Classifications

• • 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 method that is advantageously applicable to some of them to improve, in particular, their creep and tensile strengths.
• nickel-based superalloys denotes alloys in which the Ni enters for at least 50% by weight in their composition (all percentages given in this text will be weight percentages). More specifically, the invention relates to a heat treatment process applicable to alloys containing more than 2.5% in total of niobium and tantalum, and which are therefore likely to show a double precipitation:
• alloy NC19FeNb commercial designation INCONEL 718 ® (718) and alloys derived from or 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 elaboration and processing, gives it a privileged place among alloys with high characteristics used up to a close temperature. 650 ° C.
• the increase in efficiency and performance of turbomachines results in an increase in temperature at the combustion chamber outlet, and thus calls for an improvement in the creep resistance of alloy 718 to increase the possibilities of use extended to 650 ° C.
• thermomechanical treatment processes Two different thermomechanical treatment processes are known and today implemented to improve the fatigue properties of the alloy 718.
• a first option as described in FR-A-2 089 069 it has been chosen to provide a thermomechanical treatment for precipitating the Ni 3 Nb- ⁇ phase at the grain boundaries, then conducting a recrystallization treatment of the alloy at a temperature below the dissolution temperature of the Ni 3 Nb- ⁇ phase, the phase Ni 3 Nb- ⁇ precipitated at grain boundaries being used during recrystallization to prevent grain growth.
• This method makes it possible to obtain very fine grain recrystallized structures of ASTM or higher. Their fatigue characteristics are improved but their creep resistance is insufficient.
• Ni 3 Nb- ⁇ , orthorhombic structure is harmful because it sets niobium and thus limits the hardening phase of training Ni 3 Nb- ⁇ "metastable and centered tetragonal structure.
• the Ni 3 Nb- ⁇ "hardening phase makes it possible to slow down the movement of the dislocations in the crystallographic network, and thus to improve the creep resistance.
• Ni 3 Ta- ⁇ phase is detrimental because it fixes the tantalum and thus limits the formation of Ni 3 Ta- ⁇ "hardening phase.
• Another known solution to improve the properties of 718 consists in carrying out direct aging after thermomechanical treatment, that is to say without the usual solution treatment between 900 and 980 ° C. carried out between the thermomechanical treatment and the aging treatment. that this option makes it possible to limit the formation of the phase Ni 3 Nb- ⁇ likely to precipitate during the dissolution treatment, and to obtain a fine grain and to improve the tensile and fatigue properties, it has drawbacks. It turns out that one obtains heterogeneous microstructures within the same room, because of important local variations in the size of the grains and the proportion of phase ⁇ formed during the thermomechanical treatments.
• EP-A-1 398 393 discloses Ni-based superalloy treatments in the form of single crystals or orientated solidified alloys.
• the alloy is a single crystal
• a possible ⁇ phase precipitation could only be heterogeneous and would not prevent the growth of the grains. These would end up at the end of treatment with a size too high.
• the compositions of the alloy preferably described in this document would not allow precipitation of the ⁇ phase, given their Ti, Ta, Nb and Al contents, since this phase would not be stable because of the high content of Al.
• US-A-4,459,160 also discloses monocrystalline Ni-based superalloys, in which therefore no ⁇ -phase precipitation at the grain boundaries can be observed.
• the object 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 subject of the invention is a process for manufacturing a Ni-based superalloy component blank containing at least 50% Ni in weight percentages, according to which an alloy of such a superalloy is produced, and thermal treatments of said alloy are carried out, characterized in that:
• said superalloy contains in weight percentages at least 2.5% in total of Nb and Ta; a heat treatment of said alloy is carried out, comprising a plurality of stages distributed as follows: a first bearing during which said alloy is maintained between 850 and
• an aging treatment comprising a third bearing and possibly one or more additional bearings, carried out at a temperature lower than that of the first bearing and making it possible to precipitate the hardening phases ⁇ 'and / or ⁇ ".
• the Al content of the alloy is less than or equal to 3%.
• the ratio (Nb + Ta + Ti) / Al of the alloy is greater than or equal to 3.
• 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.
• phase quantity ⁇ of between 2 and 4%, better still between 2.5 and 3.5%, is preferably obtained.
• the first and second bearings are preferably made without intermediate cooling.
