EP3117017B1 - Precipitation hardening nickel alloy, part made of said alloy, and manufacturing method thereof - Google Patents

Description

The invention relates to alloys based on nickel (superalloys), and more precisely those intended for the manufacture of parts to be used at high temperatures. Typically, this is the case of the elements of terrestrial, aeronautical and other turbines.

For this type of use, a NiCo20Cr20MoTi alloy (AFNOR standard) known as "C263" is known whose composition is typically Ni, Cr (19-21%), Co (19-21%), Mo (5,6- 6.1%), Ti (1.9-2.4%), Al (<0.6%). The percentages are percentages by weight, as will be the case for all the compositions indicated thereafter.

It is a structurally hardened alloy, which is provided by the presence of γ 'phase (Ni ₃ Ti, Al), and which has good forgeability and weldability properties. On this last point, this is due to the fact that, contrary to what is often encountered for alloys hardened by the γ 'phase, it is not subject to the phenomenon of cracking due to embrittlement cracking under stress. heat (in English "strain age cracking") in the welding areas. It also has good ductility in hot traction and satisfactory heat resistance. In general, its weldability / forgeability compromise is advantageous.

However, it has the disadvantage of having a microstructural instability between 700 and 900 ° C, temperature range in which phase η can be formed at the expense of the γ phase (see reference: Zhao, Metallurgical and Materials Transactions A, 2001, vol.32A, pp1271-1282 ). Ductility and resilience are degraded. It is therefore not optimally suitable for uses where the parts are brought to such temperatures.

Other alloys are known for such uses and do not exhibit this structural instability, but they have other disadvantages.

The alloy known as INCO 617 (Ni, Cr (20-24%), Co (10-15%), Mo (8-10%), Al (0.8-1.5%), Ti (0-0.6 %)) has a good compromise weldability / forgeability, but its mechanical properties hot (especially at about 750 ° C which is a frequent use temperature for the parts to which the invention is addressed in a preferred manner) are insufficient.

The alloy known as RENE 41 (Ni, Cr (18-20%), Co (10-12%), Mo (9-10.5%), Al (1.4-1.6%) , Ti (3-3.3%)), on the other hand, has good mechanical properties when hot, but its weldability / forgeability compromise is not optimal. It is the same for the alloy known under the name of WASPALOY (Ni, Cr (18-21%), Co (12-15%), Mo (3.5-5%), Al (1.2-1.6%), Ti (2.75-3.25%). These compromises weldability / forgeability unsatisfactory are probably due to a proportion of phase γ 'too important.

JP 2013 095949 discloses a heat resistant austenitic steel, rich in Co, having good weldability and good creep resistance. It contains C = 0.03-0.15%; If ≤ 1%; Mn ≤ 2%; Ni = 40-80%; Cr = 15-40%; Mo = 4-15%; Co = 15-25%; Ti ≤ 3%; Al ≤ 2%; B = 5-100 ppm; N ≤ 300 ppm; O ≤ 300 ppm, P ≤ 0.03%; S ≤ 0.03%, Fe may also be present up to an undefined level to make up the rest of the composition with the impurities. The grain size must be less than a value depending on the B, P and S contents Ca, Mg, and rare earths may also be present.

JP 61-235529 discloses an alloy for a continuous casting machine support roll, having a high resistance to heat and oxidation. It contains 14-25% Cr; 5-25% Co; 0.1-7% of W; 0.1-3% Ti; 0.1-3% Al with Ti + Al ≤ 4%; ≤ 0.01% of O; ≤ 0.03% of N; ≤ 0.06% C; ≤ 1.0% Mn; ≤ 0.5% Si; ≤ 0.03% of P; ≤ 0.03% of S; the rest being Ni.

There is therefore a need for manufacturers to have Ni base alloys for high temperature applications (typically 700-900 ° C) having both a good microstructural stability at the temperatures of use, good mechanical properties at these same temperatures , and simultaneously a good forgeability and good weldability allowing the manufacture of said parts in the desired configurations and their integration in the devices for which they are intended.

