EP0971041B1 - Single crystal nickel-based superalloy with high solvus gamma prime phase - Google Patents

Abstract

A nickel base superalloy, having specified contents of rhenium and ruthenium, is new. A nickel base superalloy has the composition (by wt.) 3.5-7.5% Cr, 0-1.5% Mo, 1.5-5.5% Re, 0-5.5% Ru, 3.5-8.5% W, 5-6.5% Al, 0-2.5% Ti, 4.5-9% Ta, 0.08-0.12% Hf, 0.08-0.12% Si, balance Ni and impurities.

Description

The present invention relates to superalloys based on nickel particularly suitable for the manufacture of monocrystalline blades stationary and mobile gas turbine engines, and showing a high creep resistance at very high temperatures while retaining good environmental resistance of combustion. These alloys are more particularly suitable to applications in aeronautical engines used to propulsion of airplanes and helicopters.

The nickel-based monocrystalline superalloys are the the most efficient materials used today for manufacture of stationary and movable turbine blades aeronautical gas turbomachines. ONERA's work in this area began in the late 1970s and are translated, among other things, by the filing of various patents of invention relating to monocrystalline superalloys intended for different fields of application: FR 2 503 188, FR 2 555 204, FR 2 557 598, FR 2 599 757, FR 2 643 085, FR 2 686 902.

Evolution of the performance of aeronautical gas turbines, translated in terms of power and efficiency specific and lifespan, requires the ability to have alloys for turbine blades showing characteristics mechanical at high temperature (650 to 1150 ° C) and a continuous corrosion and hot oxidation resistance improved. Extreme operating conditions can indeed bring the metal to temperatures above 1100 ° C. To optimize resistance to hot corrosion and to hot oxidation, the monocrystalline blades in superalloy are also generally coated with a protective deposit of the nickel aluminide or alloy type MCrAlY. To compensate for cracking and breaking possible of these protective layers under the effect of thermal cycling, which would penalize the lifespan of parts, the superalloys must however show a high intrinsic resistance to corrosion and oxidation.

In polycrystalline blades cast by conventional foundries, much of the deformation hot during service occurs at seals grains which limits the life of the parts. The development of the monocrystalline solidification process a allowed, by eliminating grain boundaries, to increase by dramatically the performance of superalloys at nickel base. In addition, the process makes it possible to select the preferred growth direction of the part monocrystalline and therefore choose an orientation <001>, optimal vis-à-vis the creep and fatigue resistance thermal which are the two modes of stress causing the greatest damage to the turbine blades.

Successive improvements in the mechanical performance, in particular in creep, of these superalloys for monocrystalline blades have been made possible by optimizations of their chemical compositions. Indeed, in addition to nickel which is the major constituent of these alloys, various elements of addition bring their specific contributions to the properties of these. The roles of these elements will be detailed later. In the monocrystalline superalloys covered by the aforementioned patents, the major addition elements (weight concentrations up to a few percent) have generally been chosen from the following list: chromium (Cr), cobalt (Co), molybdenum (Mo) , tungsten (W), aluminum (Al), titanium (Ti), tantalum (Ta), niobium (Nb). The elements Cr, Co, Mo and part of the W mainly participate in the hardening of the austenitic matrix (γ phase) where they come into solution. The elements Al, Ti, Ta and Nb favor the precipitation in the γ matrix of hardening particles of a second phase of the Ni ₃ type (Al, Ti, Ta, Nb) (γ 'phase). Minor elements (weight concentrations less than 0.5%) such as silicon (Si), hafnium (Hf), can also be added in order to optimize the resistance to the environment as demonstrated in FR 2 686 902.

Since the early 1980s, a large number of patents devoted to new chemical compositions of superalloys for monocrystalline blades have been deposited in the world. The most recent developments have consisted of particular to incorporate in these alloys the elements rhenium (Re) and ruthenium (Ru) refractories. These additions are mainly aimed at improving the creep resistance at high temperature of monocrystalline superalloys while retaining a stable microstructure at high temperature vis-à-vis of the formation of particles of intermetallic phases which are likely to cause property loss of these alloys.

