JPH0770678A - High strength cemented carbide and high strength single crystal casting - Google Patents

Abstract

PURPOSE:To provide a cemented carbide combining creep fracture strength more excellent than that of the existing single crystal alloys and oxidation resistance on a good class among the existing single crystal alloys and to provide single crystal cast parts for blades and nozzles using the same. CONSTITUTION:This high strength cemented carbide is the one contg., by weight, 4.5 to 10% Cr, 4.5 to 6.5% Al, 3 to 9% W, 3 to 9% Ta, >3 to 9% Mo, 0.1 to 3% Co, 0.1 to 4% Re, <=0.3% Hf, and the balance Ni with inevitable impurities, and the single crystal casting tar blades and nozzles using the same is produced.

Description

Detailed Description of the Invention

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention is used for blades (moving blades) and nozzles (stating blades) of gas turbines for aircraft or ground power generation, which require high creep rupture strength in a high temperature combustion gas atmosphere. The present invention relates to a novel superalloy, a single crystal cast produced by using the superalloy, and a single crystal part for a gas turbine produced by using the single crystal, which has particularly high creep rupture strength and oxidation resistance. Regarding what is required.

[0002]

2. Description of the Related Art Turbine blades that are exposed to the most severe operating environment against the increase in combustion temperature due to higher power output and higher efficiency of gas turbine engines are made of polycrystalline ordinary casting alloys, and have crystal grains in the stress loading direction. It has undergone a transition from a unidirectional columnar solidified alloy with no boundaries to a single crystal alloy with no grain boundaries.

Among these single crystal alloys, PWA1484 is an alloy having particularly high creep rupture strength.
(U.S. Pat. No. 4,719,080, JP-A-61-28
No. 4545, "Second-generation Nickel-base Single
Crystal Superalloy "; ADCetel and DNDuhl; Supera
lloys 1988, The Metall. Soc., (1988), pp235-244), C
MSX-4 (U.S. Pat. No. 4,643,782, JP-A-60-211031, "Process and Alloy Optimizati
on for CMSX-4 Superalloy Single Cristal Airfoils ";
DJ Fraisier, JR Whetstone, K. Harris, GLErickso
n, RESchwer; HighTemp. Mater. Power Eng. 1990 Pa
rt2, (1990), pp1281-1300) and SC-83K (US Pat. No. 4,976,791, JP-A-2-138431, "N").
Development of i-based single crystal super heat resistant alloys "; Takehiro Ohno, Rikizo Watanabe; Iron and Steel, vol.77, (1991), pp832-839), etc. However, the high temperature strength of these conventional alloys However, it is difficult to satisfy the demands for high output and high efficiency of the gas turbine engine, which has undergone remarkable development in recent years.

[0004]

SUMMARY OF THE INVENTION Therefore, the present invention is a single crystal component having high oxidation resistance comparable to that of a conventional single crystal alloy, and further having high creep rupture strength superior to that of the conventional single crystal alloy. To provide a high-strength superalloy suitable for, a master ingot made of the same, and a single crystal cast made by using the alloy, and a gas turbine blade and a gas turbine nozzle made by using the single crystal cast. Intended.

[0005]

Therefore, as a result of intensive investigations by the present inventors, Ni-Cr-Al-W-Mo-Ta-
In the (Ti) -Co-Hf system, by adding Re and having a higher amount of Mo than the conventional alloy, it is possible to obtain a high degree of solid solution strengthening that is not possible with the conventional alloy while maintaining the structural stability. As a result, it was possible to obtain an alloy with a high creep rupture strength, which was much higher than that of conventional alloys.

Another feature of the present invention is to find the optimum amount of Co added. In conventional alloys, Co was regarded only as an impurity level, or conversely, about 5 to 10% was added for the purpose of improving strength. It has been clarified that this excessive Co is an undesirably large amount for the alloy of the present invention in terms of oxidation resistance and strength. This carefully controlled amount of Co, together with the addition of a small amount of Hf, had a remarkable effect on the improvement of the oxidation resistance and the high temperature strength of the alloy of the present invention.