• 18% ≤ Cr ≤ 22%, de préférence 18% ≤ Cr ≤ 20% ;
• 18% ≤ Co ≤ 22%, de préférence 19% ≤ Co ≤ 21% ;
• 4% ≤ Mo + W ≤ 8%, de préférence 5,5% ≤ Mo + W ≤ 7,5% ;
• traces ≤ Zr ≤ 0,06% ;
• traces ≤ B ≤ 0,03%, de préférence traces ≤ B ≤ 0,01% ;
• traces ≤ C ≤ 0,1%, de préférence traces ≤ C ≤ 0,06% ;
• traces ≤ Fe ≤ 1% ;
• traces ≤ Nb ≤ 0,01% ;
• traces ≤ Ta ≤ 0,01% ;
• traces ≤ S ≤ 0,008% ;
• traces ≤ P ≤ 0,015% ;
• traces ≤ Mn ≤ 0,3% ;
• traces ≤ Si ≤ 0,15% ;
• traces ≤ O ≤ 0,0025% ;
• traces ≤ N ≤ 0,0030% ;

le reste étant du nickel et des impuretés résultant de l'élaboration, les teneurs en Al et Ti satisfaisant, de plus les conditions :

• (1) Ti/Al ≤ 3 ;
• (2) Al + 1,2 Ti ≥ 2% ;
• (3) (0,2 Al - 1,25)² - 0,5 Ti ≥ 0% ;
• (4) Ti + 1,5 Al ≤ 4,5%.

For this purpose, the subject of the invention is a nickel-based alloy with a structural hardening, characterized in that its composition is, in weight percentages:

• 18% ≤ Cr ≤ 22%, preferably 18% ≤ Cr ≤ 20%;
• 18% ≤ Co ≤ 22%, preferably 19% ≤ Co ≤ 21%;
• 4% ≤ Mo + W ≤ 8%, preferably 5.5% ≤ Mo + W ≤ 7.5%;
• traces ≤ Zr ≤ 0.06%;
• traces ≤ B ≤ 0.03%, preferably traces ≤ B ≤ 0.01%;
• traces ≤ C ≤ 0.1%, preferably traces ≤ C ≤ 0.06%;
• traces ≤ Fe ≤ 1%;
• traces ≤ Nb ≤ 0.01%;
• traces ≤ Ta ≤ 0.01%;
• traces ≤ S ≤ 0.008%;
• traces ≤ P ≤ 0.015%;
• traces ≤ Mn ≤ 0.3%;
• traces ≤ If ≤ 0,15%;
• traces ≤ 0 ≤ 0.0025%;
• traces ≤ N ≤ 0.0030%;

the rest being nickel and impurities resulting from the elaboration, the contents of Al and Ti satisfying, moreover the conditions:

• (1) Ti / Al ≤ 3;
• (2) Al + 1.2 Ti ≥ 2%;
• (3) (0.2 Al - 1.25) ² - 0.5 Ti ≥ 0%;
• (4) Ti + 1.5 Al ≤ 4.5%.

Its γ 'phase fraction is preferably between 5 and 20%.

The solvus temperature of its γ 'phase is preferably less than or equal to 980 ° C.

The subject of the invention is also a process for manufacturing a nickel-based alloy part, characterized in that an ingot having the previously defined composition is prepared and homogenized at a temperature of at least 1150 ° C. for 24 to 72 hours, it is hot worked by forging or rolling in a supersolvus temperature range, it is dissolved at a temperature of 1100 to 1200 ° C for 1 to 4 hours, it is cooled to at least 1 ° C / min, for example in water, is aged at a temperature of 750 to 850 ° C for 7 to 10 hours, and is cooled, for example in calm air, or in an enclosure.

The invention also relates to a turbine element land or aeronautical alloy nickel-based, characterized in that it was prepared according to the above method.

