JP2018529022A - Nickel base alloy - Google Patents

Abstract

3.5-6.5 wt% chromium, 0.0-12.0 wt% cobalt, 4.5-11.5 wt% tungsten, 0.0-0.5 wt% molybdenum, 3. 5-7.0 mass% rhenium, 1.0-3.7 mass% ruthenium, 3.7-6.8 mass% aluminum, 5.0-9.0 mass% tantalum, 0.0- 0.5 mass% hafnium, 0.0-0.5 mass% niobium, 0.0-0.5 mass% titanium, 0.0-0.5 mass% vanadium, 0.0-0. 1 wt% silicon, 0.0-0.1 wt% yttrium, 0.0-0.1 wt% lanthanum, 0.0-0.1 wt% cerium, 0.0-0.003 wt% % Sulfur, 0.0-0.05% by weight manganese, 0.0-0.05% by weight zirconium, 0.0-0.005% by weight boron, Consists .0～0.01 wt% carbon, the balance nickel-based alloy composition consisting of nickel and unavoidable impurities.

Description

  The present invention relates to nickel-based single crystal superalloy compositions designed for high performance jet propulsion applications.

  Alloy (4th generation single crystal nickel base superalloy) shows a combination of creep resistance and oxidation resistance, but this combination is an alloy of the same grade as the alloy (4th generation single crystal nickel base superalloy) Or superior to alloys of the same grade as the alloy (4th generation single crystal nickel-base superalloy). Alloy density, cost, processing and long-term stability are also considered in the design of new alloys.

  Examples of typical compositions of fourth generation nickel based single crystal superalloys are listed in Table 1. These alloys can be used in the manufacture of rotating / fixed turbine blades used in gas turbine engines.

  The object of the present invention is to provide an alloy having comparable high temperature behavior or superior high temperature behavior compared to the fourth generation alloys listed in Table 1.

  According to the present invention, 3.5 to 6.5 wt% chromium, 0.0 to 12.0 wt% cobalt, 4.5 to 11.5 wt% tungsten, 0.0 to 0.5 wt% % Molybdenum, 3.5-7.0% by weight rhenium, 1.0-3.7% by weight ruthenium, 3.7-6.8% by weight aluminum, 5.0-9.0% by weight Tantalum, 0.0-0.5 wt% hafnium, 0.0-0.5 wt% niobium, 0.0-0.5 wt% titanium, 0.0-0.5 wt% vanadium, 0.0-0.1 wt% silicon, 0.0-0.1 wt% yttrium, 0.0-0.1 wt% lanthanum, 0.0-0.1 wt% cerium, 0.0. 0 to 0.003 mass% sulfur, 0.0 to 0.05 mass% manganese, 0.0 to 0.05 mass% zirconium, 0.0 to 0.005 The amount% boron, consisting of 0.0 to 0.01 wt% carbon, the balance nickel-based alloy composition consisting of nickel and unavoidable impurities is provided. This composition provides a good balance between cost, density, creep resistance, and oxidation resistance.

  In one embodiment, chromium with which a nickel base alloy composition is provided is 4.0-5.0% by mass%. Such alloys have good oxidation resistance and are particularly resistant to the formation of TCP.

  In one embodiment, the cobalt included in the nickel-based alloy composition is at least 0.1% by mass. This alloy results in an increase in creep resistance and a decrease in the γ 'solvus temperature, thereby increasing the solution window.

  In one embodiment, the cobalt with which a nickel base alloy composition is provided is 7.0-11.0% by mass%. In such alloys, resistance to creep deformation at the observed limit level of creep anisotropy (orientation dependence) is improved. Further, in this alloy, the processing is further simplified by reducing the γ ′ solvus temperature. By setting the maximum cobalt content to 10.9%, creep anisotropy is further limited.

  In one embodiment, tungsten with which a nickel base alloy composition is provided is 7.0-9.5% by mass%. This composition provides a compromise between low cost, light weight, and creep resistance.

  In one embodiment, the tungsten included in the nickel-based alloy composition is at least 7.1% by weight. Thereby, excellent creep resistance is achieved.

  In one embodiment, the aluminum provided in the nickel-based alloy composition is 3.7 to 6.6% by mass, preferably 5.1 to 6.6%, more preferably 5.5 to 6.6. %. This composition achieves excellent creep resistance and reduced density simultaneously with increased oxidation resistance.

In one embodiment, the tantalum included in the nickel-based alloy composition is 5.0% to 9.0% by mass. This provides a balance between creep resistance, manufacturability (based on the dissolution window) and density and / or suppresses the possibility of eta (ε) phase Ni ₃ Ta being formed. The tantalum included in this alloy is preferably 5.0 to 7.3%. This reduces the tendency of the ε phase to form, further reduces the cost and density of the alloy and increases the melting window.

  In one embodiment, molybdenum with which a nickel base alloy composition is provided is 0.1% or more by mass%. Such an alloy is advantageous in that the creep resistance is improved.

  In one embodiment, rhenium included in the nickel-based alloy composition is 4.5% to 6.0% by mass%, more preferably 5.3% to 6.0%. This composition provides a good balance between creep resistance, density, resistance to TCP phase formation, and cost.

  In one embodiment, the ruthenium included in the nickel-based alloy composition is 2.0 to 3.0% by mass. This composition provides a good balance between creep resistance and cost. Further, the ruthenium content is limited to the range of 2.1 to 2.9%, so that the compromise is further improved.

  In one embodiment, the hafnium with which the nickel base alloy composition is provided is 0.0% to 0.2% by mass. This is optimal for constraining inevitable impurities in the alloy, such as carbon.

In one embodiment, when the mass% of tantalum and aluminum contained in the alloy is W _Ta and W _Al , respectively, the nickel-based alloy composition satisfies the following formula.
33 ≦ W _Ta +5.1 W _Al ≦ 39
This is advantageous because an appropriate volume fraction γ ′ can be present.

In one embodiment, when the mass% of tantalum and aluminum contained in the alloy is W _Ta and W _Al , respectively, the nickel-based alloy composition satisfies the following formula.
2.2 ≦ 5.15W _Al −0.5W _Al ² −W _Ta
Preferably, the following formula is satisfied.
2.9 ≦ 5.15W _Al −0.5W _Al ² −W _Ta
This is advantageous because the alloy melting window can be adequate to allow the heat treatment process.