The high strength superalloy of the present invention can be single crystallized by the unidirectional solidification casting method. Since single crystal parts such as blades and nozzles of gas turbines made using this single crystal superalloy have high oxidation resistance and high temperature strength, the current gas turbines whose combustion efficiency is limited by the performance of these parts are By using the blade and nozzle of the single crystal superalloy according to the present invention, it is possible to manufacture a gas turbine with high efficiency and high output which has never been achieved.

The first invention of the present invention obtained as a result of the above-mentioned examination is, by weight%, Cr4.5 to 10%, Al.
4.5-6.5%, W3-9%, Ta3-9%, Mo3
% To 9% or less, Co 0.1 to 3%, Re 0.1 to 4
%, Hf 0.3% or less and balance unavoidable impurities and Ni
Is a high strength superalloy. Furthermore, among these compositions, the relationship between W, Mo and Re is
0.40 ≦ 2Mo / (W + 2Mo + Re) ≦ 0.65
And 0.04 ≦ Re / (W + 2Mo + Re) ≦ 0.20
It is desirable to have a range that satisfies both. Among them, the preferable composition is, by weight%, Cr 5.3 to 8%, Al 5 to 6%,
W 6.5-8%, Ta 6-8.5%, Mo 3.3-6
%, Co 0.5 to 1.5%, Re 1.0 to 2.0%, H
f 0.01-0.2% and balance unavoidable impurities and Ni
Is a high strength superalloy. In addition, less than 1% Ti can be added to these compositions. If only high temperature strength is important and oxidation resistance is not important, Co and Hf may not be intentionally added.

A second aspect of the present invention is a rod-shaped master ingot made of the above alloy. A third aspect of the present invention is a high-strength single crystal casting that is formed by unidirectionally solidifying the above alloy and that has substantially no grain boundaries.
Furthermore, this high-strength single crystal cast product had a volume ratio of undissolved eutectic γ'phase existing after the solid solution treatment of 5% or less, solid solution +
The volume ratio of the γ'phase after aging is 50 to 70%, and the γ'phase precipitated after solid solution + aging treatment is adjusted to a cubic or rectangular parallelepiped shape with a side length of 1 μm or less, resulting in excellent high temperature. Strength can be obtained. Further, this high-strength single crystal cast product can obtain a creep rupture time at 1040 ° C. and 19 kgf / mm ² of 250 hours or more. A fourth invention of the present invention is a gas turbine blade and a gas turbine nozzle, which are made of the high-strength single crystal casting of the third invention.

[0010]

The reason for limiting the components of the alloy of the present invention will be described below. Since Cr has the effect of improving the oxidation resistance and corrosion resistance of the alloy, a minimum of 4.5% is required, but excessive addition causes harmful precipitation phases such as the σ phase and reduces creep rupture strength and ductility. Therefore, it is limited to 4.5 to 10%. It is preferably 5.3 to 8%.

Al is an important element for forming a film of Al ₂ O _{3 which} contributes most to the improvement of the high temperature oxidation characteristics of the Ni-base superalloy, and in this respect, it is desirable that the amount of Al is large. At the same time, Al is also a main strengthening element that forms a γ'phase, which is an intermetallic compound that precipitation strengthens the Ni-base superalloy. Although the basic composition of the γ'phase is represented by Ni ₃ Al, it is further strengthened by solid-solving elements other than Al, such as Ti, Ta, W and Mo. The action of these elements will be described in detail below. Single crystal alloy is usually 5 in volume ratio
Although it contains a large amount of 0% or more γ'phase, a coarse γ'phase called eutectic γ'phase exists in the final solidification part at the end of solidification, so this is once dissolved in the parent phase (γ phase). In order to do so, the solution treatment is performed at a high temperature. Γ'dissolved by solution treatment
The phases strengthen the alloy by uniformly and finely precipitating during aging and during cooling.

Therefore, Al needs to be at least 4.5%, but excessive addition of more than 6.5% causes too much γ'phase, and the eutectic γ'phase is completely solidified by solution treatment. Since it cannot be melted, the strength decreases. Also, A
The amount of 1 is relatively high with respect to the solid solution strengthening elements of the γ'phase such as Ta, W, Mo, and Ti described above.
It also means that the phases are not solution strengthened. Therefore, in the present invention, Al is limited to the range of 4.5 to 6.5%. It is more preferably 5 to 6%.