Various patents thus protect areas of compositions of monocrystalline superalloys containing additions of one and / or the other of the elements Re and Ru, in particular US 4,719,080 (United Technologies Corporation), US 4,935,072 (Allied-Signal Inc), US 5 151 249 (General Electric) US 5,270,123 (General Electric) and US 5,482,789 (General Electric). However, the information available regarding properties of these alloys are very limited, and do not allow not to judge the industrial interest of these additions.

In France today monocrystalline superalloys used are called "first generation", as by example the shades AM1 and MC2 both covered by the FR patent 2,557,598 and the AM3 alloy protected by the FR patent 2,599,757. Of these, the MC2 alloy is considered to be the best performing alloy in terms of resistance creep up to 1100 ° C. The future needs of engine manufacturers however, require the availability of alloys to blades more efficient than these first generation alloys. In particular, it is necessary to increase the maximum admissible temperatures by the alloys constituting the turbine blades.

The object of the invention is therefore to propose a new family of nickel-based monocrystalline superalloys showing improved creep resistance, in particular at temperatures above 1100 ° C, but also at lower temperatures of interest to various parts of blades, compared to those of industrially exploited alloys.

To this end, we sought to introduce new elements addition, without penalizing other essential characteristics good behavior of alloys, such as mass volume, corrosion and oxidation resistance to microstructural stability.

Analysis of the state of the art and the results work carried out by the inventor quickly showed that only alloys containing additions of rhenium could allow the creep resistance of the MC2 alloy beyond 1100 ° C. To counterbalance some harmful effects of rhenium (excessive density, microstructural instability), it also seems advantageous to incorporate ruthenium.

The invention relates to a nickel-based superalloy suitable for the manufacture of parts of turbomachines by monocrystalline solidification, characterized in that its mass composition is as follows: Cr 3.5 to 7.5% MB 0 to 1.5% Re 1.5 to 5.5% Ru 0 to 5.5% W 3.5 to 8.5% al 5 to 6.5% Ti 0 to 2.5% Your 4.5 to 9% Hf 0.08 to 0.12% Yes 0.08 to 0.12%, the complement to 100% being constituted by Ni and any impurities.

More particularly, the invention proposes such a superalloy having the following composition by mass: Cr 3.5 to 5.5% MB 0 to 1.5% Re 4.5 to 5.5% Ru 2.5 to 5.5% W 4.5 to 6.5% al 5 to 6.5% Ti 0 to 1.5% Your 5 to 6.2% Hf 0.08 to 0.12% Yes 0.08 to 0.12%, the complement to 100% being constituted by Ni and any impurities.

More particularly still, the mass composition of the superalloy is as follows: Cr 3.5 to 5.5% MB 0 to 1.5% Re 3.5 to 4.5% Ru 3.5 to 5.5% W 4.5 to 6.5% al 5.5 to 6.5% Ti 0 to 1% Your 4.5 to 5.5% Hf 0.08 to 0.12% Yes 0.08 to 0.12%, the complement to 100% being constituted by Ni and any impurities.

Three specific superalloy compositions according to the invention are given below: Cr 3.5 to 4.5% 4.5 to 5.5% 3.5 to 4.5% MB 0.5 to 1.5% 0.5 to 1.5% Re 3.5 to 4.5% 3.5 to 4.5% 4.5 to 5.5% Ru 3.5 to 4.5% 4.5 to 5.5% 2.5 to 3.5% W 4.5 to 5.5% 5.5 to 6.5% 5.5 to 6.5% al 5.5 to 6.5% 5.5 to 6.5% 4.8 to 5.8% Ti 0 to 1% 0 to 1% 0.5 to 1.5% Your 4.5 to 5.5% 4.5 to 5.5% 5.7 to 6.7% Hf 0.08 to 0.12% 0.08 to 0.12% 0.08 to 0.12% Yes 0.08 to 0.12% 0.08 to 0.12% 0.08 to 0.12% the complement to 100% being constituted by Ni and any impurities.

