JP2012517524A - Method for manufacturing parts made from nickel-based superalloys and corresponding parts - Google Patents

Abstract

  The present invention relates to a method for manufacturing a blank of a part made from a Ni-based superalloy, wherein the alloy is manufactured and heat-treated, wherein the superalloy is at least 2.5%. A first step of 850 to 1000 ° C. for at least 20 minutes to precipitate the phase δ on the particle assembly, with a total amount of Nb and Ta, and a partial portion of the phase δ obtained in the first step Including a second step at a temperature higher than the temperature of the first step to allow melting, a third step at a temperature lower than the temperature of the first step and optionally one or more additional steps, and curing It relates to a method characterized in that the heat treatment is carried out in several steps with an aging treatment allowing the precipitation of the phases γ ′ and γ ″. The invention further relates to corresponding parts

Description

  The present invention relates to nickel-based superalloys, and more particularly to a heat treatment method that can be advantageously applied to a portion thereof, particularly to improve their creep resistance and tensile strength.

  “Nickel-based superalloy” refers to an alloy in which Ni comprises at least 50% by weight of the composition (all percentages shown herein are percentages by weight).

More particularly, the present invention has a double precipitation, ie, a δ phase (Ni ₃ Nb-δ or Ni ₃ Ta— of −800 to 1050 ° C.) because the total content of niobium and tantalum exceeds 2.5%. δ) grain boundary precipitation;
-Γ '(Ni ₃ (Al-Ti) -γ') and / or γ "(Ni ₃ Nb-γ" or Ni ₃ Ta-γ ") type hardening during aging carried out at about 600 to 800 ° C. The present invention relates to a heat treatment method applicable to an alloy capable of causing grain boundary precipitation of a phase.

  This is especially true for the alloy NC19FeNb, commercially known as INCONEL 718®, and alloys derived therefrom or equivalent alloys such as 625, 718Plus and 725.

  Experience has shown that in the aviation and terrestrial gas turbine industry where nickel-based resistive alloys have a variety of applications, the fatigue strength of the alloys is one of the most important factors for turbine disk and shaft sizing. .

  The 718 alloy is relatively low cost due to the absence of cobalt in its composition, and know-how about its manufacture and conversion has been acquired, so this alloy is an advanced material used at temperatures up to around 650 ° C. A special rating is given among the alloys of properties. However, the improvement in turbomachinery yield and performance is interpreted as an increase in the temperature of the combustion chamber output, and thus the creep resistance of 718 alloy is increased to increase the possibility of long-term use at temperatures up to 650 degrees. Improvement is necessary. Therefore, improving the creep resistance of 718 alloy while maintaining a fine particle microstructure (> 7 ASTM) so as not to impair fatigue strength is a major industrial concern. It is recalled that the ASTM standard that defines the criteria for particle size evaluation prescribes that the smaller the particle, the larger the ASTM number to which it belongs.

  Two different thermomechanical processing methods are known and are currently used to improve the fatigue properties of 718 alloy.

According to a first option as described in French patent application No. 2089069, after performing a thermomechanical treatment to precipitate the Ni ₃ Nb-δ phase at the grain boundaries, it is below the melting temperature of the Ni ₃ Nb-δ phase. The alloy was recrystallized at a temperature of 5 and was chosen to prevent grain growth by using the Ni ₃ Nb-δ phase precipitated at the grain boundaries during recrystallization. By using this method, it is possible to obtain a recrystallized structure having a very fine grain size of ASTM 10 or more. Their fatigue properties are improved, but their creep resistance is insufficient. In fact, the presence of a Ni ₃ Nb-δ phase having an orthorhombic structure is preferred because it is metastable by fixing niobium and limits the formation of a Ni ₃ Nb-γ ″ hardened phase having a central square structure. It is known that the Ni ₃ Nb-γ "hardened phase improves creep resistance by slowing dislocations in the crystal lattice.

Similarly, the presence of Ni3Ta-δ phase, it by fixing the tantalum, it is also known that undesirable to limit the formation of Ni 3 _Ta-γ "hardening phase.

Another known solution for improving the properties of 718 is to perform a normal solution heat treatment at 900 to 980 ° C. performed immediately after thermomechanical processing, ie between thermomechanical processing and aging processing. Without aging. This option limits the formation of the Ni ₃ Nb-δ phase that can precipitate during solution heat treatment, making it possible to obtain finer particle sizes with improved tensile and fatigue properties, but has disadvantages.

  It has been found that non-uniform microstructures can be obtained within the same part due to large local variations in particle size and the proportion of δ phase formed during thermomechanical processing.

  As a result, creep resistance is reduced over a wide temperature and stress range compared to previous methods.

  The document EP 1398393 describes the treatment of Ni-based superalloys in the form of unidirectionally solidified single crystals or alloys. When the alloy is a single crystal, there is no grain boundary, so there is clearly no δ phase precipitation at the grain boundary. In unidirectional solidification, precipitation of δ phase can only occur non-uniformly and does not prevent particle growth. At the end of the treatment, the particle size becomes too large. In addition, the alloy composition described in this document preferably does not allow the precipitation of the δ phase with respect to Ti, Ta, Nb and Al because the δ phase is not stable due to the high Al content.

  The document of U.S. Pat. No. 4,459,160 also describes a single crystal Ni-based superalloy in which no δ phase precipitation can be observed at the grain boundaries.

French patent application No. 2089069 European Patent Application No. 1398393 U.S. Pat. No. 4,459,160

  It is an object of the present invention to provide creep resistance and tensile strength of a nickel-based superalloy having a niobium and / or tantalum content greater than 2.5% while avoiding the disadvantages of the prior art without reducing fatigue properties. It is to improve the strength.

For this purpose, the subject of the present invention is a method for producing a Ni-based superalloy blank containing at least 50% Ni as a weight percentage, wherein the superalloy alloy is produced and the alloy is subjected to a heat treatment. Because
The superalloy comprises Nb and Ta up to a total amount of at least 2.5% by weight,
A first holding step in which the alloy is held at 850 to 1000 ° C. for at least 20 minutes to precipitate the δ phase at the grain boundaries;
A second holding step in which the alloy is held at a temperature higher than that of the first holding step and the δ phase obtained in the first step is partially dissolved;
* Performed at a temperature lower than the temperature of the first step and distributed to a third step of precipitating the hardened phase γ ′ and / or γ ″ and optionally an aging process comprising one or more additional steps A method comprising subjecting the alloy to a heat treatment including a plurality of holding steps.

  Preferably, the Al content of the alloy is 3% or less.

  Preferably, the (Nb + Ta + Ti) / Al ratio of the alloy is 3 or more.

  Preferably, the grain size obtained at the end of the aging treatment of the alloy is in the range of 7 to 13 ASTM, more preferably 8 to 12 ASTM, and even more preferably 9 to 11 ASTM.