• the passage from the first to the second step 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 step can be carried out between 900 and 1000 0 C for at least
• the alloy may contain by weight: between 50 and 55% of nickel, between 17 and 21% of chromium, less than 0.08% of carbon, less than 0.35% of manganese, less than 1% of cobalt less than 0.35% of silicon, between 2.8 and 3.3% of molybdenum, at least one of the elements niobium or tantalum so 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 elaboration.
• the first step can then be performed between 920 and 990 ° C. for at least 30 minutes and the second step at a temperature of between 960 and 1010 ° C. for 5 to 45 minutes.
• the total content of Nb and Ta of the alloy can then be between 5.2 and 5.5%, the first stage carried out between 960 and 990 ° C. for 45 minutes to 2 hours and the second step carried out between 990 and 1010 minutes. C for 5 to 45 minutes.
• the first stage can be carried out between 920 and 960 ° C. for 45 minutes at 2 hours and the second stage carried out between 960 and 960 ° C. 990 ° C for 5 to 45 min.
• the alloy may contain by weight: between 55 and 61% of nickel, between 19 and 22.5% of chromium, between 7 and 9.5% of molybdenum, at least one of the elements niobium or tantalum so 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 % of cobalt, less than 0,03% of carbon, less than 0,35% of manganese, less than 0,2% of silicon, less than 0,006% of boron, less than 0,015% of phosphorus, less than 0.01% sulfur, the residual percentage being iron and impurities resulting from the preparation.
• the alloy may contain, by weight: between 12 and 20% of chromium, between 2 and 4% of molybdenum, at least one of the elements niobium or tantalum so 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% of cobalt, between 0.4 and 1.4% of titanium, between 0.6 and 2.6% of aluminum , between 6 and 14% of iron, less than 0.1% of carbon, less than 0.015% of boron, less than 0.03% of phosphorus, the residual percentage being nickel and impurities resulting from the preparation.
• the aforementioned alloys contain, in percentage by weight, a phosphorus content greater than 0.007%.
• the first bearing and the second bearing can be made at sub-solvus temperatures of the phase ⁇ of the alloy, the first bearing being carried out at a temperature between the solvus temperature ⁇ minus 50 0 C and the temperature of solvus ⁇ minus 20 0 C, and the second stage being carried out at a temperature between the solvus temperature ⁇ minus 20 ° C and the solvus temperature ⁇ .
• the temperature of the hot-formed workpiece blank can be kept constant during at least one of said bearings.
• Said third stage can be achieved between 700 and 750 0 C for 4 to 16h and a fourth stage is then carried out between 600 and 650 ° C between 4 and 16h, cooling at 50 ° C / h to +/- 10 ° C / h being formed between said third and fourth bearings.
• a fourth stage is then carried out between 600 and 650 ° C between 4 and 16h, cooling at 50 ° C / h to +/- 10 ° C / h being formed between said third and fourth bearings.
• Said piece blank may have been produced in the form of an ingot, then shaped hot.
• Said blank may have been produced by a powder metallurgy process.
• the invention also relates to a nickel-based superalloy part, characterized in that it was obtained from a part blank manufactured by the above method.
• the invention consists in carrying out on a Ni-base alloy containing Nb and / or Ta a heat treatment for which the structural hardening is obtained by precipitation of the gamma '(Ni 3 Ti- ⁇ ') hardening phases. and / or gamma "(Ni 3 Nb- ⁇ " and / or Ni 3 Ta- ⁇ "), these phases respectively comprising titanium and Niobium and / or Tantalum.
• the heat treatment comprises at least three stages which are chronologically :
• a first treatment stage carried out at 850-1000 ° C. which is intended to precipitate the delta phase Ni 3 Nb- ⁇ and / or Ni 3 Ta- ⁇ at the grain boundaries, with a substantially homogeneous distribution of this phase in the joints grains, and to homogenize the microstructure of the material; it also makes it possible, in the case of partially recrystallized microstructures, to complete the recrystallization and to precipitate the ⁇ phase at the joints of the new recrystallized grains;
• Ni 3 Nb- ⁇ and / or Ni 3 Ta- ⁇ while keeping the substantially homogeneous distribution obtained at the end of the first stage, and avoiding grain enlargement; the second stage ends with oil quenching or air cooling;
• the third bearing and any subsequent bearings constitute (s) an aging heat treatment, performed at a temperature lower than that of the first bearing and making it possible to precipitate the hardening phases gamma '(Ni 3 (Al-Ti) Y) and / or gamma "(Ni 3 Nb- ⁇ " or Ni 3 Ta- ⁇ ");
• One or more intermediate coolings are possible between each level, but not required.
