U.S. patent application number 12/882405 was filed with the patent office on 2011-04-21 for method of forging a nickel base superalloy.
This patent application is currently assigned to ROLLS-ROYCE PLC. Invention is credited to Robert J. MITCHELL.
Application Number | 20110088817 12/882405 |
Document ID | / |
Family ID | 41462344 |
Filed Date | 2011-04-21 |
United States Patent
Application |
20110088817 |
Kind Code |
A1 |
MITCHELL; Robert J. |
April 21, 2011 |
METHOD OF FORGING A NICKEL BASE SUPERALLOY
Abstract
A method of forging a nickel base superalloy comprising
providing a nickel base superalloy preform (40) with a first
predetermined shape, the nickel base superalloy preform having been
produced by powder metallurgy. The nickel base superalloy preform
(40) is forged to produce a nickel base superalloy forged component
(50) with a second predetermined shape. The first predetermined
shape and the second predetermined shape are arranged such that the
effective strain at the end of the forging is less than 1. The
nickel base superalloy forged component (50) is given a supersolvus
heat treatment to produce a large grain size in the nickel base
superalloy forged component (50).
Inventors: |
MITCHELL; Robert J.;
(Nottingham, GB) |
Assignee: |
ROLLS-ROYCE PLC
London
GB
|
Family ID: |
41462344 |
Appl. No.: |
12/882405 |
Filed: |
September 15, 2010 |
Current U.S.
Class: |
148/514 ;
148/428 |
Current CPC
Class: |
C22C 19/03 20130101;
C22C 19/056 20130101; C22F 1/10 20130101 |
Class at
Publication: |
148/514 ;
148/428 |
International
Class: |
C22F 1/10 20060101
C22F001/10; B22F 3/24 20060101 B22F003/24; C22C 19/05 20060101
C22C019/05; B22F 3/17 20060101 B22F003/17 |
Foreign Application Data
Date |
Code |
Application Number |
Oct 15, 2009 |
GB |
0918020.9 |
Claims
1. A method of forging a nickel base superalloy comprising the
steps of: (a) providing a nickel base superalloy preform with a
first predetermined shape, the nickel base superalloy preform
having been produced by powder metallurgy, (b) forging the nickel
base superalloy preform to produce a nickel base superalloy forged
component with a second predetermined shape, wherein the first
predetermined shape and the second predetermined shape are arranged
such that the effective strain at the end of forging is less than
1, (c) supersolvus heat treating at least a region of the nickel
base superalloy forged component to produce a large grain size in
the at least a region of the nickel base superalloy forged
component, wherein the large grain size in the nickel base
superalloy is 80 to 140 micrometers.
2. A method as claimed in claim 1 wherein in step (b) the effective
strain at the end of the forging is less than or equal to 0.75.
3. A method as claimed in claim 1 wherein in step (b) the effective
strain at the end of the forging is less than or equal to 0.5.
4. A method as claimed in claim 1 wherein step (a) comprises
providing a stepped cylindrical preform, the cylindrical preform
having a first substantially cylindrical portion and a second
substantially annular portion arranged coaxially around the first
portion, the first cylindrical portion having a first thickness,
the second annular portion having a second thickness and the second
thickness is less than the first thickness and step (b) comprises
forging the stepped cylindrical preform to a substantially disc
shaped forged component and after step (c) the second portion
having coarser grains than the first portion.
5. A method as claimed in claim 4 wherein step (a) comprises
providing a substantially cylindrical preform, the cylindrical
preform having a third substantially annular portion arranged
coaxially around the second annular portion, the third annular
portion having a third thickness and the third thickness is less
than the second thickness and after step (c) the third portion
having coarser grains than the second portion.
6. A method as claimed in claim 5 wherein in the first cylindrical
portion the effective strain is about 0.9, in the second annular
portion the effective strain is about 0.75 and in the third annular
portion the effective strain is about 0.5.
7. A method as claimed in claim 4 wherein step (c) comprises a
subsolvus heat treatment in a first region of the forged component
and a supersolvus heat treatment in a second region of the forged
component.
8. A method as claimed in claim 7 wherein the first region
comprises the first portion and the second region comprises the
second portion.