• La stabilité de la phase γ' à haute température (700-900°C, en particulier 750°C), pour éviter qu'elle ne se transforme en phase aciculaire η (de composition Ni₃Ti, donc dépourvue d'AI) ;
• La fraction de phase γ' formée à 700-900°C, en particulier à 750°C ;
• La température de solvus de la phase γ'.

As will be understood, the invention is based on an optimization of the known C263 grade, which essentially passes through a judiciously chosen balance between the contents of Al and Ti. This balance will drive:

• The stability of the γ 'phase at high temperature (700-900 ° C, in particular 750 ° C), to prevent it from becoming acicular phase η (of composition Ni ₃ Ti, thus devoid of AI);
• The γ 'phase fraction formed at 700-900 ° C, in particular at 750 ° C;
• The solvus temperature of the γ 'phase.

On the remainder of the composition of the alloy, the changes relative to the known C263 are small, and it has been verified that the optimizations of the contents of Al and Ti according to the invention do not lead to a modification of the advantageous properties of the alloy that are not directly related to the γ 'phase.

• Les figures 1 à 8 qui montrent des micrographies d'échantillons de référence (figures 1 et 5 à 8) et selon l'invention (figures 2 à 4) ;
• La figure 9 qui montre les résultats d'essais de mesure de la résistance à la traction Rm de ces échantillons en fonction de la température ;
• La figure 10 qui montre les résultats d'essais de mesure de la limite élastique conventionnelle Rp_0,2 de ces échantillons en fonction de la température ;
• La figure 11 qui montre les résultats d'essais de mesure de l'allongement à la rupture A% de ces échantillons en fonction de la température ;
• La figure 12 qui montre les résultats d'essais de mesure de la striction Z% de ces échantillons en fonction de la température ;
• La figure 13 qui montre les résultats d'essais de fluage rupture à 750°C de ces échantillons, où la contrainte à rupture est donnée en fonction du paramètre de Larson-Miller ;
• La figure 14 qui montre les résultats d'essais de résilience de deux échantillons (un échantillon de référence et un échantillon selon l'invention), réalisés après le traitement thermique final de l'échantillon et après un survieillissement à 750°C pendant 3000 h représentatif de ce que pourrait subir le métal lors d'une utilisation à laquelle il est destiné de manière privilégiée ;
• Les figures 15 à 18 qui montrent un échantillon selon l'invention et des échantillons de référence en cours de forgeage.

The invention will be better understood with the aid of the description which follows, given with reference to the following appended figures:

• The Figures 1 to 8 which show micrographs of reference samples ( figures 1 and 5 to 8 ) and according to the invention ( Figures 2 to 4 );
• The figure 9 which shows the results of tests for measuring the tensile strength Rm of these samples as a function of temperature;
• The figure 10 which shows the results of measurements of the conventional yield strength Rp _0.2 of these samples as a function of temperature;
• The figure 11 which shows the results of tests for measuring the elongation at break A% of these samples as a function of temperature;
• The figure 12 which shows the results of tests of the necking Z% of these samples as a function of temperature;
• The figure 13 which shows the results of creep rupture tests at 750 ° C of these samples, where the breaking stress is given according to the Larson-Miller parameter;
• The figure 14 which shows the results of resilience tests of two samples (a reference sample and a sample according to the invention), carried out after the final heat treatment of the sample and after overaging at 750 ° C. for 3000 hours representative of this sample. the metal could suffer in a use for which it is intended in a privileged manner;
• The Figures 15 to 18 which show a sample according to the invention and reference samples being forged.

A first condition for optimizing the equilibrium between Al and Ti is that the phase formation η is avoided at the temperatures of use of the alloy during its preferred uses, that is to say at temperatures of 700-900 ° C, typically of the order of 750 ° C. The formation of the η phase is directly related to the Ti and Al contents present in the alloy and to their ratio. It is thus necessary to determine the ranges of contents in these elements which make it possible to avoid it with 700-900 ° C, considering the remainder of the composition of the alloy. Thermodynamic calculations, carried out using the THERMOCALC software commonly used by metallurgists and which was also used as a first approach for the rest of the optimization, indicated that for C263, if the Ti / Al ratio was lower or equal to 3, phase formation η was avoided, regardless of the Al level in the alloy.

