CA2276154C - Nickel-based monocrystalline superalloy with a high .gamma.' solvus - Google Patents
Nickel-based monocrystalline superalloy with a high .gamma.' solvus Download PDFInfo
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Abstract
The superalloy according to the invention, suitable for the manufacture of turbo-engine parts by monocrystallization solidification, possesses the following mass composition: 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 Ni and possible impurities.
Description
Nickel-based monocrysta.lline superalloy with a high ti' solvus The present invention concerns nickel-based superalloys suitable in particular for manufacturing fixed and moving monocrystalline gas turbine blades, and exhibiting a high creep resistance at very high temperatures while retaining good resistance to the combustion gas environment. These alloys are particularly suited to applications in aeronautical engines used to propel airplanes and helicopters.
Nickel-based superalloys are materials with the highest performance used today for the manufacture of fixed and moving blades of the turbines of aeronautical gas turbine engines. The work of ONERA in this field started at the end of the 1970s and resulted, among other things, in 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 and FR 2 686 902.
The development of the performances of aeronautical gas turbines, translated in terms of specific power and yield and length of life, calls for the availability of alloys for turbine blades exhibiting high temperature mechanical properties (650 to 1150'C) and improved continuous resistance to corrosion and hot oxidation. Extreme operating conditions may in point of fact raise the metal to temperatures in excess of 1100'C. In order to optimize the resistance to hot corrosion and hot oxidation, monocrystalline blades made of superalloy are moreover generally coated with a protective deposit of the nickel aluminide or MCrAlY alloy type. In order to prevent any possible cracking or fracturing of these protective layers under the effect of thermal cycling, which would adversely effect the life of the parts, superalloys must however exhibit high intrinsic oxidation and corrosion resistance.
In polycrystalline blades cast by conventional foundry processes, a large part of the hot deformation during service is produced in the region of the grain boundaries which limits the life of the parts. The development of the monocrystalline solidification. process has, by eliminating grain boundaries, enabled the performance of nickel-based superalloys to be increased spectacularly. In addition, the process makes it possible to select the preferred growth orientation of the monocrystalline part and hence to select an optimum <001> orientation as regards creep resistance and thermal fatigue: resistance which are the two stress modes causing the greatest damage to turbine blades.
Successive improvements to the: mechanical performances, in particular creep, of these superalloys for monocrystalline blades have been made possible by optimizing their chemical compositions. Indeed, apart from nickel, which is the major constituent of these alloys, various addition elements provide specific contributions to the properties of the alloys. The functions of these: elements will be detailed hereinafter. In the monocryst.alline superalloys covered by the previously mentioned pater.its, the main addition elements (weight concentratior.Ls at a level of a few percent) have generally been selected from the following list: chromium (Cr), cobalt (Co), molybdenum (Mo), tungsten (W), aluminum (Al), titanium (Ti), tantalum (Ta) and niobium (Nb). The elements Cr, Co and Mo and part of the W
participate mainly in hardenir.Lg the austenitic (y phase) matrix where they enter into solution. The elements Al, Ti, Ta and Nb promote precipitation in the y matrix of hardening particles of a secor.Ld phase of the Ni3(Al, Ti, Ta, Nb) type (T' phase). Minor elements (weight concentrations below 0.5 %), such as silicon (Si) and hafnium (Hf), may also be added in order to opti.mize the environmental resistance as is demonstrated in FR 2 686 902.
Since the start of the 1980s, a large number of patents dedicated to novel superalloy compositions for monocrystalline blades have been filed worldwide. More recent developments have consisted in particular of incorporating the refractory elements rhenium (Re) and ruthenium (Ru) in these alloys. These additions are aimed especially at improving the high temperature creep resistance of these monocrystalline superalloys while preserving a stable microstructure at high temperatures as regards the formation of particles of intermetallic phases which are liable to bring about losses in the properties of these alloys.
Various patents thus protect the fields of monocrystalline superalloy compositions containing additions of one and/or 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 as regards the properties of these alloys is very limited and does not permit a judgement to be made as to the industrial value of these additions.
In France at the present time the monocrystalline superalloys used are referred to as "first-generation", as for example the grades AM1 and. MC2, both covered by patent FR 2 557 598, and the alloy AM3 protected by the patent FR 2 599 757. Among these, the alloy MC2 is considered as the alloy with the highest performance as regards creep resistance up to 1100 C. The future requirements of engineers will call for the ability to have alloys for blades available which have a higher performance than these first-generation alloys. It is necessary in particular to increase the maximum permissible temperatures for alloys constituting turbine blades.
The object of the invention is thus to provide a novel family of nickel-based monocrystalline superalloys exhibiting improved creep resistance compared with that of alloys exploited industrially, in particular at temperatures above 1100 C, but: equally at lower temperatures affecting various parts of the blades.
To this end, attempts have been made to introduce new addition elements without penailizing other properties essential for the good perforniance of these alloys, such as density, hot corrosion and oxidation resistance and microstructural stability.
An analysis of the state of the art as well as the results of work carried out by the inventor quickly showed that only alloys containing rhenium additions could make it possible to exceed the creep r.esistance of the alloy MC2 above 1100'C. In order to cour.Lter-balance certain harmful effects of rhenium (excessive density, microstructural instability), it seems moreover advantageous to incorporate ruthenium.
The invention concerns a nickel-based superalloy suitable for the manufacture of turbo-engine parts by monocrystalline solidification, wherein its mass composition 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 Ni and possible impurities.
More particularly, the invention provides such a superalloy having the following mass composition:
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 Ni and possible impurities.
Even more particularly, the mass composition of the superalloy 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 Ni and possible impurities.
Three specific compositions of superalloys according to the invention are given below:
Cr . 3.5 to 4.5 % 4.5 to 5.5 % 3.5 to 4.5 %
Mo . 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 %
Ta . 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 %
Si . 0.08 to 0.12 % 0.08 to 0.12 % 0.08 to 0.12 %
the complement to 100 % consisting of Ni and possible impurities.
The alloys according to the invention, prepared in the form of monocrystals with a <001> orientation, have the following properties:
- a density in all cases below 9 g.cm-3 and at best 8.8 g.cm-3, making it possible to minimize the mass of monocrystalline blades and consequently to limit the centrifugal stress acting on these blades and on the turbine disc to which they are attached;
- an ability for homogenization by returning all the ry' particles into solution, including the eutectic T/T' phases;
- a high temperature for putting the hardening ry' phase into solution, in all cases higher than that of previous alloys containing neither rhenium nor ruthenium;
- the absence of brittle intermetallic phases which can precipitate when high temperatures are maintained and which are liable to bring about a reduction in the creep resistance of the alloys and an. embrittlement of the alloys;
- resistance to hot cyclic corrosion and hot cyclic oxidation greater than that of previous alloys containing neither rhenium nor ruthenium.
Obtaining all these properties simultaneously enables the creep resistance of monocrystalline blades at very high temperatures and their environniental resistance to be optimized and therefore enables their life to be improved as well as the performance of gas turbines.