In one embodiment, when the mass% of ruthenium and rhenium contained in the alloy is W _Ru and W _Re , respectively, the nickel-based alloy composition satisfies the following formula.
4.5 ≧ W _Ru + 0.225W _Re
Preferably, the following formula is satisfied.
3.9 ≧ W _Ru + 0.225W _Re
This is advantageous because a relatively low cost alloy is obtained.

In one embodiment, when the mass% of rhenium and tungsten contained in the alloy is W _Re and W _W , the nickel-based alloy composition satisfies the following formula.
15.8 ≧ 1.13W _Re + W _W
Preferably, the following formula is satisfied.
14.4 ≧ 1.13 W _Re + W _W
This is advantageous because a relatively low density alloy can be obtained.

In one embodiment, when the mass% of rhenium, molybdenum, and tungsten contained in the alloy is W _Re , W _Mo , and W _W , respectively, the nickel-based alloy composition satisfies the following formula.
21.9 ≦ 2.92 W _Re + (W _W + W _Mo )
Preferably, the following formula is satisfied.
24.6 ≦ 2.92W _Re + (W _W + W _Mo )
This is advantageous because an alloy having excellent creep resistance can be obtained.

  In one embodiment, the sum of niobium element, titanium element and vanadium element in the nickel-base alloy composition is less than 1% by mass, and preferably 0.5% by mass or less. That is, these elements do not have a detrimental effect on the environmental resistance of the alloy.

  In one embodiment, the sum of niobium element, titanium element, vanadium element and tantalum element in the nickel-base alloy composition is 5.0 to 9.0% by mass, preferably 5.0 to 7.3% by mass. is there. Thereby, γ ′ having a preferable volume fraction and preferable APB energy are obtained.

  In one embodiment, the nickel-base alloy composition has γ ′ with a volume fraction of 60-70%.

  In one embodiment, a single crystal formed of any of the nickel-based alloy compositions of the above embodiments is obtained.

  In one embodiment, a turbine blade for a gas turbine engine formed of an alloy according to any of the above embodiments is obtained.

  In one embodiment, a gas turbine engine comprising the turbine blade of the above-described embodiment is obtained.

  The term “comprising” is used herein to indicate that the composition is 100% and that the percentage is 100% by eliminating the presence of additional ingredients.

  The present invention will be more fully described by way of example only with reference to the accompanying drawings.

FIG. 1 shows the distribution coefficient of the main component in the alloy design region. FIG. 2 is an isoline diagram showing the influence of aluminum and tantalum, which are γ′-forming elements, on the volume fraction of γ ′ in an alloy in the alloy design region. This isoline diagram is obtained by the phase equilibrium calculation performed at 900 ° C. FIG. 3 is an isoline diagram showing the influence of aluminum and tantalum elements on the antiphase boundary energy in an alloy at 900 ° C. in which the volume fraction of γ ′ is 60 to 70%. FIG. 4 is an isoline diagram showing the influence of aluminum and tantalum elements on the dissolution window in an alloy at 900 ° C. with a volume fraction of γ ′ of 60 to 70%. FIG. 5 is an isoline showing the effect of rhenium and ruthenium content on raw element costs in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum and having a volume fraction of γ ′ of 60 to 70%. FIG. FIG. 6 is an isoline diagram showing the influence of rhenium and tungsten on the density in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum and having a volume fraction of γ ′ of 60 to 70%. FIG. 7 is an isoline diagram showing the influence of rhenium and tungsten elements on creep resistance in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum and having a volume fraction of γ ′ of 60 to 70%. It is. This alloy contains 0% by weight of ruthenium. FIG. 8 is an isoline diagram showing the influence of rhenium and tungsten elements on creep resistance in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum and having a volume fraction of γ ′ of 60 to 70%. It is. This alloy contains 1% by weight of ruthenium. FIG. 9 is an isoline diagram showing the influence of rhenium and tungsten elements on creep resistance in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum and having a volume fraction of γ ′ of 60 to 70%. It is. This alloy contains 2% by weight of ruthenium. FIG. 10 is an isoline diagram showing the effect of rhenium and tungsten elements on creep resistance in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum and having a volume fraction of γ ′ of 60 to 70%. It is. This alloy contains 3% by weight of ruthenium. FIG. 11 shows that the chromium element and the tungsten element have a microstructural stability in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum, 1 to 3% by mass of ruthenium, and a volume fraction of γ ′ of 60 to 70%. It is an isoline diagram which shows the influence which acts on property. This alloy contains 4% by weight of rhenium. FIG. 12 shows that the chromium element and the tungsten element have a microstructural stability in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum, 1 to 3% by mass of ruthenium, and having a volume fraction of γ ′ of 60 to 70%. It is an isoline diagram which shows the influence which acts on property. This alloy contains 5% by weight of rhenium. FIG. 13 shows that the chromium element and the tungsten element have a microstructural stability in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum, 1 to 3% by mass of ruthenium, and a volume fraction of γ ′ of 60 to 70%. It is an isoline diagram which shows the influence which acts on property. This alloy contains 6% by weight of rhenium. FIG. 14 shows that the chromium element and the tungsten element have a microstructural stability in an alloy at 900 ° C. containing 5 to 9% by mass of tantalum, 1 to 3% by mass of ruthenium, and having a volume fraction of γ ′ of 60 to 70%. It is an isoline diagram which shows the influence which acts on property. This alloy contains 7% by weight of rhenium. FIG. 15 is an isoline diagram showing the effect of cobalt on the γ ′ solvus temperature in alloys with different ratios of aluminum to tantalum. This alloy is an alloy at 900 ° C. containing 5 to 9% by mass of tantalum and having a volume fraction of γ ′ of 60 to 70%. FIG. 16 shows the time to 1% creep strain in alloy ABD-2 (circle) of the present invention compared to 4th generation single crystal turbine blade alloy TMS-138A (triangle). FIG. 17 shows the time to break for alloy ABD-2 (circle) of the present invention compared to the fourth generation single crystal turbine blade alloy TMS-138A (triangle). FIG. 18 shows the measured mass change of the fourth generation single crystal turbine blade alloy TMS-138A (triangle) and the inventive alloy ABD-2 (circle) when air oxidized at 1000 ° C.