  Preferably, the distribution of δ phase is uniform at the grain boundaries upon completion of the aging process.

  After the second holding step, an amount of δ phase of preferably 2 to 4%, most preferably 2.5 to 3.5% is obtained.

  The first and second holding steps are preferably performed without intermediate cooling.

  Switching from the first holding step to the second holding step can occur at a rate of 4 ° C./min or less, preferably 1 to 3 ° C./min.

  The first holding step can be performed at 900 to 1000 ° C. for at least 30 minutes and the second holding step can be performed at 940 to 1020 ° C. for 5 to 90 minutes, and the temperature difference between the two temperature holdings is at least 20 ° C.

The weight content of the alloy is
50 to 55% nickel,
17 to 21% chromium less than 0.08% carbon,
Less than 0.35% manganese,
Less than 1% cobalt,
Less than 0.35% silicon,
2.8 to 3.3% molybdenum,
At least one of the niobium element and the tantalum element in which the sum of niobium and tantalum is 4.75% to 5.5% and Ta is less than 0.2%;
0.65 to 1.15% titanium,
0.20 to 0.80% aluminum,
Less than 0.006% boron,
Less than 0.015% phosphorus,
The remaining percentage of iron and impurities resulting from processing may be present.

  The first holding step can then be performed at 920 to 990 ° C. for at least 30 minutes and the second holding step can be performed at a temperature of 960 to 1010 ° C. for 5 to 45 minutes.

  The total Nb and Ta content of the alloy is 5.2 to 5.5%, the first holding step is carried out at 960 to 990 ° C. for 45 minutes to 2 hours, and the second holding step is 5 From 990 to 1010 ° C. for 45 minutes.

  When the total Nb and Ta content of the alloy is 4.8 to 5.2%, the first holding step is performed at 920 to 960 ° C. for 45 minutes to 2 hours, and the second holding step is started from 5 It can be carried out at 960-990 ° C. for 45 minutes.

The weight content of the alloy is
55 to 61% nickel,
19 to 22.5% chromium,
7 to 9.5% molybdenum,
At least one of the niobium and tantalum elements, the sum of niobium and tantalum being 2.75 to 4%, and Ta being less than 0.2%;
1 to 1.7% titanium,
Less than 0.55% aluminum,
Less than 0.5% cobalt,
Less than 0.03% carbon,
Less than 0.35% manganese,
Less than 0.2% silicon,
Less than 0.006% boron,
Less than 0.015% phosphorus,
Less than 0.01% sulfur,
The remaining percentage of iron and impurities resulting from processing may be present.

Alloy
12-20% chromium,
2 to 4% molybdenum,
At least one of niobium and tantalum elements, the sum of niobium and tantalum being 5 to 7%, and Ta being less than 0.2%;
1 to 2% tungsten,
5-10% cobalt,
0.4 to 1.4% titanium,
0.6 to 2.6% aluminum,
6-14% iron,
Less than 0.1% carbon,
Less than 0.015% boron,
Less than 0.03% phosphorus,
The remaining percentage of nickel and the weight content of impurities resulting from processing can be present.

  Preferably, the alloy has a phosphorus content of greater than 0.007% by weight.

  Generally, the first and second holding steps are performed at the δ-phase subsolvus temperature of the alloy, and the first holding step is at a temperature between 50 ° C. below the δ solvus temperature and 20 ° C. below the δ solvus temperature. Once implemented, the second holding step can be performed at a temperature between 20 ° C. below the δ solvus temperature and the δ solvus temperature.

  The temperature of the hot forming blank can be kept constant throughout at least one of these steps.

  The third step can be performed at 700 to 750 ° C. for 4 to 16 hours, and then the fourth step is performed at 600 to 650 ° C. for 4 to 16 hours. Cooling from 50 ° C./hour to 50 ± 10 ° C./hour is carried out between these steps.

  Between the first step and the second step, it is possible to keep the hot-formed alloy at an intermediate temperature between the temperature of the first step and the temperature of the second step for up to 1 hour. .

  The blank may be manufactured in the form of an ingot and then hot processed.

  The blank may be manufactured using powder metallurgy.

  A further subject matter of the present invention is a nickel-based superalloy part, characterized in that it is obtained from a blank produced by the above method.

  This may be a blank of parts for an aviation or onshore gas turbine.

As will be appreciated, the present invention provides structural hardening for Ni-based alloys containing Nb and / or Ta with gamma ′ (Ni ₃ Ti-γ ′) and / or gamma ”(Ni ₃ Nb-γ” and / or Or Ni ₃ Ta-γ ") heat treatment obtained by precipitation of a hardened phase, each of these phases comprising titanium and niobium and / or tantalum. The heat treatment is at least the following steps in chronological order: Includes three holding steps.
The delta phase Ni ₃ Nb-δ and / or Ni ₃ Ta-δ is precipitated at the grain boundaries such that the distribution of this phase at the grain boundaries is substantially uniform, to homogenize the microstructure of the material. Targeted first holding step performed at 850-1000 ° C. As for the partially recrystallized microstructure, it also completes the recrystallization and precipitates the δ phase at the grain boundaries of the new recrystallized grains.
- is carried out at a temperature higher than the temperature of the first step, maintaining a substantially uniform distribution obtained after the first step, while avoiding the expansion of the particles, the delta phase Ni 3 _Nb-δ And / or a second holding step intended to partially dissolve Ni ₃ Ta-δ. The second step is completed by oil cooling or air cooling.
The third step and optionally the subsequent steps are carried out at a temperature lower than that of the first step, and gamma ′ (Ni ₃ (Al—Ti) -γ ′) and / or gamma ”(Ni ₃ Nb -gamma "or _Ni 3 _Ta-γ") is a thermal aging step to precipitate the hardening phases.

  One or more intermediate cooling operations are possible between each step, but are not essential.

  The method of the present invention produces parts with a better compromise between yield strength under heavy loads, high fatigue strength, and long creep resistance life compared to prior art parts having the same composition. Enable.

  The invention will be better understood upon reading the following description presented with reference to the accompanying drawings in which:

FIG. 3 is a diagrammatic illustration of two first heat treatment steps according to the present invention, where the temperature on the Y axis is referenced in relation to the δ phase solvus temperature. FIG. 6 is a diagrammatic illustration of two first heat treatment steps according to the present invention, also showing an intermediate step between the first step and the second step, the temperature on the Y axis being the δ phase solvus temperature Referenced in connection with FIG. 3 is a diagrammatic illustration of two first heat treatment steps according to the present invention, where the temperature on the Y axis is referenced in relation to the δ phase solvus temperature. It is a figure which shows the microscope picture of the alloy in which the reference | standard heat processing was performed. It is a figure which shows the microscope picture of the alloy in which the reference | standard heat processing was performed. It is a figure which shows the microscope picture of the alloy in which the reference | standard heat processing was performed. It is a figure which shows the microscope picture of the alloy in which the reference | standard heat processing was performed. It is a figure which shows the microscope picture of the alloy by this invention. It is a figure which shows the microscope picture of the alloy by this invention.

