WO2010023210A1

WO2010023210A1 - Process for preparing a nickel-based superalloy part and part thus prepared

Info

Publication number: WO2010023210A1
Application number: PCT/EP2009/060954
Authority: WO
Inventors: Gérard Raisson; Benjamin Picque; Dominique Cauvin
Original assignee: Aubert & Duval
Priority date: 2008-08-26
Filing date: 2009-08-25
Publication date: 2010-03-04
Also published as: FR2935395B1; FR2935395A1

Abstract

Process for preparing a nickel-based superalloy part by powder metallurgy: - manufacture of a superalloy with the composition 18 % ≤ Cr ≤ 24 %; 6 % ≤ Mo ≤ 10 %; 2.5 % ≤ Nb ≤ 5 %; traces ≤ Fe ≤ 10 %; traces ≤ Al ≤ 1 %; 0.5 % ≤ Ti ≤ 2.5 %; traces ≤ B ≤ 0,01 %; traces ≤ Mn ≤ 0.35 %; traces ≤ Si ≤ 0.2 %; traces ≤ C ≤ 0.05 %; traces ≤ Co ≤ 2 %; traces ≤ Ta ≤ 0.5 %; traces ≤ Mg ≤ 0.05 %; traces ≤ P ≤ 0.015 %; traces ≤ S ≤ 0.01 %; the rest being nickel and impurities; - atomisation of a molten mass of said superalloy in order to obtain a powder; - screening of said powder; - introduction of the powder into a container; - closure of and application of a vacuum to the container; - densification of the powder and the container by pressurizing them both in order to obtain an ingot or a billet; - hot forming of said ingot or said billet; The densification process is carried out at 500 to 1500 bar, at a temperature at which the fraction by volume of liquid phase in the superalloy is between 0.5 and 10 %, the temperature and the pressure being maintained for 3 to 16 h, and the forging being carried out without intermediate heat treatment after the densification of the powder. A hot formed part thus produced.

Description

Process for preparing a nickel-based superalloy part and part thus prepared

The present invention relates to a process for obtaining hot formed parts, in particular by forging, from powders of a nickel-based superalloy, for example the superalloy with the trade name 725^®.

Nickel superalloys are materials currently used to produce components intended for aeronautical or terrestrial turbines, such as turbine discs. Said materials are characterised by their capacity to function under high stresses and under high fatigue loads at elevated temperatures of above 650 °C which can reach 1090 °C in the case of certain applications in aeronautical turbines. The search for high-performance materials capable of resisting increasing operating temperatures and/or to resist higher stresses at the same temperature is linked to the need to improve the thermodynamic output of the turbines.

The components for aeronautical turbines of nickel-based superalloys (that is to say containing at least 50 % by weight of nickel, the rest consisting of various alloy elements) are most conventionally obtained by a method called the "ingot route" in which the nickel based superalloy is produced, by melting and remelting, as an ingot, before being hot worked by one or more thermomechanical and thermal treatments, in order to obtain the desired microstructure and final shape. Said ingot route is not, however, optimal for obtaining large sized parts having the aforementioned enhanced properties, because of a microstructure which is not sufficiently homogeneous after melting and remelting of the alloy. A highly homogeneous microstructure of the material before hot working is in fact necessary in order to be able to work the material with higher deformation rates and at higher deformation speeds, while at the same time avoiding the formation of shrinkage cracks (that is to say surface fissures formed during cooling), and the appearance of structural defects in the material.

The method of manufacture known as the "powder route" (powder metallurgy), which enables materials with a far more homogeneous structure to be obtained, was developed some years ago now for the production of high- performance components from nickel-based superalloys, in particular for applications in aeronautical turbines. Said powder route comprises in particular the following stages: - preparation of a molten mass with the composition specified for the superalloy;

- atomisation of said molten mass in order to obtain a powder;

- screening of said powder in order to retain only the particles having the desired grain size;

- introduction of the powder into a container which is closed and placed under vacuum;

- densification of the powder and the container in order to obtain an ingot or a billet of appropriate dimensions; - thermomechanical treatments (forging, for example) and, optionally, purely thermal treatments of the ingot or billet in order to obtain a final part with dimensions and structures appropriate to the intended application.

