EP2321440A2

EP2321440A2 - Method for preparing a nickel superalloy part, and part thus obtained

Info

Publication number: EP2321440A2
Application number: EP09740490A
Authority: EP
Inventors: Gérard Raisson
Original assignee: Aubert and Duval SA
Current assignee: Aubert and Duval SA
Priority date: 2008-08-26
Filing date: 2009-08-24
Publication date: 2011-05-18
Also published as: US8889064B2; US20110150693A1; WO2010023405A2; WO2010023405A3; FR2935396A1; FR2935396B1

Abstract

The invention relates to a method for preparing a nickel superalloy part by: - developing a nickel superalloy from a composition capable of hardening through double precipitation of a gamma’ phase and a gamma” or delta phase; - spraying a molten mass of said superalloy to obtain a powder; - sifting said powder; - placing said powder in a container; - closing and placing the container in a vacuum; - densifying the powder and the container to obtain an ingot or billet; - hot forming said ingot or said billet; characterized in that, before the densifying step, the powder and the container are heated for at least 4 hours to a temperature that is both greater than 1,140°C and less than the solidus temperature of the superalloy by 10°C at a pressure causing a densification less than or equal to 15% of the powder volume. The invention also relates to a part thus produced.

Description

Process for the preparation of a nickel base superalloy piece and piece thus obtained.

The present invention relates to a process for obtaining forged parts from powders of a nickel superalloy hardened by double precipitation (gamma 'and gamma' or delta), such as the 725 ^® brand superalloy.

Nickel superalloys are materials commonly used to make components for aeronautical turbines, such as turbine disks. These materials are characterized by their ability to operate under high stress and high fatigue loads at high temperatures, above 650 ° C, which can reach 1090 ° C in the case of some applications of aeronautical turbines. The search for efficient materials capable of withstanding increasingly high operating temperatures is related to the need to improve the thermodynamic efficiency of the turbines.

The components for aeronautical turbines made of superalloys based on nickel (that is to say comprising at least 50% by weight of nickel, the remainder being composed of various alloying elements) are most conventionally obtained by means of a obtained, called "ingot channel", where the nickel base superalloy is prepared by melting and remelting, then cast and put into the form of an ingot, before being hot worked by one or more thermomechanical and thermal treatments to obtain the microstructure and the desired final shape.

This ingot channel is however not optimal for producing parts having the aforementioned high properties, because of a microstructure which is not sufficiently homogeneous after melting and remelting of the alloy. Indeed, a very homogeneous microstructure of the material before hot work is necessary to be able to work the material with deformation rates and higher deformation rates, while avoiding the formation of taps (that is to say cracks). formed during cooling) during thermomechanical treatment and the appearance of structural defects in the material.

For a few years now, the so-called powder pathway (powder metallurgy) has been used to obtain structural materials much more homogeneous, developed for the realization of high-performance components in nickel base superalloys, especially for aeronautical turbine applications. This powder route comprises in particular the following steps: preparation of a melt having the target composition for the superalloy;

atomizing this melt to obtain a powder;

sieving this powder to retain only the particles having the desired particle size; - introduction of the powder into a container, which is closed and evacuated;

densification of the powder and the container to obtain an ingot or a billet of appropriate dimensions;

- Thermomechanical treatments (forging, for example) and possibly thermal ingot or billet to obtain a final piece of dimensions and structures appropriate to the intended application.

However, the parts obtained by the powder route are difficult to work by thermomechanical treatment, particularly because of the lack of ductility of the parts obtained after densification of the powder. The lack of ductility of the parts obtained from nickel base superalloy powders is explained by the characteristics of the surfaces of the original particles, which will mark the structure of the material and subsist after compaction of the powder. The surfaces of the parent particles are also known as PPB (Prior Particle Boundaries). The particles of the initial powder have surfaces which promote the formation and grouping of insoluble precipitates, such as oxides, sulphides, nitrides, sulphonitrides, carbides and / or carbonitrides, which will subsist after compaction of the powder. . This phenomenon is known as "decorations" around powder particles. During the compacting operation of the powder, the precipitates present at the PPBs form stable networks, which can not be removed by subsequent treatments.

A consequence of this phenomenon is to promote interparticle breaks during future solicitations of the room, and to make it difficult to grain magnification very substantially beyond the original particle sizes. It is classically impossible to enlarge the grain beyond three times the original particle size. This makes the billet obtained after compaction of the powder very difficult to forge and makes it impossible to obtain certain high final mechanical characteristics, such as good creep resistance.

