RU2433197C2 - Heat-resistant nickel-based alloy, part manufacturing method, and turbomachine part - Google Patents

Abstract

FIELD: metallurgy. ^ SUBSTANCE: nickel-based alloy contains the following, wt %: chrome 11.5-13.5; cobalt 11.5-16.0; molybdenum - more than 3.4 to 5.0; tungsten 3.0-5.0; aluminium 2.2-3.2; titanium 3.5-5.0; niobium 0.5-2.0; hafnium 0.25-0.35; zirconium 0-0.07; carbon 0.015-0.030; boron 0.01-0.02; nickel is the rest. Turbomachine parts are made from alloy powder by means of powder metallurgy method. ^ EFFECT: alloy can be subject to subsolvus and hypersolvus treatment and has high creep strength at high temperatures. ^ 9 cl, 2 dwg, 4 tbl

Description

The invention relates to heat-resistant alloys, or superalloys based on nickel (Ni), intended, in particular, for the manufacture of disks (impellers) of a turbine or compressor of turbomachines using powder metallurgy methods. The turbomachines in question may be aircraft (turbojet engine, turboprop engine) or ground (gas turbine for energy production).

The compressor and turbine disks located respectively upstream and downstream of the combustion chamber of a turbojet engine are subjected during operation to mechanical stresses, similar to tensile, creep and fatigue, at temperatures that can reach 800 ° C. However, it is desirable that the service life of these discs reach many thousands of hours. Therefore, these disks must be made of an alloy having at high temperatures high tensile strength, very good creep resistance, and good crack propagation resistance.

Currently, such disks can be made of Ni-based alloys using powder metallurgy methods, and these methods limit the phenomena of chemical segregation and contribute to good microstructural uniformity of the alloy.

An example of a known nickel-based alloy is described in FR 2593830. This alloy is commercially available under the brand name No. 18.

The mentioned alloy example, as well as the alloys according to the invention, relates to two-phase alloys, which contain: a phase called a gamma phase and formed by a nickel-based solid solution, which makes up the matrix of metallurgical grains, and a phase called a gamma bar phase, whose structure is based on the ordered intermetallic compound Ni ₃ Al. The gamma-ray phase forms several sets of intragranular or intergranular precipitates that appear at various stages of thermomechanical processing of the alloy and which play different roles in the mechanical properties of the alloy.

It can be stated that the totality of intergranular precipitates limits the growth of gamma-matrix grains during recrystallization heat treatment. Thus, “playing” on the recrystallization heat treatment of the alloy, they control a set of intergranular precipitates and, consequently, the size of the mentioned grains. Depending on whether the maximum temperature reached during the mentioned heat treatment is higher (hypersolvus treatment) or lower (subsolus treatment) than the dissolution temperature (or solvus temperature) of gamma-ray phase intergranular precipitates, recrystallization ends with a large grain size (at hypersolvus treatment) or with a small grain size (with subsolvus processing).

Tensile strength is usually promoted by a decrease in grain size, while creep resistance is promoted by an increase in grain size. Thus, according to the application under consideration and the intended mechanical characteristics, two-phase alloys are thermomechanically processed in order to obtain either a fine-grained microstructure (small grains), i.e. having a grain size of the order of 5-15 μm (i.e., ASTM values from 12 to 9 according to standard of the American Society for Testing Materials), or a coarse-grained microstructure, that is, having a grain size of about 20-180 microns (i.e., ASTM values from 8 to 2).

In addition, the hardening of the grains is ensured by the presence of various sets of intragranular precipitates of the gamma-ray bar phase based on Ni ₃ Al, and it is generally believed that the high-temperature mechanical tensile strength of these alloys increases with an increase in the volume fraction of the gamma bar phase, and this fraction can reach 60%

Alloy No. 18, in which the volume fraction of the gamma-ray phase is approximately 55%, is subjected mainly to subsolus treatments, since it is desired to obtain a fine-grained microstructure. The fatigue and tensile strength of a given alloy usually have an advantage with respect to its creep resistance, since the operating temperature is often below 650 ° C, i.e. relatively moderate.

