EP1840232A1 - Nickel-based alloy - Google Patents

Abstract

The invention relates to alloys, or superalloys, based on nickel (Ni) essentially comprising the following elements in the contents indicated in 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: 100% complement. Use for the production of turbine disks or turbine engine compressor, according to powder metallurgy processes.

Description

The invention relates to alloys, or superalloys, based on nickel (Ni) intended more particularly for the production of turbine disks or turbine engine compressor, according to powder metallurgy processes. The turbomachines concerned can be aeronautical (turbojet engine, turboprop engine) or terrestrial (gas turbine dedicated to the production of energy).

The compressor and turbine disks respectively located upstream and downstream of the combustion chamber of a turbojet engine are subjected in service to mechanical stresses comparable to traction, creep and fatigue, at temperatures up to 800 ° C. However, it is desired that the operating lifetimes of these disks reach several thousand hours. These discs must therefore be made of an alloy having, at high temperatures, a high resistance to tensile stresses, a very good creep resistance, as well as a good resistance to the propagation of cracks.

Today, these discs can be made of Ni-based alloys according to powder metallurgy processes, these processes limiting chemical segregation phenomena and promoting good microstructural homogeneity of the alloy.

An example of a known nickel-based alloy is described in the document FR2593830 . This alloy is marketed under the reference N18.

This alloy example, as well as the alloys of the invention, is part of the two-phase alloys which comprise: a so-called gamma phase formed by a nickel-based solid solution, which constitutes the matrix of the metallurgical grains, and a phase called gamma prime, whose structure is based on the Ni ₃ Al ordered intermetallic compound. 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 have different roles. distinct in the mechanical behavior of the alloy.

It has been found that the intergranular precipitate population limits the growth of the gamma matrix grains during a recrystallization heat treatment. Thus, by modifying the heat treatment of recrystallization of the alloy, the population of intergranular precipitates and thus the size of said grains is controlled. Depending on whether the maximum temperature reached during this heat treatment is greater (treatment said supersolvus) or lower (so-called subsolvus treatment) at the dissolution temperature (or solvus temperature) of the intergranular precipitates of gamma prime phase, the recrystallization results in a high grain size (for a supersolvus treatment) or a small grain size ( for a subsolvus treatment).

The tensile strength is generally favored by the reduction of the grain size while the resistance to creep is favored by an increase in it. Thus, according to the intended application and the mechanical characteristics referred to, the two-phase alloys are thermomechanically treated to present either a fine grain microstructure (small grains), that is to say having a grain size of about 5 at 15 .mu.m (that is to say, indices ASTM 12 to 9, according to the "American Society for Testing and Material" standard), ie a coarse-grained microstructure, that is to say having a size of grain of the order of 20 to 180 microns (that is to say, ASTM indices 8 to 2).

Moreover, the grain reinforcement is ensured by the presence of different populations of intra-granular precipitates of the gamma prime Ni ₃ Al base phase and it is generally accepted that the tensile strength, when hot, of these alloys increases with the volume fraction of the gamma prime phase, this fraction up to 60%.

The N18 alloy, whose volume fraction of gamma prime phase is about 55%, is mainly subjected to subsolvus treatments, because it is desired to obtain a fine-grained microstructure. Indeed, the fatigue and tensile strength of this alloy is generally preferred over its creep resistance, due to a temperature of use often less than 650 ° C, that is to say relatively moderate.

At temperatures above 650 ° C, high creep resistance is required and, therefore, a coarse-grained microstructure (obtained by supersolvus treatment) would be more suitable. However, a supersolvus treatment on N18 alloy large diameter discs is very difficult or impossible to achieve industrially because of the too small difference between the solvus temperature of the gamma prime phase and the burn temperature (this is ie beginning of melting) of the alloy. This temperature range for the dissolution of the gamma prime phase (that is to say to achieve a supersolvus treatment) is indeed too narrow (lower than 30 ° C) which makes hazardous the industrial practice of the heat treatment of total dissolution of the gamma prime phase.