  Traditionally, nickel-base superalloys have been designed based on empirical principles. Thus, the chemical composition of nickel-base superalloys has been identified by time-consuming and expensive experimental development including small scale processing of limited amounts of material and subsequent characterization of behavior. Thereafter, an alloy composition that is found to exhibit the best or most desirable combination of properties is employed. The large number of alloy element groups that can achieve this combination indicates that these alloys are not fully optimized and that there is a high probability that there will be more improved alloys.

  In superalloys, chromium (Cr) and aluminum (Al) are generally added to provide oxidation resistance, and cobalt (Co) is added to improve resistance to sulfidation. Molybdenum (Mo), tungsten (W), Co, rhenium (Re), and optionally ruthenium (Ru) are introduced for creep resistance. These elements have the ratio of creep deformation. This is to inhibit the thermal activation process (for example, dislocation increase) to be determined. Aluminum (Al), tantalum (Ta), and titanium (Ti) are introduced to increase the static strength and the cyclic strength, because these elements are of the precipitation hardening phase gamma prime (γ ′). This is to promote formation. This precipitated phase is coherent with a face centered cubic (FCC) matrix phase called gamma (γ).

  In this specification, the model-based approach used to identify new grades of nickel-base superalloys is described in terms of “alloy design” (ABD) methods. This approach utilizes a computational material model framework for estimating design-related properties over a very broad compositional region. In principle, this alloy design tool makes it possible to solve the so-called inverse problem. That is, it is possible to identify an optimum alloy composition that most satisfies the specified design constraint.

  The first step in the design process is to define the element table and the upper and lower limits of compositional limitations associated with the element table. In the present invention, the composition limitation for each element when adding each element, which is called “alloy design region”, is considered. This composition limitation is detailed in Table 2.

  The second step is based on thermodynamic calculations to calculate the phase diagram and thermodynamic properties of a particular alloy composition. This is often called the CALPHAD method (CALculate PHAse Diagram). By performing these calculations at the working temperature of the new alloy (900 ° C.), information about the phase equilibrium (microstructure) can be obtained.

  The third stage involves identifying an alloy composition having the desired microstructure. In the case of a single crystal superalloy that requires excellent resistance to creep deformation, the creep rupture life becomes the longest when the volume fraction of the precipitation hardening phase γ ′ is 60 to 70%. In addition, since the γ / γ ′ lattice irregularity loses coherency, it is necessary to follow one of positive and negative values. Thus, the limit depends on the absolute value of that value. The lattice mismatch δ is defined as a mismatch between the γ phase and the γ ′ phase, and is obtained by the following equation.

Here, alpha _gamma and alpha _{gamma prime} is the lattice constant of the gamma phase and gamma prime phase.

  Alloys based on inadequate microstructure are also rejected by an estimate of sensitivity to the morphological close packed (TCP) phase. By using CALPHAD modeling in this calculation, the formation of harmful TCP phase sigma (σ), wrinkles and mu (μ) is predicted.

  Therefore, this model specifies all the compositions in the design region where the calculation result of the volume fraction of γ ′ is 60 to 70%. In these compositions, the lattice irregularity of γ ′ is less than a predetermined absolute value, and the total volume fraction of the TCP phase is less than a predetermined magnitude.

  In the fourth stage, the merit index is estimated for the identified alloy composition remaining in the data set. Examples of merit indices include creep merit index (which indicates the creep resistance of an alloy based only on average composition), antiphase boundary (APB) energy, density, cost, and dissolution window.

  In the fifth stage, the calculated merit index is compared with constraints on the desired behavior, and these design constraints are considered boundary conditions for the problem. All compositions that do not meet the boundary conditions are rejected. At this stage, the size of the test data set is very small.

  The final sixth stage involves analyzing the remaining composition data set. This analysis can be done in various ways. For one, the alloy having the maximum merit index may be classified through a database. The alloys having the maximum merit index include, for example, the lightest alloy, the alloy having the highest creep resistance, the alloy having the highest oxidation resistance, and the cheapest alloy. Alternatively, a database may be used to determine the relative trade-off in performance caused by different combinations of characteristics.

  Five examples of merit indexes will be described.

  The first merit index is a creep merit index. The most important observation is that time-dependent deformation (ie, creep) of single crystal superalloys occurs due to dislocation creep with initial activity limited to the γ phase. Accordingly, since the proportion of the γ ′ phase is increased, the dislocation segment is rapidly fixed to the γ / γ ′ interface. The rate limiting step is the departure of the dislocation trapped configuration from the γ / γ ′ interface. It relies on local chemistry that causes a significant effect of the alloy composition on creep properties.

The microstructure model based on physics is used for the accumulation rate of creep strain ε ^· when the load is uniaxial and along the <001> crystallographic direction. The collective equation is the following formula.

Here, ρ _m is the movable dislocation density, φ _p is the volume fraction of the γ ′ phase, and ω is the width of the matrix channel. The terms σ and Τ are working stress and temperature, respectively. The terms b and k are Burgers vector and Boltzmann constant, respectively. The term _KCF is a constraint coefficient.

Section K _CF indicates the proximity of the cubic particles within these alloys. Equation 3 shows a dislocation multiplication process that requires estimation of the multiplication parameter C and the initial dislocation density. The term D _eff is the effective diffusivity that controls the ascending process at the particle / matrix interface.

In the above description, the composition dependency is caused by two terms φ _p and D _eff . Therefore, assuming that the microstructure is constant (most of the microstructure is controlled by heat treatment), the dependence on the chemical composition is caused by D _eff since φ _p is fixed. It can be seen that for the purposes of the alloy design modeling described herein, it is not necessary to perform the full integration of Equations 2 and 3 for each prototype alloy composition. Instead, the primary merit index M _creep that needs to be maximized is used. M _creep is obtained by the following equation.

Here, x _i is the atomic fraction of the solute i in the alloy. D _i ^~ is an appropriate interdiffusion coefficient.

