U.S. patent application number 11/694204 was filed with the patent office on 2007-10-04 for nickel-based alloy.
This patent application is currently assigned to SNECMA. Invention is credited to Isabelle Augustins Lecallier, Pierre Caron, Jean-yves Guedou, Didier Locq, Loeiz Naze.
Application Number | 20070227630 11/694204 |
Document ID | / |
Family ID | 36694327 |
Filed Date | 2007-10-04 |
United States Patent
Application |
20070227630 |
Kind Code |
A1 |
Augustins Lecallier; Isabelle ;
et al. |
October 4, 2007 |
NICKEL-BASED ALLOY
Abstract
The invention provides alloys or superalloys based on nickel
essentially comprising the following elements in the amounts
indicated as percentages by weight: Cr: 11.5% to 13.5%; Co: 11.5%
to 16.0%; Mo: 3.4% to 5.0%; W: 3.0% to 5.0%; Al: 2.2% to 3.2%; Ti:
3.5% to 5.0%; Nb: 0.5% to 2.0%; Hf: 0.25% to 0.35%; Zr: 0 to 0.07%;
C: 0.015% to 0.030%; B: 0.01% to 0.02%; and Ni: complement to 100%.
The alloy is for use in the production of turbine or compressor
disks for turbo-machines, using powder metallurgy techniques.
Inventors: |
Augustins Lecallier; Isabelle;
(Rambouillet, FR) ; Caron; Pierre; (Les Ulis,
FR) ; Guedou; Jean-yves; (Le Mee Sur Seine, FR)
; Locq; Didier; (Le Plessis Robinson, FR) ; Naze;
Loeiz; (Mennecy, FR) |
Correspondence
Address: |
OBLON, SPIVAK, MCCLELLAND, MAIER & NEUSTADT, P.C.
1940 DUKE STREET
ALEXANDRIA
VA
22314
US
|
Assignee: |
SNECMA
Paris
FR
ARMINES
Paris
FR
ONERA(OFF NAT D'ETUDES ET DE RECHERCHES AEROS)
Chatillon
FR
|
Family ID: |
36694327 |
Appl. No.: |
11/694204 |
Filed: |
March 30, 2007 |
Current U.S.
Class: |
148/428 |
Current CPC
Class: |
C22C 1/0433 20130101;
C22C 19/056 20130101; C22F 1/10 20130101; C22C 19/05 20130101 |
Class at
Publication: |
148/428 |
International
Class: |
C22C 19/05 20060101
C22C019/05 |
Foreign Application Data
Date |
Code |
Application Number |
Mar 31, 2006 |
FR |
06 51145 |
Claims
1. An alloy essentially comprising the following elements, in the
amounts indicated, as percentages by weight: Cr: 11.5% to 13.5%;
Co: 11.5% to 16.0%; Mo: 3.4% to 5.0%; W: 3.0% to 5.0%; Al: 2.2% to
3.2%; Ti: 3.5% to 5.0%; Nb: 0.5% to 2.0%; Hf: 0.25% to 0.35%; Zr: 0
to 0.07%; C: 0.015% to 0.030%; B: 0.01% to 0.02%; and Ni:
complement to 100%.
2. An alloy according to claim 1, wherein the sum of the amounts of
Al, Ti, and Nb, as an atomic percentage, is 10.5% or more and 13%
or less.
3. An alloy according to claim 1, wherein the amounts of Al, Ti,
and Nb, as an atomic percentage, are such that the ratio between
the sum of the amounts of Ti and Nb, and the amount of Al, is 0.9
or more and 1.1 or less.
4. An alloy according to claim 2, wherein the amounts of Al, Ti,
and Nb, as an atomic percentage, are such that the ratio between
the sum of the amounts of Ti and Nb, and the amount of Al, is 0.9
or more and 1.1 or less.
5. An alloy according to claim 1, wherein the amounts of W, Mo, Cr,
and Co, as atomic percentages, is such that the sum of the amounts
of W, Mo, Cr, and Co is 30% or more and 34% or less, and such that
the sum of the amounts of W and Mo is 3% or more and 4.5% or
less.
6. An alloy according to claim 2, wherein the amounts of W, Mo, Cr,
and Co, as atomic percentages, is such that the sum of the amounts
of W, Mo, Cr, and Co is 30% or more and 34% or less, and such that
the sum of the amounts of W and Mo is 3% or more and 4.5% or
less.