• une masse volumique inférieure dans tous les cas à 9 g.cm^-3, et au mieux à 8,8 g.cm^-3, permettant de minimiser la masse des aubes monocristallines et par conséquent de limiter la contrainte centrifuge agissant sur ces aubes et sur le disque de turbine sur lequel elles sont fixées;
• une aptitude à l'homogénéisation par remise en solution totale des particules de phase γ', y compris les phases eutectiques γ/γ';
• une température élevée de mise en solution de la phase durcissante γ', dans tous les cas supérieure à celle des alliages antérieurs ne contenant ni rhénium, ni ruthénium;
• l'absence de phases intermétalliques fragiles pouvant précipiter au cours de maintiens à haute température et susceptibles d'entraíner une réduction de la résistance au fluage et une fragilisation des alliages;
• une tenue à la corrosion cyclique à chaud et à l'oxydation cyclique à chaud supérieure à celle des alliages antérieurs ne contenant ni rhénium ni ruthénium.

The alloys according to the invention, produced in the form of single crystals of orientation <001>, have the following properties:

• a density lower in all cases to 9 g.cm ^-3 , and at best to 8.8 g.cm ^-3 , making it possible to minimize the mass of the monocrystalline blades and consequently to limit the centrifugal stress acting on these blades and on the turbine disk on which they are fixed;
• an aptitude for homogenization by re-dissolving the particles of phase γ ', including the eutectic phases γ / γ';
• a high dissolution temperature of the hardening phase γ ', in all cases higher than that of the previous alloys containing neither rhenium nor ruthenium;
• the absence of fragile intermetallic phases which can precipitate during maintenance at high temperature and which may cause a reduction in the creep resistance and embrittlement of the alloys;
• a resistance to hot cyclic corrosion and to hot cyclic oxidation than that of previous alloys containing neither rhenium nor ruthenium.

Obtaining all of these characteristics simultaneously optimizes creep resistance at very high temperature and environmental behavior of monocrystalline blades and therefore increase their lifespan as well as gas turbine performance.

The invention thus provides a unique combination of features alloys, which the state of the art does not not allow to predict.

The alloys of the invention are intended for the manufacture of monocrystalline parts, that is to say made up of a single metallurgical grain. This particular structure is obtained using a solidification process directed in a thermal gradient using a device for selecting grain or a monocrystalline germ at the beginning of solidification.

After solidification, the superalloys essentially consist of two phases: the austenitic matrix γ is a solid solution based on nickel in which particles of phase γ ', an intermetallic compound whose composition is based on Ni ₃ Al, precipitate during cooling to solid state. The addition elements are distributed in the two phases γ and γ 'but generally show a particular affinity for one or the other of these two phases. Thus chromium, molybdenum, rhenium and ruthenium are distributed preferentially in the matrix γ while aluminum, titanium and tantalum go preferentially in the phase γ '.

In crude monocrystalline solidification alloys, the distribution of particles of hardening phase γ 'is very heterogeneous in the volume of the single crystal due to segregation chemicals resulting from solidification conditions process specific. The microstructure is said to be dendritic. The precipitates are very fine in the heart of the dendrites which solidifies first during cooling of the alloy and get bigger in the regions then solidifying from the center of the dendrite. Of plus, at the end of solidification, eutectic phases made up of massive particles of γ 'phase containing γ phase lamellas solidify in the regions separating the dendrites.

However, experience has shown that creep resistance nickel-based superalloys was optimized when the γ 'phase particle distribution was homogeneous in the entire volume of the alloy with precipitate sizes less than 1 micrometer, the optimal size of the precipitates depending on the composition of the alloy. The γ 'phase contained in the eutectic phases does not contribute, in particular to the hardening of alloys and the potential for creep resistance of alloys is therefore not completely operated in the raw solidification state. These blocks massifs of eutectic phase γ / γ 'are also sites preferential crack initiation during stresses cyclical resulting from thermal fatigue phenomena due to gas turbine start and stop cycles.