However, the parts obtained by the powder route are difficult to work by thermomechanical treatment, due notably to the lack of ductility of the parts obtained after densification of the powder. In particular, it has proved impossible to date to forge such parts produced from alloys of the 725^® type or from an alloy of similar composition containing Ti and/or Al and/or Cr with satisfactory results. It will be recalled that a typical composition of an alloy of the 725^® type is

^* 19 % < Cr < 23 %; typically 20.75 % * 7 % < Mo < 9 %; typically 8.2 %

^* 3 % < Nb < 4 %; typically 3.4 %

^* 4 % < Fe < 6 %; typically 9 %

^* 0.3 % < Al < 0.6 %; typically 0.2 %

^* 1 % < Ti < 1.8 %; typically 1.35 % * 0.002 % < B < 0.004 %; typically 0.002 %

^* Mn < 0.35 %;

^* Si < 0.2 %;

^* C < 0.03 %; typically 0.02 %

^* P < 0.008 %; typically < 0.015 % * S < 0.002 %; typically < 0.01 % the rest being nickel and impurities resulting from manufacture.

The lack of forgeability of the parts obtained from powders of nickel-based superalloys is explained by the surface characteristics of the initial particles, which affect the structure of the material and remain after compaction of the powder. The surfaces of the initial particles exhibit a network of fine precipitates known as PPB (Prior Particle Boundaries). The particles of the initial powder therefore exhibit surfaces which promote the formation and the rearrangement of said insoluble precipitates, such as oxides, sulphides, nitrides, sulphonitrides, carbides and/or carbinitrides, which remain after compaction of the powder. Said phenomenon is known as "decoration" around the powder particles. During the powder compacting operation, the PPBs form stable networks which are impossible to remove by subsequent treatments and which form "barriers" tending to oppose the migration of the grains during the forging operations. This promotes interparticle ruptures through cohesion loss of the grains during future loading of the part, since the PPB networks give to the metal a weak forgeability when it is subject to tensile stress. Premature cracking in the region of the PPBs is then observed, in particular during the forging stages (forging cracks). It is conventionally impossible to enlarge the grain by more than three times the size of the initial particles. This makes it impossible to obtain certain enhanced final mechanical characteristics such as good creep strength.

In document US-A-5 009 704, a solution to this problem is proposed in the form of a thermomechanical treatment process applied to nickel-based alloys in general, consisting in placing a powder of the alloy in a container and applying a hot isostatic pressure to the powder-container assembly at a high temperature, which results in partial melting (so-called "burn") of the surface of the particles, but without going so far as to cause dissolution of the stable metal carbides. It is then cooled slowly to a temperature below the solidus of the alloy, at which it is kept for a sufficient time as to ensure diffusion of the alloy elements that have segregated into the liquid, before continuing the cooling to ambient temperature. The required hot deformations are then carried out, for example by forging or extrusion, said deformations being performed below the solvus temperature of the gamma' phase. Typically, the hot isostatic compaction is performed at a temperature of -4 to 10°C above the solidus at 1030 bar for 3 hours. The treatment of the UDIMET 720-type alloy, based mainly on Ni, Cr and Co and the composition of which is given in US- A-5 009 704, is a prime example of the application of said process to improve the ductility of the preforms prior to hot working. It should be noted that said alloy, unlike alloys of the 725^® type, is hardened only by simple precipitation of the gamma' phase and not by double precipitation of gamma' phases on the one hand and gamma" or delta phases on the other.

However, said process is complex in execution and does not meet the productivity criteria currently required, in particular for the production of parts of the 725^® alloy-type.

The aim of the invention is to improve the aptitude for hot straining and consequently the forgeability of parts from alloys of the 725^® type hardened by double precipitation and alloys of this type, while at the same time conferring on them enhanced final mechanical characteristics.