7. An alloy according to claim 3, wherein the amounts of W, Mo, Cr,
and Co, as atomic percentages, is such that the sum of the amounts
of W, Mo, Cr, and Co is 30% or more and 34% or less, and such that
the sum of the amounts of W and Mo is 3% or more and 4.5% or
less.
8. An alloy according to claim 4, wherein the amounts of W, Mo, Cr,
and Co, as atomic percentages, is such that the sum of the amounts
of W, Mo, Cr, and Co is 30% or more and 34% or less, and such that
the sum of the amounts of W and Mo is 3% or more and 4.5% or
less.
9. A powder of an alloy according to claim 1.
10. A method of fabricating a part, wherein a blank of said part or
the part itself is produced from a powder of an alloy according to
claim 1, using a powder metallurgy technique.
11. A method of fabricating a part according to claim 10, wherein
said blank or said part undergoes a recrystallization heat
treatment in which the blank or the part is brought to a
temperature which is higher than the solvus temperature of the
gamma-prime phase of said alloy and lower than the melting onset
temperature for said alloy.
12. A method of fabricating a part according to claim 10, wherein
said blank or said part undergoes a recrystallization heat
treatment in which the blank or the part is brought to a
temperature which is lower than the solvus temperature of the
gamma-prime phase of said alloy.
13. A method of fabricating a part, wherein a blank of said part or
the part itself is produced from a powder of an alloy according to
claim 8, using a powder metallurgy technique.
14. A method of fabricating a part according to claim 13, wherein
said blank or said part undergoes a recrystallization heat
treatment in which the blank or the part is brought to a
temperature which is higher than the solvus temperature of the
gamma-prime phase of said alloy and lower than the melting onset
temperature for said alloy.
15. A method of fabricating a part according to claim 13, wherein
said blank or said part undergoes a recrystallization heat
treatment in which the blank or the part is brought to a
temperature which is lower than the solvus temperature of the
gamma-prime phase of said alloy.
16. A turbo-machine part produced from an alloy according to claim
1.
17. A turbo-machine part according to claim 16, having a
coarse-grained structure in the zone in which it is subjected to
the highest operational temperatures and where creep plays a
significant role in damage to the part, and a small-grained
structure in the zone in which it is subjected to the lowest
operational temperatures and where damage essentially results from
tensile forces and cyclic stresses.
18. A turbo-machine part according to claim 17, consisting of a
compressor or turbine disk.
19. A turbo-machine part produced from an alloy according to claim
8.
20. A turbo-machine part according to claim 19, having a
coarse-grained structure in the zone in which it is subjected to
the highest operational temperatures and where creep plays a
significant role in damage to the part, and a small-grained
structure in the zone in which it is subjected to the lowest
operational temperatures and where damage essentially results from
tensile forces and cyclic stresses.
21. A turbo-machine part according to claim 20, consisting of a
compressor or turbine disk.
Description
[0001] The invention relates to alloys, or superalloys, based on
nickel (Ni) and more particularly intended for the production of
compressor or turbine disks for turbo-machines using powder
metallurgy processes. The turbo-machines concerned may be
aeronautical (turbojet engine, turboprop engine) or ground-based
(gas turbine for the production of energy).
BACKGROUND OF THE INVENTION
[0002] In service, compressor and turbine disks, located
respectively upstream and downstream of the combustion chamber of a
turbojet engine, are subjected to mechanical stresses that can be
attributed to tension, creep, and fatigue, at temperatures that can
reach 800.degree. C. Said disks should nevertheless have
operational service lives of several thousand hours. Thus, said
disks must be produced from an alloy which, at high temperatures,
has high resistance to traction forces, very good creep strength,
and crack propagation resistance.
[0003] Currently, said disks can be produced from nickel-based
alloys using powder metallurgy processes, said processes limiting
chemical segregation phenomena and encouraging good microstructural
homogeneity of the alloy.
[0004] One known example of a nickel-based alloy is described in
French document FR-A-2 593 830. Said alloy is sold with reference
number N18.
[0005] That example of an alloy, along with the alloys of the
invention, falls into the category of two-phase alloys that
comprise: a phase termed the gamma phase formed by a nickel-based
solid solution that constitutes the matrix for the metallurgy
grains, and a phase termed the gamma-prime phase, of structure that
is based on the coherent intermetallic compound Ni.sub.3Al. The
gamma-prime phase forms several populations of inter- or
intra-granular precipitates that appear at different stages of the
thermomechanical history of the alloy and that play distinct roles
in the mechanical behavior of the alloy.
