries, so as to
achieve the particular mechanical properties desired. The component
may also be aged using known techniques with a short stress relief
cycle at a temperature above the aging temperature of the alloy if
desirable to reduce residual stresses.
[0024] In the case of the nickel-base superalloy R104, a
supersolvus heat treatment of a type described above has typically
yielded an acceptable but not wholly optimal average grain size
range of about ASTM 5 to 7, with the result that the low cycle
fatigue behavior of the resulting turbine disk is less than
optimal, particularly at temperatures of about 400.degree. F. to
about 800.degree. F. (about 200.degree. C. to about 425.degree.
C.). The present invention provides modifications to the chemistry
of R104 to control and limit grain growth during supersolvus heat
treatment to achieve and maintain a finer grain size following
supersolvus heat treatment, as well as avoid critical grain growth.
According to one aspect of the invention, a finer and more
controllable average grain size can be achieved by modifying the
R104 alloy to have a relatively high carbon content, for example,
greater than 0.05 weight percent carbon and in some cases greater
than 0.1 weight percent carbon. According to a second aspect of the
invention, improved high temperature dwell behavior can be achieved
by modifying the R104 alloy to contain at least 0.1 weight percent
hafnium. According to additional aspects of the invention, grain
refinement can be further promoted by utilizing relatively high
strain rates and relatively low temperatures during forging. The
teachings of U.S. Pat. Nos. 4,957,567 to Krueger et al., 5,529,643
to Yoon et al., and 5,584,947 to Raymond et al. are incorporated
herein by reference, particularly regarding the use of high strain
rates during forging and the placement of an upper limit on the
strain rate (critical strain rate) to avoid critical grain growth
during supersolvus heat treatment.
[0025] In an investigation leading to the present invention, a
series of targeted alloy compositions were defined (by weight
percent) as set forth in a table in FIG. 3. For reference, the
first two compositions listed in the table fall within the
disclosed range for R104. The targeted compositions reflect the
intent to evaluate alloys with carbon contents at and above the
maximum carbon content of 0.1 weight percent for R104, as well as
additions of hafnium. On the basis of these targeted compositions,
nine alloys were procured whose actual chemistries are indicated in
a table in FIG. 4. Processing of the alloys included consolidating
a powder of the alloy compositions to produce multiple billets of
each alloy, which were then hot worked (forged) followed by a
supersolvus heat treatment. Two sets of forging conditions were
used. A first, referred to as "Hot/Slow" in FIG. 5, entailed
forging conditions that included a maximum strain rate of about
0.003/sec at a forging temperature of about 2060.degree. F. (about
1130.degree. C.). The second, referred to as "Conventional" in FIG.
5, entailed forging conditions that included a conventional maximum
strain rate of about 0.03/sec at a forging temperature of about
1925.degree. F. (about 1050.degree. C.). The supersolvus heat
treatments were performed at a temperature of about 2140.degree. F.
(about 1170.degree. C.), which is above the gamma prime solvus
temperature (but below the incipient melting temperature) of R104.
During the heat treatment, the worked grain structures of the
forged specimens were recrystallized and the gamma prime
precipitates were dissolved (solutioned).
[0026] Following the supersolvus heat treatment, the specimens were
cooled at rates that ensured re-precipitation of gamma prime within
the gamma matrix or at grain boundaries. A controlled air cooling
was employed to yield an approximately constant cooling rate of
about 200.degree. F./minute for all specimens. Finally, the
specimens were aged at about 1550.degree. F. (about 845.degree. C.)
for about four hours, followed by about eight hours at about
1400.degree. F. (about 760.degree. C.).
[0027] As noted above and well known in the art, in addition to
grain recrystallization and solutioning gamma prime precipitates,
the supersolvus heat treatment also resulted in grain growth
(coarsening), typically resulting in grain sizes coarser than the
original billet grain size. FIG. 5 indicates the average ASTM grain
size observed for each alloy composition. From FIG. 5, it can be
seen that the "Hot/Slow" forging method produced significantly
coarser grains than the "Conventional" forging method. The finer
average grain sizes observed in the latter, which were typically
ASTM 8 or finer, would be expected to promote improved mechanical
properties of the forged specimens, including low cycle fatigue
resistance, tensile strength, fatigue strength, and other
mechanical properties desired for a turbine or compressor disk. In
addition, uniform average grain sizes within a range of about two
or three ASTM units were obtained, which would be further expected
to promote the low cycle fatigue resistance and other mechanical
properties of the specimens. The absence of excessively large
grains caused by critical grain growth was attributed to
maintaining strain rates during forging of the specimens below a
critical (maximum) strain rate for the superalloy compositions,
though at rates higher than those taught by Krueger et al.