W is an element that forms a solid solution in the γ phase and γ'phase to strengthen both phases, and requires at least 3%. However, excessive addition causes precipitation of the α-W phase and the Re-W phase and lowers the strength, and further lowers the corrosion resistance at high temperatures and increases the specific gravity. Therefore, W is limited to the range of 3 to 9%. It is preferably in the range of 6.5 to 8%. Mo is also an element similar to W, is an element that forms a solid solution in the γ phase and the γ ′ phase and strengthens both phases, and is an essential additional element. In particular, Mo is
Since Ni, which is the base, has a high solid solubility in the austenite phase, it is more difficult to precipitate a different phase than when W is added. Further, regarding the distribution ratio between the γ phase and the γ ′ phase, since it is more solid-solved on the γ-phase side than W, it works more effectively for solid solution strengthening of the γ-phase. Mo is also more advantageous than W in terms of specific gravity. From the above points, in the present invention, Mo is a more important element than W, and it is necessary to add at least more than 3%.

However, excessive addition of more than 9% causes precipitation of α-Mo phase and Re-Mo phase and lowers high temperature strength. Therefore, Mo is limited to the range of more than 3% and 9% or less. More preferably, it is in the range of 3.3 to 6%. Re
Is an element that is distributed at a higher concentration in the γ phase than in the γ ′ phase, much more than W, let alone W, and is the most effective element for solid solution strengthening of the γ phase. In addition to the Mo content that preferentially solid-solution strengthens the γ phase in the range of more than 3%,
Furthermore, this combination of adding Re in an appropriate amount is a completely new idea that has not been found in the above-mentioned high-strength Ni-based single crystal alloys disclosed in the literatures so far. Therefore, the required Re is at least 0.1%. On the other hand, Re is a very expensive element, and excessive addition of more than 4% unnecessarily raises the price of the alloy, and Re-W, Re-M
Re causes the precipitation of harmful phases such as o and Re-Ta, and therefore Re is set to the range of 0.1 to 4.0%. More preferably 1.0
The range is up to 2.0%.

In order to increase the strength of the Ni-based single crystal superalloy, it is necessary to maintain the consistency of the γ phase and the γ'phase and to strengthen each phase to the maximum extent by solid solution strengthening. Γ was sufficiently solid-solution strengthened by adding Mo content and an appropriate amount of Re
By obtaining the phase, the γ'phase having a basic composition of Ni ₃ Al, which is a precipitation strengthening element, is sufficiently solid-solution strengthened by Ta and Ti to the extent commensurate with it, and the γ phase and the γ'phase are matched. It is possible to obtain a high-strength single crystal superalloy having high properties and high solid solution strengthening degree of both phases. FIG. 1 shows how Mo and Re among W, Mo, and Re mainly distributed to the γ phase are contained in novel proportions in the ratio with respect to the sum of atomic weights of these three elements. In Fig. 1, the horizontal axis is Re / (W +
2Mo + Re) and the vertical axis is 2Mo / (W + 2Mo +
Re) is taken to indicate the scope of claims 1 to 3 in the present invention and the position of the alloy in the examples.

Re / (W + 2Mo + Re) and 2Mo / (W + 2Mo) derived from the W and Mo and Re amounts of claim 1.
The range of + Re is (0.0066, 0.397),
(0.211, 0.316), (0.308, 0.46
2), (0.16, 0.72), (0.0047, 0.
853) and (0.0037, 0.664) are the areas within the solid line surrounded by 6 points. Conventional alloy No. 31 (S
C-83K), No. 32 (PWA1484) and N
o. 33 (CMSX-4), it is clear how the alloy of the present invention contains Re and has a higher Mo composition.