• the method according to the invention makes it possible to produce parts which, compared with those of the prior art having the same composition, have a better compromise between a high tensile yield strength, a high fatigue strength and a service life. in high creep.
• FIG. 1 to 3 schematically three examples of the first two heat treatment levels according to the invention, Figure 2 also showing an intermediate bearing between the first and second bearing; the ordinate temperatures are referenced with respect to the solvus temperature of the ⁇ phase.
• FIGS. 4 to 9 show micrographs of alloys which have undergone reference heat treatments ( Figures 4 to 7) and according to the invention ( Figure 8, 9).
• the process for manufacturing a Ni superalloying part according to the invention may begin with the elaboration and casting of an ingot of said superalloy by conventional methods such as a double melting process (VIM Vacuum Induction Melting, fusion under induction vacuum - VAR Vacuum Arc Remelting, vacuum arc remelting) or triple fusion (VIM - ESR Electroslag remelting, electroslag remelting - VAR).
• the process according to the invention can also be applied to a blank of a piece derived from powder metallurgy.
• the initial microstructure of a product (it being understood that by the term “product” is meant a half-product or a part blank) before the typical treatment of the invention may vary depending on the thermomechanical deformation treatments carried out upstream, for example, forging, stamping or hot rolling:
• the delta phase Ni 3 Nb- ⁇ and / or Ni 3 Ta- ⁇ may be present at the grain boundaries but in a non-uniform manner distributed between the grains following a deformation performed at a temperature below the solvus of the ⁇ phase;
• the delta phase Ni 3 Nb- ⁇ and / or Ni 3 Ta- ⁇ can be absent or almost absent ( ⁇ 1%) from the microstructure following a deformation carried out for example at a temperature above the solvus of the ⁇ phase.
• 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 of the ⁇ phase fraction present after the thermomechanical treatments because of differences in temperature more or less important after deformation.
• the first stage of the treatment according to the invention makes it possible to precipitate
• phase ⁇ grain boundaries that were free after the thermomechanical treatment.
• Those skilled in the art may also, by routine tests, adjust if necessary the execution parameters of the first stage 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 second stage of the treatment according to the invention carried out at a temperature close to the solvus of the ⁇ phase, the delta phase Ni 3 Nb- ⁇ and / or Ni 3 Ta- ⁇ is partially dissolved.
• the dissolution of the ⁇ phase is substantially uniform, since the microstructure obtained after the first step is preferably homogeneous.
• the so-called residual phase ⁇ that is to say the undissolved ⁇ phase, retains the same distribution as that obtained after the first plateau.
• the residual ⁇ phase remains substantially uniformly distributed around the grains, serves to slow the growth of all the grains and makes it possible to limit or even prevent the appearance of coarse grains during the second stage, 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 the homogeneity of the grain size in the microstructure of the alloy at the end of treatment.
• the second step thus makes it possible to reduce the quantity of phase ⁇ obtained after the first step up to a residual quantity optimally lower than 4%, or even below 3.5%, while avoiding grain enlargement.
• the greater dissolution of the ⁇ phase on a homogeneous fine-grained microstructure makes it possible to release more niobium for the precipitation of the gamma and / or gamma hardening phases during a third stage, or even other subsequent stages, constituting a aging treatment of the alloy.
• the grains which are not surrounded by phase ⁇ or which have little phase ⁇ at the grain boundaries, or a phase ⁇ not uniformly distributed will increase uncontrolled up to a grain size of approximately 5-6 ASTM.
• the presence, even very localized, of 5-6 ASTM grains reduces the fatigue life by a factor of 10 compared to a homogeneous ASTM grain microstructure.
• the combination of the first and second bearings according to the invention thus makes it possible (see FIGS. 8 and 9) to partially and homogeneously dissolve the ⁇ phase, avoiding the presence of these large grains 5-6 ASTM, which is unacceptable to guarantee high fatigue properties.
• phase ⁇ (state 1)
• the absence of the first stage therefore does not make it possible to obtain the desired microstructure, that is to say a residual phase current.
• ⁇ homogeneous and preferentially less than 4% and a homogeneous and acceptable grain size.