The compositions of the alloys of the invention were chosen so that two-phase microstructures can be obtained γ / γ ', consisting of a homogeneous precipitation of particles γ 'in a matrix γ at the end of the solidification steps monocrystalline and heat treatments detailed by the following. In order to achieve this optimized microstructure, it is necessary to apply treatment first thermal intended to dissolve the precipitates of phase γ ' contained in dendrites and eliminate eutectic phases solidified between the dendrites. The dissolution of precipitated γ 'is performed when the treatment temperature thermal reaches the solvent temperature γ '(temperature of dissolution of the precipitates of phase γ ') characteristic of the chemical composition of the alloy. In practice the value of the solvus γ 'varies periodically in the alloy monocrystalline solidification crude in relation to the local alloy chemistry. Thus the solvus γ 'increases by the heart of the dendrite towards the interdendritic regions of the makes chemical segregations, until reaching the temperature of onset of fusion of the eutectic γ 'phase which is the last solid formed during the cooling of the alloy from the liquid state. This start temperature of eutectic fusion is in practice assimilated to solidus temperature (melting start temperature) of the alloy. The temperature of the homogenization treatment must therefore remain below the solidus temperature.

Practically the complete dissolution of the γ 'precipitates and the γ / γ' eutectics has been obtained in the alloys of the invention thanks to the application of heat treatment sequences including a prior homogenization of the dendritic structures. This heat treatment sequence includes a first 3-hour pre-homogenization treatment at a temperature between 1300 and 1310 ° C, then a progressive increase of 30 ° C at the speed of 3 ° Ch ^-1 , before a new level of 3 hours at a temperature between 1330 and 1340 ° C, the final cooling to be carried out at a speed such that the final size of the precipitates of phase γ 'is less than 300 nm. All of the γ / γ 'eutectic phases are thus eliminated. This result could be obtained for all the alloys of the invention. The sequence of heat treatments which has just been described is an example making it possible to obtain the expected result. This does not preclude the possibility of obtaining a similar result using another sequence of heat treatments, the result of the treatment being more important than the manner of achieving it. The important thing is to have demonstrated the possibility of obtaining such a result in the case of the alloys of the invention.

The alloys of the invention were tested after being subjected to a sequence of homogenization treatments and dissolution of the γ 'phase as described above, then two heat treatments for income allowing fix the size and the volume fraction of the precipitates of γ 'phase. A first income consists of a treatment of 4 to 16 hours at a temperature between 1050 and 1150 ° C allowing the size of the γ 'phase precipitates to be fixed between 300 and 500 nm. A second income treatment consists of in a treatment of 15 to 25 hours at a temperature included between 850 and 870 ° C to optimize the fraction of precipitated γ 'phase. These income treatments are compatible with coating diffusion treatments protectors and soldering treatments generally applied to monocrystalline turbine blades during their manufacturing. Micrographic examination shows that the precipitates of phase γ 'are roughly cubic in shape and represent a volume fraction of at least 70% in the alloy. They are contained in the matrix γ which appears in the form of fine corridors between these precipitates.

The resistance to creep at high temperature is all the more greater than the volume fraction of the hardening phase γ ' precipitated in the alloy is high. At the temperature ambient, the alloys of the invention contain a volume fraction close to 70%. When the temperature increases from room temperature, the γ 'phase gradually dissolves in the γ matrix, slowly until around 1000 ° C, then more quickly above 1000 ° C. When the solvent temperature γ 'is exceeded, the γ 'precipitates are then completely dissolved. The reduction of the volume fraction of the γ 'phase when the temperature increases is one of the causes of the decrease in resistance creep of superalloys.

One of the major contributions of the invention is to increase by significantly the temperature of γ 'solvus in order to conserve a high volume fraction of γ 'phase at temperatures higher than 1100 ° C and therefore obtain resistance very high creep at these temperatures. The invention therefore relates to so-called "high solvus γ '" alloys showing very high creep resistance above 1100 ° C. The experience acquired by the inventor in the field has shown that increases in concentrations of Al, Ti, Ta, Mo and W resulted in an increase in the γ 'solvus. On the other hand additions of the elements Cr and Co led to a decrease of the temperature of the solvent γ '. Regarding rhenium and ruthenium, previous work has not concluded explicitly about their specific role on temperature of γ 'solvus.

Increases in concentrations of increasing elements the γ 'solvus can however lead to effects that can harm the properties of alloys. So concentrations too high in elements Al, Ti and Ta lead to the formation an excessive amount of γ / γ 'eutectic phases during solidification of alloys; these phases can no longer then be totally removed by heat treatments which adversely affects the homogeneity of the alloy and therefore its creep resistance. Plus the concentration in Ta must be limited because this element has a mass high atomic and penalizes alloys from the point of view of the density.