  The second merit index relates to antiphase boundary (APB) energy. Defect energy in the γ ′ phase, such as APB energy, has a significant effect on the deformation behavior of the nickel-base superalloy. It has been found that increasing APB energy improves mechanical properties including tensile strength and resistance to creep deformation. APB energy studies have been performed on a number of Ni-Al-X systems using density functional theory. This study calculated the effect of ternary elements on the APB energy of the γ 'phase and assumed a linear superposition of the effects of each ternary element addition when considering a complex multicomponent system. As a result, the following equation was derived.

_Here, representing the _{x Cr, x Mo, x W} , x Ta, respectively _{x Nb} and _{x Ti,} Cr in γ'-phase, Mo, W, Ta, atomic percent concentration of Nb and Ti. The composition in the γ ′ phase can be obtained by phase equilibrium calculation.

The third merit index is density. The density ρ was calculated using simple rules and correction factors for the mixture. Here, ρ _i is the density of a given element, and x _i is the atomic fraction of the alloy element.

The fourth merit index is cost. To estimate the cost of each alloy, a simple rule for the mixture was applied. Here, the cost of each alloy was obtained by multiplying the mass fraction x _i of the alloy element by the current raw material cost c _i of the alloy element (2015).

  This estimate assumes that the processing costs are the same for all alloys. That is, the product yield is not affected by the composition.

  The fifth merit index is the dissolution window. By performing thermodynamic modeling (CALPHAD) calculations over a range of temperatures, the melting window can be calculated for each alloy. This value (measured in degrees Celsius) can be used to determine whether a given alloy is suitable for conventional manufacturing processes used to manufacture single crystal turbine blades. In general, in order to enable melt heat treatment, the melting window needs to be higher than 50 degrees. The melt heat treatment is performed in the single layer region, but at this point, the alloy exists only in the γ phase region. This melt heat treatment is necessary to homogenize the composition of the cast alloy that can be highly segregated. In order to determine the melt heat treatment window, it is necessary to determine the phase equilibrium (more specifically, the phase transition) over a certain temperature range. The temperature at which complete dissolution of the γ ′ phase (known as the γ ′ solvus temperature) needs to be known, as is the solidus temperature. The difference between the solidus temperature and the γ 'solvus temperature gives a dissolution window. Therefore, the dissolution window index is calculated as the difference between the solidus temperature and the γ 'solvus temperature.

  The alloy composition of the present invention was specified using the ABD method described above. The design intent of this alloy was to specify the composition of the fourth generation single crystal nickel-base superalloy. This fourth generation single crystal nickel-base superalloy has a combination of creep and oxidation resistance that is comparable to or superior to equivalent grade alloys. Alloy density, cost, processing and long-term stability are also considered in the design of new alloys.

  The material properties of commercially used 4th generation single crystal turbine blade alloys are listed in Table 3. This material property was determined using the ABD method. New alloy designs have been considered in light of the expected properties listed for these alloys. Table 3 also shows the calculated material properties for alloy ABD-2. Alloy ABD-2 is an alloy according to the present invention and has the nominal composition shown in Table 4.

In order to maximize creep resistance, it was necessary to optimize the microstructure of the alloy. The microstructure consists mainly austenitic face-centered cubic (FCC) gamma phase (gamma) and regularized the L1 ₂ precipitation phase (gamma prime). Since it is known that this microstructure provides the maximum level of creep resistance for single crystal blade alloys, it is generally considered optimal that the volume fraction of the γ 'phase is 60-70%. It is. The goal of this alloy was to set the volume fraction of γ ′ to 60-70%. However, the alloys of the present invention can deviate from this goal.

  As shown in FIG. 1, the distribution coefficient of each element included in the alloy design region was obtained by phase equilibrium calculation performed at 900 ° C. When the distribution coefficient is 1, it means that the element is distributed with the same priority as the γ phase or the γ ′ phase. When the distribution coefficient is less than 1, it indicates that the element has a priority with respect to the γ ′ phase, and as the value of the distribution coefficient approaches 0, the priority increases. The larger the value of the partition coefficient, the more the element is preferentially present in the γ phase. The distribution coefficients of aluminum and tantalum indicate that these elements are strong γ 'forming elements. It is preferable that the chromium element, the cobalt element, the rhenium element, the ruthenium element, and the tungsten element are distributed in the γ phase. Among the elements considered within the alloy design area, aluminum and tantalum are most strongly distributed in the γ ′ phase. Accordingly, the levels of aluminum and tantalum were controlled to produce the desired volume fraction γ '.

  FIG. 2 shows the effect of the elements (mainly aluminum and tantalum) added to form the γ ′ phase on the proportion of γ ′ phase in the alloy at a certain operating temperature (900 ° C. in this case). In order to design an alloy, an alloy composition in which the volume fraction of γ ′ is 60 to 70% was studied. Therefore, 3.7 to 6.8% by mass of aluminum was necessary.

  According to the following equation, the γ ′ volume fraction changes as the contents of aluminum and tantalum change.

Here, f (γ ′) is a numerical value in an alloy having γ ′ in a desired ratio (in this case, 0.6 to 0.7). This numerical value is a value in the range of 33-39. W _Ta and W _Al are mass percentages of tantalum and aluminum contained in the alloy, respectively.

In order to increase the antiphase boundary (APB) energy of the γ 'phase, optimization of the aluminum and tantalum levels was also necessary. APB energy strongly depends on the chemical nature of the γ 'phase. FIG. 3 shows the effect of aluminum and tantalum on APB energy. The composition which has APB energy (-270mJ / m < ² >) or more which the present 4th generation single crystal alloy has was specified. By making the level of tantalum in the alloy higher than 5.0% by mass (where Al = 3.7 to 6.8% by mass, that is, the volume of γ ′ is 60 to 70%), an acceptable high APB Modeling calculations have shown that an alloy with energy and very good creep resistance is produced, and at the same time γ 'with a sufficiently high volume fraction is produced.