The invention thus provides a unique combination of properties for the alloys, which the state of the art was not able to provide.
The alloys of the invention are! intended for the manufacture of monocrystalline parts, i.e consisting of a single metallurgical grain. Thi.s special structure is obtained with the aid of a process for directed solidification in a thermal gradient using a grain selection device or a monocryst.alline seed at the start of solidification.
After solidification, the superalloys consist essentially of two phases: the austenitic y matrix is a nickel-based solid solution in which the particles of the y' phase, an intermetallic compound of which. the composition is based on Ni3Al, precipitates during cooling in the solid state. The addition elements are distributed in the two y and y' phases but generally exhibit a particular affinity for one or other of these two phases. Thus chromium, molybdenum, rhenium and ruthenium are preferably distributed in the y matrix, while aluminum, titanium and tantalum go preferentially into the y' phase.
F!
In the crude monocrystalline solidification alloys, the distribution of the particles of the hardening y' phase is very heterogeneous in the volume of the monocrystal on account of chemical segregations resulting from solidification conditions specific to the process. The microstructure is said to be d.endritic. The precipitates are very fine in the core of the dendrites which solidify first of all during cooling of the alloy and become larger in the regions which then solidify from the center of the dendrite. Moreover, at the end. of solidification, eutectic phases consisting of massive particles of the y' phase containing laminae of the y phase solidify in the regions separating the dendrites.
Experiments however showed that the creep resistance of nickel-based superalloys was optimized when the distribution of the particles of the T' phase were homogeneous in all the volume of the alloy with precipitate sizes less than 1 micrometre, the optimum size of the precipitates depending on the composition of the alloy.
The T' phase contained in the eutectic phases did not contribute, in particular, to the hardening of alloys, and the potential of the alloys for creep resistance is therefore not totally exploited in the crude solidification state. These massive blocks of the y/y' eutectic phase are moreover preferred sites for the start of cracks during cyclic stresses resulting from. thermal fatigue phenomena due to the starting and stopping cycles of gas turbines.
The compositions of the alloys of the invention have been selected so as to be able to obtain biphase -y/T' structures, consisting of a hcmogeneous precipitation of y' particles in aT matrix resulting from the monocrystalline solidification steps and the heat treatments detailed below. In order to achieve th.is optimized microstructure, it is first of all necessary to apply a heat treatment designed to dissolve the -y' phase precipitates contained in the dendrites and to eliminate the solidified eutectic phase between the dendrites. Solution of the y' precipitates is achieved when the heat treatment temperature reaches the temperature of the y' solvus (the temperature at which the y' phase precipitates are put into solution) which is a characteristic of the chemical composition of the alloy. In practice, the value of the y' solvus varies periodically in the crude monocrystalline solidification alloy in relation to the local chemistry of the alloy. Accordingly, the 1' solvus increases within the core of the dendrite towards the inter-dendritic regions on account of chemical segregations, until the initial fusion temperature of the eutectic y' phase is reached, this phase being the last solid formed during cooling of the alloy from the liquid state. This initial fusion temperature is in practice similar to the solidus temperature (initial fusion temperature) of the alloy. The homogenization treatment temperature must therefore remain below the solidus temperature.
In practice it has been possible to obtain complete solution of the y' precipitates and the eutectics y/y' in the alloys of the invention by virtue of the application of sequences of heat treatments including prior homogenization of the dendritic structures. This sequence of heat treatments comprises a first pre-homogenization treatment of 3 hours at a temperature of' between 1300 and 1310 C, then a progressive increase of' 3 0 C at a rate of 3 C. h-1, before a new step of 3 hours at a temperature of between 1330 and 1340 C, it being necessary to carry out final cooling at a rate such that tY:Le final size of the y' phase precipitates is less than 300 nm. The entirety of the eutectic y/y' phases is thus eliminated. It was possible to obtain this result for all the alloys of the invention.
The sequence of heat treatments which has just been described is an example enabling the anticipated result to be obtained. This does not exclude the possibility of obtaining a similar result by using another sequence of heat treatments, the result of the treatment being more important than the manner in which it is arrived at. The important thing is to have demonstrated the possibility of 5 obtaining such a result in the case of the alloys of the invention.
The alloys of the invention were tested after they had been subjected to a sequence of homogenization treatments and 10 treatments for putting the T' phase into solution as described above, and then to two annealing heat treatments enabling the size and volume fraction of the y' phase precipitates to be fixed. A first annealing consists of a treatment of 4 to 16 hours at a temperature of between 1050 and 1150 C enabling the size of the y' phase precipitates to be fixed at between 300 and. 500 nm. A second annealing treatment consists of a treatment of 15 to 25 hours at a temperature of between 850 and. 870'C enabling the fraction of the precipitated T' phase to be optimized. These annealing treatments are compatible with the diffusion treatments of protective coatings and brazing treatments generally applied to monocrystalline turbine blades during their manufacture. Micrographic examination shows that the ry' phase precipitates have a roughly cubic shape and represent a volume fraction of at least 70 % in the alloy.
They are contained in the y matrix which appears in the form of fine channels between these precipitates.
The high temperature resistance to creep is greater the higher the volume fraction of the hardening ry' phase precipitated in the alloy. At ambient temperature, the alloys of the invention contain a volume fraction close to 70 %. When the temperature increases from ambient temperature, the y' phase dissolves progressively in the y matrix, slowly up to about 1000'C and then more rapidly above 1000 C. When the T' solvus temperature is exceeded the ry' precipitates are then totally dissolved. The reduction in the volume fraction of the ti' phase when the temperature increases is one of the causes of the reduction of the creep resistance of superalloys.
One of the major benefits of the invention is to increase substantially the y' solvus temperature so as to preserve a high volume fraction of the y' phase at temperatures above 1100 C and thus to obtain a very high creep resistance at these temperatures. The invention therefore concerns so-called "high -y' solvus" alloys exhibiting a very high creep resistance above 1100'C. Experience acquired by the inventor in the field has shown that increases in the concentrations of Al, Ti, Ta, Mo and W bring about an increase in the ry' solvus. On the other hand, additions of the elements Cr and Co lead to a reduction in the temperature of the ry' solvus. As regards rhenium and ruthenium, previous work has not explicitly drawn any conclusions as to their specific action on the ry' solvus temperature.
Increases in the concentration of elements increasing the T' solvus may however lead to effects which can harm the properties of the alloys. Thus too high a concentration of the elements Al, Ti and Ta leads to the formation of an excessive quantity of y/y' eutectic phases during solidification of these alloys. These phases can then no longer be totally eliminated by subsequent heat treatments, which harms the homogeneity of the alloy and consequently its creep resistance. Moreover, the concentration of Ta should be limited since this element has a high atomic mass and penalizes the alloys from the point of view of density.