At temperatures above 650 ° C, a high creep resistance is required, and therefore, a coarse-grained microstructure (obtained by hypersolus treatment) will be more suitable. However, the implementation of hypersolvous processing of large-diameter disks of alloy No. 18 is very difficult or even impossible for industrial implementation due to too small a gap between the solvus temperature of the gamma-ray phase and the burn-out temperature (i.e., the melting start) of the alloy. This temperature range of re-dissolution of the gamma-ray phase (that is, for hyper-solus processing) is in fact too narrow (less than 30 ° C), which makes the practical industrial use of heat treatment with complete dissolution of the gamma-ray phase phase risky.

In addition, in the discs during rapid cooling (of the order of 100 ° C / min) following heat treatment with complete dissolution, high internal stresses arise that cause cracking (shrinkage during quenching).

The purpose of the invention is to provide Ni-based alloys for which it is possible to carry out not only subsolus processing, but also hypersolus processing on an industrial scale and which preferably have high temperature mechanical characteristics, in particular creep resistance, at least equivalent to the mechanical characteristics of alloy No. 18, and preferably exceeding them.

This goal is achieved in the invention using alloys that contain essentially (that is, in addition to possible impurities) the following elements in the indicated amounts, in mass percent:

- chromium (Cr): 11.5-13.5%;

- cobalt (Co): 11.5-16.0%;

- molybdenum (Mo): from more than 3.9 to 5.0%;

- tungsten (W): 3.0-5.0%;

- aluminum (Al): 2.2-3.2%;

- titanium (Ti): 3.5-5.0%;

- niobium (Nb): 0.5-2.0%;

- hafnium (Hf): 0.25-0.35%;

- zirconium (Zr): 0-0.07%;

- carbon (C): 0.015-0.030%;

- boron (B): 0.01-0.02%; and

- nickel (Ni): the rest is up to 100%.

In his studies leading to the invention, the applicant found that the problems encountered in the case of alloy No. 18 are partially related to the large volume fraction (55%) of the gamma-ray phase in this alloy.

In fact, on the one hand, the applicant found that such a large volume fraction tended to narrow the gap between the solvus temperature of the gamma-bar phase and the burn-out temperature of alloy No. 18, making this gap too narrow to carry out hypersolvus processing on an industrial scale.

On the other hand, the applicant has determined that the internal stresses arising in the part during rapid cooling following heat treatment with complete re-dissolution are partly the result of the separation of a large volume fraction of the gamma-ray phase.

Finally, the applicant established that, while maintaining the temperature above 650 ° C for sufficiently long periods of time, the elemental composition of alloy No. 18 caused the development of topologically compact phases, usually denoted by sigma and mu phases, detrimental to the heat resistance of the disk during operation.

Thus, the composition of the alloys according to the invention is selected in such a way as to cause the release of a limited volume fraction of gamma-ray phase.

Thus, although the alloys according to the invention will be less rich in gamma-bar phase than alloy No. 18, they have, contrary to all expectations, in their version with a fine-grained microstructure tensile characteristics and creep resistance exceeding these characteristics of this comparative alloy. It also seems that these alloys have crack propagation velocities under fatigue-creep equivalent or even lower than alloy No. 18.

For compressor disks or turbomachine turbines, high tensile strength is especially favorable for the tear resistance of said disks, which can occur suddenly during a random overspeed mode. Such high resistance is also an indicator of the good properties of multi-cycle fatigue and adequate service life.

In addition, a decrease in the volume fraction of the gamma-ray phase compared to alloy No. 18 favors the manufacture of disks having a coarse-grained microstructure and, therefore, high creep resistance at high temperature (i.e., at temperatures greater than or equal to 700 ° C). Such creep resistance, combined with very good mechanical tensile and crack propagation properties during fatigue-creep, allows the use of these disks at higher temperatures than in modern turbomachines, which allows to achieve better thermal efficiency and reduce the specific consumption of turbomachines.

The preparation of said coarse-grained microstructure is also facilitated by a convenient temperature interval between the temperature of the solvus of the gamma-bar phase and the temperature of the onset of melting of the alloy. Advantageously, the alloy compositions according to the invention are such that the size of said interval is greater than or equal to 35 ° C. This means that heat treatments above the solvus temperature can be carried out on an industrial scale without the risk of burning the alloy.