[0006] It has been shown that the inter-granular precipitate
population limits the growth of gamma matrix grains during
recrystallization heat treatment. Hence, by adjusting the
recrystallization heat treatment of the alloy, the inter-granular
precipitate population, and thus the size of said grains, are
controlled. Depending on whether the maximum temperature reached
during said heat treatment is higher (supersolvus treatment) or
lower (subsolvus treatment) than the solution temperature (or
solvus temperature) of the inter-granular precipitates of the
gamma-prime phase, recrystallization finishes with a large grain
size (for a supersolvus treatment) or a low grain size (for
subsolvus treatment).
[0007] Tensile strength is generally favored by a reduction in
grain size, while creep strength is favored by an increase. Hence,
depending on the envisaged application and the envisaged mechanical
characteristics, two-phase alloys are thermomechanically treated to
produce either a fine-grained microstructure (small grains), i.e.
with a grain size of the order of 5 .mu.m [micrometer] to 15 .mu.m
(i.e. ASTM [American Society for Testing and Materials] indices 12
to 9), or a microstructure with coarse grains, i.e. with a grain
size of the order of 20 .mu.m to 180 .mu.m (i.e. ASTM indices 8 to
2).
[0008] Further, the grain strength is ensured by the presence of
different populations of intra-granular precipitates of the
gamma-prime Ni.sub.3Al base phase and it is generally accepted that
the high temperature tensile strength of said alloys increases with
the volume fraction of the gamma-prime phase, said fraction
possibly reaching 60%.
[0009] The N18 alloy, with a volume fraction of the gamma-prime
phase of about 55%, principally undergoes subsolvus treatments
since a fine-grained microstructure is desirable. The fatigue
strength and tensile strength of said alloy are generally favored
over its creep strength, because the service temperature is often
less than 650.degree. C., i.e. relatively moderate.
[0010] At temperatures of more than 650.degree. C., high creep
strength is necessary and, as a result, a coarse-grained
microstructure (obtained by supersolvus treatment) would be better
suited. However, carrying out an industrial scale supersolvus
treatment on large diameter disks of N18 alloy is very difficult or
even impossible because the difference between the solvus
temperature of the gamma-prime phase and the melting temperature
(i.e. the onset of melting) of the alloy is too small. This
temperature range for solution of the gamma-prime phase (i.e. to
carry out a supersolvus treatment) is too narrow (less than
30.degree. C.), which renders industrial application of the total
gamma-prime phase solution heat treatment uncertain.
[0011] Further, high internal stresses arise in the disks during
rapid cooling (of the order of 100.degree. C./min) consecutive to
total solution heat treatment, and they cause cracks (quench
cracks) to appear.
OBJECT AND SUMMARY OF THE INVENTION
[0012] The invention aims to provide Ni-based alloys for which it
is possible to carry out not only a subsolvus treatment, but also a
supersolvus treatment on an industrial scale and which preferably
has high-temperature mechanical characteristics, especially creep
strength, that are at least equivalent to, and preferably better
than those of N18 alloy.
[0013] This is achieved by alloys that essentially comprise (i.e.
apart from any impurities) the following elements, in the amounts
indicated as percentages by weight: [0014] chromium (Cr): 11.5% to
13.5%; [0015] cobalt (Co): 11.5% to 16.0%; [0016] molybdenum (Mo):
3.4% to 5.0%; [0017] tungsten (W): 3.0% to 5.0%; [0018] aluminum
(Al): 2.2% to 3.2%; [0019] titanium (Ti): 3.5% to 5.0%; [0020]
niobium (Nb): 0.5% to 2.0%; [0021] hafnium (Hf): 0.25% to 0.35%;
[0022] zirconium (Zr): 0 to 0.07%; [0023] carbon (C): 0.015% to
0.030%; [0024] boron (B): 0.01% to 0.02%; and [0025] nickel (Ni):
complement to 100%.
[0026] The Applicant's research that led to the invention shows
that the problems encountered with N18 alloy are linked in part to
the high volume fraction (55%) of the gamma-prime phase in that
alloy.
[0027] In fact, the Applicant has shown firstly, that said high
volume fraction tends to reduce the difference between the solvus
temperature of the gamma-prime phase and the melting temperature of
the N18 alloy, rendering that difference too small to carry out a
supersolvus treatment on an industrial scale.