The elements Mo and W also have a beneficial effect on the y' solvus but these elements are heavy, in particular W, and their concentration must be controlled so as not to increase the density of the alloys excessively.
1.2 In addition, the solubility oi: these elements in the y matrix is limited, in the same way as that of rhenium and to a lesser extent those of cobalt and chromium, which may lead to the precipitation of brittle intermetallic phases of the a, , P or Laves phase type. The presence of these phases, referred to as topoloclically close-packed (T.C.P.) may bring about a loss of inechanical properties in superalloys where they precipitate. Obtaining alloys that are not liable to form these brittle intermetallic phases is one of the main arguments of prior patents on monocrystalline superalloys.
A reduction in the concentrations of the elements Cr and Co brings about a reduction in the temperature of the y' solvus. Thus, one of the maiii ideas of the invention is to refrain from any addition of Co whose effect on the creep resistance of the superalloys is small compared with that of the other addition elements. On the other hand, chromium has been kept since its presence is indispensable for maintaining good hot corrosion resistance.
Examples of the invention detailed below show that the objective of obtaining high-solvus alloys has been achieved by virtue of a judicious choice of chemical compositions taking into account the consicierations which have just been expounded.
Apart from the optimization of the volume fraction and the solvus temperature of the y' phase, an improvement to the creep resistance of monocrystalline superalloys can be obtained by increasing the concentrations of the refractory elements Mo, W, Re and Ta which play an important role in the solid solution hardening of the -y and y' phases. These heavy elements moreover retarci all the elementary mechanisms that are controlled by the diffusion of atoms, which has beneficial consequences on the creep resistance 1.3 of the alloys. The addition of rhenium in particular limits the growth of y' phase particles during periods where high temperatures are maintained, a phenomenon which participates in the degradation of the mechanical properties of superalloys with time. Moreover, an increase in the concentrations of refractory elements retards the heat-activated movement of dislocations which propagate deformation in superalloys, which has the effect of reducing the creep rate.
The concentrations of refractory elements should however be carefully balanced so as not t:o increase the density of the alloys excessively.
When the elements W and Mo are present in too high a concentration, they have a harmful effect on the oxidation and corrosion resistance of monocrystalline superalloys while the presence of rhenium does not penalize the environmental resistance of these alloys.
Moreover, the refractory element Ru, within the context of the invention, is of value in that it has a density half that of rhenium. Work carried out by the inventor in this field shows that Ru promotes t;he precipitation of brittle intermetallic phases to a lesser extent than does rhenium.
The alloys according to the irivention also include simultaneous additions of silicon and hafnium. Such additions make it possible to optimize the hot oxidation resistance of the alloys by inlproving the adhesion of the protective alumina layer formed at high temperatures.
Alloys according to the invent:ion were prepared and solidified in the form of monocrystals with a <001>
crystallographic orientation and were tested. This crystallographic orientation was that usually selected for the directed solidification of: monocrystalline turbine blades. It confers on these parts an optimum combination of creep resistance, heat fatigue resistance and mechanical fatigue resistance.
By way of example, the nominal chemical compositions (% by weight) of a few alloys of the: invention are collected in table I, together with that of the reference alloy MC2 described in FR 2 557 598. This alloy serves as a reference since it is, to the inventor's knowledge, the one with the highest creep performance among the alloys containing neither rhenium nor ruthenium.
Table I
Alloy Ni Co Cr Mo W Re Ru Al TTi Ta Si Hf MC2 Base 5 8 2 8 - - 5 1.5 6 - -MC820 Base - 5 1 F 2 - 5.5 1 6 0.1 0.1 - - -T-MC533 Base - 7 - L5F3 3 6 - 6 0.1 0.1 MC440 Base - 5 1 4 4 - 5.5 - 9 0.1 0.1 MC722 Base - 4.5 1 7 2.5 2.5 5.8 - 6 0.1 0.1 MC623 Base - 6 1 F 2 3 5.7 0.5 5.5 0.1 0.1 MC622 Base - 5.5 1 761 2.5 2 5.9 0.5 5 0.1 0.1 MC544 Base - 4 1 5 4 4 6 0.5 5 0.1 0.1 MC645 Base - 5 - 6 4 5 6 0.5 5 0.1 0.1 MC653 Base - 4 1 6 F5-3 5.3 1 6.2 0.1 0.1 The values of the densities of these alloys were measured and are given in Table II. These values were in all cases less than 8.95, and for the most part were less than 8.8.
They thus satisfied the set objective.
Table II
Alloy Density Ty.SOIVõS
(g. crri ) ( C) MC2 8.62 1266 5 MC820 8.78 1300 MC533 8.64 1292 MC440 8.85 1304 MC722 8.82 1300 MC623 8.71 1294 10 MC622 8.68 1298 MC544 8.75 1292 MC645 8.75 1320 MC653 8.93 1308 15 In the raw monocrystallization state, these alloys exhibit variable -y/-y' eutectic fractions but the application of homogenization treatments such as previously described makes it possible to put the T' phase precipitates back into solution completely and to eliminate the T/T' eutectic phases without bringing about local melting of the alloys.
The T' solvus temperatures were measured by dilatometric thermal analysis on specimens of previously homogenized alloys. The values of the y' solvus have been reported in Table II. The value of the y' of solvus of the alloy MC2 measured under similar conditions is also given for comparison in table II. The ry' solvus temperatures of the alloys of the invention were always greater than that of the reference alloy MC2, differences varying between 26 and 54 C according to the alloys.
Tensile creep tests were carried out on specimens machined from various alloys of the invention into monocrystalline bars with a <001> orientation. The bars had been previously homogenized and then annealed according to the procedures described previousl.y. The values of the times to rupture 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 monocrystalline reference alloy MC2.
Table III
Alloy Creep conditions/life in hours T=760 C T=950 C T=1050 C T=1100 C T=1150 C
a=840 MPa Q=300 MPa Q=150 MPa Q=130 MPa v=100 MPa MC2 369 198 485 156 5.6 All the alloys of the examples showed a life in creep at 1150'C very much greater than that of the reference alloy MC2. The ratio between the lives varied between approximately 9 and 33. This result was in agreement with the main set objective. The clain in life at this temperature was spectacular anLd is attributed, at least in part, to the significant increase in the ry' solvus temperature in the alloys of the invention compared with the reference alloy MC2.
For the other test conditions, the alloys of the invention showed variable lives which could be greater than those of the reference alloy MC2 according to the alloy and the temperature considered. Remarkable results were obtained in particular at 950 C and at 760 C in the case of certain alloys of the invention.
The highest performance alloys were the alloys MC544, MC645 and MC653. They exhibited lives in creep at least equal to, and generally greater than, that of the alloy MC2 within all the temperature interval considered except the alloy MC544 at 760 C. The larcfest gains in life were obtained at 950 and 1150 C.
Cyclic oxidation tests at 1100 C were carried out in air on specimens of superalloys of the invention homogenized and annealed according to the proce:dures previously described.