The ability to develop one or another of coarse-grained and fine-grained microstructures, as well as good mechanical properties corresponding to each of these two microstructures, is a definite advantage of the alloys according to the invention with respect to the alloys currently used, in particular, alloy No. 18.

In addition, this ability allows the manufacture of discs with a double structure. In fact, by performing heat treatment with a temperature gradient, a coarse-grained structure is developed in the peripheral zone of the disk, where operating temperatures are the highest and where creep plays a significant role in damage (destruction) of the material, and a fine-grained structure in the colder, central zone of the disk (close to the hub) where damage occurs essentially due to tensile forces and cyclic stresses.

Despite the fact that the concentration of aluminum is lower than the concentration of aluminum in alloy No. 18 (which is directly related to the lower volume fraction of the gamma-ray phase), the alloys according to the invention have a fairly low density (mass per unit volume), preferably less than or equal to 8, 3 kg / dm ³ , which allows you to limit the mass of the disk and the stress arising due to centrifugal force.

Finally, the elemental compositions of the alloys of the invention provide them with good microstructural stability with respect to the appearance of sigma and mu phases, which is slowed down by holding for 500 hours at 750 ° C.

In order to limit the risk of cracking during hardening, in particular during treatments above the solvus temperature of the gamma bar phase, alloy compositions of the invention have a limited volume fraction of gamma bar phase, and preferably less than or equal to 50%. However, sufficient gamma-ray phase must be present, therefore, the volume fraction of the gamma-ray phase is preferably in the range of 40% to 50%.

Advantageously, in order to obtain such a volume fraction of the gamma-ray phase in the alloys of the invention, the sum of the contents of Al, Ti and Nb in atomic percentages is greater than or equal to 10.5% and less than or equal to 13%, or 10.5% ≤Al + Ti + Nb≤13%.

Although the release of the gamma-ray phase in Ni-based alloys depends solely on the presence of Al in a sufficient concentration, the elements Ti and Nb, which, replacing Al, are part of this phase, are considered as elements favoring the formation of the gamma-ray phase in that the same amount, and are called gamma-ray generators. The volume fraction of the gamma-ray phase is, therefore, a function of the sum of the atomic concentrations of Al, Ti and Nb.

It should be noted that tantalum (Ta) is also a gamma-ray forming element, but that it was not left in the alloys of the invention. In fact, Ta is an element with a high atomic mass, which requires complex compositional adjustments to keep the alloy density below a reasonable limit (preferably less than or equal to 8.3 kg / dm ³ ). In addition, the cost of Ta is high, and its beneficial effect on crack resistance has not been clearly established. Finally, its strengthening effect on the gamma-ray phase was not greater than that of the elements Ti and Nb. It can even be stated that the strength of the alloys of the invention was at least equivalent to the strength of alloys containing Ta.

More preferably, the atomic percentages of Al, Ti and Nb in the alloys of the invention are such that the ratio between the sum of the contents of Ti and Nb and the Al content is greater than or equal to 0.9 and less than or equal to 1.1, i.e. 0.9 ≤ [(Ti + Nb) / Al] ≤1.1.

The Ti and Nb atoms, which replace Al in the gamma-ray phase based on Ni ₃ Al, strengthen it by mechanisms similar to the solidification solidification mechanisms. This hardening is the stronger, the greater the ratio of [(Ti + Nb) / Al]. Meanwhile, on the one hand, a coherent eta phase of Ni ₃ Ti stands out above a certain concentration of Ti in the form of elongated plates, which have a detrimental effect on the mechanical properties, in particular, on the ductility of the alloys containing it. On the other hand, the concentration of Nb should be limited, since an excessive content of Nb is harmful to the resistance to crack propagation in this type of alloys.

According to another aspect of the invention, the contents of W, Mo, Cr and Co in atomic percentages are such that the sum of the contents of W, Mo, Cr and Co is greater than or equal to 30% and less than or equal to 34%, and such that the sum of the contents of W and Mo is greater than or equal to 3% and less than or equal to 4.5%, i.e.: 30% ≤W + Mo + Cr + Co≤34% and 3% ≤W + Mo≤4.5%.