It is therefore necessary to respect the condition: $$\mathrm{Ti}/\mathrm{al}\le 3$$

Another condition is that to guarantee the tensile and creep properties at 700-900 ° C, the atomic percentage of γ 'phase present at these temperatures in the alloy should be at least 5%. Below this value, there is not sufficient structural hardening. It is considered that this condition is fulfilled when the weight percentages of Al and Ti respect the relation: $$\mathrm{al}+1.2\mathrm{Ti}\ge 2\%.$$

As regards the forgeability properties (or hot deformability in general, for example by rolling) and weldability, we can say the following.

Under standard forging conditions at high temperature, the forging is carried out in a temperature range where there is no γ 'phase precipitation which would make the metal too hard and subject to the appearance of defects, such as cracks, during deformations. It is therefore performed at a temperature above the solvus temperature of this phase. This temperature is therefore advantageous not to be too high, for a forging is possible in industrial conditions. More precisely, the solvus temperature of the γ 'phase must be as low as possible in order to avoid the precipitation of this phase during the inevitable cooling of the product during the forging.

It is also necessary to take into account the fraction of phase γ 'that can precipitate at high temperature. Indeed, the higher the hardening phase fraction precipitated at high temperature, the more the alloy is likely to harden during temperature variations that may occur during the forging, which can complicate the execution of the operation. This undesired γ 'phase precipitation at this point in product preparation also has an influence on the weldability because of the possibility of cracking due to embrittlement under heat stress. Indeed, the greater the γ 'phase fraction precipitated in the welded zone, the greater the stresses generated by the precipitation of the γ' phase in the same zone during cooling are high and promote cracking after welding.

In order for the good conditions of hot formability and weldability to be simultaneously satisfied, it is therefore necessary to maintain a solvus temperature of the γ 'phase of at most 980 ° C, and to limit the fraction of the phase γ' present at 700-900 ° C at 20% (in atomic%), in particular at 750 ° C.

• (3) (0,2 Al - 1,25)² - 0,5 Ti ≥ 0% ;
• (4) Ti + 1,5 Al ≤ 4,5%

These conditions are respected if the weight contents of Ti and Al comply with both conditions:

• (3) (0.2 Al - 1.25) ² - 0.5 Ti ≥ 0%;
• (4) Ti + 1.5 Al ≤ 4.5%

Regarding the other elements that must or may be present, either as mandatory or optional alloying elements, or as impurities to be limited, we can say the following. The preferred ranges are those where one is most assured of obtaining the cited advantages of each element without having the disadvantages.

The Cr content is between 18 and 22%, preferably 18 to 20%. Cr is important to ensure resistance to corrosion and oxidation, and to establish the resistance of the alloy to the effects of the environment at high temperatures. An excessively high content favors the obtaining of undesirable fragile phases, such as the σ phase, and the limit of 22% by weight is set accordingly.

The content of Co is between 18 and 22%, preferably 19 to 21%. A high Co content is necessary in order to improve the forgeability of the grade by decreasing the solvus temperature of the γ 'phase, however, it must be limited, mainly, for cost reasons.

The sum of the contents in Mo and W must be between 4 and 8%, preferably 5.5 to 7.5%. These two elements are substitutable for each other. The lower limit of 4% guarantees structural hardening and good creep resistance, and the upper limit of 8% prevents the formation of harmful phases.

The Zr content is between traces (in other words, a lack of voluntary addition, the residual content of possible Zr resulting only from the melting of the raw materials and the elaboration, with the associated impurities) and 0.06%. .

The content of B is between traces and 0.03%, preferably 0.003 to 0.01%.

The content of C is between traces and 0.1%, preferably 0.04 to 0.06%.

These last three elements form segregations at grain boundaries that contribute to heat resistance and ductility by trapping any harmful elements present, such as S. They promote creep resistance under low stress and high temperature conditions. However, if they are present in excess, they decrease the melting temperature of the segregated zones and strongly alter the forgeability. Their eventual presence must therefore be well controlled.

