EP2393951B1 - Verfahren zur herstellung eines teils aus einer superlegierung auf basis von nickel und entsprechendes teil - Google Patents

Verfahren zur herstellung eines teils aus einer superlegierung auf basis von nickel und entsprechendes teil Download PDF

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EP2393951B1
EP2393951B1 EP10708291.9A EP10708291A EP2393951B1 EP 2393951 B1 EP2393951 B1 EP 2393951B1 EP 10708291 A EP10708291 A EP 10708291A EP 2393951 B1 EP2393951 B1 EP 2393951B1
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phase
alloy
temperature
conducted
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EP2393951A2 (de
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Alexandre Devaux
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Aubert and Duval SA
Safran Aircraft Engines SAS
Safran Helicopter Engines SAS
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Aubert and Duval SA
Safran Aircraft Engines SAS
Safran Helicopter Engines SAS
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/10—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of nickel or cobalt or alloys based thereon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/03—Alloys based on nickel or cobalt based on nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/03—Alloys based on nickel or cobalt based on nickel
    • C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
    • C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
    • C22C19/055—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 20% but less than 30%

Definitions

  • the invention relates to nickel-based superalloys, and more particularly to a heat treatment process applicable with profit to some of them to improve, in particular, their creep and tensile strength.
  • nickel-based superalloys denotes the alloys in which Ni enters for at least 50% by weight in their composition (all the percentages given in this text will be weight percentages).
  • NC19FeNb alloy trade designation INCONEL 718® (718) and alloys derived from it or are comparable to it, such as 625, 718Plus and 725.
  • alloy 718 due to the absence of cobalt in its composition and the experience acquired for its production and processing, gives it a privileged place among the alloys with high characteristics used up to a temperature close to 650 ° C.
  • the increase in the efficiency and performance of turbomachines results in an increase in the temperature at the outlet of the combustion chamber, and thus calls for an improvement in the creep resistance of alloy 718 to increase the possibilities of use. extended up to 650 ° C. Improvement of the resistance in creep of alloy 718, while retaining a fine-grained microstructure (> 7 ASTM) so as not to compromise the resistance to fatigue, is therefore of great industrial interest. It is recalled that the ASTM standards governing the estimation of the size of the grains define the grains as being all the finer the higher the ASTM figure given.
  • thermomechanical treatment making it possible to precipitate the Ni 3 Nb- ⁇ phase at the grain boundaries
  • recrystallization treatment of the alloy at a temperature below the dissolution temperature of the phase Ni 3 Nb- ⁇ , the Ni 3 Nb- ⁇ phase precipitated at the grain boundaries being used during recrystallization to prevent grain growth.
  • This process provides very fine grain recrystallized structures of 10 ASTM or greater. Their fatigue characteristics are improved but their creep resistance is insufficient.
  • Ni 3 Nb- ⁇ phase of orthorhombic structure
  • the hardening phase Ni 3 Nb- ⁇ makes it possible to slow down the movement of dislocations in the crystallographic network, and therefore to improve creep resistance.
  • Ni 3 Ta- ⁇ phase is detrimental, because it binds tantalum and thus limits the formation of the hardening phase Ni 3 Ta- ⁇ "
  • thermomechanical treatment that is to say without the usual treatment of dissolving between 900 and 980 ° C carried out between the treatment.
  • thermomechanics and aging treatment Although this option makes it possible to limit the formation of the Ni 3 Nb- ⁇ phase liable to precipitate during the solution treatment, and to obtain a fine grain and to improve the tensile and fatigue properties, it has drawbacks. .
  • EP-A-1 398 393 describes treatments of Ni-based superalloys in the form of single crystals or oriented solidified alloys.
  • the alloy is a single crystal
  • any precipitation of the ⁇ phase could only take place heterogeneously and would not prevent the growth of the grains. These would be found at the end of treatment with a size too large.
  • the alloy compositions described in a privileged manner in this document would not make it possible to precipitate ⁇ phase, given their Ti, Ta, Nb and Al contents, because this phase would not be stable because of the high Al content.
  • the aim of the invention is to improve the creep resistance and the tensile strength of nickel-based superalloys having a niobium and / or tantalum content greater than 2.5% without deteriorating the fatigue properties and while avoiding disadvantages of the aforementioned prior art.
  • the grain size obtained at the end of the aging treatment of the alloy is between 7 and 13 ASTM, preferably between 8 and 12 ASTM, better still between 9 and 11 ASTM.
  • the distribution of the ⁇ phase is homogeneous at the grain boundaries at the end of the aging treatment.
  • the passage from the first to the second level can then be carried out at a speed less than or equal to 4 ° C / min, preferably between 1 and 3 ° C / min.
  • the first level can then be carried out between 920 and 990 ° C for at least 30 min and the second level at a temperature between 960 and 1010 ° C for 5 to 45 min.
  • the total content of Nb and Ta of the alloy can then be between 5.2 and 5.5%, the first level carried out between 960 and 990 ° C for 45min to 2h and the second level carried out between 990 and 1010 ° C for 5 to 45 min.
  • the first level can be carried out between 920 and 960 ° C for 45 min to 2 hours and the second level carried out between 960 and 990 ° C for 5 to 45 min.
  • the aforementioned alloys contain, in weight percent, a phosphorus content greater than 0.007%.
  • the first stage and the second stage can be carried out at sub-solvus temperatures of the ⁇ phase of the alloy, the first stage being carried out at a temperature between the solvus temperature ⁇ minus 50 ° C and the temperature of solvus ⁇ minus 20 ° C, and the second stage being carried out at a temperature between the temperature of solvus ⁇ minus 20 ° C and the temperature of solvus ⁇ .