The elements Mo and W also have a beneficial effect on the γ 'solvus but these elements are heavy, in particular W, and their content must be controlled so as not to increase excessively the density of the alloys.

Furthermore, the solubility of these elements in the matrix γ is limited, like that of rhenium and to a lower levels of cobalt and chromium, which can lead to precipitation of intermetallic phases fragile of type σ, µ, P or Laves phase. The presence of these phases called topologically compact (in English T.C.P .: topologically close-packed) may cause loss mechanical properties in superalloys where they rush. Obtaining alloys which are unlikely to forming these fragile intermetallic phases is one of the main arguments of previous patents on superalloys Monocrystalline.

The decrease in concentrations of Cr and Co elements causes a reduction in the temperature of the γ 'solvus. So one of the major ideas of the invention is to refrain from any addition of Co whose role on creep resistance super alloys is low compared to that of others addition items. On the other hand, the chromium was maintained because its presence is essential for maintaining good performance to hot corrosion.

The examples of the invention detailed below show that the objective of obtaining high-solvent alloys has been achieved thanks to a judicious choice of chemical compositions taking into account the considerations which have just been set out.

In addition to optimizing the volume fraction and the γ 'phase solvus temperature, improving the creep resistance of monocrystalline superalloys can be obtained by increasing the concentrations of refractory elements Mo, W, Re and Ta which play a role important in the solid solution hardening of the phases γ and γ '. These heavy elements also slow down the set of elementary mechanisms which are controlled by the diffusion of atoms, which has beneficial consequences on the creep resistance of alloys. The addition of rhenium, in particular, limits the magnification of particles γ 'phase during high temperature maintenance, phenomenon which contributes to degradation over time mechanical properties of superalloys. On the other hand increasing concentrations of refractory elements slows the thermally activated movement of dislocations which propagate strain in superalloys, which has the effect of reducing the creep rate.

The refractory element concentrations must be however carefully balanced so as not to excessively increase the density of the alloys.

The elements W and Mo present at too high contents are harmful with respect to oxidation and corrosion resistance monocrystalline superalloys while the presence of rhenium does not penalize the environmental behavior of these alloys

In addition, the refractory element Ru, in the context of the invention, has the advantage of having a density two times weaker than that of rhenium. Works of the inventor in this field show that the Ru favors less that rhenium the precipitation of intermetallic phases fragile.

The alloys according to the invention also include simultaneous additions of silicon and hafnium. Such additions make it possible to optimize the resistance to oxidation at hot alloys by improving the adhesion of the layer protective alumina formed at high temperature.

Alloys according to the invention have been developed, solidified in the form of crystal-oriented single crystals <001> and tested. This crystallographic orientation is that usually used for solidification directed monocrystalline blades of turbines. It gives to these parts an optimal combination of creep resistance and resistance to thermal fatigue and fatigue mechanical.

By way of example, the nominal chemical compositions (% by weight) of a few alloys of the invention are collated in Table I, together with that of the reference alloy MC2 described in FR 2,557,598. This alloy serves as reference because it is, to the knowledge of the inventor, the most efficient in creep among the alloys containing neither rhenium nor ruthenium. Alloy Or Co Cr MB W Re Ru al Ti Your Yes Hf MC2 Based 5 8 2 8 - - 5 1.5 6 - - MC820 Based - 5 1 8 2 - 5.5 1 6 0.1 0.1 MC533 Based - 7 - 5 3 3 6 - 6 0.1 0.1 MC440 Based - 5 1 4 4 - 5.5 - 9 0.1 0.1 MC722 Based - 4.5 1 7 2.5 2.5 5.8 - 6 0.1 0.1 MC623 Based - 6 1 6 2 3 5.7 0.5 5.5 0.1 0.1 MC622 Based - 5.5 1 6 2.5 2 5.9 0.5 5 0.1 0.1 MC544 Based - 4 1 5 4 4 6 0.5 5 0.1 0.1 MC645 Based - 5 - 6 4 5 6 0.5 5 0.1 0.1 MC653 Based - 4 1 6 5 3 5.3 1 6.2 0.1 0.1