In EP-A-0 438 338 a solution has been proposed for mitigating the harmful effects of the precipitates or decorations at the PPBs for nickel superalloys of the type with a hardening by gamma phase precipitation, such as, in particular, the alloys known in the art. the trade names ASTROLOY ^® , UDIMET 720 ^® or N18 ^® . This document specifies the typical compositions of ASTROLOY ^® and N18 ^® alloys. The typical composition of I ¹ UDIMET 720 ^® is

- 15.5% <Cr <16.5% - 14% <Co <15.5%

- 4.75% <Ti <5.25%

- 2.25% <Al <2.75%

- 2.75% <Mo <3.25%

- 1% <W ≤ 1, 5% - 0.025% <Zr <0.05%

- 0.01% ≤ C <0.02%

- 0.01% ≤ B <0.02%

- Ni = the rest

This solution consists in carrying out a pretreatment of the superalloy, before densification, at a temperature below the solvus temperature or close to the solvus temperature, of the gamma phase of the alloy (1 195 ° C. for ASTROLOY ^® and 1180 ° C for N18 ^®). This method makes it possible to mitigate the harmful effect of PPBs for superalloys cured by gamma phase precipitation, by precipitating the segregated elements inside the powder particles and not at their surface. Thanks to this pretreatment, uncoupled from the densification itself, the grains can grow beyond the size of the initial particles, which improves the forgeability of the alloy. However, it turns out that this solution, while providing remarkable technological advantages for nickel base-hardened alloys by simple gamma phase precipitation, is not applicable to nickel base superalloys for which the structural hardening is achieved by double precipitation of a gamma phase and a gamma phase or delta.

Indeed, a pretreatment carried out under the solvus temperature of the gamma phase or in the vicinity of this solvus temperature does not allow, in their case, to eliminate or mitigate the harmful effect of SCH and decorations at SCH. Nickel base superalloys hardened by double precipitation, because of their mechanical properties (mechanical resistance, creep and fatigue resistance at high temperatures), would be of great interest for aeronautical applications, especially for turbine components such as disks. or the blades. It would therefore be very important to find a mode of development by the powder route for using these superalloys for these applications, in particular the known superalloy known commercially 725 ^® , because of its mechanical properties and its corrosion resistance.

The object of the invention is to improve the ductility, and hence the forgeability, of double-precipitation hardened nickel superalloy parts by allowing the grains to grow substantially beyond the original powder particle sizes.

To this end, the subject of the invention is a process for the preparation of a nickel base superalloy part by powder metallurgy, comprising the following steps: - elaboration of a nickel base superalloy of a composition capable of providing a hardening by double precipitation of a gamma phase and a gamma phase or delta;

atomizing a melt of said superalloy to obtain a powder; sieving said powder to extract particles having a predetermined particle size;

introduction of the powder into a container, possibly under vacuum;

- closing and evacuation of the container; densification of the powder and the container by pressurizing the assembly to obtain an ingot or a billet;

- Hot forming and possibly heat treatment of said ingot or said billet; characterized in that prior to the densification step of the powder and the container, a heat treatment step is carried out consisting in heating the powder and the container for at least 4 hours, preferably for 12 to 30 hours, at a temperature of pressure resulting in a densification of the powder less than or equal to 15% of the initial volume, preferably less than or equal to 10% of the initial volume, this treatment taking place at a temperature both greater than 1140 ° C and lower at least 10 ° C at the solidus temperature of the superalloy.

The composition of the superalloy may be, in percentages by weight:

- 19% <Cr <23%;

- 7% <Mo <9.5%; - 2.75% <Nb <4%;

- traces <Fe <9%;

- traces <Al <0.6%; ; - 0.001% ≤ B <0.005% - traces <Mn <0.35%;

- traces <If <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 the elaboration.

Said heat treatment before densification is then preferably performed at a temperature of 10 to 50 ° C below the solidus temperature of the superalloy. For these alloys, the heat treatment before densification is carried out between 1140.degree. C. and 1180.degree. For an alloy of the above type, the heat treatment before densification is preferably carried out at a temperature both greater than 1140 ° C and 30 to 50 ° C lower than the solidus temperature of the superalloy.

The heat treatment before densification is, in this case, optimally carried out between 1160 and 1180 ° C for 12 hours to 30 hours at a pressure less than 50 bar.

Densification can be performed by hot isostatic compaction.

The hot shaping may include a forging blow.

The invention also relates to a nickel-base superalloy forging, characterized in that it was prepared by the above method.

This part can be an aeronautical or land gas turbine component.

