Nickel-based superalloy

Abstract

The present invention relates to a nickel-based superalloy, characterized in that its composition comprises or advantageously consists essentially of, in percent by weight of the total composition: Al 2.5 - 3.8, Co 7.9 - 16.9, Cr 9.7 - 13.1, Mo 2.6 - 4.1, Nb 0 - 0.41, Ta 0 - 1.9, Ti 4.4 - 6.4, W 1.9 - 4.2, B 0.010 - 0.030, C 0.010 - 0.040, Hf 0.20 - 0.40, Zr 0.040 - 0.070, the remainder being Ni together with unavoidable impurities. The invention also relates to the powder of a superalloy according to the invention, to the method for manufacturing parts made of superalloy according to the invention, and to the parts that can be obtained by this method.

Description

[0001] Nickel-based superalloy

[0002] Technical Field

[0003] The present invention relates to the field of nickel-based superalloys for high-temperature applications and developed for turbine disk applications with a dual coarse-grain / fine-grain structure. Increasing engine performance and reducing their specific consumption requires increasing their operating temperature. This results in the need for new materials that are ever more resistant to heat. For turbine disks, the target temperatures are around 800°C in nominal operation with temporary peaks of up to 850°C in the disk rim, i.e. near the blades. In addition, the alloy must also be very resistant to traction and fatigue in the area close to the disk bore, as it is subjected to high stresses when the complete disk + turbine blade system is rotating, but at lower temperatures (below 700°C).This specification can be achieved by combining two routes: on the one hand, the process route, which includes, among other things, the development of the alloy by powder metallurgy and the application of a dual-structure heat treatment (patent FR3043410B1) to optimize the grain size both in the rim and in the disc bore, and on the other hand, the material route, which provides a chemical composition directly adapted to this process and achieves a significant gain in mechanical properties.

[0004] Prior art

[0005] The specifications indicated above can be achieved by combining two routes:

[0006] - on the one hand the process route, which includes among other things the development of the alloy by powder metallurgy and the application of a dual structure heat treatment (patent FR3043410B1) making it possible to optimize the grain size both in the rim and in the bore of the disc;

[0007] - and on the other hand the material route which makes it possible to have a chemical composition directly adapted to this process and to achieve a significant gain in mechanical properties.

[0008] Thus, the heat treatment process for obtaining alloys having such a dual structure is known, as described in patent application FR3043410. This application presents a process for producing structure gradients on a disc-type part via a gradient heat treatment itself. This is an induction heat treatment which makes it possible to carry out a solution treatment at a staged temperature in the part in such a way that:

[0009] - in the hottest zone, the temperature is higher than the dissolution temperature of the phase blocking the grain boundaries, also called the solvus temperature (for nickel-based gamma-gamma prime alloys, the phase in question is the gamma prime phase).

[0010] - in the less hot zone, the temperature is lower than this solvus temperature.

[0011] Thus, in areas where the heat treatment temperature exceeds the solvus temperature of the gamma prime phase, the grains will grow to form a structure favorable to creep and cracking properties, whereas in areas where the heat treatment temperature remains below the solvus temperature, the structure will retain the grain size resulting from forging, which is generally relatively fine and favorable to tensile and fatigue properties.

[0012] Typically, gradient treatment is applied to an existing "classic" alloy whose chemical composition has been optimized to achieve, with a homogeneous structure across the entire part, the best compromise of mechanical properties required. Thus, the most efficient disc superalloys have a target operating temperature of 760°C, with possible peaks at 800°C. However, existing chemical compositions are not optimal for configuring gradient treatment. Improving part performance therefore requires defining a chemical composition specific to gradient treatment of the part.

[0013] A nickel-based alloy composition was jointly defined by Safran Tech and ONERA to provide an initial response to this problem (patent application FR3104613A1). This alloy was specifically designed for a dual-structure turbine disk application, where the grain size is optimized according to the mechanical characteristics required in each area of ​​the part: the bore has a fine-grained microstructure to optimize its tensile strength and fatigue life, while the rim has a coarser-grained microstructure to optimize its resistance to creep and crack propagation. The coarse-grained microstructure is obtained through a local supersolvus y' (gamma prime) heat treatment. This alloy has a density of 8.24 g / cm 3 , which is rather low compared to other known alloys (8.34 g / cm 3for the N19 powder metallurgy alloy) and therefore very interesting for this type of application. However, this composition has a relatively narrow y' (gamma prime) supersolvus heat treatment window: there is only about ten degrees of difference between the y' (gamma prime) solvus and the solidus of the alloy. From an industrial point of view, this heat treatment is feasible but the margin of error is small. In addition, even if the creep resistance at 850 °C of this alloy is very interesting, there are levers to optimize it in order to further increase the service life of the parts at this extreme temperature.