In order to achieve the desired γ 'volume fraction (60-70%) with a minimum concentration of tantalum, the Al addition should be limited to a maximum of 6.6% by weight (FIG. 2). Therefore, in order to achieve both the desired γ ′ volume fraction and high APB energy, the Al concentration is preferably 3.7 to 6.6 mass%. The maximum content of tantalum will be described below with reference to FIG. 4. The range of tantalum is 5.0 to 9.0% by mass, preferably 5.0 to 7.3% by mass. This is due to the preferred combination of APB energy and dissolution window (discussed below). That is, the preferred minimum level of tantalum ensures a higher APB energy for a given amount of aluminum and a level of at least 270 mJ / m ² within the range of aluminum in the alloy. It can be seen from FIG. 2 that the aluminum concentration, particularly for higher lower levels of tantalum, is 5.1% by mass or more, preferably 5.5% by mass or more. Thereby, γ ′ having a desired volume fraction is obtained. Accordingly, the range of the ratio of aluminum to tantalum is preferably 0.41 (Al = 3.7 mass%, Ta = 9.0 mass%) to 1.36 (Al = 6.8 mass%) in mass%. Ta = 5.0% by mass), more preferably 0.70 (Al = 5.1% by mass, Ta = 7.3% by mass) to 1.32 (Al = 6.6% by mass, Ta = 5.0). % By mass), and even more preferably 0.75 (Al = 5.5% by mass, Ta = 7.3% by mass) to 1.32 (Al = 6.6% by mass, Ta = 5.0% by mass). is there.

  Niobium element, titanium element, and vanadium element behave similarly to tantalum. That is, these elements are gamma prime forming elements that increase antiphase boundary energy. These elements can be added to the alloy as necessary. This provides an advantage compared to adding tantalum, which may include the ability to reduce cost and density. On the other hand, these elements can adversely affect the environmental resistance of the alloy, so their addition must be limited. Therefore, each of these elements can be contained only up to 0.5% by mass. These elements are preferably used instead of tantalum. That is, the total of the element group consisting of niobium, titanium, vanadium, and tantalum is preferably limited to 5.0 to 9.0% by mass, more preferably 5.0 to 7.3% by mass. . These ranges are preferred ranges for tantalum. Regardless of this, in one embodiment, in order to avoid a reduction in the environmental resistance of the alloy, the total of the element group consisting of niobium, titanium and vanadium is preferably less than 1.0% by weight, preferably less than 0.1%. It is limited to less than 5% by mass.

  By adjusting the balance between aluminum and tantalum, it is possible to balance γ ′ having a desired target volume fraction and sufficiently high APB energy. However, alloy processing must also be considered. One consideration is the dissolution window. There should be a window of sufficient temperature range below the melting temperature of the alloy and where only the γ phase is stable. Since the dissolution window depends on the dissolution of the γ ′ phase, the dissolution window is strongly influenced by the γ ′ chemistry, ie, the aluminum and tantalum content. This melt heat treatment is used to remove rich residual fine segregation and eutectic mixture in γ ′. This residual fine segregation and eutectic mixture can occur in the casting process used to produce single crystal alloys. In order to enable conventional processing methods, the melting window is preferably higher than 50 ° C. FIG. 4 shows the size (° C.) of the dissolution window when the volume percentage of γ ′ is 60 to 70% and the mass% of Al and the mass% of Ta are changed. FIG. 4 shows that an appropriate melting window of the alloy is secured by limiting the tantalum content to 9.0 mass%. The content of tantalum is preferably limited to 7.3% by mass. This produces an alloy with a melting window above 60 ° C., further improving the processing of the alloy.

  According to the following equation, the melting window changes with changing contents of aluminum and tantalum.

Here, in order to manufacture an alloy having a melting window of 50 ° C. or higher, the value of f (T _sol. ) Is set to be larger than 2.2. Preferably, the value of f (T _sol. ) Is greater than 2.9 in order to produce an alloy having a melting window with a temperature above 60 ° C.

In an alloy that satisfies the above requirements (γ ′ having a volume fraction of 60 to 70%, APB energy exceeding 270 mJ / m ² , melting window exceeding 50 ° C.), it is hardly volatile with respect to creep resistance and oxidation performance. The level of refractory elements was measured. In the fourth generation single crystal turbine blade, the creep performance is greatly imparted by adding ruthenium element, rhenium element and tungsten element. This will be described later with reference to FIGS. On the other hand, rhenium element and ruthenium element strongly influence the cost (FIG. 5). Tungsten and rhenium elements greatly increase the alloy density (FIG. 6). In addition, elements such as rhenium, tungsten, chromium (chromium is added for oxidation resistance) need to be properly balanced, but this tends to form a harmful TCP phase, This is in order to balance creep resistance and oxidation resistance without producing a microstructurally unstable alloy (FIGS. 11 to 14). Therefore, a complex balance of trade-offs in cost, density, creep resistance, oxidation resistance and microstructure stability needs to be managed. The process of optimizing these tradeoffs will be described later with reference to FIGS.

  Ruthenium and rhenium elements are currently (2015) at a considerable raw material cost. Thus, to optimize the alloy design, the ruthenium and rhenium levels that best manage the cost / creep resistance tradeoff in the present invention are selected. The isoline diagram in FIG. 5 shows the effect of rhenium and ruthenium levels on the alloy cost in an alloy at 900 ° C. with a γ ′ volume fraction of 60-70%. FIG. 5 shows that ruthenium has the strongest influence on the alloy cost. Therefore, the ruthenium content in the alloy is limited to 3.7% by mass. Thereby, the cost concerning this invention becomes below the cost concerning the 4th generation alloy of the present grade. Preferably, the ruthenium content is limited to 3.0% by mass or less. This ensures an optimal balance between cost and creep resistance.

  In order to limit the cost of the alloy, ruthenium and rhenium are preferably added according to the following formula:

Here, in order to produce an alloy having a cost of 300 $ / lb or less, the value of f (Cost) is set to 4.5 or less. W _Ru and W _Re are mass% of ruthenium and rhenium contained in the alloy, respectively. Preferably, the value of f (Cost) is set to 3.9 or less in order to produce a low-cost alloy of 260 $ / lb or less.