[0028] Secondly, the Applicant has shown that the internal stresses
arising in the part during rapid cooling consecutive upon total
solution heat treatment result in part from precipitation of a high
volume fraction of the gamma-prime phase.
[0029] Finally, the Applicant has shown that when the temperature
is held at over 650.degree. C. for a sufficiently long period, the
elemental composition of the N18 alloy causes the development of
topologically compact phases, generally denoted sigma and mu
phases, which are deleterious to the high temperature behavior of a
disk in operation.
[0030] Thus, the composition of the alloys of the invention is
selected so as to cause a limited volume fraction of gamma-prime
phase to precipitate.
[0031] While the alloys of the invention are less rich than N18
alloy in gamma-prime phase, against all expectations, their
small-grained microstructure has tensile and creep strength
characteristics that are better than those of the reference alloy.
It also appears that these alloys have equivalent fatigue-creep
crack propagation rates which are equivalent to or even better than
those of N18 alloy.
[0032] For turbo-machine compressor disks or turbine disks, high
tensile strength is particularly favorable to the rupture behavior
of said disks as may occur during accidental overspeeding. This
high strength is also an indicator of good oligocyclic fatigue
properties and adequate service lives.
[0033] Further, the reduction in the volume fraction of the
gamma-prime phase relative to the N18 alloy is favorable to the
production of disks having a coarse-grained microstructure and thus
high creep strength at high temperature (i.e. for temperatures of
700.degree. C. or more). This creep strength associated with very
good tensile and fatigue-creep crack propagation properties allows
these disks to be used at temperatures that are higher than in
current turbo-machines, providing access to better thermal
efficiencies and a reduction in the specific consumption of the
turbo-machines.
[0034] Production of said coarse-grained microstructure is further
facilitated by the comfortable range of temperatures between the
solvus temperature of the gamma-prime phase and the melting onset
temperature for the alloy. Advantageously, the compositions of the
alloys of the invention are such that this range spans 35.degree.
C. or more. This means that heat treatments above the solvus
temperature can be carried out on an industrial scale, without
risking melting the alloy.
[0035] The capability of developing one or the other
microstructure, coarse-grained and small-grained, as well as the
good mechanical properties corresponding to each of said
microstructures, is a distinct advantage in alloys of the invention
compared with those in current use, especially N18 alloy.
[0036] Further, this capability allows dual-structured disks to be
produced. By carrying out heat treatment at a temperature gradient,
a coarse-grained structure is developed in the peripheral zone of
the disk where the service temperatures are the highest and where
creep plays a significant role in material damage, and a small
grain structure is developed in the central zone of the disk (close
to the hub), which is cooler, where damage essentially results from
traction forces and cyclic stresses.
[0037] Despite an aluminum concentration that is lower than that of
the N18 alloy (which is directly correlated to a smaller volume
fraction of gamma-prime phase), the alloys of the invention have
relatively low density, preferably 8.3 kg/dm.sup.3 [kilograms/cubic
decimeter] or less, which means that the mass of the disk and
stresses resulting from centrifugal force are limited.
[0038] Finally, the elemental compositions of alloys of the
invention provide them with good microstructural stability as
regards the appearance of sigma and mu phases, which is retarded to
more than 500 hours maintained at 750.degree. C.
[0039] To limit the risk of quench cracking, in particular during
treatments at a temperature that is higher than the solvus
temperature of the gamma-prime phase, the compositions of the
alloys of the invention have a limited gamma-prime phase volume
fraction, preferably of 50% or less. Sufficient gamma-prime phase
must nevertheless be present, so the gamma-prime phase volume
fraction is preferably in the range 40% to 50%.
[0040] Advantageously, to obtain said volume fraction of the
gamma-prime phase in alloys of the invention, the sum of the Al,
Ti, and Nb contents, as atomic percentages, is 10.5% or more, and
13% or less, i.e. 10.5%.ltoreq.Al+Ti+Nb.ltoreq.13%.
[0041] Although precipitation of the gamma-prime phase in Ni-based
alloys occurs exclusively due to the presence of Al in sufficient
concentration, the elements Ti and Nb which, by being substituted
for Al, are constituents of that phase, are considered to be
elements that are favorable to the formation of the gamma-prime
phase in the same amount and they are termed gamma-prime-genic. The
value of the volume fraction of the gamma-prime phase is thus a
function of the sum of the atomic concentrations of Al, Ti, and
Nb.