[0014] Another nickel-based alloy composition was already known from the prior art (application EP1840232) with higher niobium contents. However, their higher Nb content does not allow the precipitation of M23C6 carbides at a temperature of 850 °C which remain stable up to 900 °C or even 920 °C because Nb lowers the solvus temperature of these carbides below 900 °C, or even below 800 °C. However, such carbides improve the creep resistance of the alloy.

[0015] Statement of the invention

[0016] The inventors surprisingly discovered that it was possible to obtain such properties using a nickel-based superalloy which could possibly contain tantalum but in a limited content and containing lower niobium contents than the alloy described in FR3104613 and in EP1840232 and in particular than the best performing alloy on the market (N19 / SMO43).

[0017] The present invention therefore relates to a nickel-based superalloy, characterized in that its composition comprises, advantageously consists essentially of, in particular consists of, in percentages by weight of the total composition:

[0018] Aluminum: 2.5 - 3.8, preferably 2.8 - 3.5;

[0019] Cobalt: 7.9 - 16.9, preferably 8.2 - 16.6;

[0020] Chromium: 9.7 - 13.1, preferably 10.0 - 12.8;

[0021] Molybdenum: 2.6 - 4.1, preferably 2.6 - 3.8;

[0022] Niobium: 0 - 0.41;

[0023] Tantalum: 0 - 1.9, preferably 0 - 1.8;

[0024] Titanium: 4.4 - 6.4, preferably 4.6 - 6.1;

[0025] Tungsten: 1.9 - 4.2, preferably 2.2 - 4.0;

[0026] Boron: 0.010 - 0.030, preferably 0.010 - 0.020;

[0027] Carbon: 0.010 - 0.040, preferably 0.015 - 0.035;

[0028] Hafnium: 0.20 - 0.40, preferably 0.20 - 0.35;

[0029] Zirconium: 0.040 - 0.070, preferably 0.045 - 0.065;

[0030] Nickel: balance plus unavoidable impurities; with 12.5 < Al + Ti + Nb + Ta < 14 at.%, the contents of these elements being expressed in atomic percentages; with 0.85 < Al / (Ti + Nb + Ta) < 1.2, the contents of these elements being expressed in atomic percentages; and with Mo + W > 2.5 at.%, the contents of these elements being expressed in atomic percentages.

[0031] The composition of the nickel-based superalloy according to the invention therefore contains the following elements:

[0032] - Major addition elements: Co, Cr, Mo, W, Al, Ti, Ta;

[0033] - Minor addition elements (concentration less than one mass percent): Hf, C, B, Zr, Nb.

[0034] The composition according to the invention thus comprises, in % by weight relative to the total weight of the composition, aluminium (Al) in a content within the range 2.5 - 3.8, advantageously 2.8 - 3.5, in particular 2.82 - 3.43, more particularly 2.9 - 3.3. The aluminium content is measured with an uncertainty of ±0.10.

[0035] The composition according to the invention further comprises, in % by weight relative to the total weight of the composition, cobalt (Co) in a content in the range 7.9 - 16.9, advantageously 8.2 - 16.6, in particular 8.2 - 15.6, more particularly 12.0 - 14.0. The cobalt content is measured with an uncertainty of ±0.4.

[0036] The composition according to the invention further comprises, in % by weight relative to the total weight of the composition, chromium (Cr) in a content within the range 9.7 - 13.1, advantageously 10.0 - 12.8, in particular 10.8 - 12.8, more particularly 10.9 - 12.7, even more particularly 11.65 - 12.7. The chromium content is measured with an uncertainty of ±0.35.

[0037] The composition according to the invention further comprises, in % by weight relative to the total weight of the composition, molybdenum (Mo) in a content in the range 2.6 - 4.1, advantageously 2.6 - 3.8, in particular 2.7 - 3.25. The molybdenum content is measured with an uncertainty of ±0.11.

[0038] The composition according to the invention further comprises, in % by weight relative to the total weight of the composition, niobium (Nb) in a content in the range 0 - 0.41. In an advantageous embodiment, the composition is free of niobium. In another advantageous embodiment, the composition contains niobium in an amount of at most 0.41% (in % by weight relative to the total weight of the composition), in particular in the range 0.39 - 0.41%, more particularly 0.40 - 0.41%. The niobium content is measured with an uncertainty of ±0.10.

[0039] The composition according to the invention further comprises, in % by weight relative to the total weight of the composition, tantalum (Ta) in a content in the range 0 - 1.9, advantageously 0 - 1.8. In an advantageous embodiment, the composition is free of tantalum. Indeed, tantalum (Ta) participates in the strengthening of the gamma prime phase but has the effect of increasing the density of the alloy. In another advantageous embodiment, the composition contains tantalum in an amount of at most 1.9% (in % by weight relative to the total weight of the composition), in particular at most 1.6%, more particularly in the range 1 - 1.6%. The tantalum content is measured with an uncertainty of ±0.15.