Elements that substantially replace Ni in a solid gamma solution are Cr, Co, Mo, and W.

Chromium (Cr) is necessary for the stability of the alloy to oxidation and corrosion and is involved, due to the effect on the solid solution, in the hardening of the gamma matrix.

Cobalt (Co) improves the resistance of these creep alloys at high temperature. In addition, an increase in the concentration of Co within the stability of the structure of the gamma phase allows one to lower the temperature of the solvus of the gamma-bar phase and thus facilitate the heat treatment with its partial or complete re-dissolution.

Molybdenum (Mo) and tungsten (W) give strong hardening to the gamma matrix due to the effect on the solid solution. However, these elements have high atomic masses, and their substitution of Ni (in particular, the substitution of Ni by W) is expressed in a noticeable increase in the density of the alloy.

Thus, the contents of Cr, Mo, Co and W in the alloys according to the invention must be carefully selected relative to one another in order to obtain the desired effects, in particular the optimum hardening of the gamma matrix, but without danger of causing premature phases of brittle intermetallic compounds sigma and mu phases. These phases, when they develop in excess, can actually entail a significant decrease in the ductility and mechanical strength of the alloys.

Finally, it is noted that elements called minor, such as C, B, and Zr, form segregates, mainly at grain boundaries, for example, in the form of borides or carbides. They also contribute to increasing the strength and ductility of alloys by changing the chemical composition of grain boundaries, and their absence will be harmful. However, the excessive content of these elements entails a decrease in the temperature of the onset of melting, as well as excessive precipitation of carbides and borides, which consume these elements of the alloy and which no longer participate in the hardening of the alloy. Thus, the concentrations of carbon, boron and zirconium are selected, in particular, with minimal non-zero contents of carbon and boron, so as to obtain optimal mechanical strength and ductility for the alloys of the invention at high temperature. Hafnium (Hf) is also present in moderation, as this element improves resistance to intergranular cracking at high temperature.

The invention also relates to a method for manufacturing a part, more specifically a turbomachine part, such as a disk (impeller) of a compressor or turbine, in which a blank of said part or the part itself is obtained from the alloy powder of the invention using the powder metallurgy method.

Advantageously, said billet or said part is subjected to recrystallization heat treatment at which the billet or part is brought either to a temperature which is lower than the solvus temperature of the gamma-bar phase of said alloy, or to a temperature which is higher than the solvus temperature of the gamma-bar phase of said alloy and lower than the temperature the beginning of the melting of said alloy, in order to thereby contribute to the development of a microstructure with a grain size suitable for loading conditions.

The invention, its applications and its advantages will be better understood when studying the following detailed description. This description is given with reference to the accompanying figures, in which:

Figure 1 is a photograph obtained by scanning electron microscopy and showing the microstructure of alloy A, described later, and

Figure 2 is a photograph obtained by scanning electron microscopy and showing the microstructure of the alloy C1, described later.

Parts obtained from the alloys of the invention are preferably fabricated using powder metallurgy methods.

For example, manufacturing a compressor disk or turbine using the powder metallurgy method comprises the following steps:

- manufacture of an ingot of the initial alloy by mixing and fusion of pure metal elements (up to possible impurities);

- remelting the ingot and spraying it with an inert gas, or remelting the ingot and centrifugally spraying it using a known rotating electrode method to produce a pre-fused powder;

- compaction of this pre-fused powder by hot isostatic pressing and / or extrusion;

- forming a blank of a disk by isothermal forging;

- heat treatment of the obtained workpiece and

- final machining of the disc. At the exit from the isothermal forging, various stages of heat treatment can be selected to obtain a microstructure that is better suited to the intended application. The processing temperature with the dissolution of the gamma-bar phase allows you to control the size of the metallurgical grain:

- in the case of processing at a temperature lower than the solvus temperature of the gamma-bar phase, a microstructure with small grains (from 5 to 15 μm) is obtained, and

- in the case of processing at a temperature in the range between the solvus temperature of the gamma-bar phase and the melting temperature of the alloy, a microstructure with large grains (more than 15 μm) is obtained. This last treatment can be carried out on an industrial scale only if the gap between the two above-mentioned temperatures, called the “dissolution interval”, is sufficiently significant: it is believed that for industrial alloys it should be more than 30 ° C, preferably more than 35 ° FROM.