Each test cycle comprised a constant temperature at 1100 C
followed bv cooling to ambient temperature. The behaviour in cyclic oxidation of the various alloys are illustrated in the graphs of figures la and. lb where variations in the density are given (loss in mass per unit area) of the samples as a function of the number of one hour oxidation cycles. Tests were carried out under the same conditions on the reference alloy MC2. The oxidation resistance of a superalloy improved as its density variation became lower.
All the alloys of the invention thus showed a cyclic oxidation resistance greater than that of reference alloy MC2.
Cyclic corrosion tests were carried out at 850 C on specimens of alloys of the invention and the reference alloy MC2. The specimens had been previously homogenized and annealed according to the procedures previously described. Each cycle comprised a constant temperature for one hour at 850 C followed by cooling to ambient temperature. The samples were contaminated with Na2SOd (0.5 mg.cm'Z) every 50 hours. 'Variations in the density of the alloy specimens are given as a function of the number of cycles in the graphs of figures 2a and 2b. The behaviour in corrosion was considered as satisfactory when the mass of the specimen hardly varied at all. This was the incubation period. An accelerated corrosion stage took place at the end of the incubation stage. This accelerated corrosion resulted more often in a rapid increase in mass corresponding to the formation of corrosion products. The graphs show a mediocre behaviou.r for the reference alloy MC2 for which the accelerated corrosion stage took place rapidly. The alloys of the invention showed incubation stages with variable durations, but in all cases were longer than that characterizing the reference alloy MC2, which demonstrated a better resistance to cyclic corrosion.
The microstructures of the alloys of the invention were checked at the end of the isothermal ageing treatments of 200 hours at 1050 C and at the end of creep tests taken to rupture at 760, 950, 1050, 1100 and 1150 C so as to check their microstructural stability as regards precipitation of undesirable intermetallic phases of the a, or Laves phase type. Only the alloy MC820 showed needle-like particles of a phase rich in rhenium at the end of the ageing treatment of 200 hours at 1050 C as well as at the end of creep tests to rupture at 1050 and 1100 C. These particles were localized in the cores of the dendrites, where rhenium separated preferentially during the process of directed solidification. All the other alloys of the invention cited in table I were free from particles of undesirable phases rich in rhenium at the e:nd of the ageing treatments and creep tests.
Nickel-based superalloys are materials with the highest performance used today for the manufacture of fixed and moving blades of the turbines of aeronautical gas turbine engines. The work of ONERA in this field started at the end of the 1970s and resulted, among other things, in 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 and FR 2 686 902.
The development of the performances of aeronautical gas turbines, translated in terms of specific power and yield and length of life, calls for the availability of alloys for turbine blades exhibiting high temperature mechanical properties (650 to 1150'C) and improved continuous resistance to corrosion and hot oxidation. Extreme operating conditions may in point of fact raise the metal to temperatures in excess of 1100'C. In order to optimize the resistance to hot corrosion and hot oxidation, monocrystalline blades made of superalloy are moreover generally coated with a protective deposit of the nickel aluminide or MCrAlY alloy type. In order to prevent any possible cracking or fracturing of these protective layers under the effect of thermal cycling, which would adversely effect the life of the parts, superalloys must however exhibit high intrinsic oxidation and corrosion resistance.
In polycrystalline blades cast by conventional foundry processes, a large part of the hot deformation during service is produced in the region of the grain boundaries which limits the life of the parts. The development of the monocrystalline solidification. process has, by eliminating grain boundaries, enabled the performance of nickel-based superalloys to be increased spectacularly. In addition, the process makes it possible to select the preferred growth orientation of the monocrystalline part and hence to select an optimum <001> orientation as regards creep resistance and thermal fatigue: resistance which are the two stress modes causing the greatest damage to turbine blades.
Successive improvements to the: mechanical performances, in particular creep, of these superalloys for monocrystalline blades have been made possible by optimizing their chemical compositions. Indeed, apart from nickel, which is the major constituent of these alloys, various addition elements provide specific contributions to the properties of the alloys. The functions of these: elements will be detailed hereinafter. In the monocryst.alline superalloys covered by the previously mentioned pater.its, the main addition elements (weight concentratior.Ls at a level of a few percent) have generally been selected from the following list: chromium (Cr), cobalt (Co), molybdenum (Mo), tungsten (W), aluminum (Al), titanium (Ti), tantalum (Ta) and niobium (Nb). The elements Cr, Co and Mo and part of the W
participate mainly in hardenir.Lg the austenitic (y phase) matrix where they enter into solution. The elements Al, Ti, Ta and Nb promote precipitation in the y matrix of hardening particles of a secor.Ld phase of the Ni3(Al, Ti, Ta, Nb) type (T' phase). Minor elements (weight concentrations below 0.5 %), such as silicon (Si) and hafnium (Hf), may also be added in order to opti.mize the environmental resistance as is demonstrated in FR 2 686 902.
Since the start of the 1980s, a large number of patents dedicated to novel superalloy compositions for monocrystalline blades have been filed worldwide. More recent developments have consisted in particular of incorporating the refractory elements rhenium (Re) and ruthenium (Ru) in these alloys. These additions are aimed especially at improving the high temperature creep resistance of these monocrystalline superalloys while preserving a stable microstructure at high temperatures as regards the formation of particles of intermetallic phases which are liable to bring about losses in the properties of these alloys.
Various patents thus protect the fields of monocrystalline superalloy compositions containing additions of one and/or 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 as regards the properties of these alloys is very limited and does not permit a judgement to be made as to the industrial value of these additions.
In France at the present time the monocrystalline superalloys used are referred to as "first-generation", as for example the grades AM1 and. MC2, both covered by patent FR 2 557 598, and the alloy AM3 protected by the patent FR 2 599 757. Among these, the alloy MC2 is considered as the alloy with the highest performance as regards creep resistance up to 1100 C. The future requirements of engineers will call for the ability to have alloys for blades available which have a higher performance than these first-generation alloys. It is necessary in particular to increase the maximum permissible temperatures for alloys constituting turbine blades.
The object of the invention is thus to provide a novel family of nickel-based monocrystalline superalloys exhibiting improved creep resistance compared with that of alloys exploited industrially, in particular at temperatures above 1100 C, but: equally at lower temperatures affecting various parts of the blades.
To this end, attempts have been made to introduce new addition elements without penailizing other properties essential for the good perforniance of these alloys, such as density, hot corrosion and oxidation resistance and microstructural stability.
An analysis of the state of the art as well as the results of work carried out by the inventor quickly showed that only alloys containing rhenium additions could make it possible to exceed the creep r.esistance of the alloy MC2 above 1100'C. In order to cour.Lter-balance certain harmful effects of rhenium (excessive density, microstructural instability), it seems moreover advantageous to incorporate ruthenium.
The invention concerns a nickel-based superalloy suitable for the manufacture of turbo-engine parts by monocrystalline solidification, wherein its mass composition 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 Ni and possible impurities.