To this end, the invention relates to a process for preparing a nickel-based superalloy part by powder metallurgy, comprising the following stages: - manufacture of a nickel-based superalloy with the composition: * 18 % < Cr < 24 %; * 6 % < Mo < 10 %;

^* 2.5 % < Nb < 5 %;

^* traces < Fe < 10 %;

^* traces < Al < 1 %;

^* 0.5 % < Ti < 2.5 %; * traces < B < 0.01 %;

^* traces < Mn < 0.35 %;

^* traces < Si < 0.2 %;

^* traces < C < 0.05 %;

^* trace < Co < 2 %; * traces < Ta < 0.5 %;

^* traces < Mg < 0.05 %;

^* traces < P < 0.015 %;

^* traces < S < 0.01 %; the rest being nickel and impurities resulting from manufacture. - atomisation of a molten mass of said superalloy in order to obtain a powder; - screening of said powder in order to extract the particles having a predetermined grain size;

- introduction of the powder into a container, optionally under vacuum; - closure of and application of a vacuum to the container; - densification of the powder and the container by pressurising them in order to obtain an ingot or a billet;

- hot forming and optionally heat treatment of said ingot or billet; characterised in that the densification is carried out at a pressure of 500 to

1500 bar, at a temperature at which the fraction by volume of liquid phase in the superalloy lies between 0.5 and 10 %, the densification temperature and pressure being maintained for 3 to 16 h, and in that the forging is carried out without intermediate heat treatment after the densification of the powder, optionally after returning the ingot or the billet to ambient temperature and reheating it to the hot forming temperature. Preferably, the densification temperature lies between 10 and 30 °C above the burn temperature of the alloy measured in the conditions of its manufacture according to the stages preceding densification and the chosen densification conditions.

Preferably, the alloy has the composition * 19 % < Cr < 23 %;

^* 7 % < Mo < 9.5 %;

^* 2.75 % < Nb < 4 %;

^* traces < Fe < 9 %;

^* traces < Al < 0.6 %; * 1 % < Ti < 1.8 %;

^* 0.001 % < B < 0.005 %;

^* traces < Mn < 0.35 %;

^* traces < Si < 0.2 %;

^* traces < C < 0.03 %; * traces < Mg < 0.05 %;

^* traces < P < 0.015 %;

^* traces < S < 0,01 %; the rest being nickel and impurities resulting from manufacture. Preferably, the densification pressure thus lies between 800 and 1400 bar, the densification temperature lies between 1210 and 1230 °C and the period of maintaining said temperature and said pressure lies between 3 and 16 hours.

Preferably, the temperature and the period of densification are chosen in order to obtain a fraction by volume of liquid phase in the superalloy of 0.5 to 10 % over a depth of 100 to 200 mm measured from the periphery of the container.

Preferably, at least the first stage of the hot forming process is performed at a temperature greater than the delta solvus temperature of the superalloy.

Preferably, said first hot forming stage is performed at a temperature of between 1010 and 1050 °C.

Preferably, the densification is performed by hot isostatic compaction. Preferably, the first stage of said hot forming is a potting die forging. Preferably, the hot forming initially comprises two potting die forging stages. The invention also relates to a hot formed nickel-based superalloy part, characterised in that it has been obtained by the above process.

Said part may be a component of an aeronautical turbine or a terrestrial turbine.

As will have been understood, the invention consists first of all in carrying out, on a powder of a superalloy of the 725^® or a similar type, a densification stage (for example by hot isostatic compaction) at a temperature at which the fraction by volume of liquid phase in the superalloy is from 0.5 to 10 %, and preferably able to range between 10 and 30°C above the burn temperature, hence generally, for an alloy with a composition as indicated in Table 1 below, at a temperature of between 1210 and 1230°C, said temperature being maintained for 3 to 16 h, and at a pressure of 500 to 1500 bar, preferably 800 to 1 100 bar. The period of maintaining the temperature of 3 to 16 hours depends first of all on the size of the slug. Typically, if the diameter of the slug is less than 250 mm, 4 hours suffice, and if the diameter of the slug is more than 1000 mm, the temperature will have to be maintained for at least 10 hours. Simple models or tests enable the skilled man to determine the optimum maintenance period in terms of cost and metallurgical effectiveness for a slug of given composition and dimensions.

The hot forming stage, comprising for example forging that can be followed by die stamping, is then performed directly after densification, that is to say without specific intermediate heat treatment between densification and forming (but a reduction to ambient temperature after densification, followed by reheating to the forming temperature, is possible).