In the first aspect of the invention, the 0.
40 ≦ 2Mo / (W + 2Mo + Re) ≦ 0.65,
In order to satisfy 0.04 ≦ Re / (W + 2Mo + Re) ≦ 0.20, (0.04, 0.40), (0.20,
0.40), (0.20, 0.65) and (0.0
4, 0.65) is a more desirable range in terms of easiness of solution treatment, high temperature strength and alloy price. More suitable Re / (W + 2Mo + R
The ranges of e) and 2Mo / (W + 2Mo + Re) are derived from the amounts of W, Mo and Re of claim 3 (0.064,
0.42), (0.12, 0.40), (0.13,
0.44), (0.0975, 0.59), (0.05
1, 0.62) and 6 of (0.0497, 0.57)
This is the area within the broken line surrounded by dots.

Co is an element that plays an important role in the alloy of the present invention, and there is an optimum amount of addition that clearly improves the oxidation resistance with respect to the alloy of the present invention. In terms of strength, the addition of Co lowers the stacking fault energy of the alloy and improves the creep strength in a relatively low temperature range, while it also increases the solid solubility of the γ'phase in a high temperature range, resulting in precipitation strengthening. And weakens the creep strength in the high temperature range. Due to their contradictory effects, Co has an optimum addition amount in terms of strength. Due to such effects, Co requires a minimum addition of 0.1%. However, the addition of more than 3% is no longer effective for the oxidation resistance, and the high temperature strength also decreases. Also, T
Since it is easy to generate a harmful phase called CP phase (topologically close-packed phase), Co is 0.1 to 3.
Limited to 0%. Desirably, it is 0.5 to 1.5%.

Ta is an essential strengthening element mainly for solid-solution strengthening the γ'phase and constituting a γ'phase having a consistency with the γ phase solid-solution strengthened with W, Mo, and Re described above. . Therefore, Ta needs to be at least 3%, but excessive addition exceeding 9% causes an increase in the solid solution temperature of the eutectic γ'phase and precipitation of the Re-Ta phase, and conversely decreases the high temperature strength. . Therefore, Ta is limited to the range of 3 to 9%. The preferred range is 6
~ 8.5%.

Hf is an important element for improving the oxidation resistance and high temperature strength of the alloy, and is an essential additional element. The effect appears from a very small amount of addition. But,
Excessive addition of Hf lowers the melting point of the alloy, lowers the solution treatment temperature, and the eutectic γ'phase cannot be sufficiently dissolved. Therefore, the addition amount should be as small as possible. Therefore, Hf is added at 0.3% or less, but a more preferable range is 0.01 to 0.2%.

Since Ti forms a solid solution in the γ'phase and serves to strengthen the solid solution of the γ'phase, Ti can be added as a selective element also in the alloy of the present invention. However, excessive addition of more than 1% easily forms a eutectic γ'phase and lowers the melting point of the alloy, so that the difference between the initial melting temperature and the complete solution temperature of the γ'phase,
That is, the heat treatment window is narrowed, and the solid solution of the γ'phase by the solution treatment becomes insufficient. Therefore, in the alloy of the present invention, Ti can be added within the range of 1% or less. Among the elements other than the above, C, Si, Mn, P,
If S, B, Zr, Y, REM, and Cu are within the ranges shown below, there is no particular problem in terms of characteristics, but it is desirable that they are as low as possible. C ≦ 0.015% Si ≦ 0.05% Mn
≤0.5% P ≤0.005% S ≤0.003% B
≤0.003% Zr≤0.02% Y ≤0.2% REM
≤0.2% Cu ≤0.1%

The above alloy composition group is formed into a master ingot by the method described below, and further becomes a single crystal cast product. Here, the matrix γ phase (austenite phase) and γ '
Although it is a phase different from the phase, it is usually called a single crystal because it is a matched phase with the same crystal orientation. First of all, the master ingot of the above alloy composition has individual alloy elements in advance,
Or, after refining the scrap of the single crystal casting described below and using a return material with a reduced impurity level to a level where it can be reused, it is vacuum melted into a master ingot, then remelted in vacuum and then unidirectional. It can be solidified to obtain a single crystal casting. At this time, a master ingot having a purity as high as possible is suitable for manufacturing a single crystal cast.