  • the temperature of the hot-shaped part blank can be kept constant during at least one of said stages.
  • Said third level can be carried out between 700 and 750 ° C for 4 to 16 h and a fourth level is then carried out between 600 and 650 ° C between 4 and 16 h, cooling to 50 ° C / h at +/- 10 ° C / h being carried out between said third and fourth stages.
  • Said part blank may have been produced in the form of an ingot, then hot-shaped.
  • Said part blank may have been produced by a powder metallurgy process.
  • the subject of the invention is also a part made of a nickel-based superalloy, characterized in that it has been obtained by finishing a blank of a part produced by the preceding process.
  • It may be an aeronautical or land-based gas turbine element.
  • One or more intermediate cooling is possible between each stage, but not compulsory.
  • the method according to the invention makes it possible to produce parts which, compared to those of the prior art having the same composition, exhibit a better compromise between a high elastic limit in traction, high fatigue resistance and a service life. in high creep.
  • the process for manufacturing a Ni superalloy part according to the invention can begin with the production and casting of an ingot of said superalloy by conventional processes such as a double fusion process (VIM Vacuum Induction Melting, fusion under vacuum by induction - VAR Vacuum Arc Remelting, remelting with vacuum arc) or triple melting (VIM - ESR Electroslag remelting, remelting in electroconductive slag - VAR).
  • VIM Vacuum Induction Melting fusion under vacuum by induction - VAR Vacuum Arc Remelting, remelting with vacuum arc
  • VIM - ESR Electroslag remelting, remelting in electroconductive slag - VAR triple melting
  • the method according to the invention can also be applied to a part blank resulting from powder metallurgy.
  • the first stage of the treatment according to the invention makes it possible to homogenize the distribution of the ⁇ phase within the microstructure and to reduce the variations local areas of the phase fraction ⁇ present after the thermomechanical treatments due to more or less significant temperature differences after deformation.
  • Those skilled in the art can easily, by means of routine tests, if necessary adjust the execution parameters of the first stage in order to optimize this homogenization of the distribution of the ⁇ phase.
  • the first stage of the treatment according to the invention makes it possible to precipitate in a (substantially) homogeneous manner the ⁇ phase at the grain boundaries which were therein. free after thermomechanical treatment.
  • a person skilled in the art can also, by routine tests, if necessary adjust the execution parameters of the first stage in order to optimize this homogenization of the distribution of the ⁇ phase.
  • the first stage also makes it possible to complete the recrystallization in the zones where the recrystallization would not have been complete during the thermomechanical treatment, and thus to homogenize the overall structure of the alloy. .
  • the delta phase Ni 3 Nb- ⁇ and / or Ni 3 Ta- ⁇ is partially dissolved.
  • the dissolution of the ⁇ phase takes place in a substantially uniform manner.
  • the so-called residual ⁇ phase that is to say the undissolved ⁇ phase, retains the same distribution as that obtained after the first level.
  • the residual reste phase remains substantially uniformly distributed around the grains, makes it possible to slow down the growth of all the grains and makes it possible to limit or even avoid the appearance of large grains during the second level, which is carried out at a temperature higher than that of the first stage.
  • the homogeneous distribution of the ⁇ phase at the grain boundaries promotes homogeneity of the grain size in the microstructure of the alloy at the end of the treatment.
  • the second stage therefore makes it possible to reduce the quantity of phase ⁇ obtained after the first stage to a residual quantity which is optimally less than 4%, or even below 3.5%, while avoiding an enlargement of the grain.
  • the greater dissolution of the ⁇ phase on a homogeneous fine-grained microstructure makes it possible to release more niobium for the precipitation of the hardening phases gamma 'and / or gamma "during a third stage, or even other subsequent stages, constituting a aging treatment of the alloy.
  • phase ⁇ In the absence of the first stage, when the initial microstructure results from a subsolvus deformation which led to the precipitation of phase ⁇ (state 1), the distribution of phase ⁇ is heterogeneous (see figures 4 and 5 ). Consequently, certain grains may have a large amount of ⁇ phase at the grain boundaries, or little or no ⁇ phase at the grain boundaries, or else a heterogeneous ⁇ phase distribution at the grain boundaries.
  • the grains which are not surrounded by phase ⁇ or which have little phase ⁇ at the grain boundaries, or a phase ⁇ not uniformly distributed will grow uncontrollably to a grain size of more than about 5-6 ASTM.
  • the presence, even very localized, of grains 5-6 ASTM reduces fatigue life by a factor of 10 compared to an ASTM 10 grain homogeneous microstructure.
  • the combination of the first and second bearings according to the invention thus makes it possible (see figures 8 and 9 ) partially and homogeneously dissolve the ⁇ phase, avoiding the presence of these large grains 5-6 ASTM, which is prohibitive to ensure high fatigue properties.
  • the absence of the first stage therefore does not make it possible to obtain the desired microstructure, that is to say a residual content of homogeneous ⁇ phase and preferably less than 4% and a homogeneous and acceptable grain size.
  • the grain size preferred for the products resulting from the process according to the invention results from the desire to achieve a good compromise between contradictory properties with regard to their requirements on the size of the grains.
  • the fatigue strength and the tensile strength are favored by fine grains, while the creep resistance and the cracking resistance are favored by coarse grains.
  • the preferred grain sizes are 7 to 13 ASTM, preferably 8 to 12 ASTM, more preferably 9 to 11 ASTM.
  • the inventors were able to demonstrate that the presence of the ⁇ phase between preferably 2 and 4%, and optimally between 2.5 and 3.5%, makes it possible to improve the properties of the material without weaken it.