[0040] The composition according to the invention further comprises, in % by weight relative to the total weight of the composition, titanium (Ti) in a content in the range 4.4 - 6.4, advantageously 4.6 - 6.1. The titanium content is measured with an uncertainty of ±0.15.

[0041] The composition according to the invention further comprises, in % by weight relative to the total weight of the composition, tungsten (W) in a content in the range 1.9 - 4.2, advantageously 2.2 - 4.0. The tungsten content is measured with an uncertainty of ±0.16. The composition according to the invention also comprises, in % by weight relative to the total weight of the composition, boron (B) in a content in the range 0.010 - 0.030, in particular 0.010 - 0.020. The boron content is measured with an uncertainty of ±0.003.

[0042] The composition according to the invention also comprises, in % by weight relative to the total weight of the composition, carbon (C) in a content in the range 0.010 - 0.040, advantageously 0.015 - 0.035, in particular 0.020 - 0.035. The carbon content is measured with an uncertainty of ±0.003.

[0043] The composition according to the invention further comprises, in % by weight relative to the total weight of the composition, hafnium (Hf) in a content in the range 0.20 - 0.40, advantageously 0.20 - 0.35. The hafnium content is measured with an uncertainty of ±0.002.

[0044] The composition according to the invention further comprises, in % by weight relative to the total weight of the composition, zirconium (Zr) in a content in the range 0.040 - 0.070, advantageously 0.045 - 0.065, in particular 0.055 - 0.060. The zirconium content is measured with an uncertainty of ±0.003.

[0045] The nickel-based alloy according to the invention contains the elements cobalt, aluminum, titanium, niobium and tantalum intended to form a hardening Y' (gamma prime) precipitation of ordered Ll2 structure and composition (Ni,Co)3(AI,Ti,Nb,Ta). The tantalum content is however limited in order to avoid excessively increasing the density of the alloy, the target of which is less than or equal to 8.30 g / cm 3 .

[0046] The niobium content is limited to 0.41% by mass because this element stabilizes the MC carbides (with M=Ti, Nb) to the detriment of the M23C6 carbides (with M=Cr, Mo or W). The presence of too much niobium results in a lowering of the solvus temperature of the M carbides 23 C6 below 900 °C. The minimum solvus value of carbides M 23 CÔ is thus advantageously fixed at 900 °C (solvus M 23CÔ > 900 °C), preferably at 920 °C (solvus M23C6> 920 °C), in order to ensure their stability at 850 °C and to maintain a margin in the event of future increases in the temperature peaks of the disks in operation.

[0047] Carbon, boron, zirconium, and hafnium help strengthen grain boundaries at high temperatures. Carbon also helps form M23C6 carbides (with M=Cr, Mo, or W).

[0048] Molybdenum and tungsten contribute to the hot mechanical strengthening of the alloy. Their contents are optimized to maximize this strengthening while limiting the precipitation of TCP (Topologically Close Packed) phases of type o or p, which can degrade the mechanical resistance. These elements also participate in the formation of M23C6 carbides (with M=Cr, Mo or W). In particular, the presence of W participates in the formation and stabilization of M23C6 carbides. The sum of their contents expressed in atomic percentages must respect the rule Mo + W > 2.5 at.% (the contents of these elements being expressed in atomic percentages) in order to ensure a minimum of solid solution strengthening of the y (gamma) matrix.

[0049] Chromium is present because it participates in the formation of M23C6 carbides (with M=Cr, Mo or W) and helps improve the alloy's resistance to oxidation by forming a continuous passivation layer on the surface. However, its content is limited to avoid promoting the precipitation of the TCP phases of which it is a component.

[0050] The hot mechanical strength of the alloy is favored by a molar fraction of y' (gamma prime) precipitates between 50% and 56%. The elements Al, Ti, Nb and Ta expressed in atomic percentages meet the criterion 12.5 < Al + Ti + Nb + Ta < 14%at. In order to favor the precipitation of the y' (gamma prime) phase compared to the q-Ni3Ti phase, which is undesirable from the point of view of mechanical properties, the elements Al, Ti, Nb and Ta meet the criterion 0.85 < Al / (Ti + Nb + Ta) < 1.2, the contents of these elements being expressed in atomic percentages. These criteria mean that the titanium, niobium and tantalum contents in the y' (gamma prime) phase are to be optimized to maximize the hot mechanical strengthening of the alloy, while avoiding favoring the formation of the q phase to the detriment of y' (gamma prime).