More particularly, the invention provides such a superalloy having the following mass composition:
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 Ni and possible impurities.
Even more particularly, the mass composition of the superalloy 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 Ni and possible impurities.
Three specific compositions of superalloys according to the invention are given below:
Cr . 3.5 to 4.5 % 4.5 to 5.5 % 3.5 to 4.5 %
Mo . 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 %
Ta . 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 %
Si . 0.08 to 0.12 % 0.08 to 0.12 % 0.08 to 0.12 %
the complement to 100 % consisting of Ni and possible impurities.
The alloys according to the invention, prepared in the form of monocrystals with a <001> orientation, have the following properties:
- a density in all cases below 9 g.cm-3 and at best 8.8 g.cm-3, making it possible to minimize the mass of monocrystalline blades and consequently to limit the centrifugal stress acting on these blades and on the turbine disc to which they are attached;
- an ability for homogenization by returning all the ry' particles into solution, including the eutectic T/T' phases;
- a high temperature for putting the hardening ry' phase into solution, in all cases higher than that of previous alloys containing neither rhenium nor ruthenium;
- the absence of brittle intermetallic phases which can precipitate when high temperatures are maintained and which are liable to bring about a reduction in the creep resistance of the alloys and an. embrittlement of the alloys;
- resistance to hot cyclic corrosion and hot cyclic oxidation greater than that of previous alloys containing neither rhenium nor ruthenium.
Obtaining all these properties simultaneously enables the creep resistance of monocrystalline blades at very high temperatures and their environniental resistance to be optimized and therefore enables their life to be improved as well as the performance of gas turbines.
The invention thus provides a unique combination of properties for the alloys, which the state of the art was not able to provide.
The alloys of the invention are! intended for the manufacture of monocrystalline parts, i.e consisting of a single metallurgical grain. Thi.s special structure is obtained with the aid of a process for directed solidification in a thermal gradient using a grain selection device or a monocryst.alline seed at the start of solidification.
After solidification, the superalloys consist essentially of two phases: the austenitic y matrix is a nickel-based solid solution in which the particles of the y' phase, an intermetallic compound of which. the composition is based on Ni3Al, precipitates during cooling in the solid state. The addition elements are distributed in the two y and y' phases but generally exhibit a particular affinity for one or other of these two phases. Thus chromium, molybdenum, rhenium and ruthenium are preferably distributed in the y matrix, while aluminum, titanium and tantalum go preferentially into the y' phase.
F!
In the crude monocrystalline solidification alloys, the distribution of the particles of the hardening y' phase is very heterogeneous in the volume of the monocrystal on account of chemical segregations resulting from solidification conditions specific to the process. The microstructure is said to be d.endritic. The precipitates are very fine in the core of the dendrites which solidify first of all during cooling of the alloy and become larger in the regions which then solidify from the center of the dendrite. Moreover, at the end. of solidification, eutectic phases consisting of massive particles of the y' phase containing laminae of the y phase solidify in the regions separating the dendrites.
Experiments however showed that the creep resistance of nickel-based superalloys was optimized when the distribution of the particles of the T' phase were homogeneous in all the volume of the alloy with precipitate sizes less than 1 micrometre, the optimum size of the precipitates depending on the composition of the alloy.
The T' phase contained in the eutectic phases did not contribute, in particular, to the hardening of alloys, and the potential of the alloys for creep resistance is therefore not totally exploited in the crude solidification state. These massive blocks of the y/y' eutectic phase are moreover preferred sites for the start of cracks during cyclic stresses resulting from. thermal fatigue phenomena due to the starting and stopping cycles of gas turbines.
The compositions of the alloys of the invention have been selected so as to be able to obtain biphase -y/T' structures, consisting of a hcmogeneous precipitation of y' particles in aT matrix resulting from the monocrystalline solidification steps and the heat treatments detailed below. In order to achieve th.is optimized microstructure, it is first of all necessary to apply a heat treatment designed to dissolve the -y' phase precipitates contained in the dendrites and to eliminate the solidified eutectic phase between the dendrites. Solution of the y' precipitates is achieved when the heat treatment temperature reaches the temperature of the y' solvus (the temperature at which the y' phase precipitates are put into solution) which is a characteristic of the chemical composition of the alloy. In practice, the value of the y' solvus varies periodically in the crude monocrystalline solidification alloy in relation to the local chemistry of the alloy. Accordingly, the 1' solvus increases within the core of the dendrite towards the inter-dendritic regions on account of chemical segregations, until the initial fusion temperature of the eutectic y' phase is reached, this phase being the last solid formed during cooling of the alloy from the liquid state. This initial fusion temperature is in practice similar to the solidus temperature (initial fusion temperature) of the alloy. The homogenization treatment temperature must therefore remain below the solidus temperature.
In practice it has been possible to obtain complete solution of the y' precipitates and the eutectics y/y' in the alloys of the invention by virtue of the application of sequences of heat treatments including prior homogenization of the dendritic structures. This sequence of heat treatments comprises a first pre-homogenization treatment of 3 hours at a temperature of' between 1300 and 1310 C, then a progressive increase of' 3 0 C at a rate of 3 C. h-1, before a new step of 3 hours at a temperature of between 1330 and 1340 C, it being necessary to carry out final cooling at a rate such that tY:Le final size of the y' phase precipitates is less than 300 nm. The entirety of the eutectic y/y' phases is thus eliminated. It was possible to obtain this result for all the alloys of the invention.
The sequence of heat treatments which has just been described is an example enabling the anticipated result to be obtained. This does not exclude the possibility of obtaining a similar result by using another sequence of heat treatments, the result of the treatment being more important than the manner in which it is arrived at. The important thing is to have demonstrated the possibility of 5 obtaining such a result in the case of the alloys of the invention.
The alloys of the invention were tested after they had been subjected to a sequence of homogenization treatments and 10 treatments for putting the T' phase into solution as described above, and then to two annealing heat treatments enabling the size and volume fraction of the y' phase precipitates to be fixed. A first annealing consists of a treatment of 4 to 16 hours at a temperature of between 1050 and 1150 C enabling the size of the y' phase precipitates to be fixed at between 300 and. 500 nm. A second annealing treatment consists of a treatment of 15 to 25 hours at a temperature of between 850 and. 870'C enabling the fraction of the precipitated T' phase to be optimized. These annealing treatments are compatible with the diffusion treatments of protective coatings and brazing treatments generally applied to monocrystalline turbine blades during their manufacture. Micrographic examination shows that the ry' phase precipitates have a roughly cubic shape and represent a volume fraction of at least 70 % in the alloy.
They are contained in the y matrix which appears in the form of fine channels between these precipitates.
The high temperature resistance to creep is greater the higher the volume fraction of the hardening ry' phase precipitated in the alloy. At ambient temperature, the alloys of the invention contain a volume fraction close to 70 %. When the temperature increases from ambient temperature, the y' phase dissolves progressively in the y matrix, slowly up to about 1000'C and then more rapidly above 1000 C. When the T' solvus temperature is exceeded the ry' precipitates are then totally dissolved. The reduction in the volume fraction of the ti' phase when the temperature increases is one of the causes of the reduction of the creep resistance of superalloys.