• the preferred grain size on the products resulting from the process according to the invention results from the desire to achieve a good compromise between contradictory properties as regards their requirements on the grain size. Indeed, the fatigue strength and the tensile strength are favored by fine grains, while the creep resistance and the crack resistance are favored by coarse grains. In this perspective, the preferred grain sizes are 7 to 13 ASTM, preferably 8 to 12 ASTM, better 9 to 1 1 ASTM.
• phase-free microstructures ⁇ are, in general, more susceptible to intergranular embrittlement which considerably reduces the ductility at high temperature and greatly increases the sensitivity of the alloy to the notch effect (eg premature breaks in notched creep notch). Consequently, when the ⁇ phase is absent after the thermomechanical treatment, the first step is also necessary to create a minimum phase ⁇ distributed homogeneously at the grain boundaries and to homogenize the overall structure of the material.
• the duration of maintenance of the alloy at the first stage is greater than or equal to 20 minutes.
• the temperature of the first stage is between 850 and 1000 0 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 in order to maintain after the second stage a quantity of phase ⁇ greater than the minimum required for the hot ductility.
• the second step, carried out at a temperature above the first step, is therefore necessary to allow the amount of the ⁇ phase to be lowered by dissolution to the desired level, preferably at a content of between 2 and 4%, and optimally between 2 and 4%.
• the temperature and the duration of the second stage are adjusted as a function of the fraction of the phase ⁇ obtained at the end of the first stage to obtain the desired phase residual ⁇ fraction, while avoiding a grain magnification.
• 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 shorter as the temperature of the latter is high.
• the first two stages of treatment are successive ( Figures 1 and 2).
• successive stages of treatment it is meant that the transition from the first stage to the second stage of treatment is done by gradually increasing the temperature to go from the first stage to the second without go through an intermediate temperature that would be lower than that of the first landing.
• the succession of the first two stages without falling to a temperature below the first step, for example up to room temperature, makes it possible to avoid excessive temperature gradients within the treated sample, and to avoid a heterogeneous dissolution of the ⁇ phase which could cause in some zones a grain enlargement. It is thus preferable to adopt a rise rate between the bearings that is sufficiently low ( ⁇ 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 in a cylindrical sample of 1000 cm 3 after a rise rate of 2 ° C / min from the first stage.
• any passage between the two stages at a temperature below the first stage may increase the time required for the homogenization of the temperature within the sample during the second stage, and may promote a heterogeneous dissolution of the phase ⁇ .
• a passage at a temperature below the first step is not excluded by the invention (FIG 3) if, in particular as a function of the dimensions of the treated part, the parameters of the second step are adjusted, possibly adding a intermediate bearing, so as to avoid the possible disadvantages just mentioned.
• the first treatment stage is carried out at a temperature of between approximately 900 and 1000 ° C. for a duration of at least 30 minutes and the second treatment stage is carried out at a temperature above the first plateau between 940 and 1020 ° C. 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 ranges of temperatures and durations thus defined make it possible to obtain a homogeneous microstructure with a suitable 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 phase fraction ⁇ of between 2% and 4%.
• the invention is based firstly on a synergistic effect between the two first stages, and optimized balancing between these two first stages makes it possible to best meet the desired objectives of the invention.
• the solvus temperature of the ⁇ phase depends directly on the niobium + tantalum content of the alloy.
• the amount of niobium and / or tantalum present in the composition of the alloy therefore has a direct influence on the temperature and duration of each stage.
• the optimal durations of the treatments also depend on the massiveness of the part to be treated, and can be determined by means of modelizations or experiments customary for the person skilled in the art.
• the first stage is preferably carried out at a temperature of between about 960.degree. ° C and 990 0 C for a time between about 45 minutes and 2 hours and the second step is preferably carried out at a temperature between about 990 ° C and 1010 ° C for a period between about 5 and 45 minutes.
• the first stage is preferably carried out at a temperature between about 920 0 C and 960 ° C for a time between about 45 minutes and 2 hours and the second step is preferably carried out at a temperature between about 960 ° C and 990 0 C for a period between about 5 and 45 minutes.
• the duration of treatment also depends on the massiveness of the workpiece.
• the temperatures at the treatment stages are generally kept substantially constant during the duration of the plateau.
• the rise rate from the first to the second step is preferably less than 4 ° C / min, to avoid excessive temperature gradients, especially in the case where large parts are processed.