The invention will be best understood upon reading the following description, which is given with reference to the following attached figures:

- Fig. 1 shows a die-stamped part after hot isostatic compaction carried out in standard conditions at 1120 °C;

- Fig. 2 shows a part similar to that of Fig. 1 for which, according to the invention, hot isostatic compaction has been carried out at 1230 °C. The inventors found that a 725^® type alloy is a relatively "clean" alloy, that is namely unlikely to cause major segregations when it passes above the solidus during densification. Under said conditions, they concluded that a diffusion heat treatment below the solidus after densification, as practised in US-A-5 009 704, was not worthwhile in its case. Specifically, the segregations remain limited in scale and localised, and are found not to impair the properties of the parts produced thereafter.

Furthermore, the ranges of the characteristics of the treatments according to the invention are specially adapted to obtain optimal characteristics for the parts hot formed from ingots or billets of a 725® type alloy or comparable alloys hardened by double precipitation of gamma' and gamma" phases and obtained by a powder metallurgy process. The alloys to which the invention relates contain, after the double precipitation process causing them to harden, from 5 to 30 % of gamma', gamma" and delta phases (the delta phase often accompanying the precipitation of the gamma" phase and being difficult to distinguish from the gamma" phase), preferably from 15 to 20 %. This percentage of gamma' + gamma" + delta phases depends on only the chemical composition of the metal and not its treatment conditions. It is inevitably obtained within the scope of the invention for the compositions to which the invention relates. Said treatment conditions make it possible to perform a "burn" of the alloy, in particular at the surface of the powder grains during densification, which permits the at least partial dissolution of the precipitates grouped at the PPBs, and hence to reduce their harmfulness in that they no longer hamper the mobility of the grain boundaries from which they are dissociated. The hot ductility, and consequently the forgeability, of said alloys, is therefore improved. The other mechanical properties of the alloy also remain suitable for its proposed main uses, in particular for aeronautical and terrestrial gas turbines.

A clear distinction must be drawn between the burn temperature and the solidus temperature.

The solidus temperature is the temperature at which liquid appears in an alloy that is perfectly homogeneous (in thermochemical equilibrium) at said temperature. The burn is the temperature at which liquid appears within a metal that is in a particular state. The burn is produced at a temperature lower than the solidus, owing to the local differences in the particular metal compared with the equilibrium. For nickel-based alloys, the differences from the equilibrium are created for example by the presence of carbides (MC type) and borides, which are normally redissolved before reaching the solidus. In reality, the borides and carbides not in equilibrium dissolve progressively as soon as a certain temperature level is reached, which leads to a local over-concentration of elements such as Ti or Nb, which lower the solidus temperature at points where they are concentrated. The appearance of liquid on said local over-concentrations is then observed.

While the solidus temperature depends only on the composition of the alloy, the burn temperature depends on the state of the metal and on the heating cycle to which it is subjected. It therefore has to be determined experimentally for an alloy manufactured in particular conditions.

A consequence of the appearance of liquid due to burning is the creation of defects (porosities or heterogeneities) that exert a harmful influence on forgeability. The appearance of porosities is prevented by the pressure of the HIC (or of another method of densification). The heterogeneities can be eliminated by a sub-solvus heat treatment, which is what US-A-5 009 704 proposes.

The solution proposed by the inventors consists, in contrast, in adjusting the HIC cycle in order to limit the size and the density of the heterogeneities generated by the appearance of liquid, so as to obtain the desired effect on the PPBs without needing to carry out a subsequent homogenisation treatment. In order to determine the burn temperature, and to deduce from it the HIC stage temperature to be imposed during the manufacture of the ingot or the billet, the following procedure can be adopted. HICs are performed on powder specimens with the stage temperatures staggered for example every 10°C (between 1200 and 1250°C in the case of the alloy of which the composition is given in Table 1 ), the HIC cycle (rise in temperature and pressure and stage) corresponding to that which will be used for the powders and the containers intended for industrial parts.

Characterisation of the microstructure by light microscopy of the specimens corresponding to each test is used to determine the optimum temperature range corresponding to a pronounced removal of the PPBs and the appearance of defects of limited size and density. The burn temperature is that from which said effects start to be observed.