In addition, high internal stresses arise in the disks, during rapid cooling (of the order of 100 ° C / min) consecutive to the heat treatment of total solution dissolution, and cause the appearance of cracks (quenching taps) .

The purpose of the invention is to propose Ni-based alloys for which it is possible to carry out a subsolvus treatment, but also a supersolvus treatment on an industrial scale and which preferably have high temperature mechanical characteristics, in particular a creep resistance, at least equivalent, and preferably greater than that of the alloy N18.

• chrome (Cr) : 11,5 à 13,5 %;
• cobalt (Co) : 11,5 à 16,0 %;
• molybdène (Mo) : 3,4 à 5,0 %;
• tungstène (W) : 3,0 à 5,0 %;
• aluminium (Al) : 2,2 à 3,2 %;
• titane (Ti): 3,5 à 5,0 %;
• niobium (Nb): 0,5 à 2,0 %;
• hafnium (Hf) : 0,25 à 0,35 %;
• zirconium (Zr) : 0 à 0,07 %;
• carbone (C) : 0,015 à 0,030 %;
• bore (B) : 0,01 à 0,02 %; et
• nickel (Ni) : complément à 100 %.

To achieve this object, the subject of the invention is alloys characterized in that they essentially comprise (that is to say, any impurities) the following elements, in the contents indicated in 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): 100% complement.

In his research leading to the invention, the Applicant has established that the problems encountered with the alloy N18 were related, in part, to the high volume fraction (55%) of gamma prime phase in this alloy.

Indeed, on the one hand, the plaintiff has established that this high volume fraction tends to reduce the gap between the solvation temperature of the gamma prime phase and the burn temperature of the N18 alloy, making this gap too narrow. to realize industrially a supersolvus treatment.

On the other hand, the applicant has established that the internal stresses arising in the room during the rapid cooling following the total solution heat treatment, resulted in part from the precipitation of a high volume fraction of gamma prime phase.

Finally, the Applicant has established that the elemental composition of alloy N18 allowed, during the maintenance of temperature at more than 650 ° C for sufficiently long periods, the development of topologically compact phases, generally referred to as sigma and mu phases, which are harmful. the high temperature behavior of a disc in operation.

Thus, the composition of the alloys of the invention is chosen so as to precipitate a limited volume fraction of gamma prime phase.

Although the alloys of the invention are thus less rich than the alloy N18 gamma prime phase, they have against all expectations, in their microstructural small grain, tensile characteristics and creep resistance greater than those of this reference alloy. It also appears that these alloys exhibit fatigue cracking rates equivalent to or even lower than those of alloy N18.

For turbomachine compressor or turbine discs, the high resistance to tensile stresses is particularly favorable to the bursting behavior of these discs that may occur during an accidental overspeed regime. This high resistance also makes it possible to anticipate good properties in oligocyclic fatigue and adequate lifetimes.

In addition, the reduction, relative to the N18 alloy, of the gamma prime phase volume fraction is favorable for producing discs having a coarse-grained microstructure and therefore a high resistance to creep at high temperature (ie ie for temperatures greater than or equal to 700 ° C). This creep resistance associated with very good mechanical characteristics in traction and in propagation of cracks in fatigue-creep, allows the use of these discs at higher temperatures than in the current turbomachines, which allows access to better thermal efficiency and to reduce the specific consumption of turbomachines.

The realization of this coarse-grained microstructure is further facilitated by the comfortable temperature range between the solvus temperature of the gamma prime phase and the melting start temperature of the alloy. Advantageously, the compositions of the alloys of the invention are such that the width of this range is greater than or equal to 35 ° C. This allows the industrial production of heat treatments beyond the solvus temperature, without the risk of burning of the alloy.

The ability to develop one or other of the large and small grains microstructures and the good mechanical properties corresponding to each of these two microstructures is a definite advantage of the alloys of the invention compared to those used for this purpose. day and, in particular, alloy N18.