[0051] The unavoidable impurities of the composition according to the invention originate from the manufacturing steps of the superalloy or from impurities present in the raw materials used for the manufacture of the superalloy. All the conventional impurities encountered in nickel-based superalloys are found. In particular, they are chosen from the group consisting of manganese, silicon, vanadium, sulfur, phosphorus, copper, lead, iron, bismuth, nitrogen, oxygen, hydrogen and mixtures thereof. They can constitute up to 1% by mass of the alloy and each does not represent more than 0.5% by weight of the total composition. Generally, the content of impurities in the alloy is measured with an uncertainty of 10%.

[0052] In an advantageous embodiment, the nickel-based superalloy according to the invention is characterized in that its composition comprises, advantageously consists essentially of, in particular consists of, in percentages by weight of the total composition: Aluminum: 2.8 - 3.5;

[0053] Cobalt: 8.2 - 15.6;

[0054] Chromium: 10.8 - 12.8;

[0055] Molybdenum: 2.7 - 3.25;

[0056] Niobium: 0 - 0.41;

[0057] Tantalum: 0 - 1.9;

[0058] Titanium: 4.6 - 6.1;

[0059] Tungsten: 2.2 - 4.0;

[0060] Boron: 0.010 - 0.030, preferably 0.010 - 0.020;

[0061] Carbon: 0.010 - 0.040, preferably 0.015 - 0.035; Hafnium: 0.20 - 0.40, preferably 0.20 - 0.35;

[0062] Zirconium: 0.040 - 0.070, preferably 0.045 - 0.065;

[0063] Nickel: balance plus unavoidable impurities; with 12.5 < Al + Ti + Nb + Ta < 14 at.%, the contents of these elements being expressed in atomic percentages; with 0.85 < Al / (Ti + Nb + Ta) < 1.2, the contents of these elements being expressed in atomic percentages; and with Mo + W > 2.5 at.%, the contents of these elements being expressed in atomic percentages.

[0064] In particular, the composition of the superalloy according to the present invention may be as indicated in the following Table 1.

[0065] [Table 1]

[0066] Hot creep strength is enhanced by the precipitation of M carbides 23C6 in the grain boundaries, with M=Cr, Mo or W. These carbides precipitate discreetly during tempering carried out at temperatures lower than or equal to 870 °C, preferably lower than or equal to 850 °C for a duration of between 2 h and 16 h, and have a size lower than 5 pm, preferably lower than 1 pm. The quantity of M23C6 carbides (with M=Cr, Mo or W) calculated using the CALPHAD (CALculation of PHAse Diagrams) method is advantageously between 0.4% and 1% mol at 850 °C (0.4% < M 23 C6 < 1 mol% at 850 °C), preferably between 0.5% and 0.8 mol% at 850 °C (0.5% < M 23 C6 < 0.8 mol% at 850 °C), this is to obtain a sufficient population of carbides to ensure the desired hardening while avoiding saturating the grain boundaries, which would then promote unwanted intragranular precipitation and weaken the alloy. The solvus temperature of the carbides M 23C6 (with M=Cr, Mo or W) is advantageously greater than 900°C in order to avoid putting the carbides back into solution in the event of temperature peaks above 850°C during operation, and preferably greater than 920°C in order to compensate for a possible future increase in temperature at the time of these peaks.

[0067] In an advantageous embodiment, the superalloy according to the invention has a density of less than 8.50 g / cm 3 , advantageously less than 8.30 g / cm 3The room temperature density of each superalloy was estimated using a modified version of Hull's formula. This empirical equation was proposed by Hull (FC Hull, Metal Progress, November 1969, ppl39-140). The empirical equation is based on the law of mixtures and includes correction terms derived from a linear regression analysis of experimental data (chemical compositions and measured densities) for 235 superalloys and stainless steels. This Hull formula was modified from data for 272 nickel-based, cobalt-based, and iron-based superalloys. The modified Hull formula is:

[0068] [Math. 1]

[0069] D-=-100- / [S ■(%X / DX)]-+- 1 -AX-x-%X where DX are the densities of the elements X (Cr, Ni, etc...), and D is the density of the superalloy, the densities being expressed in g / cm 3 , where AX is a coefficient expressed in g / cm 3of the elements X (Cr, Ni, etc...), and are as follows: ANi = -0.0011; AAI = 0.0622; ATa = 0.0121; ATi = 0.0317; ANb = 0.011; ACo= -0.0001; ACr = -0.0034; AMo = 0.0033; AW = 0.0033; AHf = 0.0156, and where %X are the contents, expressed in mass percentages, of the elements X of the superalloy (Cr, Ni, etc...).

[0070] In another advantageous embodiment, the superalloy according to the invention has metallurgical stability (i.e. absence of TCP phases - Topological Compact Phases) up to 800 - 850 °C.

[0071] In yet another advantageous embodiment, the superalloy according to the invention has good oxidation resistance. In yet another advantageous embodiment, the superalloy according to the invention has a difference between the solidus temperature y (gamma) and the solvus temperature y' (gamma prime) sufficiently large for carrying out the heat treatment, advantageously a difference of at least 20 °C ([solidus y - solvus y'] - 20 °C).