As will be understood, the invention consists in producing, on a powder of a superalloy capable of being hardened by double precipitation of a gamma phase and a gamma or delta phase, a particular heat treatment of the powder and its container before densification, within a given temperature range.This treatment is intended to dissociate the grain boundaries of SCH networks, which can not therefore, in the following treatments, oppose the growth of the grain boundaries, and finally we obtain more ductile structures, therefore more suitable for hot forming such as forging.

A particular variant of the invention is referred to as the superalloy ARA 725 ^® or 725 ^® whose composition is preferably the one mentioned above, and offers a range of densification before processing temperatures which is specially adapted to it. The invention will be better understood on reading the description which follows, given with reference to the following appended figures:

- Figure 1 shows a micrograph of an ingot obtained after CIC of a 725 ^® alloy powder at 1160 ° C for 3 hours at 1000 bar according to the conventional method; FIG. 2, which similarly shows a micrograph of an ingot obtained according to the process of the invention from an ARA 725 ^® powder of the same composition as for FIG. 1, also obtained by CIC at 1 160 ° C for 3 hours at 1000 bar, but after the powder has undergone a heat treatment before densification at 1160 ° C for 6 h at atmospheric pressure.

The ARA 725 ^® alloys, having the above compositions, have solidus temperatures of about 1210 ° C ranging from 1200 to 1230 ° C depending on the precise composition.

To illustrate the advantages of the invention over treatments which deviate from its precise conditions, the results of a series of experiments carried out on powder samples of two different compositions, both of which are nevertheless usual requirements for alloy 725 ^® with a solidus temperature of approximately 1210 ° C + 5 ° C. They are presented in Table 1, expressed in% by weight.

Table 1: Compositions of the samples tested

The two samples tested differ essentially in their Fe contents, their contents of hardening elements Al, Ti, Nb and especially their contents in B, which are higher in the sample 2. Beforehand, the inventors carried out a study of phases likely to be present in alloy 725 ^® , for the compositions of the invention or close to them. This study was carried out on the one hand using the THERMOCALC software, commonly used by metallurgists, which makes it possible to establish the phase diagrams of metal alloys, and on the other hand by analysis tests. differential thermal and dilatometric optical microscope and scanning electron microscope examinations after different heat treatments.

The conclusions of this study are that 725 ^® can, in fact, be mainly hardened by the intergranular gamma and delta phases that accompany the gamma phase. The inventors have therefore concluded that it is the obtaining of this gamma phase or intergranular delta that should be preferred during treatments aimed at precipitating the gamma and gamma hardening phases before densification.

The experiments carried out on the samples 1 and 2 defined above consisted in making a piece of dimensions of diameter 70 mm and height 500 mm by hot isostatic compaction (CIC) of the powder and its container according to various modalities that we are going to specify.

Powder of these alloys, having a particle size allowing them to pass through the meshes of a 100 μm sieve, were first prepared and screened in a conventional manner.

In a first series of experiments, a CIC was carried out according to "standard" modalities, namely a simple isothermal maintenance of 3 hours between 1000 and 1400 bar, at temperatures of 1025, 1120 and 1160 ° C.

In a second series of experiments, the CIC densification was preceded by a heat treatment of the powder and the container for 6 hours at 1025, 1120 and 16O ⁰ C. Then the CIC was held at 1000 bar for 3 hours at the same temperature as the heat treatment. This cycle has been called the "decoupled cycle". It will be seen that this decoupled cycle is in accordance with the invention when the temperature of the heat treatment is 1160 ° C. Microscopic observations and mechanical tests on the plots resulting from these tests were then carried out in order to assess, on the one hand, the effect of the treatments on the morphology of the grains and the grain boundaries, and on the other hand the effect of these same treatments on the forgeability of the material.

The micrographic observations were performed under an optical microscope after electrolytic chemical etching.

The forgeability tests were conducted on specimens 6.35 mm in diameter and 35 mm in length, which were deformed in tension at 1025 ° C. at low speed. The pulling speed of the machine was minimal at 1.9 mm / s. The The strain rate of the sample was 5.4 × 10 ^-2 / s, so in conditions quite close to those referred to in typical stampings of the parts to which the alloy is intended.

The influences of the various treatments on the properties of the materials can be summarized as follows.

The microstructures obtained with standard CIC cycles have a grain size that normally increases substantially with the temperature at which the CIC is performed. However, it is not possible under the operating conditions chosen to obtain a grain size having an ASTM index of less than 8, because of the presence of the PPB at the grain boundaries, which limits the growth of the grains (it is recalled that the ASTM index indicating the grain size is higher the smaller the grain size).