The density values of these alloys were measured and are reported in Table II. These values are in all cases less than 8.95, and for the most part less than 8.8. They therefore meet the objective set. Alloy Density (g.cm ^-3 ) T _{solvus γ '} (° C) MC2 8.62 1266 MC820 8.78 1300 MC533 8.64 1292 MC440 8.85 1304 MC722 8.82 1300 MC623 8.71 1294 MC622 8.68 1298 MC544 8.75 1292 MC645 8.75 1320 MC653 8.93 1308

In the raw state of monocrystalline solidification, these alloys show variable γ / γ 'eutectic fractions but the application of homogenization treatments such as those described above allow us to completely restore solution of the γ 'phase precipitates and eliminate the phases γ / γ 'eutectics without causing local alloy melting.

Solvent temperatures γ 'were measured by analysis thermal expansion on samples of alloys previously homogenized. The values of the γ 'solvus have been reported in Table II. The value of the γ 'solvus of the MC2 alloy measured under similar conditions is also reported for comparison in Table II. The Solvent temperatures γ 'of the alloys of the invention are always higher than that of the reference alloy MC2, the differences varying between 26 and 54 ° C depending on the alloys.

Creep tests in tension were carried out on test pieces machined in monocrystalline bars of orientation <001> of various alloys of the invention. The bars were previously homogenized and then returned according to the procedures described above. The values of the failure times for different creep conditions and for various alloys of the invention are compared in Table III with the values obtained under the same conditions on the reference monocrystalline alloy MC2. Alloy Creep conditions / Lifetime in hours T = 760 ° C σ = 840 MPa T = 950 ° C σ = 300 MPa T = 1050 ° C σ = 150 MPa T = 1100 ° C σ = 130 MPa T = 1150 ° C σ = 100 MPa MC2 369 198 485 156 5.6 MC820 386 205 439 168 105 MC533 561 298 401 151 52 MC440 154 162 198 102 52 MC722 118 274 248 87 109 MC623 455 222 289 126 62 MC622 175 232 257 129 117 MC544 162 458 486 199 151 MC645 2105 404 499 171 185 MC653 1153 456 726 216 194

All the alloys in the examples show a lifetime in creep at 1150 ° C much higher than that of the alloy of reference MC2. The relationship between lifetimes varies between 9 and 33 approximately. This result is consistent with the main target set. Gain in service life at this temperature is spectacular and is attributed, at least in part, to significant increase in the temperature of solvus γ ' in the alloys of the invention compared to the alloy of reference MC2.

For the other test conditions, the alloys of the invention show varying lifetimes which can be higher than those of the reference alloy MC2 according to the alloy and the temperature considered. Remarkable results are obtained in particular at 950 ° C and at 760 ° C in the case of certain alloys of the invention.

The best performing alloys are the MC544 alloys, MC645 and MC653. They show lifetimes in creep at less equal and generally greater than that of the alloy MC2 over the entire temperature range except alloy MC544 at 760 ° C. The most lifespan gains important are obtained at 950 and 1150 ° C.

Cyclic oxidation tests at 1100 ° C were carried out in the air on samples of superalloys of the invention homogenized and returned according to the procedures described before. Each test cycle includes maintaining one hour at 1100 ° C followed by cooling to temperature room. Cyclic oxidation behaviors of different alloys are illustrated in the graphs of Figures 1a and 1b where the variations in mass are reported specific (loss of mass per unit area) of samples as a function of the number of oxidation cycles of a hour. Tests were conducted under the same conditions on the reference alloy MC2. Oxidation resistance of a superalloy is all the better as its variation in specific mass is low. All the alloys of the invention thus show superior resistance to cyclic oxidation to that of the reference alloy MC2.