It should be understood that the preferred contents of the elements just mentioned are independent of each other. In other words, an alloy which has a preferential content on one or more of them only, but not on the others, must nevertheless be considered as an advantageous variant of the invention.

Concerning the elements whose contents have interest to be minimized as much as possible, one can say the following.

The Fe content is limited to 1% maximum. Beyond, it may form phases harmful to the properties of the alloy.

The contents of Nb and Ta are both limited to 0.01% maximum. These elements are expensive and have a strong tendency to segregate without these segregations having advantages that could offset their disadvantages (contrary to what can happen for Zr, B and C).

The contents of S, P, Mn and Si must also be limited so as not to reduce the hot ductility. An excess of Si would also cause a precipitation of Laves phases during solidification, and it will be difficult to put them back in solution during subsequent heat treatments. Resilience would be degraded.

The maximum levels allowed for these elements are therefore 0.008% for S, 0.015% for P, 0.3% for Mn, and 0.15% for Si.

To guarantee good mechanical properties of the alloy, it is necessary to limit the content of O to 25 ppm at the maximum and the content of N to 30 ppm at most. For this purpose, evacuation under vacuum and also involving a process such as electroslag remelting (ESR) or vacuum arc remelting (VAR) is particularly recommended. But from these points of view, the alloys of the invention are not particularly distinguished from the usual C263 to which they are called to substitute.

With regard to the process for preparing the parts, typically an ingot having the above composition is prepared by double or triple melting, thus involving at least one of the ESR and VAR processes, and homogenized at a temperature of at least 1150. For 24 to 72 hours, it is hot-worked by forging or rolling in a supersolvus temperature range, dissolved at a temperature of 1100 to 1200 ° C for 1 to 4 hours, cooled rapidly to room temperature. minus 1 ° C / min, for example in water, it is aged at 750 to 850 ° C for 7 to 10 hours, and is cooled, for example in calm air, or in an enclosure. Depending on the intended applications, variations can be made to this process, by not performing some of these steps or by adding others. They can be followed in particular by machining or any other operation of final dimensioning of the part.

An elaboration of the part using a powder metallurgy process and resulting in a product having the required compositional properties would also be conceivable.

Tests were performed on samples whose compositions are listed in Table 1. Table 1: Compositions of the samples tested Ech. Or% Cr% Co% Mo% W% B% C% % Zr al% Ti% O ppm N ppm (1) (2) (3) (4) AT 51,60 19.71 20.15 5.98 traces 0.005 0,051 0.02 0.77 1.50 3.5 17 1.06 2.57 0.45 2.66 B 47.50 20.86 20.49 5.96 1.43 0,010 0,050 0.02 1.95 1.13 3.1 18 0.58 3.31 0.17 4.06 C 51,00 19.79 20,12 6.13 traces 0,010 0,050 0.01 2.64 0.22 3.4 15 0.08 2.90 0.41 4.18 D 51.50 19.74 20.00 6.20 traces traces 0,052 0.01 0.42 2.24 3.1 22 5.33 3.11 0.24 2.87 E 50,40 19.60 20.00 5.97 traces 0,002 0.049 0,003 3.00 0.252 3 16 0.08 3.30 0.30 4.75 F 48,20 19.52 20.60 4.22 3.48 0,010 0,050 0.02 3.62 0.15 4 17 0.04 3.80 0.20 5.58 G 49.70 19.97 18.50 7.50 traces 0,010 0,060 0.02 2.20 1.95 3.3 14 0.89 4.54 -0.32 5.25 H 52.10 20.00 18.20 8.00 traces traces 0,060 0.01 1.10 0.48 3.2 16 0.44 1.68 0.82 2.13

Samples A, B and C correspond to the invention, the other samples are reference alloys which do not comply with at least one of the conditions (1) to (4) previously defined because of their Al and Ti contents. Sample B corresponds to the version of the invention considered optimal, where the contents of all the elements are in the preferred ranges. The reference sample D corresponds to a conventional C263 type alloy which does not respect the relation (1). Sample E and sample F do not respect relationship (3). Sample G does not respect relationships (3) and (4). Sample H does not respect relationship (2). This shows that the respect of all relations (1) to (4) is necessary to obtain the desired results.