• the rate of rise in temperature from the first to the second step is preferably between 1 ° C./min and 3 ° C./min.
• the invention applies to superalloys based on nickel, thus containing at least 50% of Ni, in which the sum Nb + Ta exceeds by weight 2.5%.
• the alloy is a nickel-based superalloy type 718 also called NC19FeNb (Standard AFNOR), containing by weight, between 50 and 55% of nickel, between 17 and 21% of chromium, less than 0, 08% of carbon, less than 0.35% of manganese, less than 0.35% of silicon, less than 1% of cobalt between 2.8 and 3.3% of molybdenum, at least one of the elements niobium or tantalum so 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 preparation.
• a phosphorus addition makes it possible to reinforce the resistance of the grain boundaries, particularly with regard to stresses such as creep and notched creep.
• the application of the invention to such an alloy with a phosphorus content of greater than 0.007% and less than 0.015% is of particular interest since the gain obtained in creep is then significantly greater. It is thus easy to improve creep lifetimes by a factor of 4 while maintaining the same grain size. This presence of phosphorus can also, for the same reasons, be recommended on the other alloy examples below.
• the alloy is a nickel-based superalloy type 725, containing by weight, between 55 and 61% of nickel, between 19 and 22.5% of chromium, between 7 and 9.5% of molybdenum, at least one of the elements niobium or tantalum so that the sum of niobium and tantalum is between 2.75 and 4% with Ta less than
• the alloy is a nickel-based superalloy type 718PLUS containing, by weight, between 12 and 20% of chromium, between 2 and 4% of 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% of cobalt, between 0.4 and 1, 4 % of titanium, between 0.6 and 2.6% of aluminum, between 6 and 14% of iron, less than 0.1% of carbon, less than 0.015% of boron, less than 0.03% of phosphorus on residual percentage being nickel and impurities resulting from the elaboration.
• the alloy is a nickel-based superalloy characterized by a niobium + tantalum content of greater than 2.5% and by the presence of an intergranular phase of Ni 3 Nb-Ta type ( ⁇ phase) between 800 ° C. and
• Ni 3 (Al-Ti) - ( ⁇ ') and / or Ni 3 Nb-Ta ( ⁇ ") type between 600 and 800 ° C.
• Ni 3 (Al-Ti) - ( ⁇ ') and / or Ni 3 Nb-Ta ( ⁇ ") type between 600 and 800 ° C.
• Ni 3 (Al-Ti) - ( ⁇ ') and / or Ni 3 Nb-Ta ( ⁇ ") type between 600 and 800 ° C.
• Ni 3 Nb-Ta Ni 3 Nb-Ta
• Ni 3 Nb-Ta delta type The greater dissolution of the intergranular phase of the Ni 3 Nb-Ta delta type then releases niobium ( ⁇ '-gene element) which is inserted in solid solution in the hardening phase ⁇ '- Ni 3 (Al, Ti) and hardens the latter.
• the treatment according to the invention may comprise a fourth stage which makes it possible to complete the precipitation of the gamma "(Ni 3 Nb-Ta- gamma") and / or gamma "(Ni 3 (Al-Ti) Y) curing phases at a temperature lower than the third landing.
• a fourth stage which makes it possible to complete the precipitation of the gamma "(Ni 3 Nb-Ta- gamma") and / or gamma "(Ni 3 (Al-Ti) Y) curing phases at a temperature lower than the third landing.
• a third level between 700 and 750 0 C from 4h to
• the treatment of the invention may also comprise at least one intermediate stage of short duration (maximum 1 h, see Figure 2) between the first bearing and the second bearing to facilitate the homogenization of the temperature in large parts during the rise in temperature between the two first stages.
• the content of (Ta + Nb) of the alloy is at least 2.5%, it is recommended that the Al content not exceed 3%, so as not to cause the precipitation of the ⁇ 'phase at the grain boundaries. Beyond 3% AI, the ⁇ 'phase tends to be stabilized at the expense of the ⁇ phase and the Nb is inserted into the ⁇ ' phase.
• the ratio (Nb + Ta + Ti) / Al is greater than or equal to 3.
• the first examples of implementation of the method according to the invention are applied to alloy products 718 obtained after thermomechanical treatment on an alloy obtained by conventional VIM + VAR + forging, but could also have been obtained by powder metallurgy, and typically intended for producing aerospace turbine disks.