  • ⁇ phase-free microstructures are, in general, more subject to intergranular embrittlement which considerably reduces ductility at high temperature and greatly increases the sensitivity of the alloy to the notch effect (for example to premature ruptures in the notch creep notch). Consequently, when the ⁇ phase is absent after the thermomechanical treatment, the first step is also necessary to create a minimum of ⁇ phase distributed homogeneously at the grain boundaries and to homogenize the overall structure of the material.
  • the alloy holding time at the first stage is greater than or equal to 20 minutes.
  • the temperature of the first stage is between 850 and 1000 ° C to precipitate the ⁇ phase.
  • the temperature and the holding time are adjusted as a function of the heterogeneity of the microstructure after deformation, and with a view to maintaining, after the second stage, an amount of phase ⁇ greater than the minimum required for hot ductility.
  • the second level carried out at a temperature above the first level, is therefore necessary to allow the quantity of phase ⁇ to be reduced by dissolution to the desired level, preferably at a content between 2 and 4%, and optimally between 2 , 5 and 3.5%, to release the Nb and / or Ta necessary for the precipitation of the ⁇ 'and / or ⁇ "phase while keeping a sufficient quantity of Nb and / or Ta in the form of a distributed ⁇ phase homogeneously around the grains for the hot ductility of the material.
  • the temperature and the duration of the second stage are adjusted as a function of the fraction of phase ⁇ obtained at the end of the first stage in order to obtain the desired residual fraction of phase tout, while avoiding an enlargement of the grain.
  • the duration of the second stage is also a function of the temperature determined for this stage. In general, the duration of the second stage is all the shorter the higher the temperature thereof.
  • the first two stages of treatment are successive ( fig. 1 and 2 ).
  • Successessive treatment stages it is meant that the passage from the first stage to the second treatment stage is effected by gradually increasing the temperature to pass from the first stage to the second, without go through an intermediate temperature which would be lower than that of the first stage.
  • the succession of the first two stages without falling to a temperature below the first stage, for example down to ambient temperature, makes it possible to avoid excessively large temperature gradients inside the treated sample, and to avoid a heterogeneous dissolution of the ⁇ phase which could cause grain enlargement in certain areas. It is thus preferable to adopt a sufficiently low rate of rise between the stages ( ⁇ 4 ° C./min) so that the temperature remains homogeneous within the sample treated during the second stage. It was verified during the second stage that the temperature was homogeneous after 5 minutes within a cylindrical sample of 1000 cm 3 after an increase rate of 2 ° C./min from the first stage.
  • any passage between the two stages at a temperature below the first stage risks increasing the time necessary for the homogenization of the temperature within the sample during the second stage, and risks promoting heterogeneous dissolution of the phase. ⁇ .
  • Such a change to a temperature below the first level is excluded by the invention ( fig. 3 ).
  • the first treatment level is carried out at a temperature between approximately 900 and 1000 ° C for a period of at least 30 minutes and the second treatment level is carried out at a temperature above the first level between 940 and 1020 ° C for a period of between about 5 and 90 minutes.
  • the temperature difference between the two bearings must then be at least 20 ° C.
  • the temperature and time ranges thus defined make it possible to obtain a homogeneous microstructure with an adequate grain size, that is to say between 7 and 13 ASTM, preferably between 8 and 12 ASTM, better still between 9 and 11 ASTM, and a residual fraction of ⁇ phase of between 2% and 4%.
  • the invention is based first of all on a synergistic effect between the first two stages, and an optimized balancing between these two first stages makes it possible to best meet the desired aims of the invention.
  • the solvus temperature of the ⁇ phase depends directly on the niobium + tantalum content of the alloy.
  • the quantity of niobium and / or tantalum present in the composition of the alloy therefore has a direct influence on the temperature and the duration of each stage.
  • the first level between 920 and 990 ° C for at least 30 min, and the second level between 960 and 1010 ° C for 5 to 45 min.
  • the optimal durations of the treatments also depend on the massiveness of the part to be treated, and can be determined by means of models or of experiments usual for those skilled in the art.
  • the first stage is preferably carried out at a temperature of between approximately 960 ° C and 990 ° C for a period of between about 45 minutes and 2 hours and the second level is preferably carried out at a temperature of between about 990 ° C and 1010 ° C for a period of between about 5 and 45 minutes.
  • the first stage is preferably carried out at a temperature of between approximately 920 ° C and 960 ° C for a period of between about 45 minutes and 2 hours and the second plateau is preferably carried out at a temperature of between about 960 ° C and 990 ° C for a period of between about 5 and 45 minutes.
  • the duration of treatment also depends on the massiveness of the part to be treated.
  • the temperatures at the treatment stages are generally kept substantially constant during the duration of the stage.
  • the rate of rise from the first to the second level is preferably less than 4 ° C./min, in order to avoid excessively large temperature gradients, especially in the case where large parts are being treated.
  • the rate of temperature rise from the first to the second level is preferably between 1 ° C / min and 3 ° C / min.
  • the invention applies to nickel-based superalloys, therefore containing at least 50% Ni, in which the sum Nb + Ta exceeds 2.5% by weight.
  • the alloy is a nickel-based superalloy of type 718 also called NC19FeNb (AFNOR standard), containing by weight, between 50 and 55% nickel, between 17 and 21% chromium, less than 0.08% carbon, less than 0.35% manganese, less than 0.35% silicon, less than 1% cobalt between 2.8 and 3.3% molybdenum, at least one of the elements niobium or tantalum such that the sum of niobium and tantalum is between 4.75 and 5.5% with Ta less than 0.2%, between 0.65 and 1.15% titanium, between 0.20 and 0.80% aluminum, less than 0.006% boron, less than 0.015% phosphorus, the residual percentage being iron and impurities resulting from the production.