  Methods for producing Ni superalloy parts according to the invention include double melting methods (VIM vacuum induction melting-VAR vacuum arc remelting) or triple melting (VIM-ESR (electroslag remelting) -VAR). It can begin by preparing and casting the superalloy ingot using conventional methods. However, the method according to the invention can also be applied to blanks produced by powder metallurgy. In the remainder of this document, the examples of applications described are examples where the starting product is obtained by a conventional route called "Ingot Metallurgy", but it will be clear to those skilled in the art that it can be replaced by powder metallurgy. is there. The processing that follows the high temperature processing that is a feature of the present invention is the same in both cases.

The initial microstructure of the product (understanding that the term “product” refers to a semi-finished product or blank of a part) prior to the processing typical of the present invention is a modified thermomechanical process performed upstream, For example, it may differ in connection with forging, stamping or hot rolling.
Metallurgical state 1 (or “State 1”): Delta phase Ni ₃ Nb-δ and / or Ni ₃ Ta-δ may be present at grain boundaries, but after deformation performed at a temperature below the δ phase solvus temperature Cannot be uniformly distributed among the particles.
Metallurgical state 2 (or “state 2”): delta phase Ni ₃ Nb-δ and / or Ni ₃ Ta-δ is absent from a deformed microstructure performed at a temperature higher than, for example, a δ-phase solvus Or may be substantially absent (<1%).

  In the first case, i.e. starting from metallurgical state 1, the first processing step according to the present invention is to make the distribution of the δ phase uniform within the microstructure and to thermomechanical processing due to some temperature fluctuations after deformation. It is possible to reduce the local variation of the fraction of the δ phase existing later. The person skilled in the art can easily adjust the parameters for performing the first step, if necessary, through routine testing, in order to optimize this homogenization of the δ phase distribution.

  In the second case, i.e. starting from the metallurgical state 2, the first treatment step according to the invention deposits the δ phase (substantially) uniformly at the grain boundaries lacking the δ phase after the thermomechanical treatment. The person skilled in the art can adjust the parameters for performing the first step through routine tests, if necessary, so as to optimize the homogenization of this δ phase distribution.

  Regardless of whether it is the first case or the second case, the first step is also to complete the recrystallization in a region where the recrystallization was not completed during the thermomechanical process. Uniform the overall structure of the.

In the second step of the process according to the invention carried out at a temperature close to the δ phase solvus, the delta phase Ni ₃ Nb-δ and / or Ni ₃ Ta-δ is partially dissolved.

  In the second step, dissolution of the δ phase occurs in a substantially uniform form. The so-called residual δ phase, ie the undissolved δ phase, maintains the same distribution as that obtained after the first step. For this reason, the residual δ phase is distributed substantially uniformly around the particles, slowing the growth of all particles in a second step carried out at a temperature higher than that of the first step, Limit or even avoid the generation of large particles. The uniform distribution of the δ phase at the grain boundary facilitates the uniform grain size of the microstructure of the alloy at the end of processing.

  Thus, the second step reduces the amount of δ phase obtained after the first step, optimally below 4% and even below 3.5% while avoiding particle expansion. Let

  As the solubility of the δ phase in the fine grained microstructure increases, it becomes more due to the precipitation of gamma ′ and / or gamma ”hardened phases during the third step and other subsequent steps that constitute the alloy aging process. Of niobium can be released.

  Surprisingly, the inventors have found that without the first processing step, it is not possible to obtain these effects regardless of the initial microstructure after thermomechanical processing.

  Obviously, in the initial microstructure lacking the δ phase (state 2), if the first step does not exist, the entire structure of the material is homogenized, and the δ phase is precipitated at the grain boundaries, and through the second step It becomes impossible to limit the growth of the subsequent particles.

  If the first step does not exist, the distribution of the δ phase becomes non-uniform when the initial microstructure is generated from the sub-solvus deformation that results in the δ phase precipitation (state 1) (see FIGS. 4 and 5). Thus, some particles may contain a large amount of δ phase at the grain boundaries, little or no δ phase at the grain boundaries, or even a non-uniform distribution of δ phases at the grain boundaries.

  By carrying out the heat treatment as it is at the temperature of the second step without maintaining the temperature at the temperature of the first step, the particles that are not surrounded by the δ phase, or have almost no δ phase at the grain boundary, or Particles having a non-uniformly distributed δ phase will probably uncontrollably expand to particle sizes greater than about ASTM 5-6. The presence of ASTM 5-6 particles, even though they are highly localized (see FIGS. 6 and 7), has a fatigue life of 10 minutes compared to a uniform microstructure with ASTM particle size 10 particles. Decrease to 1. Thus, the combination of the first step and the second step according to the present invention (see FIGS. 8 and 9) avoids the presence of these large ASTM 5-6 particles that hinder the guarantee of high fatigue properties, Allows partial dissolution of the δ phase in a uniform form.

  Thus, in the initial microstructure including δ phase (state 1), if the first step is not present, the desired microstructure, ie, the residual uniform δ phase content, is preferably less than 4% and is uniform and acceptable. It is not possible to obtain a microstructure having a particle size that can be obtained.

  The preferred particle size of products derived from the method of the present invention is derived from the desire to achieve a good compromise between conflicting properties with respect to their particle size requirements. Fatigue strength and tensile strength substantially benefit from small particle size, while creep resistance and crack resistance benefit from coarse particle size. In this regard, suitable particle sizes are ASTM 7 to 13, preferably ASTM 8 to 12, and most preferably ASTM 9 to 11.

  The absence of the second step after performing the first step is performed on the superalloy product to which the present invention is applied and is the normal type recognized above that they are not satisfactory. Corresponds to processing.

  In addition, for the initial microstructure lacking the δ phase (state 2), neither of the two first steps required by the present invention is performed, so that after the high temperature processing at the δ phase supersolvus temperature, When the aging treatment is performed on the alloy as it is (so-called “direct aging” treatment is performed) (state 2), the δ phase is not obtained at all in the final structure, which is not desirable.

  Surprisingly, the inventors have shown that the presence of a δ phase in practice of 2 to 4%, optimally 2.5 to 3.5%, improves the properties of a material without weakening the material. I could not.

  On the other hand, microstructures lacking the δ phase generally reduce intergranular weakness that significantly reduces hot ductility and greatly increases the alloy's susceptibility to notch effects (eg, premature creep failure at the notch point). More easy to wear. Thus, even when no δ phase is present after thermomechanical processing, a first step is necessary to produce a minimal amount of δ phase uniformly distributed at the grain boundaries and to homogenize the overall structure of the material. .