• thermomechanical treatments see Table 2
• TTM thermomechanical treatments
• thermomechanical treatment range No. 1 is a rolling carried out according to different passes at a temperature above the solvus of the ⁇ phase of the alloy.
• the products formed according to the thermomechanical treatment range No. 1 are bars whose metallurgical structure is free of Delta phase (metallurgical state 2).
• Table 2 the samples F, K, L, N were made from bars obtained according to this first thermomechanical treatment range.
• thermomechanical treatment range N ° 2 is a range of conventional forging in two hot (by “hot” is meant a maintenance in the oven followed by a deformation, “two hot” means two stages of deformation, each being preceded by maintenance in the oven) at a temperature below the solvus of the ⁇ 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 No. 2 are pancakes (by “pancake” is meant a product having generally a disc or roller form resulting from deformation by forging), the metallurgical structure of which contains ⁇ phase. distributed heterogeneously at the grain boundaries (metallurgical state 1, see FIGS. 4 and 5). In Table 2, samples C, E and H were made from pancakes obtained according to this second thermomechanical treatment range.
• thermomechanical treatment range No. 3 is a conventional hot-dip matrix at a temperature below the solvus of the ⁇ phase of the alloy.
• the products formed according to the thermomechanical treatment range No. 3 are disk blanks whose metallurgical structure contains ⁇ phase distributed very heterogeneously grain boundaries (metallurgical state 1, see Figures 4 and 5).
• Table 2 the samples A, B, D, G, I, J, M, O and P were made from blanks of turbine disks obtained according to this third thermomechanical treatment range.
• TTH heat treatment
• the type of "a” treatment series consist of a so-called isothermal dissolution stage and two aging stages.
• the dissolution stage consisted, for samples A, B, C, D, F and P, in keeping the alloy at a constant temperature between 955 and 1010 ° C. for 40 to 90 minutes.
• the two levels of aging consisted of a plateau at 720 ° C. for 8 hours followed by controlled cooling at 50 ° C./h up to a plateau at 620 ° C. for 8 hours.
• the range of heat treatments of type "b" known under the name of "Direct Aged” range does not involve dissolution and consists only of two stages of aging according to treatments of type "a". Only sample E has undergone the "b" type range
• the heat treatment series of types "c" are in accordance with the invention and comprise two so-called dissolution stages, respectively indicated 1 st step and 2 nd step, and one or two aging stages, respectively indicated 3 th level and 4 th step.
• 1 solutionizing bearing consisted of holding the alloy at a constant temperature between 940 0 C and 980 0 C. for about 50 to 60 minutes.
• the 2nd solution-bearing consisted of a holding the alloy at a constant temperature between 980 ° C and 1005 0 C for approximately 15 to 40 minutes.
• the passage of 1 to 2 nd level was carried out by controlled heating at a rate of about 2 ° C / min.
• the 3 rd and 4 th aging stages were in accordance with the corresponding aging stages of the "a" type reference ranges except for samples H and J.
• sample H the temperature of the 3 rd stage of aging treatment was raised to 750 ° C instead of 720 0 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 aging stages such as those used in the field of nickel-based superalloys.
• Sample J for its part, underwent only one aging processing step at 720 ° C. for 10 hours.
• the aging treatment undergone by the sample J shows that the invention is also applicable when the alloy only undergoes a single stage of aging treatment.
• the type "d" 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 because of a second bearing carried out at a temperature that is too high or for a length that is too long. In fact, the conditions of the 2 nd stage lead to an excessive dissolution of the ⁇ phase, and the growth of the grains is no longer controlled, which causes uncontrolled and significant magnification of the grains during the second plateau for samples I and L.
• the "e” type heat treatment range includes a single solution solution at 1005 ° C for 15 minutes and two aging stages. Only the sample O was obtained according to this range of heat treatment which does not conform to the invention as explained below.
• Samples A to L and O were alloys of type 718 at 5.3% Nb and 40 ppm P.
• Sample N was a 718 alloy at 5.0% Nb and 40 ppm P
• Samples M and P were alloys of type 718 at 5.3% Nb and 80 ppm P.
• Table 2 summarizes the processing conditions of the different samples, and the ASTM grain sizes and surface ⁇ phase percentages visible on a micrograph.