Cyclic corrosion tests were carried out at 850 ° C. on samples of alloys of the invention and of the reference alloy MC2. The samples were previously homogenized and returned according to the procedures described above. Each cycle includes a one hour hold at 850 ° C followed by cooling to room temperature. The samples are contaminated with Na ₂ SO ₄ (0.5 mg.cm ^-2 ) every 50 hours. The variations in the specific mass of the alloy samples are plotted as a function of the number of cycles in the graphs of FIGS. 2a and 2b. Corrosion behavior is considered satisfactory when the mass of the sample varies little: this is the incubation period. An accelerated corrosion stage occurs after the incubation stage. This accelerated corrosion most often results in rapid mass gain corresponding to the formation of corrosion products. The graphs show poor behavior for the reference alloy MC2 for which the accelerated corrosion stage occurs quickly. The alloys of the invention show incubation stages of variable durations, but in any case longer than that characterizing the reference alloy MC2, which demonstrates better resistance to cyclic corrosion.

The microstructures of the alloys of the invention have been checked after isothermal aging treatments 200 hours at 1050 ° C and at the end of the creep tests breaking at 760, 950, 1050, 1100 and 1150 ° C in order to control their microstructural stability vis-à-vis the unwanted intermetallic phase precipitation of the type σ, µ or Laves phase. Only the MC820 alloy shows rhenium-rich phase needle particles at the end of the 200 hour aging treatment at 1050 ° C as well at the end of the rupture creep tests at 1050 and 1100 ° C. These particles are located in the hearts of dendrites, where the rhenium preferentially separates during the directed solidification process. All other alloys of the invention cited in Table I are free of particles of undesirable phases rich in rhenium at term aging treatments and creep tests.

Claims (6)

1. Nickel-based superalloy suitable for the manufacture of turbo-engine parts by microcrystalline solidification,
   characterized in that its composition by mass is as follows: Cr 3.5 to 7.5 % Mo 0 to 1.5 % Re 1.5 to 5.5 % Ru 0 to 5.5 % W 3.5 to 8.5 % Al 5 to 6.5 % Ti 0 to 2.5 % Ta 4.5 to 9 % Hf 0.08 to 0.12 % Si 0.08 to 0.12 % the complement to 100 % consisting of nickel and possible impurities.
2. Superalloy according to claim 1, characterized in that its composition by mass is as follows: Cr 3.5 to 5.5 % Mo 0 to 1.5 % Re 4.5 to 5.5 % Ru 2.5 to 5.5 % W 4.5 to 6.5 % Al 5 to 6.5 % Ti 0 to 1.5 % Ta 5 to 6.2 % Hf 0.08 to 0.12 % Si 0.08 to 0.12 % the complement to 100 % consisting of nickel and possible impurities.
3. Superalloy according to claim 1, characterized in that its composition by mass is as follows: Cr 3.5 to 5.5 % Mo 0 to 1.5 % Re 3.5 to 4.5 % Ru 3.5 to 5.5 % W 4.5 to 6.5 % Al 5.5 to 6.5 % Ti 0 to 1 % Ta 4.5 to 5.5 % Hf 0.08 to 0.12 % Si 0.08 to 0.12 % the complement to 100 % consisting of nickel and possible impurities.
4. Superalloy according to claim 1, characterized in that its composition by mass is as follows: Cr 3.5 to 4.5 % Mo 0.5 to 1.5 % Re 3.5 to 4.5 % Ru 3.5 to 4.5 % W 4.5 to 5.5 % Al 5.5 to 6.5 % Ti 0 to 1 % Ta 4.5 to 5.5 % Hf 0.08 to 0.12 % Si 0.08 to 0.12 % the complement to 100 % consisting of nickel and possible impurities.
5. Superalloy according to claim 1, characterized in that its composition by mass is as follows: Cr 4.5 to 5.5 % Re 3.5 to 4.5 % Ru 4.5 to 5.5 % W 5.5 to 6.5 % Al 5.5 to 6.5 % Ti 0 to 1 % Ta 4.5 to 5.5 % Hf 0.08 to 0.12 % Si 0.08 to 0.12 % the complement to 100 % consisting of nickel and possible impurities.
6. Superalloy according to claim 1, characterized in that its composition by mass is as follows: Cr 3.5 to 4.5 % Mo 0.5 to 1.5 % Re 4.5 to 5.5 % Ru 2.5 to 3.5 % W 5.5 to 6.5 % Al 4.8 to 5.8 % Ti 0.5 to 1.5 % Ta 5.7 to 6.7 % Hf 0.08 to 0.12 % Si 0.08 to 0.12 % the complement to 100 % consisting of nickel and possible impurities.

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