The single crystal cast product can be industrially used by being subjected to heat treatment such as solution treatment, aging treatment and surface coating treatment. This single crystal casting is
The heat treatment and the alloy composition are preferably adjusted so as to have the structure shown below. First of all, if the eutectic γ'phase generated during solidification cannot be sufficiently dissolved by the solution treatment, this undissolved eutectic γ'phase portion becomes the starting point of creep fracture.
Therefore, the undissolved eutectic γ ′ existing after the solution treatment
The phase volume ratio is preferably 5% or less. Secondly, the amount of γ'phase present after solid solution + aging treatment also has a great influence on the strength and ductility of the casting. If the volume ratio of the γ'phase is less than 50%, sufficient high temperature strength cannot be obtained. On the contrary, if it exceeds 70%, undissolved eutectic γ'phase becomes excessively left in the solution treatment. . Therefore, the volume ratio of the γ ′ phase after solid solution + aging treatment is limited to 50 to 70%. Therefore, a more preferable volume ratio of the γ'phase is 55 to 65%.

Thirdly, it is desirable that the γ'phase precipitated during this aging treatment has a lattice constant that is sufficiently matched with the austenite phase that is the matrix, and that it is finely precipitated in a regular cubic or rectangular parallelepiped shape. If the austenite phase and the γ'phase are not sufficiently matched, the length of one side exceeds 1 μm, the corners of a cube or a rectangular parallelepiped are deformed, and spherical precipitation occurs, resulting in sufficient high temperature strength. You won't get it. Therefore, it is necessary that the γ'phase precipitated during the aging treatment has a cubic shape or a rectangular parallelepiped shape with a side length of 1 μm or less. Therefore, the desirable side length of the γ ′ phase is 0.02 to 0.7 μm. Further, this single crystal cast product has characteristics of 1040 ° C., 1
It is desirable that the creep rupture time at 9 kgf / mm ² is 250 hours or more.

The high-strength single crystal casting obtained by unidirectionally solidifying the above-mentioned novel material is suitable for an article used in a harsh environment where high creep rupture strength and excellent oxidation resistance are required. Since the gas turbine blade and the gas turbine nozzle made of the above-mentioned high-strength single crystal casting have high creep rupture strength and excellent oxidation resistance, it is possible to raise the combustion gas temperature of the gas turbine higher than at present. It becomes possible, and as a result, the thermal efficiency of the gas turbine is significantly improved.

[0026]

【Example】

(Example 1) Table 1 shows the chemical compositions of the samples used for comparing the properties of the alloy of the present invention, the comparative alloy and the conventional alloy, and 2M.
o / (W + 2Mo + Re) value and Re / (W + 2M
o + Re) value is shown. Inventive alloy No. 1-10, comparative alloy No. 21-22 and conventional alloy No. 31-3
As for No. 3, a 5 kg master ingot was prepared by vacuum induction melting. All conventional alloys were melted aiming at the same composition as the published composition. Of the conventional alloys, No. 31 is SC-83K, N
o. 32 is PWA 1484, No. 32. 33 is CMSX-4
Indicates.

[0027]

[Table 1]

Table 2 shows 1100 ° C. for each alloy in the crucible.
Oxidation weight loss after 10 cycles of air-cooling heat cycle after heating for 16 hours, alloy structural stability, and 1040
The creep rupture time at −19 kgf / mm ² and the elongation at that time are shown. Further, regarding the structural stability of the alloy, from the microstructure after each heat treatment, for the alloy composed of only the γ phase and the γ ′ phase, a mark “◯” is described in the property column.

[0029]

[Table 2]

In each of the various tests shown in Table 2, the above master ingot was single-crystallized in a pull-down type directional solidification furnace, and then the following heat treatment was carried out. Regarding the creep rupture test, the conventional alloy No. For all other alloys except 31 and 33, the parallel part diameter is 6.35 mm and the distance between the scores is 2
The test piece was processed into a 5.4 mm test piece, and the test was carried out under the above conditions based on the ASTM method. Among the conventional alloys, alloy No. 1 whose creep rupture curve organized by Larson Miller parameters is publicly known. For 31 and 33,
The rupture time corresponding to 1040 ° C.-19 kgf / mm ² was read from the creep rupture curve and is also shown in Table 2. Further, the crucible oxidation-resistant test piece used was a disk-shaped test piece having a diameter of 7 mm and a thickness of 4 mm.