9. A method as claimed in claim 5 wherein step (c) comprises a
supersolvus heat treatment in a first region of the forged
component and a supersolvus heat treatment in a second region of
the forged component.
10. A method as claimed in claim 9 wherein the first region
comprises the first portion and the second region comprises the
second portion and the third portion.
11. A method as claimed in claim 1 wherein step (c) comprises
supersolvus heat treating all of the nickel base superalloy forged
component to produce a large grain size in all of the nickel base
superalloy forged component.
12. A method as claimed in claim 1 wherein the forged component
comprises a forged component for a turbine disc or a compressor
disc.
13. A method as claims in claim 1 wherein the nickel base
superalloy consists of 18.5 wt % cobalt, 15 wt % chromium, 5 wt %
molybdenum, 2 wt % tantalum, 3.6 wt % titanium, 3 wt % aluminium,
0.5 wt % hafnium, 0.015 wt % boron, 0.06 wt % zirconium, 0.027 wt %
carbon and the balance nickel plus incidental impurities.
14. A method as claimed in claim 1 wherein step (b) comprises
forging at a temperature in the range of 1050.degree. C. to
1150.degree. C. and at a strain rate in the range of 0.001 to 0.1
per second.
15. A method as claimed in claim 1 wherein step (c) comprises
supersovlus heat treating at 20.degree. C. to 50.degree. C. above
the gamma prime solvus temperature for up to 4 hours.
16. A method as claimed in claim 1 wherein step (c) comprises
supersolvus heat treating at a temperature of 1110.degree. C. to
1150.degree. C. for up to 4 hours.
17. A method of forging a nickel base superalloy comprising the
steps of: (a) providing a nickel base superalloy preform with a
first predetermined shape, the nickel base superalloy preform
having been produced by powder metallurgy. (b) forging the nickel
base superalloy preform to produce a nickel base superalloy forged
component with a second predetermined shape, wherein the first
predetermined shape and the second predetermined shape are arranged
such that the effective strain at the end of forging is less than
1, (c) supersolvus heat treating at least a region of the nickel
base superalloy forged component to produce a large grain size in
the at least a region of the nickel base superalloy forged
component.
18. A nickel base superalloy forged component with a large grain
size in at least a region, wherein the large grain size is 80 to
140 micrometers.
Description
[0001] The present invention relates to a method of forging a
nickel base superalloy and in particular to a method of forging a
nickel base superalloy turbine disc or a nickel base superalloy
compressor disc.
[0002] Currently high strength, powder processed nickel base
superalloys such as RR1000, Rene 104, Alloy 10, LSHR, IN100, Rene
88DT and Rene 95 etc may be processed to produce one of three
different microstructures or combinations thereof. These different
microstructures are a) fine grain size microstructure ASTM 10-12
ALA 7, b) midi gain size microstructure ASTM 7.5-9 ALA 4 and c)
coarse grain size microstructure ASTM 5-7 ALA 2.
[0003] In addition a single component may be processed to produce a
range of grain sizes, with different grain sizes at different
locations, e.g. a dual microstructure turbine disc.
[0004] The grain size achieved in these nickel base superalloys is
influenced by three factors. The first factor is alloy chemistry,
in particular the levels of boron and carbon. The levels of boron
and carbon dictate the levels of borides or carbides present in the
nickel base superalloy and these borides or carbides pin grain
boundaries and restrict grain growth. The second factor is
processing parameters, principally forging temperature, strain rate
and final strain. The third factor is heat treatment. Fine grain
nickel base superalloys are heat treated below the gamma-prime
solvus temperature, a subsolvus heat treatment, the gamma-prime
phase acts to pin the grain boundaries. Coarse grain nickel
superalloys are heat treated above the gamma-prime solvus
temperature, a supersolvus heat treatment. Midi grain nickel base
superalloys are forged at a lower than usual temperature and then
heat treated above the gamma-prime solvus temperature, supersolvus
heat treatment. In any heat treatment that involves going above the
gamma-prime solvus temperature it is necessary to have an
understanding of both the chemistry of the nickel base superalloy
and prior mechanical working history to avoid critical grain
growth.
[0005] There is a need within the industry for a nickel base
superalloy which has a large grain size in excess of coarse grain
without changing the chemistry of the nickel base superalloy.