  The retention time of the alloy at the temperature of the first step is 20 minutes or more. The temperature of the first step for precipitating the δ phase is 850 to 1000 ° C. The temperature and holding time are adjusted for microstructure non-uniformity after deformation and to maintain the amount of δ phase after the second step greater than the minimum amount required for hot ductility.

  Thus, the second step, performed at a higher temperature than the first step, maintains a sufficient amount of Nb and / or Ta in the form of a δ phase uniformly distributed around the particles towards the hot ductility of the material. However, in order to release Nb and / or Ta necessary for the precipitation of the γ ′ phase and / or the γ ″ phase, the amount of the δ phase is dissolved to a desired amount, preferably 2 to 4%, optimally 2 Required to reduce content from 5 to 3.5%.

  The temperature and duration of the second step is adjusted with respect to the fraction of δ phase obtained after the first step to obtain the desired residual fraction of δ phase while avoiding particle expansion. The duration of the second step is also related to the temperature determined for this step. In general, the shorter the duration of the second step, the higher its temperature.

  According to a preferred variant of the invention, the two first processing steps are sequential steps (FIGS. 1 and 2).

  “Continuous process step” means that the switching from the first process step to the second process step gradually increases the temperature without going through an intermediate temperature lower than the temperature of the first step. It means that it occurs by moving from the first step to the second step.

Performing the two first steps sequentially without lowering the temperature of the first step, for example, to ambient temperature, avoids large temperature gradients inside the processed sample, and in some areas the particles It is possible to avoid inhomogeneous dissolution of the δ phase, which can cause expansion. Therefore, it is preferable to employ a sufficiently low rate of temperature increase (<4 ° C./min) between steps so that the temperature is maintained uniformly in the treated sample throughout the second step. It was confirmed in the second step that the temperature became uniform after 5 minutes in a 1000 cm ³ cylindrical sample after a temperature rise rate of 2 ° C./min from the first step. Therefore, switching between two steps at a temperature lower than the first step can increase the time required to equalize the temperature in the sample in the second stage, resulting in non-uniform dissolution of the δ phase. There is a risk of promoting. However, especially with regard to the size of the processing parts, if the parameters of the second step are adjusted by optionally adding intermediate steps to avoid the possible disadvantages mentioned above, Said switching to a temperature lower than the temperature is not excluded by the present invention (FIG. 3).

  Preferably, the first treatment step is performed at a temperature of about 900 to 1000 ° C. for at least 30 minutes, and the second treatment step is a temperature of the first step of 940 to 1020 ° C. for a time of 5 to 90 minutes. Performed at higher temperatures. Therefore, the temperature difference between the two steps must be at least 20 ° C. The temperature range and duration so obtained is sufficient particle size, ie ASTM 7 to 13, preferably ASTM 8 to 12, most preferably ASTM 9 to 11, and 2% to 4% residual δ phase. It is possible to obtain a uniform microstructure with a fraction.

  As will be appreciated, the present invention is primarily based on a synergistic effect between two first steps and is set by the present invention by optimizing the balance of these two first steps. It is possible to fully meet the objectives.

  The δ phase solvus temperature is directly dependent on the niobium + tantalum content of the alloy. Thus, the amount of niobium and / or tantalum present in the alloy composition has a direct effect on the temperature and duration of each step.

  If a type 718 alloy (whose standardized composition is detailed below) is used, a first temperature hold of 920 to 990 ° C. for at least 30 minutes is performed and from 960 for 5 to 45 minutes It is pointed out that a second temperature hold of 1010 ° C. is carried out. The optimal duration of processing depends on the mass of the part to be processed and can be determined using models or experiments commonly used by those skilled in the art.

  If the total Nb and Ta content of 718 alloy is 5.2 to 5.5% (Ta less than 0.2%), the first step is preferably about 960 over a period of about 45 minutes to 2 hours. The second step is preferably carried out at a temperature of about 990 ° C. to 1010 ° C. for a time of about 5 to 45 minutes.

  If the Nb + Ta content of the 718 alloy is about 4.8 to 5.2% (Ta less than 0.2%), the first step is preferably about 920 ° C. to 960 ° C. over a period of about 45 minutes to 2 hours. The second step is preferably performed at a temperature of about 960 ° C. to 990 ° C. for a time of about 5 to 45 minutes. The duration of processing also depends on the mass of the part to be processed.

  The temperature of the process step is generally kept substantially constant throughout the duration of the temperature hold.

  The rate of temperature increase between the first step and the second step is preferably 4 in order to avoid excessive temperature gradients, especially when the part being processed is a large part. It is less than ℃ / min.

  The rate of temperature increase from the first step to the second step is preferably 1 ° C./min to 3 ° C./min.

  Thus, the present invention applies to nickel-based superalloys that contain at least 50% Ni and the total Nb + Ta exceeds 2.5% by weight.

In one particular case, the alloy is
50 to 55% nickel,
17 to 21% chromium less than 0.08% carbon,
Less than 0.35% manganese,
Less than 0.35% silicon,
Less than 1% cobalt,
2.8 to 3.3% molybdenum,
At least one of the niobium element and the tantalum element, wherein the sum of niobium and tantalum is 4.75 to 5.5% and Ta is less than 0.2%;
0.65 to 1.15% titanium,
0.20 to 0.80% aluminum,
Less than 0.006% boron,
Less than 0.015% phosphorus,
A 718 type nickel-based alloy, also called NC19FeNb (AFNOR standard), with the remaining percentage of iron and the weight content of impurities resulting from processing.

  Elements whose minimum content is not present are present in a trace form, in other words a content that can be zero, and in any case may be sufficiently low to have no metallurgical effect (this is This is true for the composition described).

  Advantageously, the addition of phosphorus makes it possible to enhance the grain boundary strength against stresses such as creep and notch creep. The application of the present invention to such alloys with a phosphorus content greater than 0.007% and less than 0.015% is of particular interest since the resulting creep gain is clearly increased. Therefore, it is possible to easily improve the creep life by a factor of 4 while maintaining the same particle size. The presence of this phosphorus can be recommended for other examples of alloys shown below for the same reason.

In another specific case, the alloy is
55 to 61% nickel,
19 to 22.5% chromium,
7 to 9.5% molybdenum,
At least one of the niobium and tantalum elements, the sum of niobium and tantalum being 2.75 to 4%, and Ta being less than 0.2%;
1 to 1.7% titanium,
Less than 0.55% aluminum,
Less than 0.5% cobalt,
Less than 0.03% carbon,
Less than 0.35% manganese,
Less than 0.2% silicon,
Less than 0.006% boron,
Less than 0.015% phosphorus,
Less than 0.01% sulfur,
725 type nickel-based superalloy with the remaining percentage of iron and the weight content of impurities resulting from processing.