  To design an alloy with excellent resistance to creep deformation, the addition of tungsten, rhenium and ruthenium elements is optimized. Creep resistance was determined using a creep merit index model. Maximizing the creep merit index is desirable because it is associated with improved creep resistance. The influence of tungsten, rhenium and ruthenium on creep resistance is shown in FIGS. This shows that creep resistance is improved by increasing the levels of tungsten, rhenium and ruthenium. On the other hand, the desired amounts of tungsten and rhenium have a strong influence on the alloy density (FIG. 6). Calculations for creating the graphs of FIGS. 6 to 10 are performed so that the volume fraction of γ ′ is 60 to 70% at 900 ° C. Therefore, it is necessary to balance the trade-off between creep resistance and alloy density.

  In order to limit the alloy density, tungsten and rhenium are preferably added according to the following formula:

Here, in order to manufacture an alloy having a density of 9.0 g / cm ³ or less, the value of f (Density) is set to 15.8 or less. W _W is mass% of tungsten contained in the alloy. Preferably, in order to produce an alloy having a density of 8.9 g / cm ³ or less, the value of f (Density) is set to 14.4 or less.

The current fourth generation single crystal alloy has a creep merit index of 12 × 10 ⁻¹⁵ m ⁻² s or more (see Table 3). This level of creep resistance is desirably achieved in combination with a low density of less than 9.0 g / cm ³ , preferably a low density of less than 8.9 g / cm ³ . The isolines in FIG. 6 showing the effect of rhenium and tungsten on the density overlap with the effect of rhenium, tungsten and ruthenium on the creep merit index as broken lines (FIGS. 7-10).

In order to set the minimum creep merit index to 12 × 10 ⁻¹⁵ m ⁻² s, the ruthenium contained in the alloy is at least 1.0 mass%. In order to obtain even higher creep resistance, the ruthenium content is preferably 2.0% by mass or more, more preferably 2.1% by mass or more. Ruthenium is preferably limited to 3.0% by mass. This provides a favorable balance between cost and creep resistance. More preferably, the maximum level of ruthenium is 2.9% by mass. This provides the benefits of a high creep merit index while further reducing costs. By limiting the tungsten content to 11.5 mass% or less, the alloy density can be reduced to 9.0 g / cm ³ or less. The tungsten content is preferably limited to 9.5% by mass. This produces an even lower density alloy (FIGS. 6 and 10). Microstructural stability is also ensured by making tungsten low (FIGS. 11-14).

  7 to 10 show that the creep merit index is increased by setting the minimum content of rhenium to 3.5% by mass or more. The rhenium content is preferably higher than 4.5% by mass. This produces an alloy with a better balance between density (FIG. 6) and creep resistance (FIG. 10). More preferably, the rhenium contained in the alloy is at least 5.3% by weight. This composition produces an alloy with even better balance between creep resistance and density. In such alloys, the level of ruthenium that can be required is low, so the cost can also be reduced (FIG. 9).

  Molybdenum behaves similarly to tungsten. That is, this slow diffusion element can improve the creep resistance. Accordingly, it is preferred that the molybdenum be present in an amount of at least 0.1% by weight. However, molybdenum increases the tendency to form harmful TCP phases in the alloy. Therefore, it is necessary to control the addition of molybdenum. Therefore, molybdenum is limited to 0.5 mass% or less.

  7 to 10 and the knowledge that tungsten can be replaced with molybdenum, it can be seen that when tungsten, rhenium, and molybdenum are added according to the following formula, a good level of creep resistance can be obtained.

Here, the numerical value of f (Creep) is 21.9 or more. W _Mo is the mass% of molybdenum contained in the alloy. As a result, an alloy with a creep merit index calculation result of 12 × 10 ⁻¹⁵ m ⁻² s or more is produced. Preferably, the value of f (Creep) is greater than 24.6. This results in an alloy having better creep resistance, as well as a lower level of ruthenium that provides equivalent creep resistance, i.e., reduces cost.

  The addition of cobalt is optional. However, model calculations have shown that cobalt increases the creep merit index. It is also known that the addition of cobalt reduces the stacking fault energy in the gamma matrix, which also improves creep resistance. Furthermore, the addition of cobalt can lower the γ ′ solvus temperature, thereby improving the ease of processing and contributing to an increase in the dissolution window. Therefore, it is desirable that cobalt is present at least 0.1% by mass. FIG. 15 shows that the γ ′ solvus temperature is decreased by adding cobalt when the ratio of aluminum to tantalum is 0.75 to 1.36, which is the most preferable range. The lower limit of cobalt is preferably 7.0% by mass. Thereby, an alloy having excellent creep resistance and a reduced γ ′ solvus temperature is produced. This is beneficial in the heat treatment process. On the other hand, if the level of cobalt is increased, the creep anisotropy of the alloy increases, particularly in the primary creep, so it is necessary to limit the addition of cobalt. As a result, the creep rate strongly depends on the orientation of the single crystal. In order to control the amount of creep anisotropy to an acceptable level, the upper limit of cobalt is set to 12.0% by mass. The upper limit of cobalt is preferably 11.0% by mass. Thereby, creep anisotropy is further reduced. More preferably, the upper limit of cobalt is 10.9 mass%. This further reduces the possibility of creep anisotropy.

  In order to maintain the resistance to creep over a long period of time, it is necessary to add rhenium, tungsten and ruthenium which are slow diffusion elements. Chromium addition is also necessary to improve resistance to oxidation / corrosion damage. On the other hand, it has been clarified that the addition of tungsten, rhenium and chromium at a high level increases the tendency to form undesirable TCP phases, mainly σ, Ρ and μ phases. FIGS. 11 to 14 show the effect of addition of chromium, tungsten and rhenium on the total proportion of the TCP phase (σ + μ + Ρ). By controlling the addition of these elements, the level of the TCP phase is preferably set to be equal to or lower than the level of the TCP phase of the current fourth generation superalloy (Table 3).