[0072] In an advantageous embodiment, the nickel-based superalloy according to the invention is free of tantalum and / or niobium, advantageously tantalum and niobium.

[0073] In another advantageous embodiment, the nickel-based superalloy according to the invention comprises tantalum and / or niobium.

[0074] The present invention further relates to the superalloy powder according to the invention. Indeed, the superalloy according to the invention may be in the form of a powder with a particle size of between 10 μm and 100 μm. The present invention further relates to a method for manufacturing a nickel-based superalloy powder according to the invention comprising the following steps:

[0075] A- mixture of elementary or pre-alloyed raw materials;

[0076] B- melting of the mixture obtained in step A), advantageously in a vacuum induction furnace (VIM);

[0077] C- gas atomization, advantageously with argon, of the product obtained in step B) so as to obtain a powder, advantageously predominantly spherical (i.e. without sharp angles);

[0078] D- sieving of the powder obtained in step C), advantageously under an inert atmosphere, so as to obtain the targeted particle size;

[0079] E- recovery of the powder obtained.

[0080] The method may further have the following successive steps, after step E:

[0081] F - containerization of the powder (in particular a cylindrical container, more particularly under vacuum); G - hot compaction of the filled container, for example by Hot Uniaxial Compaction (HUC) in a press, or by Hot Isostatic Compaction (HIP), in particular in an autoclave;

[0082] H - hot spinning of the powder container (the container passes through a die which reduces its diameter. This allows us to obtain a cylindrical bar with a 100% dense, fine and homogeneous structure),

[0083] I - peeling (to remove the peripheral sheath corresponding to the walls of the initial container) and controls, in particular by ultrasound;

[0084] J - cutting into sections.

[0085] The particle size of the powder is thus adapted according to the manufacturing technology of the parts based on superalloy powder envisaged. The particle size ranges used for the different manufacturing processes vary according to the technology, the equipment and the intended applications. In general, if all the applications are combined, the powder used for these processes will have more or less wide particle size distributions in number between 10 μm and 100 μm. The present invention further relates to a method for manufacturing a part, in particular turbines, made of superalloy according to the invention or of superalloy powder according to the invention, characterized in that it comprises the following steps: a- forging, b- gradient heat treatment of the part obtained in step a), c- final heat treatment of the entire part with dual microstructure obtained in step b) d- recovery of the part obtained in step c).

[0086] Forging step a) can be carried out by methods well known to those skilled in the art, in particular on the sections obtained in step J. This can be, for example, die-forging (such as, for example, isothermal forging). This technique is well known to those skilled in the art. This step a) makes it possible to obtain a superalloy part.

[0087] Step b) of gradient heat treatment can be implemented using the method and device described in patent application FR3043410.

[0088] Advantageously, step b) of gradient heat treatment of the part obtained in step a) comprises: bl- a first heating of a zone of the part to a first temperature (Tl) higher than the solvus temperature of the gamma prime phase of said superalloy and lower than the melting temperature of said superalloy.

[0089] It may thus comprise heating an area of ​​the part (for example the rim of the disc) to a first temperature (Tl) at least 5°C higher than the solvus temperature of the gamma prime phase of said superalloy (advantageously between +5°C and +15°C relative to the solvus temperature of the gamma prime phase of said superalloy) and lower than the melting temperature of said superalloy (this is therefore a supersolvus treatment, i.e. a supersolvus solution treatment). The duration of this treatment may be between 1 hour and 8 hours.

[0090] Gradient heat treatment may for example be carried out by local induction heating or by any method or device described in FR3043410.

[0091] In the case where the part is a turbine disc, the area of ​​the part undergoing the first temperature (Tl) consists of the rim area of ​​the disc, the rest of the part not being impacted by this treatment.

[0092] Thus, the supersolvus treatment (temperature Tl) makes it possible to use 100% of the hardening potential associated with the gamma prime phase to maintain hardening that is still effective at high temperatures (800 °C and can even withstand peaks at 850 °C), while increasing the grain size to improve the alloy's resistance to hot creep and crack propagation. The grain size is thus advantageously greater than or equal to 15 pm (measured by the intercept method). The grain size is advantageously 40 pm on average to maintain good fatigue resistance.

[0093] Step b) therefore makes it possible to obtain a part with a dual or gradient microstructure, i.e. one that does not have a homogeneous microstructure, in particular one whose grain size is not the same, depending on whether one is in the area of ​​the part that has undergone the supersolvus treatment or that which has not been impacted by the supersolvus treatment. Thus the area of ​​the part that has undergone the supersolvus treatment contains large grains and that which has not been impacted by the supersolvus treatment contains fine grains.The part with dual or gradient microstructure according to the invention therefore contains large grains and fine grains, advantageously large grains having a size greater than or equal to 15 μm, in particular in the area of ​​the part having undergone the supersolvus treatment, for example in the rim of the disc, and fine grains, in particular resulting from forging, having a size less than 15 μm, in particular in the area of ​​the part not having been impacted by the supersolvus treatment, for example in the bore of the disc. The grain size is measured by the intercept method.