Regarding the heat treatment conditions, with respect to the alloys of the present invention and the comparative alloys, the structures which were heated in the range of 1250 to 1350 ° C. for 4 hours and then air-cooled were examined in advance. Select the temperature to be used as the solution treatment temperature,
After holding at that temperature for 4 hours, an air-cooled solution treatment was performed.
Regarding the aging conditions after the solution treatment, a two-stage aging treatment of heating at 1080 ° C. for 4 hours and then air cooling followed by heating at 870 ° C. for 20 hours and air cooling was performed. For conventional alloys, N
o. 31 (SC-83K) was heated at 1320 ° C. for 4 hours and then air-cooled, heated at 1080 ° C. for 5 hours and then air-cooled, and heated at 870 ° C. for 20 hours and then air-cooled. No. 32 (P
WA1484) was heated at 1316 ° C. for 4 hours, air-cooled, heated at 1080 ° C. for 4 hours and air-cooled, and further heated at 870 ° C. for 20 hours and air-cooled.

No. 33 (CMSX-4) is Cannon-M
Recommended heat treatment conditions by uskegon (Source: “Single crystallization of latest nickel-base superalloy and its high temperature strength characteristics”, Yoshio Ota)
Others; 12 according to iron and steel, vol.76, (1990), pp940-947)
Hold at 72 ° C for 2 hours, raise temperature at 1288 ° C for 2 hours, raise temperature at 1296 ° C for 3 hours, raise temperature at 1304 ° C for 3 hours
After 6 hours of continuous solution treatment of temperature rise after holding for 3 hours, hold at 1313 ° C. for 3 hours, hold at 1316 ° C. for 2 hours and air cooling, hold for 4 hours at 1080 ° C. and air cooling and 871 ° C.
After maintaining at 20 ° C. for 20 hours, air cooling aging treatment was performed.

From Table 1 and FIG. 1, it can be seen that the Re and Mo contents of the alloy of the present invention are both higher than those of the comparative alloy and the conventional alloy. From Table 2, the alloy No. of the present invention. It can be seen that all of 1 to 10 have good oxidation resistance, creep rupture life, creep rupture elongation and microstructure stability. In particular, the creep rupture life of the alloy of the present invention is No. 1, which is considered to have the highest strength among conventional alloys. 31
It can be seen that a life longer than (SC-83K) is obtained, and that it is an optimal alloy as a blade material and a nozzle material for a high efficiency and high output gas turbine engine. On the other hand, comparative alloy No. Nos. 21 and 22 are inferior to the alloys of the present invention in creep rupture life regardless of whether or not Ti is added. Compared with the alloys of the present invention, this is due to the low Mo content and 2Mo / (W + 2Mo + R
e) by having a quantity.

Of the conventional alloys, No. 31 (SC-83
K) has a good oxidation resistance because it contains a small amount of Co and Hf, but since it does not contain Re, the creep rupture life is inferior to that of the alloy of the present invention. No. 32 (PWA1484) and N
o. Although 33 (CMSX-4) contains Re, its creep rupture life is inferior to that of the alloy of the present invention because the amount of Mo and the amount of 2Mo / (W + 2Mo + Re) are low.

(Example 2) The alloy N of the present invention in Example 1
o. 2 was used to manufacture the gas turbine blade shown in FIG. 2 and the gas turbine nozzle shown in FIG. 3 with and without a core. FIG. 4 shows a front view of a core for a gas turbine blade, and FIG. 5 shows a front view of a core for a gas turbine nozzle. In gas turbines of recent years, it is common to provide cooling holes of complicated shapes inside the blades and nozzles in order to lower the temperature on the metal surface and inside as the combustion gas temperature rises. In order to manufacture such a hollow structure blade and nozzle, FIG. 4 and FIG.
A core made of a refractory containing silica as the main component was used. Create a wax model around this core,
Further, a ceramic shell was formed on the outside with a refractory material such as alumina, zircon, and yttria, which was dewaxed and fired to obtain a mold.