  In order to achieve superior oxidation resistance compared to current fourth generation single crystal alloys containing Cr in the range of 2.0-3.2 wt%, the minimum chromium content in the present invention is: It is 3.5% by mass or more, preferably 4.0% by mass or more. That is, on the premise of achieving oxidation resistance superior to that of current fourth generation alloys, higher mass% chromium is included than these alloys. In order to weaken the tendency for harmful TCP phases to form in the alloy, the chromium content is limited to 6.5% by weight (FIGS. 11-14). The chromium content in the alloy is preferably limited to 5.0% by weight. This produces an alloy with the best balance between oxidation resistance and microstructure stability. The rhenium content in the alloys in the present invention is limited to 7.0% by weight or less (to ensure acceptable microstructure stability) (FIG. 14). By setting the rhenium level to 4.5-6.0% by mass, the density, creep resistance and microstructure stability are well balanced, so the rhenium content should be limited to 6.0% by mass or less. Is more preferable. Based on the level of rhenium when the microstructure stability, density and creep resistance are acceptablely balanced, the minimum level of tungsten required for the present invention is 4.5% by weight or more. This balances creep resistance (FIGS. 7-10), cost and microstructure stability (FIGS. 11-14). In order to achieve high creep resistance (FIGS. 7-10), the minimum level of tungsten is preferably 7.0% by weight and desirably at least 7.1% by weight.

  When producing an alloy, it is beneficial for the alloy to be substantially free of inevitable impurities. The impurities may include carbon (C), boron (B), sulfur (S), zirconium (Zr) and manganese (Mn) elements. If the concentration of carbon is maintained below 100 PPM (mass basis), no undesirable carbide phase is formed. The boron content is preferably limited to 50 PPM or less (mass basis) to prevent the formation of undesirable boride phases. The carbide phase and boride phase constrain elements such as tungsten and tantalum added to give strength to the γ phase and γ ′ phase. Therefore, if carbon and boron are present in a large amount, mechanical properties including creep resistance deteriorate. It is preferred that the sulfur (S) element and the zirconium (Zr) element be maintained at less than 30 PPM and less than 500 PPM (mass basis), respectively. Manganese (Mn) is an unavoidable impurity and is preferably limited to 0.05% by mass (500 PPM on a mass basis). When sulfur is present in an amount of more than 0.003% by mass, the alloy becomes brittle and segregates at the alloy / oxide interface formed during oxidation. This segregation can increase the peeling of the protective oxide scale. Zirconium and manganese can cause casting defects, such as segregation, in the casting process. Therefore, it is necessary to control the levels of zirconium and manganese. When the concentration of these inevitable impurities exceeds a predetermined level, problems surrounding the product yield occur, and deterioration of the material properties of the alloy is expected.

  The addition of hafnium (Hf) up to 0.5% by weight, more preferably up to 0.2% by weight, is beneficial for binding inevitable impurities in the alloy, especially carbon. Because hafnium is a strong carbide former, the addition of this element is beneficial in constraining residual carbon impurities that may be included in the alloy. Also, the addition of hafnium can result in further grain boundary strengthening, which is beneficial when introducing low-angle grain boundaries into the alloy.

So-called “reactive elements” (silicon (Si), yttrium (Y), lanthanum (La) and cerium (Ce)) are added at a level of up to 0.1% by mass. This is beneficial for improving the adhesion of a protective oxide layer such as Al ₂ O ₃ . These reactive elements can “clean up” harmful elements such as sulfur. Sulfur segregates at the alloy oxide interface, weakens the bond between the oxide and the substrate, and causes oxide separation. In particular, it has been shown that the addition of silicon at levels up to 0.1% by weight to nickel-base superalloys is beneficial for oxidation properties. In particular, silicon segregates at the alloy / oxide interface and improves the bond strength of the oxide to the substrate. Thereby, exfoliation of the oxide is suppressed, and as a result, the oxidation resistance is improved.

  Based on the description of the invention in this section, a wide range and preferred range of each element addition was defined. These ranges are listed in Table 4. The composition of the example—alloy ABD-2— was selected from the preferred composition range, the composition of which is specified in Table 4. Alloy ABD-2 has been found to follow standard methods used in the manufacture of single crystal turbine blade components. In this manufacturing method, preparation of an alloy having the composition of ABD-2, preparation of a mold for casting the alloy using the investment casting method, directionality in which a “grain selector” for manufacturing a single crystal alloy is used Includes casting of the alloy using solidification techniques followed by multi-step heat treatment of single crystal casting.

  Experiments on alloy ABD-2 verified the important material properties that are the main objective of the alloys of the present invention. The material properties were verified mainly for sufficient creep resistance and improved oxidation performance compared to current single crystal alloys used for IGT applications. The behavior of alloy ABD-2 was compared to that of alloy TMS-138-A tested under the same experimental conditions.

  Single crystal castings of alloy ABD-2 with nominal composition according to Table 4 were produced using conventional methods of producing single crystal compositions. The casting was in the form of a cylindrical bar having a diameter of 10 mm and a length of 160 mm. The cast rod was confirmed to be a single crystal oriented within 10 ° from the <001> direction.

  The cast material was subjected to a series of additional heat treatments to produce the desired γ / γ ′ microstructure. When a melt heat treatment was performed at 1325 ° C. for 6 hours, it was revealed that the remaining fine segregation and eutectic mixture were removed. It has been found that the heat treatment window of the alloy is sufficient to avoid initial melting during the melt heat treatment. After the melt heat treatment, the alloy was subjected to two-stage aging heat treatment. The aging heat treatment was performed at 1120 ° C. for 2 hours in the first stage and at 870 ° C. for 16 hours in the second stage.

  A creep test piece having a gauge length of 20 mm and a diameter of 4 mm was machined from a completely heat-treated single crystal rod. The specimen was oriented within 10 ° from the <001> direction. The creep performance of the ABD-2 alloy was evaluated using an experimental temperature range of 800-1100 ° C. Repeated oxidation tests were performed on fully heat treated materials. The repeated oxidation test was performed at 1000 ° C. in a 2 hour cycle over 50 hours.

  In order to compare the creep resistance of alloy ABD-2 with alloy TMS-138A, a Larson-Miller diagram was used. In FIG. 16, the time to 1% creep strain in both alloys is compared and shown. The time to 1% strain is important because most gas turbine components are manufactured with strict accuracy to achieve maximum engine performance. Even at low distortion levels—on the order of a few percent—parts are often replaced. It can be seen that the time to 1% creep strain in alloy ABD-2 is comparable to TMS-138A. In FIG. 17, the time to creep rupture in both alloys is compared and shown. FIG. 17 shows that the fracture life of alloy ABD-2 is comparable to TMS-138A.