An advantage of said type of treatment according to the invention is that the proportion of liquid (which generates defects) can grow slowly with the temperature, whereas with a treatment above the solidus the appearance of liquid is far more rapid and it is difficult to keep it to less than 10 %. The fraction of liquid appearing during said treatment lies, according to the invention, between 0.5 and 10 %. In the case of 725^®, the measured temperature range is ± 15°C, which makes the process applicable to industrial conditions.

It should be noted that specimens of small size are treated in a homogeneous manner in the HIC cycle with said characteristics. In the case of industrial containers whose diameter can exceed 1 meter, only a surface layer with an estimated thickness of 100 to 200 mm will have exceeded the burn temperature. Since only said surface layer undergoes a tensile deformation during the subsequent thermomechanical transformations, the indicated treatment is sufficient to confer to the blanks the forgeability required for their transformation. A temperature more than 40 °C higher than the burn temperature of the superalloy is not desirable during densification, since the defects generated are large and this affects the forgeability of the metal.

In a preferred embodiment of the invention, the densification pressure is from 800 to 1400 bar, the densification temperature lies between 10 and 30 °C above the burn temperature, i.e. about 1210 and 1230 °C for 725^®, and its duration lies between 3 and 16 hours. In said conditions, a liquid phase fraction of 0.5 to 10 % by volume is obtained in the alloy, which enables the results sought to be obtained in an optimum manner.

The type 725^® alloys and the similar alloys to which the invention applies exhibit a relatively small increase in the proportion of liquid phase as a function of temperature, which makes it possible, in the prescribed parameter ranges, to effectively control the optimum proportion of liquid phase without causing the appearance of heterogeneities harmful to forgeability.

As has been stated, the fraction by volume of liquid phase during densification is from 0.5 to 10 %. It is preferably measured over a depth which, starting from the periphery of the container, is between 100 and 200 mm, preferably about 150 mm. At greater depths, the fraction by volume of liquid may be zero.

The inventors found that, by means of an adapted forging range, only the peripheral zone of the aforementioned slug over said depth is subject to tensile stress and is therefore capable of cracking prematurely. The densification temperature and the period of maintaining said temperature are therefore, in fact, independent of the diameter and the volume of the container. A consequence of said finding is that the productivity of the process can be optimised by not necessarily aiming to obtain the burn and the optimum proportion of liquid phase over the whole of the powder, including at the centre of the ingot or the billet undergoing formation, but only at the aforementioned depth of 200 mm, if not less.

Before densification, it is preferable to apply a vacuum to the container before its actual filling. The inventors found that the scale of the PPB phenomenon is reduced and its homogeneity improved by said procedure. The vacuum is maintained after the closure of the filled container, since it is necessary for densification.

Preferably, the densification is performed by hot isostatic compaction, said process being particularly suited to the production of ingots that can reach 20 t in conditions of perfect control of the densification parameters.

After densification, hot transformation is performed, by forging or by another process. As has been stated, said hot transformation can take place immediately after densification, or else after cooling conducted without special requirements, for example a natural cooling to ambient temperature, followed by reheating to the hot forming temperature. In any case, no special heat treatment is required between densification and hot transformation.

The forging temperature, or the hot forming temperature in general, typically lies between 900 and 1100 °C. However, it was found that a hot forming at a temperature slightly higher than the solvus of the delta phase, i.e. between 1010 and 1050 °C, for example 1025 °C, corresponded to an optimum in forgeability and control of the microstructure leading to the desired mechanical properties. Preferably, at least the first hot forming stage takes place at a temperature higher than the delta solvus of the alloy.

The forging (or the hot forming in general) is preferably carried out with a first stay in the furnace following by at least two further stays in the furnace (in other words "in three heats") enabling the part to be brought back to the nominal forging temperature in order to compensate for its cooling, with a rational deformation defined by LogN (initial height/final height of the slug) of preferably more than 0.30 in order to obtain the microstructure leading to the desired mechanical properties.