• Table 3 summarizes the main mechanical properties of some of these same samples, namely:
• the number of cycles before failure during a fatigue test at 450 ° C comprising, in a sinusoidal cycle with a maximum stress of 1050 MPa, a frequency of 10 Hz and a load ratio R of 0.05;
• the grain size is defined according to the ASTM standard, and it is also specified, in cases where the grain size is relatively inhomogeneous, the maximum grain size (ALA).
• the alloy products 718 F, K, L, N have therefore been transformed according to the thermomechanical range No. 1 which does not allow to precipitate phase. ⁇ .
• the product F is a reference sample which after the thermomechanical range No. 1 has been treated according to a standard heat treatment range "a" of the alloy 718 (treatment comprising a single step of dissolution in solution subsolvus phase ⁇ ) .
• the product L was treated with two-stage solution dissolution but with a second bearing made at a temperature and time too high, outside the field of the invention for an alloy 718.
• Products K and N do not have the same niobium content, but both have undergone a heat treatment range "c" according to the invention.
• the alloy products 718 identified C, E and H have been 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 has been treated according to a standard "a" type heat treatment range of alloy 718 (treatment comprising a single solution dissolution stage).
• the product E is also a reference sample which after the thermomechanical range No. 2 has been treated according to the heat treatment range of type "b” and has therefore been directly aged after forging ("direct aged”), and has not therefore not undergoing dissolution treatment before aging.
• 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 alloy products 718 identified A, B, D, G, I, J, M, O and P have been transformed according to the thermomechanical range No. 3 which makes it possible to precipitate the ⁇ phase in a very heterogeneous manner.
• the products A, B and P were treated according to a standard treatment range of the alloy 718 (type "a" treatment comprising a single solution dissolution stage subsolvus).
• the product D was treated with a treatment comprising a single dissolution stage but at a higher temperature than the products A, B and P, that is to say at a temperature close to the solvus of the ⁇ phase.
• the product I was treated with two-stage solution dissolution but with a duration, for the second stage, too high with respect to the temperature.
• the heat treatment undergone by I is therefore outside the scope of the invention.
• thermomechanical treatment No. 3 After the thermomechanical treatment No. 3, the product G was treated with two-stage solution in the field of the invention (heat treatment "c").
• Product J was also treated with two-stage solution in the field of the invention, but was not treated with a fourth step.
• the product M has been treated with two-stage solution in the field of the invention, but has a phosphorus content equal to 0.008% which is twice as high as that of the products A-L and N-O.
• Product O has undergone a heat treatment "e” with a dissolution at a single stage; this treatment is outside the scope of the invention.
• the product P is a reference sample with a phosphorus content of 0.008%. It has been treated according to a standard treatment range of alloy 718 (type "a" treatment comprising a single solubilization step).
• Products A, B, C that have been treated with a standard subsolvus heat treatment have a fine grain microstructure (> 9 ASTM) but have a higher ⁇ (> 4.5%) phase fraction the phase fraction ⁇ sought preferably in the context of the invention.
• the mechanical properties obtained by these products constitute the reference to appreciate the properties in traction, fatigue and creep obtained on the thermomechanical ranges (TTM) 2 and 3.
• Product D has been processed at a higher temperature than products A, B, and C, it has ASTM grains and a ⁇ phase which is heterogeneously distributed ( ⁇ 2.5%) and is less than the fraction. phase ⁇ sought preferably in the context of the invention. It can be seen 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 the products A, B, and C. The considerable reduction in fatigue life is attributable to the presence of large ASTM grains which constitute the fatigue initiation sites.
• the product E which was directly aged after the thermomechanical treatment No. 2 has a very heterogeneous grain size (10 to 14 ASTM) and significant variations in the phase rate ⁇ , this rate being found in most areas of the room (particularly those biased in creep) greater than the desired phase fraction ⁇ .
• the tensile and fatigue properties of the product E are superior to those of the products A, B, C, notes that the creep lifetimes obtained with product E are lower 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 phase fraction ⁇ (2.9% and 3.5%) included in the invention.
• preferred phase fraction ⁇ interval namely 4% maximum and 2.5% minimum. It is found that the tensile properties are significantly higher than those of the products A, B, C and of the same level as those of the product E. It is also noted that the creep properties of the products G, H, M are significantly higher than those of products A, B, C, E while the grain size is similar in these products.
• the fine-grained microstructure of the products G, H, M makes it possible to preserve the fatigue properties obtained with the products A, B, C, E and the weaker phase fraction ⁇ of the products G, H, M makes it possible to improve the creep.