[0006] Accordingly the present invention seeks to provide a novel
method of forging a nickel base superalloy which has been produced
by powder metallurgy, which reduces, preferably overcomes, the
above mentioned problem.
[0007] Accordingly the present invention provides a method of
forging a nickel base superalloy comprising the steps of:
(a) providing a nickel base superalloy preform with a first
predetermined shape, the nickel base superalloy preform having been
produced by powder metallurgy, (b) forging the nickel base
superalloy preform to produce a nickel base superalloy forged
component with a second predetermined shape, wherein the first
predetermined shape and the second predetermined shape are arranged
such that the effective strain at the end of the forging is less
than 1, (c) supersolvus heat treating at least a region of the
nickel base superalloy forged component to produce a large grain
size in the at least a region of the nickel base superalloy forged
component, wherein the large grain size in the nickel base
superalloy is 80 to 140 micrometers.
[0008] Preferably in step (b) the effective strain at the end of
forging is less than or equal to 0.75.
[0009] More preferably in step (b) the effective strain at the end
of forging is less than or equal to 0.5.
[0010] Preferably step (a) comprises providing a stepped
cylindrical preform, the cylindrical preform having a first
substantially cylindrical portion and a second substantially
annular portion arranged coaxially around the first portion, the
first cylindrical portion having a first thickness, the second
annular portion having a second thickness and the second thickness
is less than the first thickness and step (b) comprises forging the
stepped cylindrical preform to a substantially disc shaped forged
component and after step (c) the second portion having coarser
grains than the first portion.
[0011] Preferably step (a) comprises providing a substantially
cylindrical preform, the cylindrical preform having a third
substantially annular portion arranged coaxially around the second
annular portion, the third annular portion having a third thickness
and the third thickness is less than the second thickness and after
step (c) the third portion having coarser grains than the second
portion.
[0012] Preferably in the first cylindrical portion the effective
strain is about 0.9, in the second annular portion the effective
strain is about 0.75 and in the third annular portion the effective
strain is about 0.5.
[0013] Step (c) may comprise a subsolvus heat treatment in a first
region of the forged component and a supersolvus heat treatment in
a second region of the forged component. The first region may
comprise the first portion and the second region may comprise the
second portion or the second region comprises the second portion
and the third portion.
[0014] Step (c) comprise supersolvus heat treating all of the
nickel base superalloy forged component to produce a large grain
size in all of the nickel base superalloy forged component.
[0015] The forged component may comprise a forged component for a
turbine disc or a compressor disc.
[0016] The present invention will be more fully described by way of
example with reference to the accompanying drawings in which:
[0017] FIG. 1 shows a gas turbine engine having a nickel base
superalloy turbine disc manufactured according to the present
invention.
[0018] FIG. 2 shows a nickel base superalloy turbine disc
manufactured according to the present invention.
[0019] FIG. 3 shows a micrograph of a nickel base superalloy forged
according to the prior art.
[0020] FIG. 4 shows a micrograph of a nickel base superalloy forged
according to the present invention.
[0021] FIG. 5 is a graph of grain size after heat treatment against
effective strain.
[0022] FIG. 6 is a graph of grain size after heat treatment against
time above gamma prime solvus
[0023] FIG. 7 is a schematic of a nickel base superalloy preform
for forging.
[0024] FIG. 8 is a schematic of a forged component produced from
the nickel base superalloy preform.
[0025] FIG. 9 is a graph showing the grain size against position in
the forged component of FIG. 8 following a supersolvus heat
treatment.
[0026] FIG. 10 is a graph showing the grain size against position
in the forged component of FIG. 8 following a dual microstructure
heat treatment.
[0027] A turbofan gas turbine engine 10, as shown in FIG. 1,
comprises in axial flow series an intake 12, a fan section 14, a
compressor section 16, a combustion section 18, a turbine section
20 and an exhaust 22. The turbine section 20 comprises a turbine
disc 24, which carries a plurality of circumferentially space
turbine blades 26. The gas turbine engine 10 is quite conventional
and its construction and operation will not be described
further.