The decoupled cycles for which the preliminary heat treatment has been carried out at 1025 and 1120 ° C make it possible to obtain, after CIC at the same temperature, a microstructure product close to that obtained after the standard CIC cycles carried out at the same temperatures. In contrast, a temperature of

1 160 ° C makes it possible to increase the size of the grains to 6 or 7 ASTM, and there is a partial decoupling between the surfaces of the powder particles and the grain boundaries. This is shown by the comparison between FIGS. 1 and 2, which show the microstructures of two ingots made from the sample powder.

1:

one (FIG. 1) by direct CIC of the powder at 1160 ° C. for 3 hours at 1000 bar; the other (Figure 2) by CIC performed under these same conditions, but preceded, according to the invention, by a heat treatment of the powder and its container, exposed to 1160 ° C for 6 hours at atmospheric pressure.

It is clearly seen that on the ingot produced according to the invention, the size of the grains is much more homogeneous than for the reference, and that the very small grains have disappeared, sign that the SCH has not constituted obstacles to their growth.

Table 2 shows the main results of the mechanical tests carried out on the different samples as a function of the treatments undergone. Rm is the tensile strength, elongation at break and Z necking of the specimen.

Table 2: Mechanical test results on the different samples

The influence of CIC cycle type on forgeability can be commented on as well.

There is an improvement in the forgeability of standard CIC blanks when the standard CIC cycle temperature increases.

Sample 1 has a very poor ductility when the CIC was at 1025 ° C (A =

5.4%). With a CIC at 1160 ° C, the ductility of this alloy 1 is higher (A =

13%). However, the differences between the 1025 and 1120 ° C CICs are not significant, reflecting the weak influence of CIC temperature on forgeability in this temperature range due to the similarity of the microstructures obtained.

With regard to decoupled cycles, the improvement of the forgeability of alloy 1 is only very marked for the cycle at 1160 ° C. In this case, we obtain a value of A of 28%. At lower temperatures, forgeability remains the same as for standard CIC cycles at equivalent temperature. It is thought that this finding may be attributed to the lack of free decoupling between the surfaces of the powder particles and the grain boundaries for these temperatures. The comparison between the results obtained on alloys 1 and 2 makes it possible to evaluate in particular the influence of boron on forgeability. It is particularly pronounced in the case where a standard CIC cycle at 1025 and 1120 ° C is used, if one goes from a trace boron content to a content of 30 ppm. But for the decoupled cycle at 1160 ° C, the effect of boron is not significant.

Above all, the beneficial influence of the decoupled CIC cycle is also evident in the pattern of damage observed after the samples were broken. The low ductilities of the samples obtained with CIC cycles lead to intergranular fracture facies, corresponding substantially to the grains of the original powder. On the contrary, the decoupled CIC cycle according to the invention, provided that it is carried out at 1160 ° C., provides high ductility samples which exhibit a transgranular fracture facies, thanks to the partial decoupling between the powder grains and the particles. grain boundaries. This decoupling partially offsets the damage that normally occurs at the grain boundaries and is a fundamental factor in improving the forgeability of the material.

In general, the inventors were able to extrapolate these results, obtained on the 725 ^® , to the other nickel-base superalloys that can be hardened by double precipitation of gamma and gamma or delta phases.The known alloys of the IN706, IN718 type, IN725 fall into this category.

Their conclusions are that in order to be effective on the forgeability of the material by causing a decoupling between the grains of the initial powder and the grain boundaries of the product after the densification of the powder and its container, the heat treatment prior to the densification must bear the powder for at least 4 hours at a temperature greater than 1140 ^° C., and also less than 10 ° C. or more at the solidus temperature of the superalloy, preferably between 10 and 50 ° C. below the solidus in order to substantially change the PPB without risk of generating defects caused by a local burn. In the case of 725 ^® , this may correspond to a temperature of 1,160 to 1,180 ° C, depending on the precise solidus temperature of the alloy, within the compositional limits set by the usual specifications. temperature being maintained for 12 hours to 30 hours. The optimum is between 30 and 50 ° C below the solidus. It is under these conditions that a sufficient modification of the PPBs is obtained, which significantly decreases their capacity to prevent the growth of the grain during densification of the powder.

The duration of the heat treatment before densification can be up to 30 h depending on the dimensions of the workpiece. Of course, one of the parameters to be taken into account for the optimization of the treatment time is the size of the part to be produced, the processing time being all the greater as the workpiece is thick so that the treatment can relate homogeneously throughout its thickness. Optimally, this treatment time is from 15 to 17 hours so that a depth of treatment of the order of 150 mm, corresponding to what is desirable to achieve on the components of aeronautical turbines of conventional dimensions, is certainly achieved. to which the invention applies in a privileged but not exclusive manner.