In another specific case, the alloy is
12-20% chromium,
2 to 4% molybdenum,
At least one of niobium and tantalum elements, the sum of niobium and tantalum being 5 to 7%, and Ta being less than 0.2%;
1 to 2% tungsten,
5-10% cobalt,
0.4 to 1.4% titanium,
0.6 to 2.6% aluminum,
6-14% iron,
Less than 0.1% carbon,
Less than 0.015% boron,
Less than 0.03% phosphorus,
718 PLUS type nickel-based superalloy with the remaining percentage nickel and the weight content of impurities due to processing.

Generally, the alloy, the content of niobium + tantalum is greater than 2.5%, 1050 ° C. of _Ni 3 Nb-Ta type intergranular phase from 800 ° C. ([delta] phase) is present and 600 from 800 ° C. Ni ₃ It is a nickel-based superalloy characterized by the presence of (Al—Ti)-(γ ′) type and / or Ni ₃ Nb—Ta (γ ″) type intergranular phase. More than 2.5% In a nickel-based superalloy comprising niobium and / or tantalum, containing niobium and / or tantalum and having an intergranular phase of the Ni ₃ Nb—Ta type, the γ ”hardened phase Ni ₃ Nb— Even if Ta does not exist, the effect of the present invention is found. Therefore, when the solubility of the intergranular phase of the delta Ni ₃ Nb—Ta type increases, niobium (γ ′ generating element that is inserted into the γ ′ hardened phase—Ni ₃ (Al, Ti) in the solid solution and hardens the phase. ) Is released.

Treatment of the present invention, the gamma of the third temperature lower than the temperature in step _{"(Ni 3 Nb-Ta-} γ") and / or gamma _{'(Ni 3 (Al-Ti} ) -γ') of hardening phase precipitation A fourth step that allows completion can be included.

  For example, after a third step of 700 to 750 ° C. over 4 to 16 hours, to a temperature of the fourth step of 600 to 650 ° C. held for 4 to 16 hours at 50 ° C./hour±10° C./hour Can be cooled.

  The process of the present invention has a short duration between the first step and the second step in order to promote temperature uniformity in the large part during the temperature rise between the two first steps. It can also include at least one intermediate step of time (within 1 hour; see FIG. 2).

  According to the present invention, when the (Ta + Nb) content of the alloy is at least 2.5%, it is recommended that the Al content is 3% or less so as not to cause precipitation of the γ 'phase at the grain boundaries. The When Al is 3% or more, the γ ′ phase tends to be stable enough to damage the δ phase, and Nb is inserted into the γ ′ phase.

  In order to give priority to the precipitation of the δ phase at the grain boundary, the (Nb + Ta + Ti) / Al ratio is preferably 3 or more.

  The invention will now be illustrated by means of some examples of heat treatment embodiments according to the invention which are given in a non-limiting manner.

  A first example of an embodiment of the method according to the invention is obtained after thermomechanical processing of an alloy, which is obtained via the conventional path of VIM + VAR + forging, but can also be obtained by powder metallurgy, Applies to 718 alloy articles intended for the manufacture of aviation turbine discs.

  At the experimental level, the VIM method was used, then the VAR method was used to prepare a 718 alloy remelted ingot and then three different thermomechanical treatments numbered 1 to 3 in Table 2 (Table 2) They were hot processed according to the schedule (see TTM, Table 2 (Table 2)). The product obtained after thermomechanical treatment was cut into samples (represented from A to P in Table 1 (Table 1)). The samples were then subjected to a different heat treatment (TTH) comprising 2 to 4 steps depending on the different cases (see Table 2 (Table 2)).

  The schedule for thermomechanical treatment N ° 1 included rolling performed in different passes at a temperature above the δ phase solvus temperature of the alloy. The product formed by the thermomechanical treatment schedule N ° 1 is a bar whose metallurgical structure lacks a delta phase (metallurgical state 2). In Table 2 (Table 2), Samples F, K, L, and N were made from the bars obtained by this first thermomechanical processing schedule.

  The thermomechanical treatment schedule N ° 2 refers to reheating at a temperature below the δ phase solvus temperature of the alloy (“subsolvus temperature”) (“heat” refers to deforming after holding in the furnace, thus “reheating”. Was a conventional forging schedule including two deformation steps each preceded by holding in a furnace. This schedule allows the precipitation of the δ phase in the alloy. The product formed according to the thermomechanical treatment schedule N ° 2 is a so-called pancake (generally with some δ phase (see metallurgical state 1, FIGS. 4 and 5)) whose metallurgical structure is unevenly distributed at grain boundaries. , Products with rough shape of flat disk due to deformation by forging). In Table 2 (Table 2), Samples C, E and H were made from pancakes obtained by this second thermomechanical processing schedule.

  The thermomechanical processing range N ° 3 was a conventional punching schedule with a single thermal step at a lower temperature than the alloy δ phase solvus. The product formed by the thermomechanical treatment schedule N ° 3 is a disc blank containing some δ phase (see Metallurgical State 1, FIGS. 4 and 5) whose metallurgical structure is very unevenly distributed at grain boundaries. In Table 2 (Table 2), Samples A, B, D, G, I, J, M, O, and P were made from turbine disk blanks obtained by this third thermomechanical processing schedule.

  Samples A to P were then subjected to five different schedules of heat treatment (“TTH”) represented by a, b, c, d, e, including 2 to 4 steps for each case (Table 2 (Table 2 ( TTH column in Table 2).

  The heat treatment schedule of type “a” or “b” is a reference heat treatment schedule indicating advanced technology.

  A type “a” heat treatment schedule includes one step of so-called isothermal solution treatment and two aging steps. In these schedules, the solution heat treatment for samples A, B, C, D, F and P consisted of holding the alloy at a constant temperature of 955 to 1010 ° C. for 40 to 90 minutes. The two aging steps consisted of holding at 720 ° C. once for 8 hours and then controlled cooling at 50 ° C./hour to a holding temperature of 620 ° C. over 8 hours.

  A type “b” heat treatment schedule known as “direct aging” does not include solution heat treatment and consists of only two aging steps according to type “a” treatment. Only sample E was given a type “b” schedule.

  The type “c” heat treatment schedule is in accordance with the present invention and includes two so-called solution heat treatment steps, referred to as the first step and the second step, respectively, and one or two, respectively referred to as the third step and the fourth step. Includes an aging step.

  In these schedules relating to samples G, H, J, K, M and N, the first solution heat treatment step consists of holding the alloy at a constant temperature of 940 ° C. to 980 ° C. for about 50 to 60 minutes. It was. The second solution treatment step consisted of holding the alloy at a constant temperature of 980 ° C. to 1005 ° C. for about 15 to 40 minutes. Switching from the first temperature hold to the second temperature hold was performed by controlled reheating at a rate of about 2 ° C./min. The third and fourth aging steps followed the corresponding aging step of type “a” reference schedule, except for samples H and J.

  For sample H, the temperature of the third aging step was 750 ° C. instead of 720 ° C. used for the other samples. Due to this difference, the scope of the present invention is not limited to the limited temperature conditions and duration of the aging step, in contrast to the temperature and duration of the aging step such as those used in the nickel-based superalloy field. It was possible to prove that the present invention can be applied to the above.