In addition, this ability allows for dual structure disks. Indeed, by carrying out a heat treatment under a temperature gradient, a coarse grain structure is developed in the peripheral zone of the disk, where the temperatures in service are the highest and where creep plays a significant role in the damage of the material. , and a small-grain structure in the central zone of the disk (near the hub), colder, where the damage results essentially from the traction forces and cyclic stresses.

Despite an aluminum concentration lower than that of alloy N18 (which is directly correlated to a lower volume fraction of gamma prime phase), the alloys of the invention have a relatively low density, preferably less than or equal to 8. , 3 kg / dm ³ , which makes it possible to limit the mass of the disk and the stresses resulting from the centrifugal force.

Finally, the elemental compositions of the alloys of the invention provide them with good microstructural stability with regard to the appearance of the sigma and mu phases, which is delayed beyond 500 hours of maintenance at 750 ° C.

In order to limit the risks of quenching cracking, in particular during treatments at a temperature above the solvus temperature of the gamma prime phase, the compositions of the alloys of the invention have a limited gamma prime phase volume fraction and, preferably, less than or equal to 50%. The gamma prime phase must nevertheless be in sufficient proportion, the volume fraction of gamma prime phase is preferably between 40 and 50%.

Advantageously, to obtain such a volume fraction of gamma prime phase in the alloys of the invention, the sum of the contents of Al, Ti and Nb, in atomic percentages, is greater than or equal to 10.5% and less than or equal to 13%, ie: 10.5% ≤ Al + Ti + Nb ≤ 13%.

Although the precipitation of the gamma prime phase in Ni-based alloys depends exclusively on the presence of Al in sufficient concentration, the elements Ti and Nb, which, by substituting for Al, constitute this phase, are, like this one, considered as favorable elements for the formation of the gamma prime phase and are called gamma prime-genes. The value of the volume fraction of the gamma prime phase is therefore a function of the sum of the atomic concentrations of Al, Ti and Nb.

It will be noted that tantalum (Ta) is also a gamma primegen element but that it has not been retained in the composition of the alloys of the invention. Indeed, Ta is an element of high atomic mass, which necessitates complex composition adjustments in order to maintain the density of the alloy under a reasonable limit (preferably less than or equal to 8.3 kg / dm ³ ). In addition, the cost of Ta is high and its beneficial role on crack resistance could not be clearly established. Finally, its enhancement effect of the gamma prime phase does not appear to be greater than that of the elements Ti and Nb. It has even been found that the resistance of the alloys of the invention was at least equivalent to that of the alloys containing Ta.

Advantageously, the Al, Ti and Nb contents, in atomic percentages in the alloys of the invention, are such that the ratio between the sum of the contents of Ti and Nb, and the Al content, is greater than or equal to 0.9 and less than or equal to 1.1, ie: 0.9 ≤ [(Ti + Nb) / Al] ≤ 1.1.

The Al-substituted Ti and Nb atoms in the Ni ₃ Al base gamma prime phase have the effect of reinforcing it by mechanisms similar to those of solid solution hardening. This hardening is even higher than the ratio [(Ti + Nb) / Al] is high. However, on the one hand, beyond a certain value of the Ti concentration, the ordered phase eta Ni ₃ Ti precipitates in the form of elongated platelets which have a detrimental effect on the mechanical behavior, especially on the ductility of the alloys which contain them. On the other hand, the concentration of Nb must be limited because an excessive content of Nb is detrimental to the resistance to the propagation of cracks in this type of alloys.

According to another aspect of the invention, the contents in W, Mo, Cr and Co, in atomic percentages, are such that the sum of the contents in W, Mo, Cr and Co is greater than or equal to 30% and less than or equal to at 34%, and such that the sum of the contents in W and Mo is greater than or equal to 3% and less than or equal to 4.5%, namely: 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 the resistance to oxidation and corrosion of the alloy and participates, as a solid solution, in the hardening of the gamma matrix.

Co improves the resistance of these alloys to creep at high temperatures. In addition, the increase in Co concentration, within the limits of the stability of the gamma phase structure, makes it possible to lower the solvus temperature of the gamma prime phase and thus to facilitate the implementation of heat treatments. partial or total solution of solution thereof.