[0042] It should be noted that tantalum (Ta) is also a
gamma-prime-genic element, but it does not appear in the
composition of the alloys of the invention. Ta is a high atomic
mass element, which means that complex compositional adjustments
have to be made to maintain the density of the alloy within
reasonable limits (preferably 8.3 kg/dm.sup.3 or less). Further, Ta
is expensive and it has not been possible to establish clearly that
it has any beneficial role in crack resistance. Finally, its
strengthening effect on the gamma-prime phase does not appear to be
greater than that of the elements Ti and Nb. It has even been shown
that the strength of the alloys of the invention is at least
equivalent to that of alloys containing Ta.
[0043] Also advantageously, the amounts of Al, Ti, and Nb, as an
atomic percentage in the alloys of the invention, are such that the
ratio between the sum of the amounts of Ti and Nb and the amount of
Al is 0.9 or more and 1.1 or less, i.e. 0.9 ([(Ti+Nb)/Al] (1.1.
[0044] The Ti and Nb atoms substituting for Al in the gamma-prime
phase Ni3Al base strengthen it by mechanisms analogous to those of
solid solution hardening. Said hardening is greater as the ratio
[(Ti+Nb)/Al] rises. However, beyond a certain value of the
concentration of Ti, the coherent Ni3Ti eta phase precipitates in
the form of elongate plates that have a deleterious effect on the
mechanical behavior, especially on the ductility, of alloys
containing it. Further, the concentration of Nb must be limited,
since an excessive Nb content is prejudicial to the crack
propagation resistance in this type of alloy.
[0045] In accordance with a further aspect of the invention, the
amounts of W, Mo, Cr, and Co, as an atomic percentage, are such
that the sum of the amounts of W, Mo, Cr, and Co is 30% or more and
34% or less, and such that the sum of the amounts of W and Mo is 3%
or more and 4.5% or less, i.e.: 30%.ltoreq.W+Mo+Cr+Co .ltoreq.34%;
and 3%.ltoreq.W+Mo .ltoreq.4.5%.
[0046] The elements which essentially substitute for Ni in the
gamma solid solution are Cr, Co, Mo, and W.
[0047] Cr is essential for oxidation and corrosion properties of
the alloy, and it participates in hardening the gamma matrix by the
solid solution effect.
[0048] Co improves the high-temperature creep strength of these
alloys. Further, an increase in the concentration of Co within the
stability limits of the structure of the gamma phase can reduce the
solvus temperature of the gamma-prime phase and hence facilitate
carrying out the partial or complete solution heat treatments
thereof.
[0049] Mo and W greatly harden the gamma matrix by the solid
solution effect. However, those elements have high atomic masses
and their substitution for Ni (in particular substitution of W for
Ni) results in a substantial increase in the density of the
alloy.
[0050] The amounts of Cr, Mo, Co, and W in the alloys of the
invention must thus be carefully adjusted relative to one another
in order to obtain the desired effects, in particular optimum
hardening of the gamma matrix, without in any way risking causing
the premature appearance of fragile intermetallic compound phases,
namely sigma and mu. Said phases, when they develop in excessive
quantities, can cause a significant reduction in the ductility and
mechanical strength of the alloys.
[0051] Finally, it should be noted that the minor elements, which
are C, B, and Zr, form segregations principally at the grain
boundaries, for example in the form of carbides or borides. They
thus contribute to increasing the strength and ductility of alloys
by modifying the chemistry of the grain boundaries, and their
absence would be prejudicial. However, an excess of those elements
causes a reduction in the temperature of melting onset and causes
excessive precipitation of carbides and borides, which consume the
elements of the alloy and which no longer participate in hardening
the alloy. The concentrations of carbon, boron, and zircon are thus
adjusted, in particular with non-zero minimum amounts of carbon and
boron, so as to obtain good high-temperature strength and optimum
ductility for alloys of the invention. Hf is also present in
moderate quantities, since that element improves the
high-temperature inter-granular cracking resistance.
[0052] The invention also provides a method of fabricating a part,
more particularly a turbo-machine part such as a compressor or
turbine disk, wherein a blank of said part or the part itself is
produced from a powder of an alloy of the invention, using a powder
metallurgy technique.