  • NC19FeNb AFNOR standard
  • the elements for which a minimum content is not given may be present only in trace amounts, in other words at a content which may be zero, in any case sufficiently low to be without metallurgical effect (this is true for other compositions which will be cited).
  • an addition of phosphorus makes it possible to reinforce the resistance of the grain boundaries, in particular with regard to stresses such as creep and notched creep.
  • the application of the invention to such an alloy with a phosphorus content greater than 0.007% and less than 0.015% is of very particular interest since the gain obtained in creep is then markedly greater. It is thus easily possible to improve the creep life times by a factor of 4 while maintaining the same grain size. This presence of phosphorus can also, for the same reasons, be recommended for the other alloy examples below.
  • the alloy is a nickel-based superalloy of type 725, containing by weight, between 55 and 61% nickel, between 19 and 22.5% chromium, between 7 and 9.5% molybdenum, at least one of the elements niobium or tantalum such that the sum of niobium and tantalum is between 2.75 and 4% with Ta less than 0.2%, between 1 and 1.7% titanium, less than 0.55% aluminum, less than 0.5% cobalt less than 0.03% carbon, less than 0.35% manganese, less than 0.2% silicon, less than 0.006% boron, less than 0.015% phosphorus, less than 0.01% sulfur, the residual percentage being iron and impurities resulting from the production.
  • the alloy is a nickel-based superalloy of the 718PLUS type, containing by weight, between 12 and 20% chromium, between 2 and 4% molybdenum, at least one of the elements niobium or tantalum such that the sum of niobium or tantalum is between 5 and 7% with Ta less than 0.2%, between 1 and 2% of tungsten, between 5 and 10% cobalt, between 0.4 and 1.4% titanium, between 0.6 and 2.6% aluminum, between 6 and 14% iron, less than 0.1% carbon, less than 0.015% boron, less than 0.03% phosphorus the residual percentage being nickel and impurities resulting from the production.
  • the alloy is a nickel-based superalloy characterized by a niobium + tantalum content greater than 2.5% and by the presence of an intergranular phase of Ni 3 Nb-Ta type ( ⁇ phase) between 800 ° C and 1050 ° C and by the presence of an intragranular phase of Ni 3 (Al-Ti) - ( ⁇ ') type and / or of Ni 3 Nb-Ta ( ⁇ ") type between 600 and 800 ° C.
  • Ni 3 Nb-Ta the effect of the invention is also found even in the absence of the hardening phase ⁇ "Ni 3 Nb- Ta.
  • the greater dissolution of the intergranular phase of delta type Ni 3 Nb-Ta then releases niobium ( ⁇ '-gene element) which is inserted in solid solution into the hardening phase ⁇ '- Ni 3 (Al, Ti) and hardens the latter.
  • the treatment according to the invention can comprise a fourth stage making it possible to complete the precipitation of the hardening phases gamma "(Ni 3 Nb-Ta- ⁇ ') and / or gamma' (Ni 3 (Al-Ti) - ⁇ ') to a temperature lower than that of the third stage.
  • the treatment of the invention can also include at least one short-term intermediate level (1 hour maximum; see figure 2 ) between the first level and the second level to facilitate the homogenization of the temperature within large parts during the rise in temperature between the first two levels.
  • the (Ta + Nb) content of the alloy is at least 2.5%
  • the Al content does not exceed 3%, so as not to cause the precipitation of the ⁇ 'phase at the grain boundaries.
  • the ⁇ 'phase tends to be stabilized to the detriment of the ⁇ phase and the Nb comes to be inserted into the y' phase.
  • the (Nb + Ta + Ti) / Al ratio is preferable for the (Nb + Ta + Ti) / Al ratio to be greater than or equal to 3.
  • the first examples of implementation of the process according to the invention are applied to products in alloy 718 obtained after thermomechanical treatment on an alloy obtained by the conventional VIM + VAR + route. forging, but could also have been obtained by powder metallurgy, and typically intended for the production of aeronautical turbine disks.
  • thermomechanical treatments see table 2
  • TTM thermomechanical treatments
  • the products obtained after thermomechanical treatments were debited to produce samples (designated by A to P in Table 1).
  • TTH various heat treatments
  • thermomechanical treatment N ° 1 is a rolling carried out according to various passes at a temperature higher than the solvus of the phase ⁇ of the alloy.
  • the products formed according to the thermomechanical treatment range N ° 1 are bars whose metallurgical structure is free of Delta phase (metallurgical state 2).
  • samples F, K, L, N were produced from bars obtained according to this first range of thermomechanical treatment.
  • thermomechanical treatment range No. 2 is a conventional two-hot forging range (by “hot” means holding in the oven followed by deformation; “two hot” therefore means two deformation stages, each preceded by holding in the oven) at a temperature below the solvus of phase ⁇ of the alloy (“sub-solvus” temperature). This range makes it possible to precipitate the ⁇ phase in the alloy.
  • the products formed according to the thermomechanical treatment range N ° 2 are pancakes (by “pancake” is meant a product having the overall shape of a disc or a pebble resulting from deformation by forging), the metallurgical structure of which contains phase ⁇ distributed heterogeneously at the grain boundaries (metallurgical state 1, see figures 4 and 5 ).
  • samples C, E and H were made from pancakes obtained according to this second range of thermomechanical treatment.
  • thermomechanical treatment range N ° 3 is a conventional hot forging range at a temperature below the solvus of the ⁇ phase of the alloy.