Mo and W provide a strong hardening of the gamma matrix by the effect of solid solution. However, these elements have high atomic masses and their substitution with Ni (in particular the substitution of W for Ni) results in a significant increase in the density of the alloy.

The contents of Cr, Mo, Co and W in the alloys of the invention must therefore be carefully adjusted with respect to each other to obtain the desired effects, in particular optimum curing of the gamma matrix, without thereby risking causing the premature appearance of the phases of fragile intermetallic compounds, sigma and mu. These phases, when they develop in excessive amounts, can indeed cause a significant reduction in the ductility and mechanical strength of the alloys.

Finally, it will be noted that the so-called minor elements C, B and Zr form segregations mainly at the grain boundaries, for example in the form of carbides or borides. They contribute to increase the strength and ductility of alloys by changing the chemistry of grain boundaries and their absence would be detrimental. However an excess content of these elements leads to a reduction of the melting temperature beginning and excessive precipitation of carbides and borides which consumes alloy elements that no longer participate in the hardening of the alloy. The concentrations of carbon, boron and zirconium are therefore adjusted, in particular with non-zero minimum levels of carbon and boron, in order to obtain at high temperature optimum mechanical strength and ductility for the alloys of the invention. Hf is also present in moderate amounts because this element improves resistance to intergranular hot cracking.

The subject of the invention is also a method for manufacturing a part, more particularly a turbomachine part such as a compressor or turbine disk, characterized in that a blank of said part is produced, or the part itself from a powder of an alloy according to the invention, by a powder metallurgy technique.

Advantageously, said blank or said part is subjected to a recrystallization heat treatment according to which the blank or part is brought, either at a temperature below the solvus temperature of the gamma prime phase of said alloy, or at a temperature of less than a temperature greater than the solvus temperature of the gamma prime phase of said alloy, and less than the melting start temperature of this alloy, so as to promote the development of a grain size microstructure adapted to the stress conditions.

• la figure 1 est un cliché obtenu par microscopie électronique à balayage montrant la microstructure de l'alliage A, décrit plus loin, et
• la figure 2 est un cliché obtenu par microscopie électronique à balayage montrant la microstructure de l'alliage C1, décrit plus loin.

The invention, its applications and its advantages will be better understood on reading the detailed description which follows. This description refers to the appended figures in which:

• FIG. 1 is a photograph obtained by scanning electron microscopy showing the microstructure of alloy A, described below, and
• Figure 2 is a photograph obtained by scanning electron microscopy showing the microstructure of alloy C1, described below.

The parts made from the alloys according to the invention are preferably manufactured by powder metallurgy techniques.

• la fabrication d'un lingot d'alliage mère par mélange et fusion d'éléments métalliques purs (aux éventuelles impuretés près),
• la refusion du lingot et sa pulvérisation par gaz inerte ou la refusion du lingot et sa pulvérisation centrifuge par une technique connue d'électrode tournante, afin d'obtenir une poudre pré-alliée,
• la consolidation de cette poudre pré-alliée par compactage isostatique à chaud et/ou par filage,
• la mise en forme d'une ébauche de disque par forgeage isotherme,
• le traitement thermique de cette ébauche, et
• l'usinage final du disque.

For example, the production of a compressor disk or turbine, according to a powder metallurgy technique, comprises the following steps:

• the manufacture of a parent alloy ingot by mixing and melting pure metal elements (with any impurities),
• the remelting of the ingot and its spraying by inert gas or the reflow of the ingot and its centrifugal spraying by a known rotary electrode technique, in order to obtain a pre-alloyed powder,
• the consolidation of this pre-alloyed powder by hot isostatic compaction and / or by spinning,
• the shaping of a disc blank by isothermal forging,
• the heat treatment of this blank, and
• the final machining of the disc.