  The oxidation performance of alloy ABD-2 and alloy TMS-138A was also compared. As turbine temperatures continue to rise (engine thermal efficiency improves), component failures due to corrosion wear such as oxidation become more common. Therefore, by improving the oxidation resistance / corrosion resistance, the life of the parts can be greatly advanced. Alloy ABD-2 was designed to have improved oxidation behavior compared to current second generation alloys. The results of repeated oxidation of ABD-2 and TMS-138A are shown in FIG. The suppression of mass increase over time is evidence of improved oxidation performance. This improvement in oxidation performance is due to the limited entry of oxygen into the substrate material by the formation of a protective oxide scale. Compared with TMS-138A, the ABD-2 alloy shows that the increase in mass with respect to time is greatly suppressed, indicating that the oxidation performance is improved.

  Overall, alloy ABD-2 exhibits equivalent creep behavior compared to TMS-138A. This behavior was achieved using an alloy with greatly improved oxidation performance. That is, the design goal was achieved while the alloy was made low cost and low density according to conventional manufacturing techniques.

Claims (25)

  3.5-6.5 wt% chromium, 0.0-12.0 wt% cobalt, 4.5-11.5 wt% tungsten, 0.0-0.5 wt% molybdenum, 3. 5-7.0 mass% rhenium, 1.0-3.7 mass% ruthenium, 3.7-6.8 mass% aluminum, 5.0-9.0 mass% tantalum, 0.0- 0.5 mass% hafnium, 0.0-0.5 mass% niobium, 0.0-0.5 mass% titanium, 0.0-0.5 mass% vanadium, 0.0-0. 1 wt% silicon, 0.0-0.1 wt% yttrium, 0.0-0.1 wt% lanthanum, 0.0-0.1 wt% cerium, 0.0-0.003 wt% % Sulfur, 0.0-0.05% by weight manganese, 0.0-0.05% by weight zirconium, 0.0-0.005% by weight boron, Consists .0～0.01 wt% carbon, the balance nickel-based alloy composition consisting of nickel and unavoidable impurities.   The nickel-base alloy composition according to claim 1, wherein chromium is 4.0 to 5.0% by mass.   Cobalt is 0.1 to 12.0% by weight, more preferably 0.1 to 11.0% or 0.1 to 10.9%, most preferably 7.0 to 11.0% or 7 The nickel-base alloy composition according to claim 1 or 2, which has a content of 0.01 to 10.9%.   4. The tungsten according to claim 1, wherein the tungsten content is 7.0 to 9.5% or 7.1 to 11.5%, preferably 7.1 to 9.5%. Nickel-based alloy composition.   The aluminum content is 3.7 to 6.6%, preferably 5.1 to 6.6%, more preferably 5.5 to 6.6% by mass. The nickel-base alloy composition described in 1.   The nickel-base alloy composition according to any one of claims 1 to 5, wherein tantalum is 5.0 to 7.3% by mass.   The nickel-base alloy according to any one of claims 1 to 6, wherein the molybdenum is at least 0.1% by mass.   The nickel-based alloy according to any one of claims 1 to 7, wherein rhenium is 4.5 to 6.0% by mass, more preferably 5.3 to 6.0%.   The nickel-base alloy according to any one of claims 1 to 8, wherein ruthenium is 2.0 to 3.0%, preferably 2.1 to 2.9% by mass. The nickel-base alloy composition according to any one of claims 1 to 9, wherein the following formula is satisfied when the mass% of tantalum and aluminum contained in the alloy is W _Ta and W _Al , respectively.
33 ≦ W _Ta +5.1 W _Al ≦ 39 The nickel-base alloy composition according to any one of claims 1 to 10, which satisfies the following formula when the mass% of tantalum and aluminum contained in the alloy is W _Ta and W _Al , respectively.
2.2 ≦ 5.15W _Al −0.5W _Al ² −W _Ta
Preferably, 2.9 ≦ 5.15 W _Al −0.5 W _Al ² −W _Ta The nickel-base alloy composition according to any one of claims 1 to 11, which satisfies the following formula, when the mass% of ruthenium and rhenium contained in the alloy is W _Ru and W _Re , respectively.
4.5 ≧ W _Ru + 0.225W _Re
Preferably, 3.9 ≧ W _Ru + 0.225W _Re The nickel-base alloy composition according to any one of claims 1 to 12, which satisfies the following formula when the mass% of rhenium and tungsten contained in the alloy is W _Re and W _W , respectively.
15.8 ≧ 1.13W _Re + W _W
Preferably, 14.4 ≧ 1.13 W _Re + W _W The nickel-base alloy according to any one of claims 1 to 13, which satisfies the following formula when the mass% of rhenium, molybdenum and tungsten contained in the alloy is W _Re , W _Mo and W _W , respectively. Composition.
21.9 ≦ 2.92 W _Re + (W _W + W _Mo )
Preferably, 24.6 ≦ 2.92W _Re + (W _W + W _Mo )   The nickel-base alloy composition according to any one of claims 1 to 14, wherein the sum of niobium, titanium and vanadium is less than 1%, preferably less than 0.5%, by weight.   The nickel-base alloy composition according to any one of claims 1 to 6, wherein hafnium is 0.0 to 0.2% by mass.   The nickel-base alloy composition according to any one of claims 1 to 16, wherein the volume fraction has γ 'of 60% to 70%.   The nickel according to any one of claims 1 to 17, wherein the total of niobium, titanium, vanadium and tantalum is 5.0 to 9.0%, preferably 5.0 to 7.3% by mass. Base alloy.   A single crystal formed of the nickel-base alloy composition according to claim 1.   A turbine blade for a gas turbine engine formed of an alloy according to any one of the preceding claims.   A gas turbine engine comprising the turbine blade according to claim 20.   An alloy substantially as herein described and / or substantially as described in any one of the accompanying drawings.   A single crystal article substantially as herein described and / or substantially as described in any one of the accompanying drawings.   A turbine blade substantially as herein described and / or substantially as described in any one of the accompanying drawings.   A gas turbine engine substantially as herein described and / or substantially as described in any one of the accompanying drawings.