FIG. 6 is a sectional view of a mold for a gas turbine blade, and FIG. 7 is a sectional view of a mold for a gas turbine nozzle. In both cases of FIG. 6 and FIG. 7, the mold 10 was fixed on the water-cooled copper chill 11 and set in the mold heater 13. Next, the alloy No. of the present invention melted by high frequency heating. 1
The master ingot having the composition of No. 1 is cast into the mold 10 heated above the melting point of the alloy, and the pulling speed is 30 cm / h.
Withdraw from the mold heater 13 and starter part 1
One-way solidification was sequentially carried out from No. 4. Although many columnar crystals grow in the starter portion 14, only one crystal in the columnar crystal was grown using the selector portion 15, and the portion above the selector portion 15 was a single crystal cast.

The mold heater 13 was kept at a temperature not lower than the melting point of the alloy until the mold 10 was completely drawn out and the casting was completely solidified. After the steps of setting the mold 10 in the water-cooled copper chill 11 among the above steps, the steps were performed in vacuum. After cooling, the mold 10 is taken out, the core is removed with an alkali, and the starter part, the selector part, the riser part, etc. are cut,
A gas turbine blade having the shape shown in FIG. 2 and a gas turbine nozzle having the shape shown in FIG. 3 were obtained. The gas turbine blade has a total length of about 220 mm, of which the blade part is about 130 mm,
The gas turbine nozzle has approximately 13 between the two sidewalls.
It is 0 mm. Here, the bypass portion 12 is used to single-crystallize a portion of the overhang portion, such as the seal fin 3 and the sidewall 8 where the cross-sectional area changes abruptly with respect to the crystal growth direction. Cut and remove like hot water. By using this, generation of foreign crystals in the overhang portion of the large single crystal cast was suppressed, and the yield was improved.

The crystal can be grown in the blade so that the <001> direction is the longitudinal direction of the blade (the direction from the enlarged portion 16 to the dovetail portion 5), that is, the <001> direction is the direction to which the centrifugal force is applied. Desirably, in the nozzle, the <001> direction is the lateral direction of the wing (sidewall 7
To the side wall 8), that is, the <001> direction is preferably the direction in which the thermal stress generated from the thermal cycle accompanying the start and stop of the gas turbine is applied. In each of the examples, single crystal castings were obtained in which the deviation of the crystal growth orientation from the <001> direction was within 5 degrees. Both the gas turbine blade and the gas turbine nozzle were heated at 1300 ° C. for 4 hours in a vacuum, and after air-cooled solution treatment, a blade of CoNiCrAlY having a thickness of 100 μm was formed.
An alloy layer was formed by the plasma spraying method, and in order to adapt this alloy layer to the substrate, air-cooling diffusion treatment was performed after 1080 ° C. for 4 hours. Next, ZrO ₂ − with a thickness of 300 μm was formed on the outer layer.
A 6 wt% Y ₂ O ₃ film was coated by the plasma spraying method. After that, 870 for the purpose of adjusting the shape of the γ'layer.
After heating at 20 ° C. for 20 hours, an aging treatment of air cooling was performed.

Among these gas turbine blades and gas turbine nozzles, a creep rupture test piece was indexed in the <001> direction from a heat-treated material in a solid state without a core, and the creep rupture test was conducted under the same conditions as in Example 1. Was carried out. As a result, the creep rupture lives of the blade and the nozzle were 268 h and 274 h, respectively, and the elongations were 12.4% and 11.4%, respectively, and it was confirmed that they had the same performance as that of Example 1.

[0040]

INDUSTRIAL APPLICABILITY As described above, the alloy of the present invention has a creep rupture strength superior to that of existing single crystal alloys and a good class of oxidation resistance among existing single crystal alloys. As a result, it has become possible to operate as a single crystal cast for blades and nozzles of high-power and high-efficiency gas turbines for aircraft etc., which was difficult to apply in the past, under severe oxidizing environment and high creep stress. A gas turbine with high output and high efficiency that could not be achieved can be obtained.