• avec un traitement à une température inférieure à la température de solvus de la phase gamma prime, on obtient une microstructure à petits grains (5 à 15 µm), et
• avec un traitement à une température comprise entre la température de solvus de la phase gamma prime et la température de début de fusion de l'alliage, on obtient une microstructure à gros grains (supérieure à 15 µm). Ce dernier traitement n'est réalisable industriellement que si l'écart entre les deux températures précitées, appelé "fenêtre de mise en solution ", est suffisamment important : on considère pour les alliages industriels qu'il doit être supérieur à 30°C et, de préférence, à 35°C.

At the end of the isothermal forging, different heat treatment steps can be chosen so as to obtain the microstructure best suited to the intended application. The temperature of the dissolution treatment of the gamma prime phase makes it possible to control the metallurgical grain size:

• with a treatment at a temperature lower than the solvus temperature of the gamma prime phase, a microstructure with small grains (5 to 15 μm) is obtained, and
• with a treatment at a temperature between the solvus temperature of the gamma prime phase and the melting start temperature of the alloy, a coarse-grained microstructure (greater than 15 μm) is obtained. This last treatment is industrially feasible only if the difference between the two aforementioned temperatures, called the "dissolution window", is sufficiently important: it is considered for industrial alloys that it must be greater than 30 ° C and, preferably at 35 ° C.

The rate of cooling following the dissolution treatment makes it possible to control the distribution of the intragranular precipitates of gamma prime phase.

One or more treatments of income make it possible to control the size of the tertiary precipitates of phase gamma prime and to relax the internal stresses which result from the quenching.

The nominal compositions of two alloys of the prior art and of three alloys according to the invention, given by way of example, are presented in Table I in which the contents of the elements of each alloy are indicated in atomic percentages, and in Table II in which the contents are indicated in percentages by mass. Alloys C1, C2 and C3 have a dissolution window greater than 50 ° C and can therefore be treated according to the two types of heat treatment presented above, which allows a large latitude of microstructures. Table I (contents indicated in atomic percentages) Alloy Co Cr MB W al Ti Nb Hf VS B Zr AT 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 Alloy Co Cr MB W al Ti Nb Hf VS B Zr AT 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

The alloy A is the alloy N18 and the alloy B is sold under the reference Rene-88DT.

• fabrication de lingots d'alliage mère par mélange et fusion d'éléments métalliques purs,
• pulvérisation centrifuge avec électrodes tournantes,
• consolidation des poudres pré-alliées par filage à chaud,
• traitements thermiques comprenant un traitement subsolvus ou supersolvus.

To carry out tests on these alloys, parts have been developed by means of powder metallurgy according to the following procedure:

• manufacture of mother alloy ingots by mixing and melting pure metallic elements,
• centrifugal spraying with rotating electrodes,
• consolidation of the pre-alloyed powders by hot spinning,
• heat treatments comprising a subsolvus or supersolvus treatment.

For the subsolvus treatment, a partial solution treatment of the gamma prime phase was performed at a temperature below the solvus temperature (Tsolvus) of the gamma prime phase (at about Tsolvus - 25 ° C). The cooling rate was of the order of 100 ° C./min after dissolution. This treatment was followed by an income of 24 hours at 750 ° C and air cooling.

For the supersolvus treatment, a treatment of total dissolution of the gamma prime phase was performed at a temperature above the solvus gamma prime (at about Tsolvus + 15 at 20 ° C). The cooling rate was of the order of 140 ° C./min after the dissolution. This treatment was followed by an 8 hour income at 760 ° C and air cooling.

In Tables III and IV are presented some results of mechanical tests carried out in tension, in creep and in crack propagation, respectively for alloys having received a subsolvus treatment (Table III) and a supersolvus treatment (Table IV).

The tensile tests were conducted in air at 650 ° C for the subsolvus treatment (Table III) and 700 ° C for the supersolvus treatment (Table IV) and Rm is the maximum stress measured during these tests.

The creep tests were carried out in air at 700 ° C. under an initial stress of 550 MPa (650 MPa for the C1 alloy). The data t _0.2% is the holding time in hours to reach a plastic deformation of 0.2%.