  • Products formed according to the range of thermomechanical treatment N ° 3 are blanks of disks whose metallurgical structure contains phase ⁇ distributed in a very heterogeneous way at the grain boundaries (metallurgical state 1, see figures 4 and 5 ).
  • samples A, B, D, G, I, J, M, O and P were produced from blanks of turbine disks obtained according to this third range of thermomechanical treatment.
  • TTH ranges of heat treatments
  • the “a” or “b” type heat treatment ranges are reference heat treatment ranges representative of the state of the art.
  • the “a” type treatment ranges consist of a so-called isothermal solution stage and two aging stages.
  • the solution stage consisted, for samples A, B, C, D, F and P, in maintaining the alloy at a constant temperature between 955 and 1010 ° C. for 40 to 90 minutes.
  • the two aging stages for their part, consisted of a stage at 720 ° C. for 8 hours followed by controlled cooling at 50 ° C./h up to a stage at 620 ° C. for 8 hours.
  • Heat treatment ranges types "c" are in accordance with the invention and comprise two said bearings of solution treatment, respectively indicated 1 and level 2 nd level, and one or two levels of aging, respectively indicated 3rd bearing and 4 th step.
  • the 1st stage of dissolution consisted of maintaining the alloy at a constant temperature between 940 ° C and 980 ° C. for about 50 to 60 minutes.
  • the 2 nd dissolution stage consisted of maintaining the alloy at a constant temperature of between 980 ° C. and 1005 ° C. for approximately 15 to 40 minutes.
  • the passage from the 1 st to the 2 nd level was carried out by a controlled heating at a rate of approximately 2 ° C / min.
  • the 3 rd and 4 th aging stages conformed to the corresponding aging stages of the “a” type reference ranges except for samples H and J.
  • the temperature of the 3rd aging treatment step was increased to 750 ° C instead of 720 ° C in the case of other samples.
  • This difference has made it possible to show that the field of the invention is not limited to restricted conditions of temperatures and durations of the aging stages, but that, on the contrary, the invention is also applicable for temperatures and durations of the aging stages. aging stages such as those practiced in the field of nickel-based superalloys.
  • Sample J for its part, underwent only one level of aging treatment at 720 ° C. for 10 hours.
  • the aging treatment undergone by sample J shows that the invention is also applicable when the alloy undergoes only one level of aging treatment.
  • the “d” type heat treatment ranges include two dissolution stages and two aging stages. Samples I and L were processed according to these ranges. However, these treatments are not in accordance with the invention due to a second stage carried out at too high a temperature or for too long a period. Indeed, the conditions of the 2 nd stage lead to too great a dissolution of phase ⁇ , and the growth of the grains is no longer controlled, which causes an uncontrolled and significant enlargement of the grains during the second stage for samples I and L.
  • the “e” type heat treatment range includes a single stage for dissolving at 1005 ° C. for 15 minutes and two stages for aging. Only sample O was obtained according to this range of heat treatment which is not in accordance with the invention as explained below.
  • Samples A to L and O were type 718 alloys at 5.3% Nb and 40 ppm P.
  • Sample N was type 718 alloy at 5.0% Nb and 40 ppm P
  • Samples M and P were type 718 alloys at 5.3% Nb and 80 ppm P.
  • Table 1 compositions of the samples tested Samples Ni% Fe% Cr% Al% Ti% Nb% Mo% B% VS % P% AL, O 54.2 rest 17.9 0.5 0.97 5.3 3 0.003 0.03 0.004 NOT 53.7 rest 17.9 0.49 0.98 5.0 3 0.003 0.02 0.004 M, P 54.0 rest 18.1 0.5 1.00 5.3 3 0.003 0.03 0.008
  • Table 2 summarizes the processing conditions for the various samples, and the ASTM grain sizes and percentages of ⁇ surface phase visible on a micrograph.
  • the grain size is defined according to the ASTM standard, and in cases where the grain size is relatively inhomogeneous, the maximum grain size (ALA) is also specified.
  • Table 2 Characteristics and treatments of the different test samples Solution stages Aging Stages Ech Alloy Nb% TTM TTH 1st level Ramp 2nd level 3rd tier cooling 4 th bearing Microstructure T ° (° C) Duration (min.) ° C / min T ° (° C) Duration (min.) T ° (° C) Duration (h.) ° C / h T ° (° C) Duration (h.) ASTM grain % ⁇ AT 718 5.3 3 at 955 60 - - - 720 8 50 620 8 11-12 5.9 B 718 5.3 3.
  • Product F is a reference sample which, after the thermomechanical range n ° 1, was treated according to a standard “a” heat treatment range of alloy 718 (treatment comprising a single stage of dissolution of phase subs). .
  • the product L was treated with a solution in two stages but with a second stage carried out at too high a temperature and time, outside the scope of the invention for an alloy 718.
  • Products K and N do not have the same niobium content, but have both undergone a range of heat treatment “c” according to the invention.
  • the 718 alloy products marked C, E and, H were transformed according to the thermomechanical range No. 2 which makes it possible to precipitate the ⁇ phase in a heterogeneous manner.
  • Product C is a reference sample which, after thermomechanical range No. 2, was treated according to a standard type “a” type heat treatment range of alloy 718 (treatment comprising a single level of subsolvus solution).
  • Product E is also a reference sample which, after the thermomechanical range n ° 2, was treated according to the type “b” heat treatment range and was therefore directly aged after forging (“direct aged”), and has not therefore not undergone solution treatment before aging. After the thermomechanical range n ° 2,
  • the product H has undergone a heat treatment according to the invention (type "c") with a two-stage solution in the field of the invention.