One of the major benefits of the invention is to increase substantially the y' solvus temperature so as to preserve a high volume fraction of the y' phase at temperatures above 1100 C and thus to obtain a very high creep resistance at these temperatures. The invention therefore concerns so-called "high -y' solvus" alloys exhibiting a very high creep resistance above 1100'C. Experience acquired by the inventor in the field has shown that increases in the concentrations of Al, Ti, Ta, Mo and W bring about an increase in the ry' solvus. On the other hand, additions of the elements Cr and Co lead to a reduction in the temperature of the ry' solvus. As regards rhenium and ruthenium, previous work has not explicitly drawn any conclusions as to their specific action on the ry' solvus temperature.
Increases in the concentration of elements increasing the T' solvus may however lead to effects which can harm the properties of the alloys. Thus too high a concentration of the elements Al, Ti and Ta leads to the formation of an excessive quantity of y/y' eutectic phases during solidification of these alloys. These phases can then no longer be totally eliminated by subsequent heat treatments, which harms the homogeneity of the alloy and consequently its creep resistance. Moreover, the concentration of Ta should be limited since this element has a high atomic mass and penalizes the alloys from the point of view of density.
The elements Mo and W also have a beneficial effect on the y' solvus but these elements are heavy, in particular W, and their concentration must be controlled so as not to increase the density of the alloys excessively.
1.2 In addition, the solubility oi: these elements in the y matrix is limited, in the same way as that of rhenium and to a lesser extent those of cobalt and chromium, which may lead to the precipitation of brittle intermetallic phases of the a, , P or Laves phase type. The presence of these phases, referred to as topoloclically close-packed (T.C.P.) may bring about a loss of inechanical properties in superalloys where they precipitate. Obtaining alloys that are not liable to form these brittle intermetallic phases is one of the main arguments of prior patents on monocrystalline superalloys.
A reduction in the concentrations of the elements Cr and Co brings about a reduction in the temperature of the y' solvus. Thus, one of the maiii ideas of the invention is to refrain from any addition of Co whose effect on the creep resistance of the superalloys is small compared with that of the other addition elements. On the other hand, chromium has been kept since its presence is indispensable for maintaining good hot corrosion resistance.
Examples of the invention detailed below show that the objective of obtaining high-solvus alloys has been achieved by virtue of a judicious choice of chemical compositions taking into account the consicierations which have just been expounded.
Apart from the optimization of the volume fraction and the solvus temperature of the y' phase, an improvement to the creep resistance of monocrystalline superalloys can be obtained by increasing the concentrations of the refractory elements Mo, W, Re and Ta which play an important role in the solid solution hardening of the -y and y' phases. These heavy elements moreover retarci all the elementary mechanisms that are controlled by the diffusion of atoms, which has beneficial consequences on the creep resistance 1.3 of the alloys. The addition of rhenium in particular limits the growth of y' phase particles during periods where high temperatures are maintained, a phenomenon which participates in the degradation of the mechanical properties of superalloys with time. Moreover, an increase in the concentrations of refractory elements retards the heat-activated movement of dislocations which propagate deformation in superalloys, which has the effect of reducing the creep rate.
The concentrations of refractory elements should however be carefully balanced so as not t:o increase the density of the alloys excessively.
When the elements W and Mo are present in too high a concentration, they have a harmful effect on the oxidation and corrosion resistance of monocrystalline superalloys while the presence of rhenium does not penalize the environmental resistance of these alloys.
Moreover, the refractory element Ru, within the context of the invention, is of value in that it has a density half that of rhenium. Work carried out by the inventor in this field shows that Ru promotes t;he precipitation of brittle intermetallic phases to a lesser extent than does rhenium.
The alloys according to the irivention also include simultaneous additions of silicon and hafnium. Such additions make it possible to optimize the hot oxidation resistance of the alloys by inlproving the adhesion of the protective alumina layer formed at high temperatures.
Alloys according to the invent:ion were prepared and solidified in the form of monocrystals with a <001>
crystallographic orientation and were tested. This crystallographic orientation was that usually selected for the directed solidification of: monocrystalline turbine blades. It confers on these parts an optimum combination of creep resistance, heat fatigue resistance and mechanical fatigue resistance.
By way of example, the nominal chemical compositions (% by weight) of a few alloys of the: invention are collected in table I, together with that of the reference alloy MC2 described in FR 2 557 598. This alloy serves as a reference since it is, to the inventor's knowledge, the one with the highest creep performance among the alloys containing neither rhenium nor ruthenium.
Table I
Alloy Ni Co Cr Mo W Re Ru Al TTi Ta Si Hf MC2 Base 5 8 2 8 - - 5 1.5 6 - -MC820 Base - 5 1 F 2 - 5.5 1 6 0.1 0.1 - - -T-MC533 Base - 7 - L5F3 3 6 - 6 0.1 0.1 MC440 Base - 5 1 4 4 - 5.5 - 9 0.1 0.1 MC722 Base - 4.5 1 7 2.5 2.5 5.8 - 6 0.1 0.1 MC623 Base - 6 1 F 2 3 5.7 0.5 5.5 0.1 0.1 MC622 Base - 5.5 1 761 2.5 2 5.9 0.5 5 0.1 0.1 MC544 Base - 4 1 5 4 4 6 0.5 5 0.1 0.1 MC645 Base - 5 - 6 4 5 6 0.5 5 0.1 0.1 MC653 Base - 4 1 6 F5-3 5.3 1 6.2 0.1 0.1 The values of the densities of these alloys were measured and are given in Table II. These values were in all cases less than 8.95, and for the most part were less than 8.8.
They thus satisfied the set objective.
Table II
Alloy Density Ty.SOIVõS
(g. crri ) ( C) MC2 8.62 1266 5 MC820 8.78 1300 MC533 8.64 1292 MC440 8.85 1304 MC722 8.82 1300 MC623 8.71 1294 10 MC622 8.68 1298 MC544 8.75 1292 MC645 8.75 1320 MC653 8.93 1308 15 In the raw monocrystallization state, these alloys exhibit variable -y/-y' eutectic fractions but the application of homogenization treatments such as previously described makes it possible to put the T' phase precipitates back into solution completely and to eliminate the T/T' eutectic phases without bringing about local melting of the alloys.
The T' solvus temperatures were measured by dilatometric thermal analysis on specimens of previously homogenized alloys. The values of the y' solvus have been reported in Table II. The value of the y' of solvus of the alloy MC2 measured under similar conditions is also given for comparison in table II. The ry' solvus temperatures of the alloys of the invention were always greater than that of the reference alloy MC2, differences varying between 26 and 54 C according to the alloys.