  For sample J, this sample was only subjected to one aging step at 720 ° C. for 10 hours. The aging treatment applied to sample J shows that the present invention can be applied even when only a single aging treatment step is applied to the alloy.

  The heat treatment type “d” schedule included two solution heat treatment steps and two aging steps. Samples I and L were processed according to these schedules. However, these treatments are not in accordance with the present invention because the second step is performed at a temperature that is too high or for a duration that is too long. The conditions of the second step actually cause too much dissolution of the δ phase and particle growth becomes uncontrolled, leading to large uncontrolled particle expansion through the second step for samples I and L.

  The type “e” heat treatment schedule included a single solution heat treatment step at 1005 ° C. for 15 minutes, and two aging steps. As explained below, only Sample O was obtained by this heat treatment schedule not subject to the present invention.

  Samples A to L and O were type 718 alloys containing 5.3% Nb and 40 ppm P. Sample N was a type 718 alloy containing 5.0% Nb and 40 ppm P. Samples M and P were type 718 alloys containing 5.3% Nb and 80 ppm P.

  Table 2 (Table 2) outlines the processing conditions for the different samples and the ASTM particle size and proportion of the surface δ phase that can be seen in the micrographs.

Table 3 (Table 3) shows some of the main mechanical properties of these samples: yield strength (YS) during tensile testing at -20 ° C;
Ultimate tensile strength (UTS) during a tensile test at -20 ° C;
The number of cycles before failure during fatigue testing at 450 ° C., including a frequency of 10 Hz and a load factor R of 0.05, with a sine cycle and a maximum stress of 1050 MPa;
The outline | summary of the lifetime at the time of the creep test in 650 degreeC under the stress of -550 Mpa and the stress of 690 Mpa is shown.

  The particle size is determined according to ASTM standards, and if the particle size is relatively non-uniform, the maximum particle size (ALA) is also specified.

  Thus, the products 718, F, K, L, N of 718 alloy were converted according to a thermomechanical schedule n ° 1, which does not allow δ phase precipitation.

  Product F is a reference sample that was processed using standard type “a” thermomechanical processing of alloy 718 (processing including a single solution heat treatment step of the δ phase subsolvus) after thermomechanical processing schedule n ° 1. .

  Product L was treated with a two-step solution heat treatment, but the second step was performed at too high a temperature and for a too long duration, and was outside the scope of the invention for alloy 718.

  Products K and N did not have the same niobium content, but both were subjected to a heat treatment schedule of type “c” according to the present invention.

  The product of 718 alloy, identified as C, E and H, was converted by a thermomechanical schedule n ° 2 that allowed heterogeneous precipitation of the δ phase.

  Product C is a reference sample that has been processed by a standard “a” type heat treatment of alloy 718 (a treatment that includes only one solution heat treatment at a subsolvus temperature) after thermomechanical schedule n ° 2.

  Product E is also a reference sample that was processed according to the type “b” heat treatment schedule after the thermomechanical schedule n ° 2, and thus was aged as it was after forging and was not subjected to solution heat treatment before aging.

  Product H was subjected to a heat treatment (type “c”) according to the present invention including a two-step solution heat treatment within the scope of the present invention.

  Products of 718 alloy identified as A, B, D, G, I, J, M, O, and P were converted by a thermomechanical schedule n ° 3 that allows highly non-uniform precipitation of the δ phase. .

  After thermomechanical treatment n ° 3, products A, B and P were processed by standard treatment for alloy 718 (type “a” treatment including single subsolvus solution heat treatment).

  Product D included only one solution heat treatment step, but was processed by processing at a higher temperature than products A, B, and P, ie, a temperature close to the δ phase solvus.

  After thermomechanical treatment, Product I was subjected to a two-step solution heat treatment where the duration of its second step was too long with respect to temperature. Therefore, the heat treatment applied to product I is outside the scope of the present invention.

  After thermomechanical treatment n ° 3, product G was treated by a two-step solution heat treatment (heat treatment “c”) within the scope of the present invention.

  Product J was also treated by a two-step solution heat treatment within the scope of the present invention, but not in the fourth step.

  Product M was processed by a two-step solution heat treatment within the scope of the present invention, but has a phosphorus content of 0.008%, twice that of products A-L and N-O.

  Product O was subjected to an “e” heat treatment by a one-step solution heat treatment, which is outside the scope of the present invention.

  Product P is a reference sample having a phosphorus content of 0.008%. It was processed using the standard processing schedule for alloy 718 (an “a” type process that includes only one solution heat treatment at sub-solvus temperature).

  Products A, B, C treated by standard subsolvus heat treatment (schedule type “a”) have a fine particle microstructure (> 9 ASTM), but preferably a fraction exceeding the desired δ phase fraction according to the invention ( > 4.5%) δ phase. The mechanical properties obtained for these products constitute the basis for the evaluation of the tensile, fatigue and creep properties obtained with thermomechanical schedules (TTM) 2 and 3.

  Product D is processed at a higher temperature than products A, B and C, and has a particle size of ASTM 5 and a non-uniformly distributed δ phase (<2.5), preferably less than the desired δ phase according to the present invention. %)including. This treatment does not allow the fine particle microstructure (at least ASTM 7, preferably at least ASTM 8, most preferably ASTM 9) nor to maintain the satisfactory fatigue properties found for products A, B and c. It has been confirmed that The significant decrease in fatigue life can be attributed to the presence of large particle size ASTM 5 particles that form the fatigue initiation region.

  Product E, aged as it is after thermomechanical treatment N ° 2, has a very non-uniform particle size (ASTM 10 to 14), and the amount of δ phase varies greatly, which is a large part of the region of the part. It is found in the part (especially in the area subject to creep) and is larger than the desired fraction of the δ phase. The tensile and fatigue properties of product E are higher than those of products A, B and C, but the creep life obtained for product E is confirmed to be shorter than that of products A, B and C. .

  In the absence of solution heat treatment, the microstructure cannot be homogenized, causing very fine particles (> 12 ASTM) and the presence of too large a δ phase fraction that is responsible for this degradation of creep properties. .

  Also, the absence of solution heat treatment for product E allows maintenance of residual work hardening after forging, which is beneficial for tensile properties but detrimental to creep resistance in low stress areas.

  Products G, H, M are processed within the scope of the present invention and are included in at least 2.5% of the fine particle microstructure (> 9 ASTM) and the desired range of δ phase fractions, i.e. 4% or less. It has a phase fraction (2.9% and 3.5%). The tensile properties are clearly higher than the tensile properties of products A, B and C, and are confirmed to be at the same level as the tensile properties of product E. The creep characteristics of products G, H and M are clearly better than those of products A, B, C and E, while the particle size is also confirmed to be similar within these products. Yes. The fine-grained microstructure of products G, H, M allows the maintenance of the fatigue properties obtained for products A, B, C, E, and the smaller δ phase fraction of products G, H, M is the creep resistance. To improve.