  • the 718 alloy products marked A, B, D, G, I, J, M, O and P were processed according to the thermomechanical range No. 3 which makes it possible to precipitate the ⁇ phase in a very heterogeneous manner.
  • thermomechanical treatment no. 3 After the thermomechanical treatment no. 3, the products A, B and P were treated according to a standard treatment range of alloy 718 (type “a” treatment comprising a single level of subsolvus solutionization).
  • Product D was treated with a treatment comprising a single stage of dissolution but at a higher temperature than products A, B and P, that is to say at a temperature close to the solvus of phase ⁇ .
  • thermomechanical treatment After thermomechanical treatment, product I was treated with a solution in two stages but with a duration, for the second stage, too high with regard to the temperature.
  • the heat treatment undergone by I is therefore outside the scope of the invention.
  • thermomechanical treatment No. 3 After thermomechanical treatment No. 3, product G was treated with a two-stage solution in the field of the invention (heat treatment “c”).
  • Product J has also been treated with a two-stage dissolving within the scope of the invention, but has not been treated with a fourth stage.
  • the product M has been treated with a two-stage dissolving in the field of the invention, but has a phosphorus content equal to 0.008% which is twice that of the products A-L and N-O.
  • the product O has undergone a heat treatment "e" with a solution in a single stage; this treatment is outside the scope of the invention.
  • Product P is a reference sample with a phosphorus content of 0.008%. It was treated according to a standard treatment range for alloy 718 (“a” type treatment comprising a single level of subsolvus dissolving).
  • Products A, B, C which have been treated with a standard subsolvus heat treatment have a fine-grained microstructure (> 9 ASTM) but have a higher fraction of phase phase (> 4.5%) with the fraction of phase ⁇ preferably sought within the framework of the invention.
  • the mechanical properties obtained by these products constitute the reference for assessing the tensile, fatigue and creep properties obtained on the thermomechanical ranges (TTM) 2 and 3.
  • Product D was processed at a higher temperature than Product A, B, and C, it has ASTM grains and a ⁇ phase which is heterogeneously distributed ( ⁇ 2.5%) and is less than the fraction of phase ⁇ preferably sought in the context of the invention. It is noted that this treatment did not make it possible to retain a fine-grained microstructure (at least 7 ASTM, preferably at least 8, better 9 ASTM) and the satisfactory fatigue properties observed for products A, B, and C. The considerable reduction in fatigue lives is attributable to the presence of ASTM coarse grains which constitute the fatigue initiation sites.
  • Product E which was directly aged after thermomechanical treatment N ° 2 has a very heterogeneous grain size (10 to 14 ASTM) and significant variations in the phase rate taux, this rate being found in most areas of the part. (particularly those subjected to creep) greater than the desired de phase fraction. Although the tensile and fatigue properties of product E are superior to those of products A, B, C, we notes that the creep lifetimes obtained with product E are less than the creep lifetimes of products A, B, C.
  • the products G, H, M have been treated in the field of the invention and comprise a fine-grained microstructure (> 9 ASTM) and a fraction of ⁇ phase (2.9% and 3.5%) included in the preferably desired interval of phase fraction ché, namely 4% at most and 2.5% at minimum. It is observed that the tensile properties are clearly superior to those of products A, B, C and of the same level as those of product E. It is also observed that the creep properties of products G, H, M are clearly superior to those of products. products A, B, C, E while the grain size is similar in these products.
  • the fine-grained microstructure of products G, H, M makes it possible to retain the fatigue properties obtained with products A, B, C, E and the lower ⁇ phase fraction of products G, H, M makes it possible to improve the creep resistance.
  • the combination of an addition of phosphorus and of the treatment according to the invention therefore has a synergistic effect which is positive on the creep properties of the alloy obtained.
  • the invention aims to conserve a fraction of residual ⁇ phase (preferably greater than 2.5%) which makes it possible to maintain satisfactory ductility at high temperature. Too low a ⁇ phase content has an effect on damage and tensile ductility at high temperature (650 ° C with a strain rate of 10 -5 s -1 ). It is in fact observed that product D with a ⁇ phase content close to 2% has a ductility (elongation at break of 7%) much lower than that of product G (elongation at break of 27%) which comprises a fraction phase ⁇ close to 3%. This reduction in ductility for product D results from intergranular damage caused by a too small and heterogeneously distributed ⁇ phase fraction.
  • the figures 4 and 5 present the microstructure of samples A, B, C, D, E, G, H, I, J, M, O and P (metallurgical state 1) after they have undergone a range of sub-solvus thermomechanical deformation (thermomechanical range 2 or 3). It is a microstructure which exhibits delta phase Ni 3 Nb- ⁇ and / or Ni 3 Ta- ⁇ at the grain boundaries, but in a non-uniform manner distributed between the grains.
  • the figure 4 shows that the samples have a fine grain of approximately ASTM size 11, with a heterogeneous distribution of the ⁇ phase (black spots at the grain boundaries). After the range of thermomechanical deformation, the ⁇ phase percentage is 2.8-6% and the grain size is 10-13 ASTM. We therefore have a very heterogeneous microstructure from these two points of view.
  • the figure 5 shows the microstructure of the samples with a higher magnification and shows grains whose boundaries are very largely free of phase ⁇ (the latter appearing in white on this micrograph).
  • sample B When a treatment comprising only a first stage of dissolution at 970 ° C for approximately 60 minutes is applied to a sample (sample B), a percentage of phase ⁇ of 4.7 to 5.5% is obtained and a grain size of 11 to 12 ASTM. The homogeneity of the sample is therefore improved, but a large fraction of phase ⁇ is retained, which is known (see sample B, Tables 1 & 2) to be very unfavorable to creep resistance.