The cooling rate that follows the dissolution treatment allows you to control the distribution of intragranular gamma-ray precipitates.

One or several holidays allows you to control the size of the tertiary precipitates of the gamma-ray phase and to relieve internal stresses that occur during hardening.

The nominal compositions of the two prior art alloys and the three alloys according to the invention, shown as examples, are presented in Table I, in which the content of elements in each alloy are indicated in atomic percent, and in Table II, in which these contents are indicated in mass percent. Alloys C1, C2 and C3 have a dissolution interval of more than 50 ° C and can thus be processed according to the two types of heat treatment presented above, which provides a wider range of microstructures.

Table I (contents are in atomic percent) Alloy With Cr Mo W Al Ti Nb Hf FROM AT Zr BUT 15.0 12.5 3.8 0 9.2 5.3 0 0.125 0,079 0,083 0,022 AT 12.9 18.1 2,4 1.3 4.6 4,5 0.4 0 0.190 0,077 0,027 C1 15.1 13.6 2.2 1.3 6.4 5,6 0.5 0,100 0.109 0,093 0 C2 15.4 14.1 2.5 1,5 6.0 5,0 1,0 0,093 0.128 0,080 0 C3 12.0 14.6 2.9 1,0 5.5 4.6 1,0 0,100 0,100 0,080 0,038

Table II (contents are indicated in mass percent) Alloy With Cr Mo W Al Ti Nb Hf FROM AT Zr BUT 15.9 11.7 6.6 0 4.4 4,5 0 0.400 0,017 0.016 0,036 AT 13.1 16,2 4.0 4.0 2.2 3,7 0.7 0 0,039 0.014 0,043 C1 15.4 12,2 3,7 4.0 3.0 4.6 0.8 0.310 0,023 0.018 0 C2 15,5 12.6 4.1 4.7 2,8 4.1 1,5 0.285 0,026 0.015 0 C3 12.15 13.0 4.8 3.15 2,55 3.8 1,6 0.310 0,021 0.015 0,060

Alloy A is alloy No. 18, and alloy B is marketed under the brand name Rene-88DT.

To conduct tests on these alloys, the parts were manufactured by powder metallurgy according to the following procedure:

- manufacture of ingots of the original alloy by mixing and fusion of pure metal elements;

- centrifugal atomization with rotating electrodes;

- compaction of pre-fused powders by hot extrusion;

- heat treatment, including subsolvus or hypersolvus processing.

In the case of the subsolvus treatment, the treatment with partial dissolution of the gamma-ray phase was carried out at a temperature lower than the solvus temperature (T _solvus ) of the gamma-ray phase (approximately T _solvus is 25 ° C). The cooling rate after dissolution was about 100 ° C / min. This treatment was followed by tempering for 24 hours at 750 ° C and cooling in air.

In the case of hypersolvus treatment, the treatment with complete dissolution of the gamma-bar phase was carried out at a temperature higher than the temperature of the solvus of the gamma-bar phase (approximately T _solvus + 15-20 ° C). The cooling rate after dissolution was about 140 ° C / min. This treatment was followed by tempering for 8 hours at 760 ° C and cooling in air.

Tables III and IV present several results of tensile, creep, and crack propagation tests for alloys subjected to subsolus treatment (Table III) and hypersolus treatment (Table IV), respectively.

Tensile tests were carried out in air at 650 ° C for subsolus treatment (Table III) and at 700 ° C for hypersolus treatment (Table IV), and R _m corresponds to the maximum stress measured during these tests.

Creep tests were carried out in air at 700 ° C under an initial stress of 550 MPa (650 MPa for alloy C1). The value of t of _0.2 % represents the exposure time in hours until plastic deformation of 0.2% is achieved.