  Comparing samples B and P, it can be seen that an increase in the phosphorus content of the 718 alloy that has undergone the reference treatment (a) does not substantially improve the creep resistance.

  Surprisingly, when the treatment according to the present invention is applied to product M having a higher phosphorus content (80 ppm), the phosphorus content of product M is higher than that of products A, B and C. However, the creep life can be significantly improved up to 4 times as compared with the product P which has not been treated according to the present invention.

  Therefore, the combination of phosphorus addition and the inventive treatment has a positive synergistic effect on the creep properties of the resulting alloy.

The present invention aims to maintain a residual δ phase fraction (preferably higher than 2.5%) that makes it possible to maintain satisfactory ductility at high temperatures. If the content of the δ phase is too low, this affects the damage and ductility under tensile testing at high temperature (650 ° C. at a strain rate of 10 ⁻⁵ s ⁻¹ ). Product D, whose δ phase content is close to 2%, has a ductility (7% elongation at break) that is much lower than the ductility of product G (27% break elongation), whose δ phase fraction is close to 3%. It is actually confirmed to have. The low ductility of this product D is due to intergranular damage caused by the fraction of the δ phase being too small and unevenly distributed.

  Next, the influence of the treatment of the present invention on the microstructure will be described in detail.

  Investigating Samples A, B, C, D, E, G, H, I, J, M, O and P, which are 718 alloy and converted using thermomechanical schedule N ° 2 or n ° 3 It was.

4 through 9 are
-Samples A, B, C, D, E, G, H, I, J, M, O and P in their initial state after thermomechanical treatment (FIGS. 4 and 5),
Samples D and O after being subjected to a heat treatment comprising only one solution heat treatment step (FIGS. 6 and 7),
-A photomicrograph showing the microstructure of samples G, H and M (Figs. 8 and 9) after being subjected to a heat treatment according to the invention.

FIGS. 4 and 5 show microscopic samples A, B, C, D, E, G, H, I, J, M, O, and P after sub-solvus thermomechanical deformation (thermomechanical schedule 2 or 3). The structure (metallurgical state 1) is shown. This is a microstructure that exhibits delta phase Ni ₃ Nb-δ and / or Ni ₃ Ta-δ at grain boundaries, but they are not evenly distributed among the grains.

  FIG. 4 shows that the sample has fine particles with a particle size of ASTM about 11, and the δ phase is unevenly distributed (dark shadows at grain boundaries). After the thermomechanical deformation schedule, the proportion of δ phase will be 2.8 to 6% and the particle size will be ASTM 10 to 13. This gives a very non-uniform microstructure from these two perspectives.

  FIG. 5 shows the larger microstructure of the sample, showing particles whose grain boundaries are almost completely devoid of the δ phase (this phase is shown in white in the micrograph).

  When a treatment involving only one solution heat treatment step at 970 ° C. for about 60 minutes is applied to the sample (Sample B), the resulting δ phase percentage is 4.7 to 5.5% and the particle size is ASTM From 11 to 12. Thus, the uniformity of the sample is improved, but is known to be highly detrimental to creep resistance (Table 1 (Table 1) and Table 2 (Table 2), sample B)) a large δ phase fraction Is maintained.

  A process in which the heat treatment applied to the sample (eg, sample O in Table 1 (Table 1)) includes only one solution heat treatment step at 1005 ° C. for about 15 minutes, corresponding to the “second step” of the present invention. The ratio of the obtained δ phase (see FIGS. 6 and 7) is 1.1 to 3.5% and the particle size is ASTM 5 to 9. Therefore, the amount of δ phase is reduced, which is the preferred direction for creep resistance, but a non-uniform distribution of particle sizes is observed. This can be explained by non-uniform grain growth throughout this step due to the non-uniform distribution of the δ phase inherited from the initial microstructure.

  In fact, as already explained, the initial microstructure arises from subsolvus deformation (state 1), and the distribution of the δ phase in the initial microstructure is non-uniform. As a result, some particles may contain a large amount of δ phase at the grain boundaries in the initial structure, whereas other particles have little or no δ phase at the grain boundaries (FIG. 5). reference).

  By performing the heat treatment as it is at the temperature of the second step without the intermediate temperature holding of the temperature of the first step according to the present invention, the particles not surrounded by the δ phase, or the δ phase at the grain boundary Particles that have little or no particle expand uncontrollably to a particle size that can exceed ASTM about 5 to 6, whereas other particles surrounded by the δ phase are prevented from growing and have a particle size close to ASTM 9 Will bring. Particle size non-uniformity can be clearly seen in the micrographs of FIGS. The presence of ASTM 5-6 particles significantly reduces the fatigue life, even if they are highly localized.

  On the other hand, if several samples (samples G, H and M) are subjected to the heat treatment according to the invention, ie a first step of 980 ° C. over 60 minutes is followed immediately by 1005 ° C. over 15 minutes. When heated to a second temperature hold of 2 ° C./min, a δ phase of 2.9 to 3.5% with a particle size of ASTM 10 to 12 is obtained.

The micrographs in FIGS. 8 and 9 are compared to the initial state of the sample,
The particle size is kept more uniform and very fine;
-It shows here that the δ phase is regularly distributed at the grain boundaries and its growth is efficiently prevented.

Small-scale formation of δ-phase precipitates that enable use of dissolved Nb and Ta atoms, reduction in particle size, uniform distribution of δ-phase at grain boundaries, and well-adjusted abundance of this δ-phase As a result, creep resistance and tensile strength are improved. Especially,
-Strong fatigue properties that avoid premature crack initiation in large particles and prioritize crack initiation in niobium carbide;
-Improvement in yield strength by broader hardening produced by a larger fraction of the hardening phase;
The fine particle size associated with controlled dissolution of the δ phase which makes it possible to achieve the object of the invention, which is a clear and even a significant improvement in the creep resistance of an alloy with sufficient phosphorus content (sample M) It is.

  When the alloy is processed according to the present invention, the final part can be obtained by proceeding with normal finishing operations in the prior art.

  In addition, the inventors have performed additional tests on samples of type 718Plus and 725 alloys, and the present invention shows that other nickel-based super alloys having niobium and / or tantalum content greater than 2.5%. It could be confirmed that when applied to alloys, it allowed a significant improvement in their creep resistance and tensile strength.