[0094] Step c) of the method according to the invention may comprise the following successive steps: cl- solution treatment of the entire part obtained in step b) at a temperature lower than the solvus temperature of the gamma prime phase of said superalloy (T2) (this is therefore a subsolvus treatment or subsolvus solution treatment), for example at a temperature T2 of between -80°C and -20°C relative to the solvus of the gamma prime phase; c2- cooling or quenching of the entire part obtained in step cl); advantageously the cooling rate is adapted to the mass of the part in order to obtain an optimal size and distribution of the gamma prime hardening phase, more advantageously it is of the order of 100°C / min (the rate may vary from 80 to 120°C / min depending on the location).c3- at least one tempering treatment of part or all of the part obtained in step c2), advantageously at a temperature in the range 730°C - 870°C, so as to precipitate carbides of type M23C6 with M=Cr, Mo or W.

[0095] The final heat treatment of step c) is therefore the classic heat treatment of gamma / gamma prime alloys. The objective of this treatment is to treat the structure not impacted by the gradient treatment in order to have, in these areas, a final structure and therefore mechanical properties equivalent to the desired level. Indeed, in the area treated only in subsolvus (temperature T2), the grain size remains small, advantageously less than 10 pm (measured by the intercept method), which makes it possible to obtain good tensile and fatigue properties at average temperatures, for example below 750 °C. This second solution treatment thus makes it possible to refine the size of the y' (gamma prime) precipitates throughout the part, while maintaining a fine grain size in the areas that have not undergone the first solution treatment.The resulting part thus has a grain size of 6-7 ASTM (28 - 40 pm) in the rim intended to be stressed in creep at very high temperatures, and a grain size of 10-12 ASTM (5 - 10 pm) in the rest of the part which is mainly stressed in tension and fatigue at lower temperatures.

[0096] The duration of stage cl) can be in the range 1 hour - 8 hours.

[0097] In particular, step c3) may consist of one or more tempering treatments, advantageously two tempering treatments. This may be a single-stage or double-stage tempering treatment. Thus, a relatively hot final tempering treatment (for example in the range 730°C - 870°C, in particular 730°C-850°C, for example a first tempering at a temperature around 850°C, more particularly 850°C, followed by a second tempering at a temperature between 730°C and 800°C, advantageously around 800°C, in particular 800°C, makes it possible to stabilize the microstructure of the part at high temperature. It also makes it possible to relax the residual stresses resulting from the quenching associated with the treatment at temperature T2.

[0098] The duration of step c3) may be in the range 2 hours - 24 hours (for example a first tempering for a duration of 4 hours - 8 hours, in particular at a temperature of 850°C, followed by a second tempering for a duration of 4 hours - 16 hours, in particular at a temperature of 800°C).

[0099] The tempering and quenching treatments are carried out using techniques well known to those skilled in the art.

[0100] The tempering heat treatments are intended to precipitate M23C6 type carbides with M=Cr, Mo or W and to stabilize the y' (gamma prime) precipitate populations. Thus the first tempering treatment is carried out to precipitate M23C6 carbides with M=Cr, Mo or W in nodular form and the second tempering treatment is carried out to stabilize the y' (gamma prime) precipitate population at a temperature close to the target operating temperature for the hottest part of the disc.

[0101] These tempering heat treatments can be carried out uniformly over the entire part, but it is also possible to treat only part of the part in order to optimize the tempering according to the characteristics targeted in each area. For example, only the area of ​​the part that has undergone the supersolvus treatment, such as the rim of the disc, can be treated at a tempering temperature around 850 °C (advantageously 850 °C) to precipitate the carbides M 23 C6, before treating the entire part at a temperature between 730 °C and 800 °C to stabilize the precipitation / (gamma prime) in the bore of the part. This optimizes creep resistance in the rim thanks to the precipitation of M carbides 23C6 at 850 °C, but without increasing the size of the precipitates / in the bore of the disc and thus keeping the precipitates / as fine as possible in order to optimize the tensile strength of this area.

[0102] The present invention finally relates to a part made of superalloy according to the present invention or of superalloy powder according to the present invention, having a dual microstructure, advantageously capable of being obtained by the method according to the present invention. Advantageously, it is a turbomachine part, more advantageously a turbine part, in particular a turbine disk, a compressor disk, a ring, a flange, or a turbine casing, more particularly a turbine disk, for example of aircraft and / or helicopter engines.