Tensile creep tests were carried out on specimens machined from various alloys of the invention into monocrystalline bars with a <001> orientation. The bars had been previously homogenized and then annealed according to the procedures described previousl.y. The values of the times to rupture 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 monocrystalline reference alloy MC2.
Table III
Alloy Creep conditions/life in hours T=760 C T=950 C T=1050 C T=1100 C T=1150 C
a=840 MPa Q=300 MPa Q=150 MPa Q=130 MPa v=100 MPa MC2 369 198 485 156 5.6 All the alloys of the examples showed a life in creep at 1150'C very much greater than that of the reference alloy MC2. The ratio between the lives varied between approximately 9 and 33. This result was in agreement with the main set objective. The clain in life at this temperature was spectacular anLd is attributed, at least in part, to the significant increase in the ry' solvus temperature in the alloys of the invention compared with the reference alloy MC2.
For the other test conditions, the alloys of the invention showed variable lives which could be greater than those of the reference alloy MC2 according to the alloy and the temperature considered. Remarkable results were obtained in particular at 950 C and at 760 C in the case of certain alloys of the invention.
The highest performance alloys were the alloys MC544, MC645 and MC653. They exhibited lives in creep at least equal to, and generally greater than, that of the alloy MC2 within all the temperature interval considered except the alloy MC544 at 760 C. The larcfest gains in life were obtained at 950 and 1150 C.
Cyclic oxidation tests at 1100 C were carried out in air on specimens of superalloys of the invention homogenized and annealed according to the proce:dures previously described.
Each test cycle comprised a constant temperature at 1100 C
followed bv cooling to ambient temperature. The behaviour in cyclic oxidation of the various alloys are illustrated in the graphs of figures la and. lb where variations in the density are given (loss in mass per unit area) of the samples as a function of the number of one hour oxidation cycles. Tests were carried out under the same conditions on the reference alloy MC2. The oxidation resistance of a superalloy improved as its density variation became lower.
All the alloys of the invention thus showed a cyclic oxidation resistance greater than that of reference alloy MC2.
Cyclic corrosion tests were carried out at 850 C on specimens of alloys of the invention and the reference alloy MC2. The specimens had been previously homogenized and annealed according to the procedures previously described. Each cycle comprised a constant temperature for one hour at 850 C followed by cooling to ambient temperature. The samples were contaminated with Na2SOd (0.5 mg.cm'Z) every 50 hours. 'Variations in the density of the alloy specimens are given as a function of the number of cycles in the graphs of figures 2a and 2b. The behaviour in corrosion was considered as satisfactory when the mass of the specimen hardly varied at all. This was the incubation period. An accelerated corrosion stage took place at the end of the incubation stage. This accelerated corrosion resulted more often in a rapid increase in mass corresponding to the formation of corrosion products. The graphs show a mediocre behaviou.r for the reference alloy MC2 for which the accelerated corrosion stage took place rapidly. The alloys of the invention showed incubation stages with variable durations, but in all cases were longer than that characterizing the reference alloy MC2, which demonstrated a better resistance to cyclic corrosion.
The microstructures of the alloys of the invention were checked at the end of the isothermal ageing treatments of 200 hours at 1050 C and at the end of creep tests taken to rupture at 760, 950, 1050, 1100 and 1150 C so as to check their microstructural stability as regards precipitation of undesirable intermetallic phases of the a, or Laves phase type. Only the alloy MC820 showed needle-like particles of a phase rich in rhenium at the end of the ageing treatment of 200 hours at 1050 C as well as at the end of creep tests to rupture at 1050 and 1100 C. These particles were localized in the cores of the dendrites, where rhenium separated preferentially during the process of directed solidification. All the other alloys of the invention cited in table I were free from particles of undesirable phases rich in rhenium at the e:nd of the ageing treatments and creep tests.
Claims (6)
1. A nickel-based superalloy suitable for the manufacture of turbo-engine parts by monocrystalline solidification, wherein its mass composition 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 Ni and possible impurities.
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 Ni and possible impurities.
2. The superalloy as claimed in claim 1, wherein its mass composition 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 Ni and possible impurities.
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 Ni and possible impurities.
3. The superalloy as claimed in claim 1, wherein its mass composition 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 Ni and possible impurities.
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 Ni and possible impurities.
4. The superalloy as claimed in claim 1, wherein its mass composition is as follows:
Cr : 3.5 to 5.5 %
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 Ni and possible impurities.
Cr : 3.5 to 5.5 %
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 Ni and possible impurities.
5. The superalloy as claimed in claim 1, wherein its mass composition 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 Ni and possible impurities.
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 Ni and possible impurities.
6. The superalloy as claimed in claim 1, wherein its mass composition 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 Ni and possible impurities.
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 Ni and possible impurities.
Applications Claiming Priority (2)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| FR9808693A FR2780982B1 (en) | 1998-07-07 | 1998-07-07 | HIGH SOLVUS NICKEL-BASED MONOCRYSTALLINE SUPERALLOY |
| FR98/08693 | 1998-07-07 |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| CA2276154A1 CA2276154A1 (en) | 2000-01-07 |
| CA2276154C true CA2276154C (en) | 2007-09-18 |
Family
ID=9528367