Crack propagation tests were carried out in air and at 650 ° C. The load cycle is as follows: load increase in 10 seconds, hold time of 300 seconds at maximum load and discharge in 10 seconds with a load ratio (minimum load / maximum load) equal to 0.05. The data V _f35 is the crack propagation velocity, measured at a delta K value equal to 35 MPa.m ^1/2 . Table III - (* under an initial stress of 650 MPa) Alloy Traction at 650 ° C Rm (MPa) Creep at 700 ° C at 550 MPa t _0.2% (h) Crack propagation at 650 ° CV _f35 (m / cycle) AT 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 - - Alloy Traction at 700 ° C Rm (MPa) Creep at 700 ° C at 550 MPa t _0.2% (h) Crack propagation at 650 ° CV _f35 (m / cycle) B 1320 150 9.10 ^-6 C 1 1440 1750 * 3.10 ^-6 C2 1428 > 3000 5.10 ^-6

The results of Tables III and IV show that the alloys of the invention make it possible to achieve a significant increase in the mechanical properties when hot (traction and creep) while preserving resistance to crack propagation that is identical to or better than that of known alloys. .

With reference to FIGS. 1 and 2, microstructural examinations were performed on subsolvus-treated alloys A and C1 to detect the appearance of topologically compact phases (i.e., fragile intermetallic compounds) after an aging heat treatment of 500 hours at 750 ° C. The observations were made by scanning electron microscopy, backscattered electron contrast, on untouched samples. The severe aging of 500 hours at 750 ° C causes, in alloy A, the inter and intra-granular formation of phases rich in heavy elements. These phases appear in clear contrast (white edging) at the grain boundaries in Figure 1. These phases, when formed in excessive amounts, can cause a significant reduction in the ductility and mechanical strength of the alloys. Examinations of C1 alloy having been subjected to the same treatment of 500 hours at 750 ° C. show that these phases did not form during aging. The alloys of the invention thus have a better stability than the alloy A (N18) with respect to the formation of fragile intermetallic compounds which are the topologically compact phases.

Claims (11)

Alloy characterized in that it essentially comprises the following elements in the contents indicated in 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: 100% complement. An alloy according to claim 1, wherein the sum of the Al, Ti and Nb contents, in atomic percentages, is greater than or equal to 10.5% and less than or equal to 13%. An alloy according to claim 1, wherein the Al, Ti and Nb contents, in atomic percentages, are such that the ratio of the sum of the contents of Ti and Nb, and the Al content, is greater than or equal to 0, 9 and less than or equal to 1.1. An alloy according to any one of claims 1 to 3, wherein the W, Mo, Cr and Co contents, in atomic percentages, are such that the sum of the contents of W, Mo, Cr and Co is greater than or equal to 30 % and less than or equal to 34%, and such that the sum of the contents in W and Mo is greater than or equal to 3% and less than or equal to 4.5%. An alloy powder according to any one of claims 1 to 4. Process for manufacturing a part, characterized in that a blank of said part, or the part itself, from an alloy powder according to any one of Claims 1 to 4 is produced by a powder metallurgy technique. A method of manufacturing a workpiece according to claim 6, wherein said blank or said workpiece is subjected to a recrystallization heat treatment in which the blank or work piece is brought to a temperature above the solvus temperature of the gamma prime phase of said alloy, and lower than the melting start temperature of this alloy. A method of manufacturing a workpiece according to claim 6, wherein said blank or said workpiece is subjected to a recrystallization heat treatment in which the blank or workpiece is brought to a temperature below the solvus temperature of the gamma prime phase of said alloy. Turbomachine part, characterized in that it is made of an alloy according to any one of Claims 1 to 4. A turbomachine part according to claim 9, having a coarse grain structure in the zone in which it is subjected to the highest operating temperatures and where creep plays a significant role in the damage of the workpiece, and a small structure grains in the zone in which it is subjected to the lowest operating temperatures and where the damage results mainly from the tensile stresses and cyclic stresses. Turbomachinery part according to claim 9 or 10, consisting of a compressor or turbine disk.

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