The samples tested were made by VIM-VAR double melting (that is, as is conventional, by melting the raw materials in a vacuum induction furnace, followed by casting and solidification of an electrode, the latter being refined by vacuum reflow in an arc furnace), to obtain ingots of 200 kg. This method is commonly used for the manufacture of ingots for forming forged or laminated parts of high purity inclusionary and low levels of residual elements, especially gaseous. It is however not necessarily used to develop the alloys of the invention, if they are intended for the production of parts that do not have very high requirements on these points. In these cases, less complex conventional methods of preparation can be used, provided that they make it possible to reach the necessary low levels on certain residual elements, in particular by a suitable choice of raw materials.

These ingots were homogenized at a temperature greater than 1150 ° C for 48 h, then forged into rods with a diameter of 80 mm between 1200 and 1050 ° C.

• Mise en solution à 1140°C +/- 10°C pendant 2 h, suivie d'une trempe à l'eau ;
• Vieillissement à 800°C+/-10°C pendant 8 h suivi d'un refroidissement à l'air.

The examples then underwent the following heat treatments:

• Dissolving at 1140 ° C +/- 10 ° C for 2 h, followed by quenching with water;
• Aging at 800 ° C +/- 10 ° C for 8 h followed by cooling in air.

This heat treatment is typical of the C263 alloy for its usual applications such as turbine elements.

The THERMOCALC software does not provide any phase appearance η for these samples in their treatment conditions, except for sample D.

In fact, micrographs were made on portions of said samples which had undergone overaging at 750 ° C for 3000 h to simulate a use of the corresponding alloys at high temperature. Field electron microscopy micrographs are shown on the figures 1 (sample D), 2 (sample A), 3 (sample B), 4 (sample C), 5 (sample E), 6 (sample F), 7 (sample G) and 8 (sample H).

It is confirmed that only the sample D, representative of a conventional C263 alloy, contains a significant amount of η acicular phase (in needles). The other samples, in particular those of the invention A, B and C, do not have this phase whose invention aimed in particular to prevent the appearance during use at 700-900 ° C, typically 750 ° C.

The figure 9 shows the results of mechanical tensile tests on these same samples for the measurement of Rm, carried out between ambient and 800 ° C. The figure 10 shows the measurement results of Rp _0.2 , the figure 11 shows the results of measurement of elongation at break A%, and the figure 12 shows the results of Z% necking tests carried out under the same conditions.

It turns out that the alloys B and C according to the invention have tensile results (Rm and Rp _0.2 ) similar to those of the reference alloy D. The tensile results of the alloy A according to FIG. The invention is slightly degraded with respect to those of alloy D but remains satisfactory. And the hot ductility of alloy A is the best of all, which can be a benefit for some uses. The invention therefore makes it possible to optimally optimize or preserve all of these mechanical properties with respect to the reference alloy C263.

Alloys E, F and G have very good results in traction, especially hot. But their loss of hot ductility is very important, which can be attributed to a poor balancing of the contents of Al and Ti.

Alloy H is unsatisfactory in all respects at high temperatures.

The figure 13 shows the results of breaking creep tests at 750 ° C: the breaking stress in MPa is given according to the Larson-Miller parameter (PLM) as is conventional to proceed.

The alloys A, B, C according to the invention, and the reference alloys F and G have longer rupture times than that of the reference alloy D. This shows that, from this point of view too, the invention provides an improvement in the performance of the alloy D which is closest thereto. The alloy E has a short life because of its insufficient hot ductility, and the tests could not be prolonged beyond a PLM of 23.4. Alloy H is, again, very clearly unsatisfactory.