  • the grains which are not surrounded by phase ⁇ or which have little phase ⁇ at the grain boundaries will grow larger. uncontrollably to a grain size of up to about 5-6 ASTM, while the growth of the other grains surrounded by phase ⁇ will be thwarted and will give rise to grain sizes close to 9 ASTM. This heterogeneity in grain size is evident on micrographs of the figures 6 and 7 . The presence, even very localized, of grains 5-6 ASTM considerably reduces the lifetimes in fatigue.
  • the inventors moreover carried out additional tests on samples of alloys of type 718 Plus and 725, and were thus able to confirm that the invention applied to other nickel-based superalloys having a niobium and / or tantalum content greater than 2.5% made it possible to clearly improve their creep resistance and their tensile strength.

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Claims (16)

  1. Verfahren zum Herstellen eines Rohlings eines Teils aus einer Superlegierung auf Basis von Ni, welche wenigstens 50 Gew.-% Ni enthält, wobei man eine Legierung einer solchen Superlegierung bereitstellt, diese einer Deformierung unter Wärme unterwirft und thermische Behandlungen der besagten Legierung realisiert, dadurch gekennzeichnet, dass:
    - besagte Superlegierung in Gewichtsprozent umfasst:
    • entweder
    zwischen 50 und 55% Nickel,
    zwischen 17 und 21% Chrom,
    weniger als 0,08% Kohlenstoff,
    weniger als 0,35% Mangan,
    weniger als 0,35% Silizium,
    weniger als 1% Kobalt,
    zwischen 2,8 und 3,3% Molybdän,
    wenigstens eines der Elemente Niob oder Tantal, derart, dass die Summe von Niob und Tantal zwischen 4,75 und 5,5% liegt mit Ta geringer als 0,2%, zwischen 0,65 und 1,15% Titan,
    zwischen 0,20 und 0,80% Aluminium,
    weniger als 0,006% Bor,
    weniger als 0,015% Phosphor,
    wobei der Restprozentgehalt Eisen und Verunreinigungen sind, die vom Bereitstellen stammen,
    • oder:
    zwischen 55 und 61% Nickel,
    zwischen 19 und 22,5% Chrom,
    zwischen 7 und 9,5% Molybdän,
    wenigstens eines der Elemente Niob oder Tantal, derart, dass die Summe von Niob und Tantal zwischen 2,75 und 4% liegt mit Ta geringer als 0,2%,
    zwischen 1 und 1,7% Titan,
    weniger als 0,55% Aluminium,
    weniger als 0,5% Kobalt,
    weniger als 0,03% Kohlenstoff,
    weniger als 0,35% Mangan,
    weniger als 0,2% Silizium,
    weniger als 0,006% Bor,
    weniger als 0,015% Phosphor,
    weniger als 0,01% Schwefel,
    wobei der Restprozentgehalt Eisen und Verunreinigungen sind, die vom Bereitstellen stammen,
    • oder:
    zwischen 12 und 20% Chrom,
    zwischen 2 und 4% Molybdän,
    wenigstens eines der Elemente Niob oder Tantal, derart, dass die Summe von Niob oder Tantal zwischen 5 und 7% liegt mit Ta geringer als 0,2%,
    zwischen 1 und 2% Wolfram,
    zwischen 5 und 10% Kobalt,
    zwischen 0,4 und 1,4% Titan,
    zwischen 0,6 und 2,6% Aluminium,
    zwischen 6 und 14% Eisen,
    weniger als 0,1% Kohlenstoff,
    weniger als 0,015% Bor,
    weniger als 0,03% Phosphor,
    wobei der Restprozentgehalt Nickel und Verunreinigungen sind, die vom Bereitstellen stammen,
    - man eine thermische Behandlung besagter Legierung durchführt, welche eine Mehrzahl von Schritten aufweist, die auf die nachfolgende Weise aufgeteilt sind:
    * einen ersten Schritt, während dem man besagte Legierung bei zwischen 850 und 1000 °C für 20 Minuten aufrechterhält, um δ-Phase an Korngrenzen abzuscheiden,
    * einen zweiten Schritt, während dem man besagte Legierung bei einer Temperatur aufrechterhält, welche höher ist als die des ersten Schritts, und der es ermöglicht, ein partielles Auflösen der während des ersten Schritts erlangten δ-Phase zu realisieren, und, am Ende des zweiten Schritts, eine Menge der δ-Phase zwischen 2 und 4%, vorzugsweise zwischen 2,5 und 3,5%, zu erlangen, wobei der erste und der zweite Schritt ohne zwischenzeitliches Abkühlen realisiert werden,
    * eine Alterungsbehandlung, welche einen dritten Schritt und eventuell einen oder mehrere Zusatzschritte aufweist, die bei einer Temperatur realisiert werden, welche geringer ist als die des ersten Schritts, und es ermöglicht, härtende y'- und/oder γ"-Phasen abzuscheiden,
    * wobei der erste Schritt bei zwischen 900 und 1000 °C für wenigstens 30 Minuten und der zweite Schritt bei zwischen 940 und 1020 °C für 5 bis 90 Minuten realisiert wird, wobei die Temperaturdifferenz zwischen den zwei Schritten wenigstens 20 °C beträgt.
  2. Verfahren gemäß Anspruch 1, dadurch gekennzeichnet, dass die Korngröße, welche am Ende der Behandlung der Legierung erlangt wird, zwischen 7 und 13 ASTM, vorzugsweise zwischen 8 und 12 ASTM, besser zwischen 9 und 11 ASTM, liegt.