The crack propagation tests were carried out in air and at 650 ° C. The loading cycle was as follows: lifting the load in 10 seconds, holding for 300 seconds at maximum load and unloading in 10 seconds at a load ratio (minimum load / maximum load) of 0.05. The value of V _f35 is the crack propagation rate measured at a delta K of 35 MPa · m ^1/2 .

Table III (* - under an initial voltage of 650 MPa) Alloy Tensile at 650 ° C R _m (MPa) Creep at 700 ° C under 550 MPa, t _0.2% (h) _Crack propagation at 650 ° C, V _f35 (m / cycle) BUT 1474 340 12 · 10 ^-5 AT 1445 610 3 · 10 ^-5 C1 1590 3000 * 2 · 10 ^-5 C2 1635 2300 3 · 10 ^-5 C3 1589 - -

Table IV (* - under an initial voltage of 650 MPa) Alloy Tensile at 700 ° C R _m (MPa) Creep at 700 ° C under 550 MPa, t _0.2% (h) _Crack propagation at 650 ° C, V _f35 (m / cycle) AT 1320 150 9 · 10 ^-6 C1 1440 1750 * 3 · 10 ^-6 C2 1428 > 3000 5 · 10 ^-6

The results in tables III and IV show that the alloys according to the invention can achieve a significant increase in high temperature mechanical properties (tensile and creep), while still maintaining crack propagation resistance, identical or better than that of known alloys.

Referring to Figs. 1 and 2, microstructural studies were carried out on sub-solus-treated alloys A and C1 in order to detect the appearance of topologically compact phases (i.e. brittle intermetallic compounds) after heat treatment, aging for 500 hours at 750 ° C. The observations were carried out by scanning electron microscopy with contrast in the reflected electron mode on samples not subjected to etching. In alloy A, aging under severe conditions for 500 hours at 750 ° C caused inter- and intragranular formation of phases enriched in heavy elements. These phases are contrasting (white edges) at the grain boundaries in Figure 1. These phases, when they are formed in excess, can entail a significant decrease in the ductility and strength of the alloys. Studies of alloy C1 subjected to the same treatment for 500 hours at 750 ° C show that these phases did not form during aging. Thus, the alloys of the invention were more stable than alloy A (No. 18) with respect to the formation of brittle intermetallic compounds, which are topologically compact phases.

Claims (9)

1. A heat-resistant alloy based on Nickel, essentially containing the following elements in the indicated amounts, wt.%:
Cr 11.5-13.5 With 11.5-16.0 Mo from more than 3.4 to 5.0 W 3.0-5.0 Al 2.2-3.2 Ti 3,5-5,0 Nb 0.5-2.0 Hf 0.25-0.35 Zr 0-0.07 FROM 0.015-0.030 AT 0.01-0.02 Ni the rest is up to 100 2. The alloy according to claim 1, having the following composition, wt.%:
Cr 13 With 12.15 Mo 4.8 W 3.15 Al 2,55 Ti 3.8 Nb 1,6 Hf 0.310 Zr 0,060 FROM 0,021 AT 0.015 Ni the rest is up to 100 3. The alloy powder according to claim 1 or 2. 4. A method of manufacturing a part in which a blank of said part or the part itself is obtained from an alloy powder according to claim 3 using the powder metallurgy method. 5. The method of manufacturing a part according to claim 4, wherein said preform or said part is subjected to recrystallization heat treatment, at which the preform or part is brought to a temperature that is higher than the solvus temperature of the gamma-ray phase of the alloy and lower than the melting temperature of the alloy. 6. The method for manufacturing a part according to claim 4, wherein said preform or said part is subjected to recrystallization heat treatment, at which the preform or part is brought to a temperature which is lower than the solvus temperature of the gamma-ray phase of said alloy. 7. Detail of a turbomachine obtained from an alloy according to claim 1 or 2. 8. A turbomachine part according to claim 7, having a coarse-grained structure in the zone in which it is exposed to the highest operating temperatures, and where creep plays an important role in damage to the part, and a fine-grained structure in the zone in which it is exposed to the lowest operating temperatures , and where the damage occurs, essentially due to tensile forces and cyclic loads. 9. A detail of the turbomachine of claim 8, which is a disk of a compressor or turbine.

Images