[0053] Advantageously, said blank or said part undergoes
recrystallization heat treatment during which the blank or part is
brought either to a temperature that is below the solvus
temperature of the gamma-prime phase of said alloy or to a
temperature that is above the solvus temperature of the gamma-prime
phase of said alloy, and lower than the melting onset temperature
of said alloy, to encourage the development of a microstructure
with a grain size which is adapted to the stress conditions.
BRIEF DESCRIPTION OF THE DRAWING
[0054] The invention, its applications and its advantages can be
better understood from the following detailed description. Said
description makes reference to the accompanying figures in
which:
[0055] FIG. 1 is a scanning electron microscope image showing the
microstructure of alloy A, described below; and
[0056] FIG. 2 is a scanning electron microscope image showing the
microstructure of alloy C1, described below.
MORE DETAILED DESCRIPTION
[0057] The parts produced from the alloys of the invention are
preferably fabricated using powder metallurgy techniques.
[0058] As an example, production of a compressor or turbine disk
using a powder metallurgy technique comprises the following steps:
[0059] fabricating a master alloy ingot by mixing and melting
metallic elements that are pure (apart from any impurities); [0060]
re-melting the ingot and pulverizing it with an inert gas or
remelting the ingot and centrifugal pulverization using a known
rotating electrode technique, to obtain a pre-alloyed powder;
[0061] consolidating said pre-alloyed powder by hot isostatic
pressing and/or by drawing; [0062] forming a disk blank by
isothermal forging; [0063] heat treating said blank; and [0064]
Final machining of the disk.
[0065] At the end of the isothermal forging, different heat
treatment steps may be selected to obtain the microstructure which
is best suited to the envisaged application. The temperature of the
gamma-prime phase solution heat treatment allows the metallurgy
grain size to be controlled: [0066] with a treatment at a
temperature which is below the solvus temperature of the
gamma-prime phase, to obtain a microstructure with small grains (5
.mu.m to 15 .mu.m); and [0067] with a treatment at a temperature in
the range between the solvus temperature of the gamma-prime phase
and the melting onset temperature of the alloy, to obtain a
coarse-grained microstructure (more than 15 .mu.m). Said final
treatment can be carried out industrially only if the difference
between the two said temperatures, termed the "solution window", is
sufficiently large: for industrial alloys, it is assumed that it
must be more than 30.degree. C., preferably more than 35.degree.
C.
[0068] The cooling rate which follows the solution treatment can
control the distribution of intra-granular precipitates of
gamma-prime phase.
[0069] One or more tempering treatments can control the size of the
tertiary precipitates of gamma-prime phase and relax internal
stresses which result from quenching.
[0070] The nominal compositions of two prior art alloys and three
alloys of the invention, given by way of examples, are shown in
Table I in which the amounts of the elements of each alloy are
shown as atomic percentages, and in Table II in which the amounts
are shown as percentages by weight. Alloys C1, C2 and C3 have a
solution window of more than 50.degree. C. and are thus treated
using the two types of heat treatment presented above, which
provides a great range of microstructures. TABLE-US-00001 TABLE I
Alloy Co Cr Mo W Al Ti Nb Hf C B Zr A 15.0 12.5 3.8 0 9.2 5.3 0
0.125 0.079 0.083 0.022 B 12.9 18.1 2.4 1.3 4.6 4.5 0.4 0 0.190
0.077 0.027 C1 15.1 13.6 2.2 1.3 6.4 5.6 0.5 0.100 0.109 0.093 0 C2
15.4 14.1 2.5 1.5 6.0 5.0 1.0 0.093 0.128 0.080 0 C3 12.0 14.6 2.9
1.0 5.5 4.6 1.0 0.100 0.100 0.080 0.038 (amounts shown as atomic
percentages)
[0071] TABLE-US-00002 TABLE II Alloy Co Cr Mo W Al Ti Nb Hf C B Zr
A 15.9 11.7 6.6 0 4.4 4.5 0 0.400 0.017 0.016 0.036 B 13.1 16.2 4.0
4.0 2.2 3.7 0.7 0 0.039 0.014 0.043 C1 15.4 12.2 3.7 4.0 3.0 4.6
0.8 0.310 0.023 0.018 0 C2 15.5 12.6 4.1 4.7 2.8 4.1 1.5 0.285
0.026 0.015 0 C3 12.15 13.0 4.8 3.15 2.55 3.8 1.6 0.310 0.021 0.015
0.060 (amounts shown as percentages by weight)
[0072] Alloy A is alloy N18 and alloy B is sold with reference
number Rene-88DT.