The forging process comprises a first forging stage most preferably conducted "in a potting die", that is to say conducted by placing the ingot or the billet to be forged in an annular part called a potting die which, during the forging, allows the product to be stressed radially, and hence good micro-structural homogeneity to be obtained in the radial directions of the alloy. Any subsequent forging stage or stages can be conducted in an open die. The forging process preferably comprises initially at least two potting die forging stages, to which one or two open die forging stages (without potting die) can be added.

After densification according to the invention, the PPBs no longer have harmful effects on the recrystallisation of the alloy. The potting die forging process performed in this way consequently enables good microstructural homogeneity of the alloy and improved foregeability to be achieved. Due to the improved characteristics of the alloy that are obtained after potting die forging, the last forming stage or stages of the part can be one or more forgings performed as open die forging, or can be directly die stamping stages performed in a forging tool kit (die) in order to give to the part the desired final geometry. The open die forging or die stamping operation can thus be performed in one to three stages according to the dimensions of the intended final part, with rational deformations preferably greater than 0.30 for each stage, while at the same time preventing the formation of forging cracks.

The process according to the invention makes it possible to obtain a part without cracks the material of which has a mean grain size after forging or die stamping of 8 to 10 ASTM and improved mechanical characteristics after forging and conventional heat treatment:: a yield strength of more than 1050 MPa, a tensile strength of more than 1400 MPa, and an elongation at break of 15 to 20 % before forging and 25 to 30 %, in some cases far more, after forging.

In order to illustrate the advantages of the invention compared with treatments that fail to meet its precise conditions, the results of a series of tests, conducted on powder specimens with two different compositions that nevertheless both meet the conventional specifications for the 725^® alloy, will be shown. These results are given in Table 1 , expressed in % by weight.

Specimen Ni Fe Cr A Ti Mo Nb C Co Ta

1 balance 4.87 20 .2 0. 38 1.42 7.61 3.62 0. 015 <0.01 <0. 02

2 balance 5.25 20 .7 0. 42 1.44 7.56 3.73 0. 014 0.016 <0. 003

Specimen B Si Mg Mn P S O N

1 <0.0005 0.033 <0.0005 <0.03 <0 .0049 <0.0021 0. 0085 0. 0065

2 0.0031 0.033 <0.001 0.029 <0 .003 0.0017 0. 0071 0. 0106

Table 1 : Compositions of the specimens tested

The two specimens tested differ mainly in their contents of Fe, their contents of hardening elements Al, Ti, Nb, and above all their contents of B, which are higher in specimen 2. Their burn temperatures are 1200°C ± 5°C and the delta solvus is 970 ± 1 O⁰C.

The tests conducted on specimens 1 and 2 defined above consisted in producing a slug with a diameter of 70 mm and a height of 500 mm by hot isostatic compaction (HIC) of the powder and its container according to various conditions that will be described. Firstly, powders of said alloys that possessed a grain size enabling them to pass through the mesh of a 100 μm screen were prepared and screened in the conventional manner.

In a first series of tests, HIC was conducted according to the "standard" methods, namely simple isothermal maintenance between 1000 and 1400 bar for 3 h, at temperatures of 1025, 1120 and 1 16O⁰C, i.e. all lower than the solidus temperature and the burn temperature of the alloy.

In a second series of tests, densification by hot isostatic compaction (HIC) at 1220 or 1230°C for 3 h at 1000 bar was performed, according to the invention. The mechanical characteristics of the parts obtained were measured either after HIC or after forging. Said forging process consisted, according to the cases, in:

- a rational deformation of 0.30 at 1025 °C in a heat at 5 x10^"2/s (treatment 1 ) produced by potting die forging, optionally followed by maintenance at 1025°C for 2 h or 1230°C for 1 h;

- a deformation in three heats at 1000°C (treatment 2) comprising two initial potting die forging stages, then an open die forging stage followed by maintenance at 1025°C for 2 h.

Micrographic observations and mechanical tests on the ingots resulting from said tests were then carried out, in order to evaluate on the one hand the effect of the treatments undergone on the morphology of the grains and the grain boundaries, and on the other hand the effect of said same treatments on the forgeability of the material.

The micrographic observations were carried out under the light microscope after electrolytic chemical attack.