[0103] The part according to the invention therefore has a dual or gradient microstructure, that is to say that it does not have a homogeneous microstructure. In particular, the size of the grains of the part is not the same depending on the area of ​​the part. It therefore contains coarse grains and fine grains, advantageously coarse grains having a size greater than or equal to 15 μm, advantageously 40 μm on average, and fine grains having a size less than 15 μm. Thus, advantageously, one area of ​​the part contains coarse grains and the rest of the part and / or another area of ​​the part contains fine grains. Thus, in the case where the part is a turbine disk, the rim area of ​​the disk is coarse-grained, advantageously having grains having a size greater than or equal to 15 μm, advantageously 40 μm on average, and the bore area of ​​the disk is fine-grained, advantageously having grains having a size less than 15 μm.Grain size is measured by the intercept method.

[0104] Advantageously, the coarse-grained area of ​​the part has good creep resistance according to standard NF EN ISO 204 August 2009 at a temperature of 850°C, more advantageously a duration greater than 37 hours, at 0.2% elongation under a stress of 200 MPa and a temperature of 850°C.

[0105] Advantageously, the fine-grained zone of the part has good tensile strength according to Standard NF EN 2002-001 / 06 at a temperature below 750 °C, in particular an elastic limit at 20 °C greater than 1100 MPa.

[0106] The present invention will be better understood in light of the description of the figures and examples which follow. The examples are given for informational purposes only and are not limiting.

[0107] Brief description of the drawings

[0108] Figure 1 shows the diagram of a forged disc with a gradient microstructure, with large grains of 28 - 40 pm in the rim and small grains of 5 - 10 pm in the rest of the part. The transition zone between the two microstructures is shown schematically by dotted lines.

[0109] EXAMPLE

[0110] Nickel-based superalloys according to the invention (examples 1 to 8 and 10 to 17) were manufactured according to the following process: vacuum casting of an ingot, then atomization under argon of this ingot, sieving at 53 μm, placing the powders in containers with degassing, then hot spinning of these powders in the form of a bar. The nickel-based superalloy of example 9 was produced by VIM casting, according to a process well known to those skilled in the art. They were compared with 5 superalloys: 1 superalloy according to the example of application FR3104613 (ex FR3104613A1), 1 superalloy according to the example of application EP1840232 (ex EP1840232) and 3 superalloys having excessively high niobium contents (counter-examples 18 to 20)

[0111] The alloys manufactured according to the invention have the chemical composition in % by mass indicated in Table 1 above. The 5 comparative superalloys have the composition indicated in Table 2 below.

[0112] [Table 2]

[0113] The calculated solvus / (gamma prime) and solidus temperatures for the different alloys and the solidus-solvus / (gamma prime) gap are shown in Table 3 below. The density estimated using a modified version of Hull's formula as shown above is also shown in Table 3. The value of the formulas Al + Ti + Nb + Ta in at% and Al / (Ti + Nb + Ta) in at% is also shown in Table 3.

[0114] A portion of the bar was then treated at a temperature between 1180 °C and 1200 °C for a duration of 2 h (supersolvus treatment) followed by cooling at around 30 °C / min. The entire bar was then treated at a temperature between 1145 °C and 1165 °C for a duration of 2 h (subsolvus treatment) followed by quenching at a rate of 100 °C / min and tempering at a temperature of 850 °C for a duration of 4 h to 8 h followed by a second tempering at a temperature of 800 °C for a duration of 4 h to 16 h. The quantity of M23C6 carbides calculated using the CALPHAD (CALculation of PHAse Diagrams) method is shown in Table 3. The solvus temperature of M23C6 carbides is also shown in Table 3.

[0115] [Table 3]

[0116] Compared to current solutions, in this case the FR3104613 alloy or the EP1840232 alloy, the advantages obtained by this invention are twofold: the first advantage is easier industrial implementation thanks to the widening of the grain enlargement heat treatment window which goes from 10 °C to 20 °C or more, and the second advantage is an increase in the lifetime of the parts at very high temperature (850 °C). The alloys according to the invention have an improvement in creep resistance at 850 °C to increase the lifetime of the part with respect to creep. At this temperature, the grain boundaries are considered to be the weak points of the microstructure. This is why grain size enlargement is used to reduce the density of the grain boundaries and thus limit their impact. In this new alloy, this enlargement is also associated with a precipitation of M23C6 carbides in the grain boundaries.These carbides are intended to strengthen grain boundaries in order to limit creep-diffusion mechanisms and slow down the deformation of the metal by creep, and thus allow parts to operate longer at 850°C.