[0073] To carry out tests on these alloys, parts were produced by
powder metallurgy using the following procedure: [0074] fabricating
master alloy ingots by mixing and fusing pure metallic elements;
[0075] centrifugal spraying with rotating electrodes; [0076]
consolidating pre-alloyed powders by hot drawing; [0077] heat
treatment including a subsolvus or supersolvus treatment.
[0078] For the subsolvus treatment, a partial solution treatment
for the gamma-prime phase was carried out at a temperature below
the solvus temperature (Tsolvus) of the gamma-prime phase (at about
Tsolvus -25.degree. C.). The rate of cooling was of the order of
100.degree. C./minute after solution. This treatment was followed
by tempering for 24 hours at 750.degree. C. and air cooling.
[0079] For the supersolvus treatment, a total gamma-prime phase
solution treatment was carried out at a temperature above the
gamma-prime solvus temperature (at about Tsolvus +15.degree. C. to
20.degree. C.). The rate of cooling was of the order of 140.degree.
C./min after solution. Said treatment was followed by tempering for
8 hours at 760.degree. C. and air cooling.
[0080] Tables III and IV show some results of mechanical tests
carried out in tension, creep, and crack propagation respectively
for alloys which received a subsolvus treatment (Table III) and a
supersolvus treatment (Table IV).
[0081] The tensile tests were carried out in air at 650.degree. C.
for the subsolvus treatment (Table III) and at 700.degree. C. for
the supersolvus treatment (Table IV), and Rm corresponds to the
maximum stress measured during these tests.
[0082] The creep tests were carried out in air at 700.degree. C. at
an initial stress of 550 MPa (650 MPa [megapascal] for alloy C1).
The parameter t.sub.0.2% is the time in hours to reach a plastic
deformation of 0.2%.
[0083] The crack propagation tests were carried out in air and at
650.degree. C. The stress cycle was as follows: load ramp-up for 10
seconds, hold for 300 seconds at maximum load and release in 10
seconds with a load ratio (minimum load/maximum load) of 0.05. The
parameter V.sub.f35 is the crack propagation rate, measured at a
value of delta K of 35 MPam.sup.1/2. TABLE-US-00003 TABLE III
Tension at 700.degree. C. Creep at 700.degree. C., Crack
propagation at Alloy Rm (MPa) 550 MPa t.sub.0.2% (h) 650.degree. C.
V.sub.f35 (m/cycle) A 1474 340 12.10.sup.-5 B 1445 610 3.10.sup.-5
C1 1590 3000* 2.10.sup.-5 C2 1635 2300 3.10.sup.-5 C3 1589 -- --
*under initial stress of 650 MPa
[0084] TABLE-US-00004 TABLE IV Tension at 700.degree. C. Creep at
700.degree. C., Crack propagation at Alloy Rm (MPa) 550 MPa
t.sub.0.2% (h) 650.degree. C. V.sub.f35 (m/cycle) B 1320 150
9.10.sup.-6 C1 1440 1750* 3.10.sup.-6 C2 1428 >3000 5.10.sup.-6
*under initial stress of 650 MPa
[0085] The results of Tables III and IV show that the alloys of the
invention can produce a large increase in the high-temperature
mechanical properties (tension and creep) while keeping the crack
propagation resistance identical to or better than known
alloys.
[0086] Referring to FIGS. 1 and 2, micro structural examinations
were carried out on alloys A and C1 which had undergone a subsolvus
treatment, to detect the appearance of topologically compact phases
(i.e. fragile intermetallic compounds) after an ageing heat
treatment of 500 hours at 750.degree. C. The observations were
carried out by back-diffused electron scanning electron microscopy
on non-attacked specimens. In alloy A, severe ageing of 500 hours
at 750.degree. C. caused inter- and intra-granular formation of
phases rich in heavy elements. These phases show up in clear
contrast (white borders) at the grain boundaries in FIG. 1. These
phases, when formed in excessive quantities, may cause a
significant reduction in the ductility and strength of the alloys.
Tests on alloy C1 which had undergone the same treatment of 500
hours at 750.degree. C. showed that said phases were not formed
during ageing. The alloys of the invention were thus more stable
than alloy A (N18) as regards the formation of fragile
intermetallic compounds, which are topologically compact
phases.
* * * * *