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| CA002276154A Expired - Lifetime CA2276154C (en) | 1998-07-07 | 1999-06-22 | Nickel-based monocrystalline superalloy with a high .gamma.' solvus |
Country Status (7)
| Country | Link |
|---|---|
| EP (1) | EP0971041B1 (en) |
| JP (1) | JP3902714B2 (en) |
| AT (1) | ATE225410T1 (en) |
| CA (1) | CA2276154C (en) |
| DE (1) | DE69903224T2 (en) |
| ES (1) | ES2181375T3 (en) |
| FR (1) | FR2780982B1 (en) |
Families Citing this family (25)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| FR2813392B1 (en) | 2000-08-28 | 2002-12-06 | Snecma Moteurs | STRUCTURAL REVELATION PROCESS FOR MONOCRYSTALLINE SUPERALLOYS |
| JP2004149859A (en) * | 2002-10-30 | 2004-05-27 | National Institute For Materials Science | DESIGN SUPPORT PROGRAM AND DESIGN SUPPORT SYSTEM FOR GAMMA'-PRECIPITATION-STRENGTHENED PLATINUM-GROUP-ELEMENT-ADDED Ni-BASE SUPERALLOY |
| US8968643B2 (en) | 2002-12-06 | 2015-03-03 | National Institute For Materials Science | Ni-based single crystal super alloy |
| CN100357467C (en) * | 2002-12-06 | 2007-12-26 | 独立行政法人物质·材料研究机构 | Ni-based single crystal superalloy |
| FR2881439B1 (en) | 2005-02-01 | 2007-12-07 | Onera (Off Nat Aerospatiale) | PROTECTIVE COATING FOR SINGLE CRYSTALLINE SUPERALLIAGE |
| US20080240926A1 (en) * | 2005-03-28 | 2008-10-02 | Toshiharu Kobayashi | Cobalt-Free Ni-Base Superalloy |
| US7824606B2 (en) | 2006-09-21 | 2010-11-02 | Honeywell International Inc. | Nickel-based alloys and articles made therefrom |
| JP5327442B2 (en) * | 2006-10-02 | 2013-10-30 | 昭栄化学工業株式会社 | Nickel-rhenium alloy powder and conductor paste containing the same |
| CA2663438C (en) * | 2006-10-02 | 2013-08-06 | Shoei Chemical Inc. | Nickel-rhenium alloy powder and conductor paste containing the same |
| FR2914319B1 (en) * | 2007-03-30 | 2009-06-26 | Snecma Sa | THERMAL BARRIER DEPOSITED DIRECTLY ON MONOCRYSTALLINE SUPERALLIANCES. |
| JP5467307B2 (en) * | 2008-06-26 | 2014-04-09 | 独立行政法人物質・材料研究機構 | Ni-based single crystal superalloy and alloy member obtained therefrom |
| JP5467306B2 (en) * | 2008-06-26 | 2014-04-09 | 独立行政法人物質・材料研究機構 | Ni-based single crystal superalloy and alloy member based thereon |
| JP5418589B2 (en) * | 2009-04-17 | 2014-02-19 | 株式会社Ｉｈｉ | Ni-based single crystal superalloy and turbine blade using the same |
| IT1394975B1 (en) | 2009-07-29 | 2012-08-07 | Nuovo Pignone Spa | NICKEL-BASED SUPERLEGA, MECHANICAL COMPONENT MADE WITH SUCH A SUPERLEGA, TURBOMACCHINA INCLUDING SUCH COMPONENT AND RELATIVE METHODS |
| FR2988736B1 (en) | 2012-04-02 | 2014-03-07 | Onera (Off Nat Aerospatiale) | PROCESS FOR OBTAINING A NICKEL ALUMINUM COATING ON A METALLIC SUBSTRATE, AND PART HAVING SUCH A COATING |
| GB2540964A (en) * | 2015-07-31 | 2017-02-08 | Univ Oxford Innovation Ltd | A nickel-based alloy |
| TWI595098B (en) * | 2016-06-22 | 2017-08-11 | 國立清華大學 | High-entropy superalloy |
| FR3057880B1 (en) | 2016-10-25 | 2018-11-23 | Safran | SUPERALLIAGE BASED ON NICKEL, MONOCRYSTALLINE AUBE AND TURBOMACHINE |
| KR101862059B1 (en) | 2016-11-29 | 2018-05-29 | 국방과학연구소 | Method for designing high strength ni-based powder superalloys |
| FR3073527B1 (en) | 2017-11-14 | 2019-11-29 | Safran | SUPERALLIAGE BASED ON NICKEL, MONOCRYSTALLINE AUBE AND TURBOMACHINE |
| FR3073526B1 (en) | 2017-11-14 | 2022-04-29 | Safran | NICKEL-BASED SUPERALLOY, SINGLE-CRYSTALLINE BLADE AND TURBOMACHINE |
| FR3091709B1 (en) * | 2019-01-16 | 2021-01-22 | Safran | High mechanical strength nickel-based superalloy at high temperature |
| FR3092340B1 (en) * | 2019-01-31 | 2021-02-12 | Safran | Nickel-based superalloy with high mechanical and environmental resistance at high temperature and low density |
| US20200255924A1 (en) * | 2019-02-08 | 2020-08-13 | United Technologies Corporation | High Temperature Combustor and Vane Alloy |
| FR3097879B1 (en) * | 2019-06-28 | 2021-05-28 | Safran Aircraft Engines | PROCESS FOR MANUFACTURING A PART IN MONOCRISTALLINE SUPERALLY |
Family Cites Families (10)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US3933483A (en) * | 1972-07-14 | 1976-01-20 | Kabushiki Kaisha Toyota Chuo Kenkyusho | Silicon-containing nickel-aluminum-molybdenum heat resisting alloy |
| US4222794A (en) * | 1979-07-02 | 1980-09-16 | United Technologies Corporation | Single crystal nickel superalloy |
| US4719080A (en) * | 1985-06-10 | 1988-01-12 | United Technologies Corporation | Advanced high strength single crystal superalloy compositions |
| US4781772A (en) * | 1988-02-22 | 1988-11-01 | Inco Alloys International, Inc. | ODS alloy having intermediate high temperature strength |
| FR2643085B1 (en) * | 1989-02-10 | 1991-05-10 | Onera (Off Nat Aerospatiale) | NICKEL-BASED SUPERALLOY FOR INDUSTRIAL TURBINE BLADES |
| US5151249A (en) * | 1989-12-29 | 1992-09-29 | General Electric Company | Nickel-based single crystal superalloy and method of making |
| DE69423061T2 (en) * | 1993-08-06 | 2000-10-12 | Hitachi, Ltd. | Gas turbine blade, method for producing the same and gas turbine with this blade |
| US5783318A (en) * | 1994-06-22 | 1998-07-21 | United Technologies Corporation | Repaired nickel based superalloy |
| JPH08143995A (en) * | 1994-11-17 | 1996-06-04 | Hitachi Ltd | Nickel-base single crystal alloy and gas turbine using same |
| DE69701900T2 (en) * | 1996-02-09 | 2000-12-07 | Hitachi Metals, Ltd. | High-strength nickel-based superalloy for directionally solidified castings |
-
1998
- 1998-07-07 FR FR9808693A patent/FR2780982B1/en not_active Expired - Fee Related
-
1999
- 1999-06-21 EP EP99401533A patent/EP0971041B1/en not_active Expired - Lifetime
- 1999-06-21 AT AT99401533T patent/ATE225410T1/en active
- 1999-06-21 ES ES99401533T patent/ES2181375T3/en not_active Expired - Lifetime
- 1999-06-21 DE DE69903224T patent/DE69903224T2/en not_active Expired - Lifetime
- 1999-06-22 CA CA002276154A patent/CA2276154C/en not_active Expired - Lifetime
- 1999-07-05 JP JP19070299A patent/JP3902714B2/en not_active Expired - Lifetime
Also Published As
| Publication number | Publication date |
|---|---|
| JP2000034531A (en) | 2000-02-02 |
| ES2181375T3 (en) | 2003-02-16 |
| EP0971041A1 (en) | 2000-01-12 |
| ATE225410T1 (en) | 2002-10-15 |
| FR2780982A1 (en) | 2000-01-14 |
| JP3902714B2 (en) | 2007-04-11 |
| EP0971041B1 (en) | 2002-10-02 |
| DE69903224D1 (en) | 2002-11-07 |
| CA2276154A1 (en) | 2000-01-07 |
| DE69903224T2 (en) | 2003-12-11 |
| FR2780982B1 (en) | 2000-09-08 |
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| Date | Code | Title | Description |
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| EEER | Examination request | ||
| MKEX | Expiry |
Effective date: 20190625 |