  3. Verfahren gemäß Anspruch 1 oder 2, dadurch gekennzeichnet, dass die Verteilung der δ-Phase am Ende der Alterungsbehandlung homogen an den Korngrenzen ist.
  4. Verfahren gemäß einem der Ansprüche 1 bis 3, dadurch gekennzeichnet, dass der erste Übergang vom ersten zum zweiten Schritt bei einer Geschwindigkeit geringer oder gleich 4 °C/min, vorzugsweise zwischen 1 und 3 °C/min, erfolgt.
  5. Verfahren gemäß einem der Ansprüche 1 bis 4, dadurch gekennzeichnet, dass besagte Superlegierung in Gewichtsprozent umfasst:
    zwischen 50 und 55% Nickel,
    zwischen 17 und 21% Chrom,
    weniger als 0,08% Kohlenstoff,
    weniger als 0,35% Mangan,
    weniger als 0,35% Silizium,
    weniger als 1% Kobalt,
    zwischen 2,8 und 3,3% Molybdän,
    wenigstens eines der Elemente Niob oder Tantal, derart, dass die Summe von Niob und Tantal zwischen 4,75 und 5,5% liegt mit Ta geringer als 0,2%,
    zwischen 0,65 und 1,15% Titan,
    zwischen 0,20 und 0,80% Aluminium,
    weniger als 0,006% Bor,
    weniger als 0,015% Phosphor,
    wobei der Restprozentgehalt Eisen und Verunreinigungen sind, was vom Bereitstellen stammt,
    und dadurch, dass der erste Schritt bei zwischen 920 und 990 °C für wenigstens 30 min realisiert wird und der zweite Schritt bei einer Temperatur zwischen 960 und 1010 °C für 5 bis 45 min realisiert wird.
  6. Verfahren gemäß Anspruch 5, dadurch gekennzeichnet, dass der Gesamtgehalt an Nb und Ta der Legierung zwischen 5,2 und 5,5% liegt, dass der erste Schritt bei zwischen 960 und 990 °C für 45 min bis 2 h realisiert wird und, dass der zweite Schritt bei zwischen 990 und 1010 °C für 5 bis 45 min realisiert wird.
  7. Verfahren gemäß Anspruch 5, dadurch gekennzeichnet, dass der Gesamtgehalt von Nb und Ta der Legierung zwischen 4,8 und 5,2% liegt, dass der erste Schritt bei zwischen 920 und 960 °C für 45 min bis 2 h realisiert wird und, dass der zweite Schritt bei zwischen 960 und 990 °C für 5 bis 45 min realisiert wird.
  8. Verfahren gemäß einem der Ansprüche 1 bis 7, dadurch gekennzeichnet, dass die Legierung in Gewichtsprozent einen Phosphorgehalt größer als 0,007% hat.
  9. Verfahren gemäß einem der Ansprüche 1 bis 8, dadurch gekennzeichnet, dass der erste Schritt und der zweite Schritt bei Sub-Solvus-Temperaturen der δ-Phase der Legierung realisiert, wobei der erste Schritt bei einer Temperatur zwischen der δ-Solvustemperatur minus 50 °C und der δ-Solvustemperatur minus 20 °C realisiert wird und der zweite Schritt bei einer Temperatur zwischen der δ-Solvustemperatur minus 20 °C und der δ-Solvustemperatur realisiert wird.
  10. Verfahren gemäß einem der Ansprüche 1 bis 9, dadurch gekennzeichnet, dass die Temperatur des unter Wärme in Form gebrachten Rohlings des Teils während wenigstens einem der besagten Schritte konstant aufrechterhalten wird.
  11. Verfahren gemäß einem der Ansprüche 1 bis 10, dadurch gekennzeichnet, dass besagter dritter Schritt bei zwischen 700 und 750 °C für 4 bis 16 h realisiert wird und dass ein vierter Schritt bei zwischen 600 und 650 °C für zwischen 4 bis 16 h realisiert wird, wobei ein Abkühlen bei 50 °C/h bis +/-10 °C/h zwischen besagtem dritten und vierten Schritt realisiert wird.
  12. Verfahren gemäß einem der Ansprüche 1 bis 11, dadurch gekennzeichnet, dass man zwischen dem ersten und dem zweiten Schritt wenigstens ein Aufrechterhalten der unter Wärme in Form gebrachten Legierung bei einer Zwischentemperatur zwischen den Temperaturen des ersten und des zweiten Schritts für maximal 1 h realisiert.
  13. Verfahren gemäß einem der Ansprüche 1 bis 12, dadurch gekennzeichnet, dass besagter Rohling des Teils in Form eines Gussblocks bereitgestellt wurde, anschließend unter Wärme in Form gebracht wurde.
  14. Verfahren gemäß einem der Ansprüche 1 bis 13, dadurch gekennzeichnet, dass besagter Rohling des Teils mittels eines Pulvermetallurgie-Verfahrens bereitgestellt wurde.
  15. Teil aus einer Superlegierung auf Basis von Ni, dadurch gekennzeichnet, dass es mittels Fertigstellens eines Rohlings eines Teils erlangt wurde, welcher mittels des Verfahrens gemäß einem der Ansprüche 1 bis 14 hergestellt wurde.
  16. Teil gemäß Anspruch 15, dadurch gekennzeichnet, dass es sich um ein aeronautisches oder terrestrisches Gasturbinenelement handelt.
EP10708291.9A 2009-02-06 2010-02-05 Verfahren zur herstellung eines teils aus einer superlegierung auf basis von nickel und entsprechendes teil Active EP2393951B1 (de)

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WO2010089516A2 (fr) 2010-08-12
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US20120037280A1 (en) 2012-02-16
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FR2941962B1 (fr) 2013-05-31
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