The figure 14 shows the results of resilience tests conducted on several test specimens of the alloys A according to the invention and D of reference, firstly after treatment heat treatment solution and then aging as described above, secondly after over-aging of 3000 h at 750 ° C following the previous heat treatment, again to simulate the evolution of the alloy in use. The results are clear: the resilience Kv is practically unaffected by the over-aging of the sample A, whereas it drops very substantially for the sample D. This confirms that the phase η formed during a high use The temperature of the conventional C263 alloy has a strong embrittling effect, and the invention overcomes this problem.

Forging tests were also carried out, under identical conditions (homogenization at more than 1150 ° C for 48 h and forging at 1200 ° C-1050 ° C up to 80 mm diameter), and Figures 15 to 18 present the results obtained.

The alloys A, B and C according to the invention, as well as the alloy H of reference, were forged without problems as would have been the alloy D: no crack appeared during the forging. The figure 15 shows the alloy A being forged at about 1100 ° C and no crack is actually visible. The figure 16 shows the alloy E being forged at the same temperature, and slight cracks are visible. The figure 17 shows the F alloy being forged at the same temperature, and the cracks are much deeper than in the previous cases. The figure 18 shows the G alloy being forged at the same temperature, and again deep coves are visible. The good forgeability of the alloys according to the invention is thus confirmed, and is attributed to a lower proportion of γ 'phase than for the reference samples E, F and G.

A preferred application of the invention is the manufacture of terrestrial and aeronautical turbine elements, but it is, of course, not exclusive.

Claims (5)

1. A precipitation hardened nickel-base alloy, characterized in that its composition is, in weight percentages:

   - 18% ≤ Cr ≤ 22%, preferably 18% ≤ Cr ≤ 20%;

   - 18% ≤ Co ≤ 22%, preferably 19% ≤ Co ≤ 21%;

   - 4% ≤ Mo + W ≤ 8%, preferably 5.5% ≤ Mo + W ≤ 7.5%;

   - trace amounts ≤ Zr ≤ 0.06%;

   - trace amounts ≤ B ≤ 0.03%, preferably trace amounts ≤ B ≤ 0.01%;

   - trace amounts ≤ C ≤ 0.1%, preferably trace amounts ≤ C ≤ 0.06%;

   - trace amounts ≤ Fe ≤ 1%;

   - trace amounts ≤ Nb ≤ 0.01%;

   - trace amounts ≤ Ta ≤ 0.01%;

   - trace amounts ≤ S ≤ 0.008%;

   - trace amounts ≤ P ≤ 0.015%;

   - trace amounts ≤ Mn ≤ 0.3%;

   - trace amounts ≤ Si ≤ 0.15%;

   - trace amounts ≤ O ≤ 0.0025%;

   - trace amounts ≤ N ≤ 0.0030%;

   the remainder being nickel and impurities resulting from the elaboration, the Al and Ti contents further satisfying the conditions:

   - (1) Ti/Al ≤ 3 ;

   - (2) Al + 1.2 Ti ≥ 2% ;

   - (3) (0.2 Al - 1.25)² - 0.5 Ti ≥ 0% ;

   - (4) Ti + 1.5 Al ≤ 4.5%.
2. The alloy according to claim 1, characterized in that its γ' phase fraction is comprised between 5 and 20%.
3. The alloy according to claims 1 or 2, characterized in that the solvus temperature of its γ' phase is less than or equal to 980°C.
4. A method for manufacturing a part in a nickel-base alloy, characterized in that an ingot is prepared, having the composition according to claim 1, it is homogenized at a temperature of at least 1,150°C for 24 to 72 h, it is hot-worked by forging or rolling in a supersolvus temperature range, it is solution treated at a temperature from 1,100 to 1,200°C for 1 to 4 h, it is cooled at a rate of at least 1°C/min, for example in water, it is aged at a temperature from 750 to 850°C for 7 to 10 h, and it is cooled for example in calm air or in an enclosure.
5. An element of a land or aeronautical turbine engine, characterized in that it was prepared according to the process of claim 4.

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