The forgeability tests were conducted on test specimens 6.35 mm in diameter and 35 mm in length which were deformed by traction at 1025°C at a low speed. The traction speed of the machine was minimal at 1.9 mm/s. The rate of deformation of the specimen was 5 x 4.10^"2/s, hence in conditions fairly close to those aimed at during typical die stampings of the parts for which the alloy manufactured is intended.

The influences of the various treatments on the properties of the materials can be summarised as follows. The microstructures obtained with the Standard HIC cycles exhibit a grain size which normally increases appreciably with the temperature at which HIC is carried out. It is not however possible, in the operating conditions adopted, to obtain a grain size having an ASTM index of less than 6 to 7 after HIC, because of the presence of the PPBs at the grain boundaries, which limits the growth of the grains (it should be noted that the ASTM index indicating the size of the grains is higher the smaller the size of the grains).

Since the cycles according to the invention were carried out at 1220 or

1230°C, hence at a temperature higher than the burn temperature of the alloy, they make it possible to obtain after HIC a product having a grain size greater than that obtained after the standard HIC cycles conducted at the same temperatures.

Some grains even have a size of 5 ASTM.

On the ingot produced according to the invention, the size of the grains is significantly more homogeneous than on the standard one, and the grains of a very small size have disappeared, which is a sign that the PPBs did not constitute obstacles to their growth. Thus it is clear that the PPBs no longer have harmful effects in terms of the recrystallisation of the alloy and damage during the forging.

Hot isostatic compaction is a preferred method of densification within the scope of the invention, but other methods can be considered, such as for example unidirectional compression or extrusion.

After densification the resulting ingot or the billet is, conventionally, peeled and then hot formed, generally by forging, for example in three stages separated by reheatings to the nominal forging temperature, then by die stamping if the forging process has not conferred on the part its definitive shape. Potting die forging (also called "potting die upsetting") is particularly recommended for the preferred applications considered, since it permits the semi-finished product to be calibrated for the die stamping process and its surface to be kneaded in order to obtain microstructural and forgeability characteristics at said surface which approximate as closely as possible to those that are found at the centre of the semi-finished product.

The parts obtained after standard HIC, after potting die forging and then open die forging, exhibited forging cracks, while the parts obtained after HIC according to the invention and after potting die forging and then open die forging were devoid of forging cracks.

In a third and fourth series of tests, forging tests were conducted on slugs and parts of large dimensions and weighing between 40 and 200 kg. The same batch of powder was used for the third and the fourth series of tests. Said powder batch used had a composition corresponding to that of specimen no. 2 of Table 1.

In the third series of tests, slugs weighing 44 kg were obtained after standard HIC at a densification temperature of 1120 °C for 4 h at 1000 bar, and other slugs weighing 44 kg were obtained after HIC according to the invention at a densification temperature of 1230 °C for 4 h at 1000 bar. After each HIC, peeling allowed slugs 148 mm in diameter and 31 1 mm high to be obtained. The two types of slugs (those obtained by standard HIC and those obtained by HIC according to the invention) underwent the same forging program:

1 st pass in a potting die 185 mm in diameter with a rational deformation of 0.44 at a deformation rate of between 5 and 8 mm/s;

2nd pass in a potting die 230 mm in diameter with a rational deformation of 0.43 and a deformation rate of between 5 and 8 mm/s;

3rd pass as an open die forging process between flat dies in order to bring the height to 89 mm with a rational deformation of 0.39 and a deformation rate of between 5 and 8 mm/s.

The forgings were carried out at 1025 and 1050°C. In the case of the slugs densified at 1120 °C, the forged parts systematically exhibited cracks in the corners, while the forged parts made from slugs densified at 1230 °C were devoid of any cracks. In the fourth series of tests slugs weighing 200 kg were obtained, like the slugs weighing 44 kg of the third series of tests, after standard HIC at 1 120 °C and after HIC according to the invention at 1230 °C. After each HIC, peeling enabled slugs 228 mm in diameter and 570 mm in height to be obtained. The first two forging passes were conducted with a potting die at 1025°C with rates of reduction close to those obtained for the above-mentioned 44 kg slugs. The third pass consisted in die stamping in a forging tool kit with deformation levels and rates of between 5 and 8 mm/s, i.e. greater than those obtained for the 44 kg slugs involving far greater stressing of the metal. The die stamping was also conducted at 1025°C. The part die stamped from the metal densified at 1 120 °C exhibited numerous large cracks (see Fig. 1 ), while the part die stamped from metal densified at 1230 °C was totally devoid of the latter (see Fig. 2).