Optimally, the filling of the container with the powder is carried out under vacuum, this evacuation being maintained during the densification itself.

Preferably, the heat treatment before densification according to the invention is carried out in an inert atmosphere to avoid forming calamine on the container and oxides within the powder. It can be carried out at atmospheric pressure or under low pressure. It must not cause densification of the powder, or only a low densification of the powder of less than or equal to 15%, preferably less than or equal to 10% of the initial volume, densification of the powder, or at least a portion of the powder. very majority of this operation, to be carried out during the stage which is especially dedicated to it. Beyond such densification during heat treatment, it is difficult or impossible to avoid the harmful effects of SCH and decorations at SCH. A densification greater than about 15%, therefore does not achieve the above purpose of the invention. For this purpose, a pressure of at most 50 bar is generally recommended.

The densification which follows is carried out at a temperature generally identical or comparable to that of the heat treatment for a period of the order of 4 to 16 hours, again depending in particular on the dimensions of the container-powder assembly. A modeling of the densification and the distribution of the temperature in the room during the landing makes it possible to fix the duration of the latter depending on the desired temperature homogeneity. The densification of the powder in the container is followed by a heat treatment in usual ways to achieve the final characteristics of the alloy. When argon has been used as atomizing gas, the usual heat treatment after densification is carried out at a temperature which is at least 30 ° C lower than the densification temperature to avoid the occurrence of porosities due to the presence of Argon in the mixture.

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

After densification, the ingot or billet resulting is conventionally peeled and then shaped hot. This hot shaping generally comprises in particular a forging. This is preferably carried out at a supersolvus temperature, typically for 725 ^® between about 1010 and 1030 ° C, preferably at 1025 ° C. This forging can be followed by forging in a forging tool (die) to give it the desired final geometry. This operation can be carried out in one to three stages, depending on the dimensions of the final piece targeted.

A forging in breath (also called "blowdown"), for example in three steps, is particularly recommended for the preferred applications envisaged, because it allows to calibrate the half-product for stamping and to knead its surface to obtain microstructural features that are as close as possible to those found at the heart of the semi-finished product. It is recalled that the term "forging in blow" a forging during which the billet or the ingot to be forged is placed in an annular piece called "soufle" which, during the forging, allows to constrain radially the billet or the ingot to obtain a microstructural homogeneity of the billet or ingot in the radial directions.

Claims

A process for the preparation of a nickel base superalloy part by powder metallurgy, comprising the following steps:

developing a nickel base superalloy of a composition capable of providing a double precipitation hardening of a gamma phase and a gamma or delta phase;

atomizing a melt of said superalloy to obtain a powder;

sieving said powder to extract particles having a predetermined particle size;

introduction of the powder into a container, possibly under vacuum;

- closing and evacuation of the container;

densification of the powder and the container by pressurizing the assembly to obtain an ingot or a billet; - Hot forming and possibly heat treatment of said ingot or said billet; characterized in that prior to the densification step of the powder and the container, a heat treatment step is carried out for at least 4 hours, preferably for 12 to 30 hours, at a pressure resulting in densification of the lower powder. or equal to 15% of the initial volume, preferably less than or equal to 10% of the initial volume, this treatment taking place at a temperature both greater than 1140 ° C and at least 10 ° C lower than the temperature of solidus of the superalloy.

2. Method according to claim 1, characterized in that the composition of the superalloy is, in percentages by weight:

- 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 <If <0.2%;

- traces <C <0.03%;

- traces <Mg <0.05%;

3. Method according to claim 1 or 2, characterized in that said heat treatment before densification is carried out at a temperature of 10 to 50 ° C below the solidus temperature of the superalloy.

4. Method according to claim 3, characterized in that the heat treatment before densification is carried out at a temperature both greater than 1140 ° C and 30 to 50 ° C lower than the solidus temperature of the superalloy.

5. Method according to claim 4, characterized in that the heat treatment before densification is carried out between 1160 and 1180 ° C for 12 hours to 30 hours at a pressure less than or equal to 50 bar.

6. Process according to claim 5, characterized in that the heat treatment before densification is carried out at atmospheric pressure.

7. Method according to one of claims 1 to 6, characterized in that the densification is carried out by hot isostatic compaction.

8. Method according to one of claims 1 to 7, characterized in that the hot shaping comprises a forging in sulfur.

9. Forged nickel base superalloy, characterized in that it was prepared by the method according to one of claims 1 to 8.

10. Part according to claim 9, characterized in that it is an aeronautical gas turbine component.

1. Part according to claim 9, characterized in that it is a gas turbine component land.