Claims

Demands 1. Nickel-based superalloy, characterized in that its composition comprises, advantageously consists essentially of, the following percentages by weight of the total composition: Aluminium: 2.5 - 3.8, advantageously 2.8 - 3.5; Cobalt: 7.9 - 16.9, advantageously 8.2 - 16.6; Chrome: 9.7 - 13.1, advantageously 10.0 - 12.8; Molybdenum: 2.6 - 4.1, advantageously 2.6 - 3.8; Niobium: 0 - 0.41; Tantalus: 0 - 1.9, advantageously 0 - 1.8; Titanium: 4.4 - 6.4, advantageously 4.6 - 6.1; Tungsten: 1.9 - 4.2, advantageously 2.2 - 4.0; Boron: 0.010 - 0.030, advantageously 0.010 - 0.020; Carbon: 0.010 - 0.040, advantageously 0.015 - 0.035; Hafnium: 0.20 - 0.40, advantageously 0.20 - 0.35; Zirconium: 0.040 - 0.070, advantageously 0.045 - 0.065; Nickel: balance as well as the inevitable impurities; with 12.5 < Al + Ti + Nb + Ta < 14 at.%, the contents of these elements being expressed in atomic percentages; with 0.85 < Al / (Ti + Nb + Ta) < 1.2, the contents of these elements being expressed in atomic percentages; and with Mo + W > 2.5 at.%, the contents of these elements being expressed in atomic percentages.

2. Nickel-based superalloy according to claim 1, characterized in that its composition comprises, advantageously consists essentially of, in weight percentages of the total composition: Aluminium: 2.8 - 3.5; Cobalt: 8.2 - 15.6; Chrome: 10.8 - 12.8; Molybdenum: 2.7 - 3.25; Niobium: 0 - 0.41; Tantalus: 0 - 1.9; Titanium: 4.6 - 6.1; Tungsten: 2.2 - 4.0; Boron: 0.010 - 0.030, advantageously 0.010 - 0.020; Carbon: 0.010 - 0.040, advantageously 0.015 - 0.035; Hafnium: 0.20 - 0.40, advantageously 0.20 - 0.35; Zirconium: 0.040 - 0.070, advantageously 0.045 - 0.065; Nickel: balance as well as the inevitable impurities; with 12.5 < Al + Ti + Nb + Ta < 14 at.%, the contents of these elements being expressed in atomic percentages; with 0.85 < Al / (Ti + Nb + Ta) < 1.2, the contents of these elements being expressed in atomic percentages; and with Mo + W > 2.5 at.%, the contents of these elements being expressed in atomic percentages.

3. Nickel-based superalloy according to any one of claims 1 or 2, characterized in that the temperature difference [solidus y - solvus Y'] - 20 °C.

4. Nickel-based superalloy according to any one of claims 1 to 3, characterized in that the M23C6 carbide content is 0.4% < M23C6 < 1% molar at 850 °C, with M=Cr, Mo or W, advantageously 0.5% < M23C6< 0.8% molar at 850 °C, with M=Cr, Mo or W.

5. Nickel-based superalloy according to any one of claims 1 to 4, characterized in that the solvus temperature of M23C6 carbides > 900 °C, with M=Cr, Mo or W, advantageously the solvus temperature of M23C6 carbides > 920 °C, with M=Cr, Mo or W.

6. Nickel-based superalloy according to any one of claims 1 to 5, characterized in that it is free of tantalum and / or niobium.

7. Nickel-based superalloy according to any one of claims 1 to 5, characterized in that it comprises tantalum and / or niobium.

8. Powder of a superalloy according to any one of claims 1 to 7.

9. A method for manufacturing a part in superalloy according to any one of claims 1 to 7 or in superalloy powder according to claim 8, characterized in that it comprises the following steps: a- forging, b- gradient heat treatment of the part obtained in step a), c- final heat treatment of the whole part with dual microstructure obtained in step b) d- recovery of the part obtained in step c).

10. A process according to claim 9, characterized in that step b) of gradient heat treatment of the part obtained in step a) comprises: bl- a first heating of an area of ​​the part to a first temperature (Tl) higher than the solvus temperature of the prime gamma phase of said superalloy and lower than the melting temperature of said superalloy.

11. A process according to any one of claims 9 or 10, characterized in that the final heat treatment step c) comprises the following successive steps: c1- dissolving the whole part obtained in step b) at a temperature below the solvus temperature of the gamma prime phase of said superalloy (T2); c2- cooling or quenching the whole part obtained in step c1); c3- at least one tempering treatment of part or all of the part obtained in step c2), advantageously at a temperature in the range of 730 °C - 870 °C, so as to precipitate M23C6 type carbides with M=Cr, Mo or W.

12. Part in superalloy according to any one of claims 1 to 7 or in superalloy powder according to claim 8, having a dual microstructure, advantageously capable of being obtained by the process according to any one of claims 9 to 11.

13. Part according to claim 12 characterized in that it is a turbomachine part, advantageously a turbine part, in particular a turbine disc, a compressor disc, a ring, a flange or a turbine housing.

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