According to Krueger et al., the critical strain rate of a gamma
prime nickel-base superalloy is composition, microstructure, and
temperature dependent, and can be determined for a given superalloy
by deforming test samples under various strain rate conditions, and
then performing suitable supersolvus heat treatments. The critical
strain rate is then defined as the strain rate that, if exceeded
during deformation and working of a superalloy and accompanied by a
sufficient amount of total strain, will result in critical grain
growth after supersolvus heat treatment. In the present
investigation, it was concluded that the upper strain rate limit
for the alloy specimens is greater than 0.03 per second, and
possibly as high as 0.32 per second.
[0028] FIG. 6 contains scanned images of two microphotographs of
the forged specimen identified as 101B in FIG. 5, as well as
scanned images of two microphotographs of a forged R104 specimen.
The images evidence that the carbide network within the 101B
specimen was significantly increased over that of R104. The
increased carbide network was attributed to the high carbon content
and the presence of hafnium in the 101B specimen. Without wishing
to be held to any particular theory, because hafnium is a strong
primary MC carbide former the hafnium content of the 101B specimen
may have promoted the formation of highly stable carbides,
contributing to high temperature carbide stability and aiding in
the ability to control grain size by the dispersion of primary MC
carbides in the matrix. FIG. 7 is a plot comparing ASTM average
grain size versus carbon content, and evidences the significant
influence carbon content had on average grain size in the forged
specimens. For example, at forging temperatures of about
2060.degree. F. (about 1130.degree. C.) carbon contents above 0.1
weight percent resulted in average grain sizes of finer than ASTM
7, and at forging temperatures of about 1925.degree. F. (about
1050.degree. C.) carbon contents above 0.05 weight percent and
above 0.1 weight percent resulted in average grain sizes of finer
than ASTM 8 and ASTM 8.5, respectively. On the basis of FIG. 1, the
finer average grain sizes achieved with the higher carbon contents
would be expected to correspond to improved low cycle fatigue
resistance. FIG. 7 also evidences that significantly finer average
grain sizes were obtained by forging at higher maximum strain rates
and lower forging temperatures. From these results, it was
concluded that finer average grain sizes can be achieved with
increasing carbon content above the disclosed upper limit for R104.
In part, the effect of the increased carbon content is believed to
be an increased pinning force that inhibits abnormal grain growth.
Generally, the finely dispersed carbides observed in FIGS. 6(a) and
6(b) were concluded to have restricted grain boundary motion during
supersolvus heat treatment, such that the grains are not permitted
to grow excessively and/or randomly to the extent that critical
grain growth occurs. From this investigation, another benefit
appears to be the ability to perform the forging operation at
relatively low temperatures, for example, about 1925.degree. F.
(about 1050.degree. C.) and likely in a range of about 1875 to
about 1975.degree. F. (about 1025 to about 1080.degree. C.).
[0029] A relationship between ASTM grain size and tensile behavior
of the forged specimens is evidenced in FIGS. 8 and 9, which show
tensile behavior and ductility at about 800.degree. F. (about
425.degree. C.) versus ASTM grain size. Improved tensile properties
were attributed to the presence of increased carbon and the forging
technique used, resulting in refining of the specimen grain
size.
[0030] In view of the above results, broad, narrower, and preferred
compositions and weight percent ranges were devised for the purpose
of obtaining improvements in low cycle fatigue resistance and dwell
crack growth behavior over the conventional R104 superalloy. These
compositions and ranges are set forth below in Table I.
TABLE-US-00001 TABLE I Broad Narrower Preferred Co 16.0-22.4 18 to
22 20.2 to 20.9 Cr 6.6-14.3 10 to 14 12.3 to 13.3 Al 2.6-4.8 2.5 to
4.0 3.1 to 3.7 Ti 2.4 to 4.6 3.0 to 4.2 3.4 to 3.8 W 1.9-4.0 1.9 to
3.0 1.7 to 2.2 Mo 1.9-3.9 2.5 to 3.9 3.5 to 3.9 Nb 0.9-3.0 0.9 to
2.0 0.9 to 1.0 Ta 1.4-3.5 1.7 to 3.0 2.1 to 2.6 Hf at least 0.1 0.1
to 0.6 0.2 to 0.5 C >0.05 >0.10 to 0.125 0.11 to 0.12 B
0.02-0.10 0.02 to 0.05 0.02 to 0.03 Zr 0.03-0.10 0.03 to 0.08 0.04
to 0.06 Ni Balance Balance Balance
[0031] While the invention has been described in terms of
particular processing parameters and compositions, the scope of the
invention is not so limited. Instead, modifications could be
adopted by one skilled in the art, such as by modifying the
disclosed processing by substituting other processing steps or
including additional processing steps. Accordingly, the scope of
the invention is to be limited only by the following claims.
* * * * *