According to a first option, as described in French Patent Application No. 2089069, after performing the thermomechanical treatment to precipitate Ni 3 _Nb-δ phase at the grain boundaries, the dissolution of Ni 3 _Nb-δ phase It was chosen to recrystallize the alloy at a temperature below the temperature and use the Ni ₃ Nb-δ phase precipitated at the grain boundaries during recrystallization to prevent grain growth. When this method is used, it is possible to obtain a recrystallized structure having a very fine grain size of ASTM 10 or more. Their fatigue properties are improved, but their creep resistance is insufficient. In fact, the presence of a Ni ₃ Nb-δ phase having an orthorhombic structure is preferred because it is metastable by fixing niobium and limits the formation of a Ni ₃ Nb-γ ″ hardened phase having a central square structure. It is known that the Ni ₃ Nb-γ "hardened phase improves creep resistance by slowing dislocations in the crystal lattice.

Similarly, the presence of the Ni ₃ Ta-δ phase is also known to be undesirable because it limits the formation of the Ni ₃ Ta-γ "hardened phase by fixing tantalum.

Surprisingly, the inventors have shown that the presence of a δ phase in practice of 2 to 4%, optimally 2.5 to 3.5%, improves the properties of a material without weakening the material. I was able to .

Claims (24)

A method for producing a blank of a Ni-based superalloy component comprising at least 50% Ni by weight, comprising preparing the superalloy alloy and performing a heat treatment of the alloy, comprising:
The superalloy comprises a total amount of Nb and Ta of at least 2.5% by weight;
-* Holding the alloy at 850-1000 ° C for at least 20 minutes to precipitate the δ phase at the grain boundaries;
A second step in which the alloy is maintained at a temperature higher than the temperature of the first step and the δ phase obtained in the first step is partially dissolved;
An aging process carried out at a temperature lower than that of the first step and comprising a third step and optionally one or more additional steps to precipitate the γ ′ and / or γ ″ curing phase. Applying a heat treatment comprising a plurality of steps to the alloy.   The method according to claim 1, wherein the Al content of the alloy is 3% or less.   The method according to claim 1, wherein the alloy has a (Nb + Ta + Ti) / Al ratio of 3 or more.   Method according to any one of claims 1 to 3, characterized in that the grain size obtained at the end of the alloy treatment is 7 to 13 ASTM, preferably 8 to 12 ASTM, most preferably 9 to 11 ASTM. .   The method according to any one of claims 1 to 4, wherein the distribution of the δ phase after the aging treatment is uniform at the grain boundaries.   6. A method according to any one of the preceding claims, characterized in that after the second step an amount of delta phase of 2 to 4%, preferably 2.5 to 3.5% is obtained.   The method according to any one of claims 1 to 6, wherein the first and second steps are performed without intermediate cooling.   8. A method according to claim 7, characterized in that the switching from the first step to the second step is carried out at a rate of 4 [deg.] C / min or less, preferably 1 to 3 [deg.] C / min.   The first step is performed at 900 to 1000 ° C. for at least 30 minutes, the second step is performed at 940 to 1020 ° C. for 5 to 90 minutes, and the temperature difference between the two steps is at least 20 ° C. The method according to any one of claims 1 to 8. The alloy is in weight percent
50 to 55% nickel,
17 to 21% chromium less than 0.08% carbon,
Less than 0.35% manganese,
Less than 0.35% silicon,
Less than 1% cobalt,
2.8 to 3.3% molybdenum,
At least one of the niobium element and the tantalum element in which the sum of niobium and tantalum is 4.75% to 5.5% and Ta is less than 0.2%;
0.65 to 1.15% titanium,
0.20 to 0.80% aluminum,
Less than 0.006% boron,
Less than 0.015% phosphorus,
10. A method according to any one of the preceding claims, comprising a remaining percentage of iron and impurities resulting from processing.   The method of claim 10, wherein the first step is performed at 920 to 990 ° C for at least 30 minutes and the second step is performed at a temperature of 960 to 1010 ° C for 5 to 45 minutes.   The total Nb and Ta content of the alloy is 5.2 to 5.5%, the first step is performed at 960 to 990 ° C. for 45 minutes to 2 hours, and the second step is 990 for 5 to 45 minutes. The process according to claim 11, wherein the process is carried out at from 10 to 1010 ° C.   The total Nb and Ta content of the alloy is 4.8 to 5.2%, the first step is performed at 920 to 960 ° C. for 45 minutes to 2 hours, and the second step is for 5 to 45 minutes. The process according to claim 11, wherein the process is carried out at 960 to 990 ° C. The alloy is in weight percent
55 to 61% nickel,
19 to 22.5% chromium,
7 to 9.5% molybdenum,
At least one of the niobium and tantalum elements, the sum of niobium and tantalum being 2.75 to 4%, and Ta being less than 0.2%;
1 to 1.7% titanium,
Less than 0.55% aluminum,
Less than 0.5% cobalt,
Less than 0.03% carbon,
Less than 0.35% manganese,
Less than 0.2% silicon,
Less than 0.006% boron,
Less than 0.015% phosphorus,
Less than 0.01% sulfur,
10. A method according to any one of the preceding claims, comprising a remaining percentage of iron and impurities resulting from processing. The alloy is in weight percent
12-20% chromium,
2 to 4% molybdenum,
At least one of niobium and tantalum elements, the sum of niobium and tantalum being 5 to 7%, and Ta being less than 0.2%;
1 to 2% tungsten,
5-10% cobalt,
0.4 to 1.4% titanium,
0.6 to 2.6% aluminum,
6-14% iron,
Less than 0.1% carbon,
Less than 0.015% boron,
Less than 0.03% phosphorus,
10. A method according to any one of the preceding claims, comprising a remaining percentage of nickel and impurities resulting from processing.   16. A method according to any one of claims 10, 14 or 15, characterized in that the alloy has a phosphorus content of greater than 0.007% by weight.   The first step and the second step are performed at the sub-solvus temperature of the δ phase of the alloy, and the first step is performed at a temperature between 50 ° C. below the δ solvus temperature and 20 ° C. below the δ solvus temperature. The method according to claim 1, wherein the second step is performed at a temperature between 20 ° C. below the δ solvus temperature and a δ solvus temperature.   18. A method according to any one of the preceding claims, characterized in that the temperature of the hot-worked blank part is kept constant during at least one of the steps.   The third step is performed at 700 to 750 ° C. for 4 to 16 hours, the fourth step is performed at 600 to 650 ° C. for 4 to 16 hours, and the third step and the fourth step 19. A method according to any one of the preceding claims, characterized in that cooling is performed between 50 ° C / hour and 50 ± 10 ° C / hour.   Between the first step and the second step, the hot work alloy is maintained at at least one intermediate temperature between the temperature of the first step and the temperature of the second step for not more than 1 hour. 20. A method according to any one of the preceding claims, characterized in that   21. A method according to any one of the preceding claims, characterized in that the blank part is prepared in the form of an ingot and then hot processed.   21. A method according to any one of the preceding claims, wherein the blank part has been prepared using powder metallurgy.   23. A nickel-based superalloy part obtained from a blank part manufactured using the method according to any one of the preceding claims.   24. The component of claim 23, wherein the component is an aviation or onshore gas turbine component.

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