[0028] The gas turbine engine turbine disc 24, as shown more
clearly in FIG. 2, comprises a hub, or cob, 28, a web 30 and a rim
32. The hub 28 is at the radially inner end of the turbine disc 24,
the rim 32 is at the radially outer end of the turbine disc 24 and
the web 30 extends radially between and interconnects the hub 28
and the rim 32. The rim 32, in this example, has a plurality of
circumferentially spaced slots 34 to receive the roots of turbine
blades 26, shown in FIG. 1, and circumferentially spaced posts 36
are provided on the rim 32 of the turbine disc 24 to define the
sides of the slots 34. The slots 34 may be firtree shape, or
dovetail shape. The turbine disc 24 comprises a high strength
nickel base superalloy, for example RR1000.
[0029] There is a need within the industry for a nickel base
superalloy which has a large grain size in excess of a coarse grain
size (ASTM 5-7). The range of grain sizes for these large grain
size microstructures in nickel base superalloy is ASTM 5-2 ALA 00.
A large uniform grain size such as this is desirable because
increasing the grain size increases the resistance to both creep
deformation and fatigue crack growth. However, strength and fatigue
properties are reduced.
[0030] FIGS. 3 and 4 show micrographs of high strength nickel base
superalloy RR1000 processed by a powder metallurgy route. RR1000
consists of 18.5 wt % cobalt, 15 wt % chromium, 5 wt % molybdenum,
2 wt % tantalum, 3.6 wt % titanium, 3 wt % aluminium, 0.5 wt %
hafnium, 0.015 wt % boron, 0.06 wt % zirconium, 0.027 wt % carbon
and the balance nickel plus incidental impurities. It is to be
noted that RR1000 consists of less than 0.03 wt % carbon and it is
preferred that the nickel base superalloy consists of less than
0.03 wt % carbon.
[0031] FIG. 4 shows RR1000 with a large grain size of ASTM 5-2 ALA
00 produced according to the present invention whereas FIG. 3 shows
RR1000 with a coarse grain size of ASTM 5-7 ALA 2 produced
according to the prior art.
[0032] In the prior art a nickel base superalloy turbine disc is
produced by isothermal forging a substantially cylindrical preform
to a substantially disc shaped forged component using a large
amount of plastic deformation. The effective strain at the end of
the forging process is greater than 1. Subsequently, the nickel
base superalloy forged component is given a supersolvus heat
treatment to produce a coarse grain size in the nickel base
superalloy forged component limited to the ASTM 5-7 range of
conventional coarse grain nickel base superalloys.
[0033] In the present invention the nickel base superalloy preform
is provided with a first predetermined shape, the nickel base
superalloy preform having been produced by powder metallurgy, the
nickel base superalloy preform is forged to produce a nickel base
superalloy forged component with a second predetermined shape,
wherein the first predetermined shape and the second predetermined
shape are arranged such that the effective strain at the end of the
forging process is less than 1. Subsequently, the nickel base
superalloy forged component is given a supersolvus heat treatment
to produce a large grain size in the nickel base superalloy forged
component in the ASTM 5-2 ALA 00, 80 to 140 micrometers, which is
greater in size than the conventional coarse grain in nickel base
superalloys.
[0034] FIG. 5 is a graph of grain size after heat treatment against
effective strain at the end of the forging process, or forging
step. It is seen in FIG. 5 that at effective strains above 1 the
grain size remains substantially constant. At effective strains
below 1 the grain size increases gradually from an effective strain
of 1 down to an effective strain of 0.75. The grain size increases
more rapidly from an effective strain of 0.75 to an effective
strain of 0.5 and then increases to a maximum grain size below an
effective strain of 0.5. We have found that for an effective strain
less than 1 the grain size is greater than that for an effective
strain greater than 1 for identical supersolvus heat treatments. We
have found that the grain size is relatively insensitive to
instantaneous strain rate in RR1000 compared to the prior art.
[0035] FIG. 6 is a graph of grain size against time above the gamma
prime solvus for three series of tests, M, P and D>1 where the
effective strain was greater than 1 and a single series of tests
D<1 where the effective strain was less than 1. This graph shows
the effect of time above the gamma prime solvus on the grain size.