Claims

1. Process for preparing a nickel-based superalloy part by powder metallurgy, comprising the following stages: - manufacture of a nickel-based superalloy with the composition: * 18 % < Cr < 24 %; * 6 % < Mo < 10 %;

^* 2.5 % < Nb < 5 %;

^* traces < Fe < 10 %; * traces < Al < 1 %;

^* 0.5 % < Ti < 2.5 %;

^* traces < B < 0.01 %;

^* traces < Mn < 0.35 %;

^* traces < Si < 0. 2%; * traces < C < 0.05 %;

^* traces < Co < 2 %;

^* traces < Ta < 0. 5 %;

^* traces < Mg < 0.05 %;

^* traces < P < 0.015 %; * traces < S < 0.01 %; the rest being nickel and impurities resulting from manufacture;

- atomisation of a molten mass of said superalloy in order to obtain a powder;

- screening of said powder in order to extract from it the particles having a predetermined grain size; - introduction of the powder into a container, optionally under vacuum;

- closure of and application of a vacuum to the container;

- densification of the powder and the container by pressurising them both in order to obtain an ingot or a billet;

- hot forming and optionally heat treatment of said ingot or said billet; characterised in that densification is carried out at a pressure of 500 to 1500 bar, at a temperature at which the fraction by volume of liquid phase in the superalloy is between 0.5 and 10 %, the densification temperature and pressure being maintained for 3 to 16 h, and in that forging is carried out without intermediate heat treatment after the densification of the powder, optionally after the ingot or the billet has returned to ambient temperature and been reheated to the hot forming temperature.

2. Process according to claim 1 , characterised in that the densification temperature is between 10 and 30 °C above the burn temperature of the alloy measured in the conditions of its manufacture according to the stages preceding densification and the chosen densification conditions.

3. Process according to claim 1 or claim 2, characterised in that the alloy has the composition

^* 19 % < Cr < 23 %;

^* 7 % < Mo < 9.5 %;

^* 2.75 % < Nb < 4 %;

^* traces < Fe < 9 %; * traces < Al < 0.6 %;

^* 1 % < Ti < 1 .8 %;

^* 0.001 % < B < 0.005 %;

^* traces < Mn < 0.35 %;

^* traces < Si < 0.2 %; * traces < C < 0.03 %;

^* trace < Mg < 0.05 %;

^* trace < P < 0.015 %;

^* traces < S < 0.01 %; the rest being nickel and impurities resulting from manufacture.

4. Process according to claim 3, characterised in that the densification pressure is between 800 and 1400 bar, the densification temperature is between

1210 and 1230°C and the period of maintaining said temperature and said pressure is between 3 and 16 hours.

5. Process according to any one of claims 1 to 4, characterised in that the densification temperature and the period of densification are chosen in order to obtain a fraction by volume of liquid phase in the superalloy of 0.5 to 10 % over a depth of 100 to 200 mm starting from the periphery of the container.

6. Process according to any one of claims 1 to 5, characterised in that at least the first hot forming stage is carried out at a temperature greater than the delta solvus temperature of the superalloy.

7. Process according to claim 6, characterised in that said first hot forming stage is carried out at a temperature of between 1010 and 1050 °C.

8. Process according to any one of claims 1 to 7, characterised in that the densification is carried out by hot isostatic compaction.

9. Process according to any one of claims 1 to 8, characterised in that the first stage of said hot forming is a potting die forging.

10. Process according to claim 9, characterised in that the hot forming process comprises two initial potting die forging stages.

1 1. Hot formed part of a nickel-based superalloy, characterised in that it has been obtained by the process according to any one of claims 1 to 10.

12. Part according to claim 1 1 , characterised in that it is an aeronautical gas turbine component.

13. Part according to claim 1 1 , characterised in that it is a terrestrial gas turbine component.