[Brief description of drawings]

FIG. 1 is a diagram showing a region of a high-strength superalloy according to the present invention in which the vertical axis represents the value of 2Mo / (W + 2Mo + Re) and the horizontal axis represents the value of Re / (W + 2Mo + Re).

FIG. 2 is a perspective view of a gas turbine blade according to the present invention.

FIG. 3 is a perspective view of a gas turbine nozzle according to the present invention.

FIG. 4 is a front view of a core used for manufacturing a gas turbine blade according to the present invention.

FIG. 5 is a front view of a core used for manufacturing the gas turbine nozzle according to the present invention.

FIG. 6 is a vertical cross-sectional view of a gas turbine blade mold showing a method for manufacturing a gas turbine blade according to the present invention.

FIG. 7 is a vertical cross-sectional view of a gas turbine nozzle mold showing a method for manufacturing a gas turbine nozzle according to the present invention.

[Explanation of symbols]

 1 Wing Part 2 Platform Part 3 Seal Fin 4 Shank Part 5 Dovetail Part 6 Wing Part 7, 8 Sidewall 10 Mold 11 Water-cooled Copper Chill 12 Bypass Part 13 Mold Heating Heater 14 Starter Part 15 Selector Part 16 Expanding Part

 ─────────────────────────────────────────────────── ─── Continuation of the front page (72) Hideki Tamaki, Inventor Hideki Tamaki 7-1, Omika-cho, Hitachi City, Ibaraki Hitachi Ltd. (72) Inventor Akira Yoshinari 7-chome, Omika-cho, Hitachi City, Ibaraki Prefecture No. 1 Hitachi Ltd., Hitachi Research Laboratory (72) Inventor Mitsuru Kobayashi 7-1, 1-1 Omika-cho, Hitachi City, Ibaraki Prefecture Hitachi Ltd. Hitachi Research Laboratory (72) Inventor Noriyuki Watanabe Hitachi City, Ibaraki Prefecture 7-1-1 Omika-cho, Hitachi, Ltd. Hitachi Research Laboratory

Claims (10)

[Claims] 1. By weight%, Cr 4.5-10%, Al
4.5-6.5%, W3-9%, Ta3-9%, Mo3
% To 9% or less, Co 0.1 to 3%, Re 0.1 to 4
%, Hf 0.3% or less and balance unavoidable impurities and Ni
A high strength superalloy. 2. The high-strength superalloy according to claim 1, wherein W, Mo and Re represented by weight% satisfy the following formula. 0.40 ≦ 2Mo / (W + 2Mo + Re) ≦ 0.65 0.04 ≦ Re / (W + 2Mo + Re) ≦ 0.20 3. By weight%, Cr 5.3 to 8%, Al 5 to
6%, W 6.5-8%, Ta 6-8.5%, Mo 3.3
~ 6%, Co 0.5-1.5%, Re 1.0-2.0
%, Hf 0.01 to 0.2%, and the balance unavoidable impurities and Ni, a high strength superalloy. 4. High strength superalloy according to claim 1 or 3, characterized in that it comprises less than 1% Ti by weight. 5. A polycrystalline master ingot made of the alloy according to claim 1. 6. A high-strength single crystal cast product having substantially no grain boundaries, which is obtained by unidirectionally solidifying the alloy according to any one of claims 1 to 4. 7. The volume ratio of the undissolved eutectic γ ′ phase present after the solution treatment is 5% or less, the volume ratio of the γ ′ phase after solution + aging is 50 to 70%, and the solution + The high-strength single crystal cast product according to claim 6, wherein the γ'phase precipitated by the aging treatment has a cubic or rectangular parallelepiped shape with a side length of 1 µm or less. 8. The high strength single crystal cast product according to claim 6, wherein the creep rupture time at 1040 ° C. and 19 kgf / mm ² is 250 hours or more. 9. A gas turbine blade comprising the high-strength single crystal cast product according to claim 6. 10. A gas turbine nozzle comprising the high-strength single crystal cast product according to claim 6.