US20070227630A1

US20070227630A1 - Nickel-based alloy

Info

Publication number: US20070227630A1
Application number: US11/694,204
Authority: US
Inventors: Isabelle Augustins Lecallier; Pierre Caron; Jean-yves Guedou; Didier Locq; Loeiz Naze
Original assignee: Association pour la Recherche et le Developpement des Methodes et Processus Industriels; Office National dEtudes et de Recherches Aerospatiales ONERA; SNECMA SAS
Current assignee: Association pour la Recherche et le Developpement des Methodes et Processus Industriels; Office National dEtudes et de Recherches Aerospatiales ONERA; Safran Aircraft Engines SAS
Priority date: 2006-03-31
Filing date: 2007-03-30
Publication date: 2007-10-04
Also published as: JP5398123B2; RU2007111861A; FR2899240A1; CA2583140C; EP1840232B1; FR2899240B1; EP1840232A1; DE602007001092D1; CA2583140A1; JP2007277721A; RU2433197C2

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.

Description

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

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° 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.
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.
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.
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 Ni₃Al. 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.
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).
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 μm [micrometer] to 15 μm (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 μm to 180 μm (i.e. ASTM indices 8 to 2).
Further, the grain strength is ensured by the presence of different populations of intra-granular precipitates of the gamma-prime Ni₃Al 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%.
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° C., i.e. relatively moderate.
At temperatures of more than 650° 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° C.), which renders industrial application of the total gamma-prime phase solution heat treatment uncertain.
Further, high internal stresses arise in the disks during rapid cooling (of the order of 100° C./min) consecutive to total solution heat treatment, and they cause cracks (quench cracks) to appear.

OBJECT AND SUMMARY OF THE INVENTION

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.
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:

- chromium (Cr): 11.5% to 13.5%;
- cobalt (Co): 11.5% to 16.0%;
- molybdenum (Mo): 3.4% to 5.0%;
- tungsten (W): 3.0% to 5.0%;
- aluminum (Al): 2.2% to 3.2%;
- titanium (Ti): 3.5% to 5.0%;
- niobium (Nb): 0.5% to 2.0%;
- hafnium (Hf): 0.25% to 0.35%;
- zirconium (Zr): 0 to 0.07%;
- carbon (C): 0.015% to 0.030%;
- boron (B): 0.01% to 0.02%; and
- nickel (Ni): complement to 100%.

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.
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.
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.
Finally, the Applicant has shown that when the temperature is held at over 650° 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.
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.
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.
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.
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° 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.
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° 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.
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.
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.
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/dm³[kilograms/cubic decimeter] or less, which means that the mass of the disk and stresses resulting from centrifugal force are limited.
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° C.
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%.
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%≦Al+Ti+Nb≦13%.
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.
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/dm³or 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.
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.
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.
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%≦W+Mo+Cr+Co ≦34%; and 3%≦W+Mo ≦4.5%.
The elements which essentially substitute for Ni in the gamma solid solution are Cr, Co, Mo, and W.
Cr is essential for oxidation and corrosion properties of the alloy, and it participates in hardening the gamma matrix by the solid solution effect.
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.
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.
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.
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.
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.
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

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:
FIG. 1 is a scanning electron microscope image showing the microstructure of alloy A, described below; and
FIG. 2 is a scanning electron microscope image showing the microstructure of alloy C1, described below.

MORE DETAILED DESCRIPTION

The parts produced from the alloys of the invention are preferably fabricated using powder metallurgy techniques.
As an example, production of a compressor or turbine disk using a powder metallurgy technique comprises the following steps:

- fabricating a master alloy ingot by mixing and melting metallic elements that are pure (apart from any impurities);
- 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;
- consolidating said pre-alloyed powder by hot isostatic pressing and/or by drawing;
- forming a disk blank by isothermal forging;
- heat treating said blank; and
- Final machining of the disk.

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:

- 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 μm to 15 μm); and
- 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 μm). 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° C., preferably more than 35° C.

The cooling rate which follows the solution treatment can control the distribution of intra-granular precipitates of gamma-prime phase.
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.

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° C. and are thus treated using the two types of heat treatment presented above, which provides a great range of microstructures.

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)

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)

Alloy A is alloy N18 and alloy B is sold with reference number René-88DT.
To carry out tests on these alloys, parts were produced by powder metallurgy using the following procedure:

- fabricating master alloy ingots by mixing and fusing pure metallic elements;
- centrifugal spraying with rotating electrodes;
- consolidating pre-alloyed powders by hot drawing;
- heat treatment including a subsolvus or supersolvus treatment.

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° C.). The rate of cooling was of the order of 100° C./minute after solution. This treatment was followed by tempering for 24 hours at 750° C. and air cooling.
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° C. to 20° C.). The rate of cooling was of the order of 140° C./min after solution. Said treatment was followed by tempering for 8 hours at 760° C. and air cooling.
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).
The tensile tests were carried out in air at 650° C. for the subsolvus treatment (Table III) and at 700° C. for the supersolvus treatment (Table IV), and Rm corresponds to the maximum stress measured during these tests.
The creep tests were carried out in air at 700° C. at an initial stress of 550 MPa (650 MPa [megapascal] for alloy C1). The parameter t_0.2%is the time in hours to reach a plastic deformation of 0.2%.
The crack propagation tests were carried out in air and at 650° 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_f35is the crack propagation rate, measured at a value of delta K of 35 MPa·m^1/2.

TABLE III

Tension at 700° C. Creep at 700° C., Crack propagation at

Alloy Rm (MPa) 550 MPa t_0.2%(h) 650° C. V_f35(m/cycle)

A 1474 340 12.10⁻⁵

B 1445 610 3.10⁻⁵

C1 1590 3000* 2.10⁻⁵

C2 1635 2300 3.10⁻⁵

C3 1589 — —

*under initial stress of 650 MPa

TABLE IV


	Tension at 700° C.	Creep at 700° C.,	Crack propagation at
Alloy	Rm (MPa)	550 MPa t_0.2%(h)	650° C. V_f35(m/cycle)

B	1320	150	9.10⁻⁶
C1	1440	1750*	3.10⁻⁶
C2	1428	>3000	5.10⁻⁶

*under initial stress of 650 MPa

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.
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° 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° 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° 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.

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.