It is clear from the data from these tests that the grain size
reaches a steady state after a short period of time. Irrespective
of effective strain, grain growth does not occur immediately, but
rather takes some time to occur. Grain growth occurs rapidly in the
nickel base superalloy where effective strain is less than 1 before
stabilising at the large grain size. Grain growth occurs more
slowly in the nickel base superalloy where effective strain is
greater than 1 before stabilising at the coarse grain size.
[0036] The definition of effective strain is given by the equation
{acute over (.epsilon.)}=
2/3[(.epsilon..sub.1-.epsilon..sub.2).sup.2+(.epsilon..sub.2-.epsilon..su-
b.3).sup.2+(.epsilon..sub.3-.epsilon..sub.1).sup.2].sup.1/2, where
{acute over (.epsilon.)} is the effective strain at the end of
forging, .epsilon..sub.1, .epsilon..sub.2 and .epsilon..sub.3 are
the principal strains in each of the three principal directions.
Thus for an isotropic material, knowledge of the uniaxial tensile
test behaviour together with the yield function enables the
stress-strain behaviour to be predicted for any stress system. The
effective strain equals the square root of two divided by three
multiplied by the square root of the squares of the difference in
the principal strains.
[0037] FIG. 7 shows an example of nickel base superalloy preform
40, pre-forge, which is used in a method according to the present
invention. In particular the nickel base superalloy preform 40 is a
stepped cylindrical preform and the stepped cylindrical preform 40
has a first substantially cylindrical portion 42, a second
substantially annular portion arranged 44 coaxially around the
first cylindrical portion 42 and a third substantially annular
portion 46 arranged coaxially around the second annular portion 44.
The first, second and third portions 42, 44 and 46 are integral. In
this example the first cylindrical portion 42 has a first thickness
10 h, the second annular portion 44 has a second thickness 4 h and
the third annular portion has a third thickness 2 h. The second
thickness 4 h is less than the first thickness 10 h and the third
thickness 2 h is less than the second thickness 4 h. In this
example the second annular portion 44 is twice as thick as the
third annular portion 46 and the first cylindrical portion 42 is
five times thicker than the third annular portion 46.
[0038] FIG. 8 shows a nickel base superalloy disc shaped forged
component 50 which has a thickness of h, which is produced by
forging, isothermally forging, the nickel base superalloy preform
40. The nickel base superalloy disc shaped forged component 50 is
then given a supersolvus heat treatment, such that the second
annular portion 44 has coarser grains than the first portion 42 and
the third portion 46 has coarser grains than the second portion 44.
The forging temperature is in the range of 1050.degree. C. to
1150.degree. C., preferably in the range of 1090.degree. C. to
1120.degree. C., and the strain rate is in the range of 0.001 to
0.1 per second, preferably 0.1 per second. The disc shaped forging
is produced in a close die shaped to form the nickel base
superalloy preform 40 to the desired shape of the nickel base
superalloy disc shaped forged component 50.
[0039] Considering the strain in the portions 42, 44 and 46 of the
stepped cylindrical preform 40 in FIG. 7, in a simplistic sense by
assuming there is no radial growth during forging. The simplistic
effective strain in the first portion 42 is
.epsilon..sub.eff=e/I.sub.0, where .epsilon..sub.eff is the
effective strain at the end of forging, e is the change in
dimension and I.sub.0 is the original dimension 9 h/10 h=0.9. The
effective strain in the second portion 44 is
.epsilon..sub.eff=e/I.sub.0, 3 h/4 h=0.75. The effective strain in
the third portion 46 is .epsilon..sub.eff=e/I.sub.0, h/2 h=0.5.
Thus, in the first cylindrical portion 42 the effective strain is
about 0.9, in the second annular portion 44 the effective strain is
about 0.75 and in the third annular portion 46 the effective strain
is about 0.5.
[0040] It is preferred to use an effective strain E.sub.eff of less
than or equal to 0.5 in order to produce a uniform large grain size
in the nickel base superalloy. An effective .epsilon..sub.eff of
less than, or equal, to 0.75 but greater than 0.5 would produce a
slightly finer grain size than the large grain size in the nickel
base superalloy.