• the combination of a phosphorus addition and the treatment according to the invention therefore has a synergistic effect which is positive on the creep properties of the alloy obtained.
• the object of the invention is to preserve a residual phase fraction ⁇ (preferably greater than 2.5%) which makes it possible to maintain satisfactory ductility at high temperature.
• a too low ⁇ phase content has an effect on the damage and ductility in tension at high temperature (650 ° C. with a strain rate of 10 "5 s " 1 ).
• the product D with a phase content ⁇ close to 2% has a ductility (elongation at break of 7%) much lower than that of the product G (elongation at break of 27%) which comprises a fraction phase ⁇ close to 3%.
• This ductility decrease for product D results from intergranular damage caused by a fraction of phase ⁇ that is too weak and heterogeneously distributed.
• Figures 4 to 9 are micrographs representative of the microstructures:
• FIGS. 4 and 5 show 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 thermomechanical deformation under -solvus (thermomechanical range 2 or 3). It is a microstructure which has delta phase Ni 3 Nb- ⁇ and / or Ni 3 Ta- ⁇ at the grain boundaries, but not uniformly distributed between the grains.
• Figure 4 shows that the samples have a fine grain size of about 11 ASTM, with a heterogeneous distribution of the ⁇ phase (black spots at grain boundaries). After the thermomechanical deformation range, the ⁇ phase percentage is 2.8 to 6% and the grain size is 10 to 13 ASTM.
• Figure 5 shows the microstructure of the samples with a higher magnification and shows grains whose joints are largely completely free of phase ⁇ (this appears in white on this micrograph).
• sample B When a treatment is applied to a sample (sample B) comprising only a first dissolution stage at 970 ° C. for approximately 60 minutes, a phase percentage ⁇ of 4.7 to 5.5% is obtained. a grain size of 11 to 12 ASTM. The homogeneity of the sample is thus improved, but a significant phase fraction ⁇ of which we know (see sample B Tables 1 & 2) is very disadvantageous to creep resistance.
• a heat treatment comprising only one dissolution stage at 1005 ° C. for approximately 15 minutes, corresponding to the "second stage" of the sample.
• the invention provides (see FIGS. 6 and 7) a phase percentage ⁇ of 1.1 to 3.5%, and a grain size of 5 to 9 ASTM.
• the phase rate ⁇ is therefore reduced, which goes in the right direction for creep resistance, but a heterogeneous grain size distribution is observed. This is explained by a heterogeneous grain growth during this plateau resulting from a nonhomogeneous distribution of the ⁇ phase inherited from the initial microstructure.
• the distribution of the ⁇ phase is heterogeneous in the initial microstructure.
• some grains may have a significant amount of phase ⁇ at the grain boundaries in the initial microstructure, while other grains show little or no phase ⁇ at the grain boundaries (see Figure 5).
• the grains which are not surrounded by phase ⁇ or which have little phase ⁇ at the grain boundaries will increase uncontrollably up to a grain size of approximately 5-6 ASTM, while other grains surrounded by ⁇ phase will be thwarted and give rise to grain sizes close to 9 ASTM.
• This heterogeneity of grain size is evident in the micrographs of FIGS. 6 and 7. The presence, even very localized, of 5-6 ASTM grains considerably reduces fatigue life.
• samples G, H and M are subjected to a heat treatment according to the invention, namely a first step at 980 ° C. for 60 minutes and, immediately after, a heating according to a ramp from 2 ° C / min to a second plateau at 1005 ° C. for 15 min, a phase percentage ⁇ of 2.9 to 3.5% is obtained, with a grain size of 10 to 12 ASTM.
• phase ⁇ is now distributed evenly at the grain boundaries, which effectively prevents their growth.
• phase ⁇ precipitates which leaves the Nb and Ta elements available in dissolved form
• the reduced grain size the homogeneity of the distribution of the ⁇ phase at the grain boundaries and at a single level. adjusted well for the presence of this phase ⁇ , creep and tensile strength are improved.
• the fine grain size associated with the controlled dissolution of the ⁇ phase that makes it possible to achieve the objectives of the invention which are:
• the inventors have also carried out additional tests on 718Plus and 725 type alloy samples, and have thus been able to confirm that the invention applied to other nickel-based superalloys having a niobium and / or tantalum content greater than 2.5% significantly improved their creep strength and tensile strength.

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