[0041] FIG. 9 is a graph showing the grain size against position in
the forged component of FIG. 8 following a conventional supersolvus
heat treatment. It is seen that in the third annular portion 46 the
grain size is larger than the second annular portion 44 and the
grain size in the second annular portion 44 is larger than the
grain size in the first cylindrical portion 42. Thus, the nickel
base superalloy preform of FIG. 8 is forged to produce a nickel
base superalloy disc shaped forged component 50 and the nickel base
superalloy disc shaped forged component 50 is given a supersolvus
heat treatment and ageing heat treatment and then machined to
produce a nickel base superalloy turbine disc or compressor disc.
Thus, the finished nickel base superalloy turbine disc will have an
increasing grain size in the turbine disc with a coarse grain size
in the rim 32, third annular portion 46, a finer grain size in the
web 30, second annular portion 44, and a finest grain size in the
hub 28, first cylindrical portion 42. In the case of RR1000 the
nickel base superalloy preform is isothermally forged at a
temperature of 1050.degree. C. to 1150.degree. C., preferably in
the range of 1090.degree. C. to 1120.degree. C., at a strain rate
in the range of 0.001 to 0.1 per second, preferably 0.1 per second.
The RR1000 nickel base superalloy disc shaped forging is given a
conventional subsolvus solution heat treatment at 20.degree. C. to
30.degree. C. below the gamma prime solvus temperature for 1 to 4
hours, then air cooled, followed by a supersolvus heat treatment at
20.degree. C. to 50.degree. C. above the gamma prime solvus
temperature for up to 1 hour, then air cooled, followed by an
ageing heat treatment at 760.degree. C. (1400.degree. F.) for 16
hours and air cooled.
[0042] FIG. 10 is a graph showing the grain size against position
in the forged component of FIG. 8 following a dual microstructure
heat treatment. In a dual microstructure heat treatment one region
is given a supersolvus heat treatment and one region is given a
subsolvus heat treatment. It is seen that in the third annular
portion 46 the grain size is larger than the second annular portion
44 and the grain size in the second annular portion 44 is larger
than the grain size in the first cylindrical portion 42. Thus, the
nickel base superalloy preform of FIG. 8 is forged to produce a
nickel base superalloy disc shaped forged component 50 and the
nickel base superalloy disc shaped forged component 50 is given a
supersolvus heat treatment and ageing heat treatment and then
machined to produce a nickel base superalloy turbine disc or
compressor disc. Thus, the finished nickel base superalloy turbine
disc will have an increasing grain size in the turbine disc with a
large grain size in the rim 32, a coarse grain size in the web 30
and a fine grain size in the hub 28. In this particular example the
first cylindrical portion 42 is given a subsolvus heat treatment to
retain fine grains in the hub 28, the second annular portion 44 and
the third annular portion 46 are given a supersolvus heat treatment
to produce coarse grains in the web 30 and large grains in the rim
32 respectively. In the case of RR1000 the first cylindrical
portion is given a conventional subsolvus solution heat treatment
at 20.degree. C. to 30.degree. C. below the gamma prime solvus
temperature for 1 to 4 hours, then air cooled, followed by
subsolvus solution heat treatment at 20.degree. C. to 30.degree. C.
below the gamma prime solvus temperature for times up to 1 hour,
then air cooled, followed by an ageing heat treatment at
760.degree. C. (1400.degree. F.) for 16 hours and then air cooled.
The second and third annular portions were given a conventional
subsolvus solution heat treatment at 20.degree. C. to 30.degree. C.
below the gamma prime solvus temperature for 1 to 4 hours, then air
cooled, followed by a supersolvus heat treatment at 20.degree. C.
to 50.degree. C. above the gamma prime solvus temperature for times
up to 1 hour, then air cooled, followed by an ageing heat treatment
at 760.degree. C. (1400.degree. F.) for 16 hours and air
cooled.
[0043] In an alternative example a stepped cylindrical preform may
be provided, the cylindrical preform has a first substantially
cylindrical portion and a second substantially annular portion
arranged coaxially around the first portion, the first cylindrical
portion has a first thickness, the second annular portion has a
second thickness and the second thickness is less than the first
thickness. The stepped cylindrical preform is forged to a generally
disc shaped forged component and then the nickel base superalloy
forged component is given a supersolvus heat treatment and ageing
heat treatment and then machined to produce a nickel base
superalloy turbine disc or compressor disc. The second portion has
coarser grains than the first portion.
[0044] By controlling the shape of the nickel superalloy preform
and the shape of the nickel alloy forged component the effective
strain is kept to a level below 1, such that on subsequent
supersolvus heat treatment a large grain size is produced in the
nickel base superalloy. It is the final effective strain and not
the maximum instantaneous strain rate of the forging process, or
forging step, which dictates the grain size in the nickel base
superalloy. The supersolvus solution heat treatment may be applied
directly after the forging step or a subsolvus heat treatment may
be applied after the forging step and a supersolvus heat treatment
applied after the subsolvus heat treatment. The forging step may
comprise isothermal forging and isothermal application of strain
during the isothermal forging step allows the imparted strain
levels to be accurately controlled. The present invention may be
applied to specific locations of a component. The present invention
produces a very creep resistant, fatigue crack growth resistant
nickel base superalloy with a high gamma prime volume fraction.
[0045] The advantage of the present invention is that it enables
critical rotating components to be produced with enhanced high
temperature properties, in particular creep resistance and fatigue
crack growth resistance. This provides an increase in the operating
life of the component or enables the component to operate at higher
temperatures and may decrease the weight of the component.
[0046] The present invention is applicable to all high strength
powder processed nickel base superalloys used for gas turbine
engine turbine discs, compressor discs, high-pressure compressor
cones and turbine cover plates.
[0047] Other suitable nickel base superalloys are Rene 95, Rene
88DT, Alloy 10, LSHR, Rene 104 and IN100. Rene 95 consists of 8.12
wt % Co, 12.94 wt % Cr, 3.45 wt % Mo, 3.43 wt % W, 3.42 wt % Al,
2.44 wt % Ti, 3.37 wt % Nb, 0.05 wt % Zr, 0.07 wt % C, 0.012 wt % B
and the balance Ni and incidental impurities. Rene 88DT consists of
13.1 wt % Co, 15.8 wt % Cr, 4 wt % Mo, 3.9 wt % W, 2 wt % Al, 3.7
wt % Ti, 0.7 wt % Nb, 0.045 wt % Zr, 0.05 wt % C, 0.016 wt % B and
the balance Ni and incidental impurities. Alloy 10 consists of
17.93 wt % Co, 10.46 wt % Cr, 2.52 wt % Mo, 4.74 wt % W, 3.53 wt %
Al, 3.79 wt % Ti, 1.61 wt % Ta, 0.97 wt % Nb, 0.07 wt % Zr, 0.027
wt % C, 0.028 wt % B and the balance Ni and incidental impurities.
LSHR consists of 20.8 wt % Co, 12.7 wt % Cr, 2.74 wt % Mo, 4.37 wt
% W, 3.48 wt % Al, 3.47 wt % Ti, 1.65 wt % Ta, 0.049 wt % Zr, 0.024
wt % C, 0.028 wt % B and the balance Ni and incidental impurities.
Rene 104 consists of 20.6 wt % Co, 13.0 wt % Cr, 3.80 wt % Mo, 2.1
wt % W, 3.4 wt % Al, 3.7 wt % Ti, 2.4 wt % Ta, 0.05 wt % Zr, 0.04
wt % C, 0.03 wt % B and the balance Ni and incidental impurities.
IN100 consists of 18.5 wt % Co, 12.5 wt % Cr, 3.2 wt % Mo, 5.0 wt %
Al, 4.4 wt % Ti, 0.06 wt % Zr, 0.07 wt % C, 0.02 wt % B and the
balance Ni and incidental impurities.
[0048] In general the subsolvus heat treatment for these nickel
base superalloys is at a temperature of 20.degree. C. to 40.degree.
C. below the gamma prime solvus temperature for times of 1 to 6
hours and the supersolvus heat treatment for these nickel base
superalloys is at a temperature of 20.degree. C. to 50.degree. C.
above the gamma prime solvus temperature for times of up to 4
hours. Alternatively the subsolvus heat treatment is 1110.degree.
C. to 1150.degree. C. for times of 1 to 6 hours and the supersolvus
heat treatment is 1160.degree. C. to 1210.degree. C. for times of
up to 4 hours.
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