JP2000063969A - Nickel base superalloy, its production and gas turbine part - Google Patents

Abstract

PROBLEM TO BE SOLVED: To produce an Ni base superalloy capable of easily executing heat treatment and provided with excellent creep strength and high temp. corrosion resistance, to provide a method for producing it and to provide a gas turbine parts improved in reliability by applying the obtd. Ni base superalloy to gas turbine parts. SOLUTION: This superalloy compsn. contg., by weight, 10 to 14% Co, 10 to 13% Cr, 0.1 to 3% Mo, 4 to 8% W, 4 to 6% Al, 1 to 4% Ti, 4 to 8% Ta, 0.1 to 0.5% Hf, <=2% Re, and the balance Ni base with inevitable impurities.

Description

Detailed Description of the Invention

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a Ni-base superalloy which is used as a component having large characteristic dimensions in turbine blades and vanes of a gas turbine for power generation and which is difficult to heat treat,
The present invention relates to a method for producing the same and a gas turbine component, and in particular, has excellent creep rupture properties and high temperature corrosion resistance,
The present invention relates to a Ni-base superalloy that facilitates heat treatment.

[0002]

2. Description of the Related Art As the efficiency of a gas turbine increases, the combustion temperature rises, and the moving blades and stationary blades of the gas turbine are required to have a longer creep life. Along with this, the materials for turbine rotor blades and vanes are unidirectionally solidified alloys that eliminate the grain boundaries in the stress axis direction from conventionally used cast alloys and improve creep strength at high temperatures. By eliminating the grain boundaries, the grain boundary strengthening elements that had caused the deterioration of the heat treatment characteristics were removed, and the optimal heat treatment increased the precipitation rate of the γ'phase, further improving the creep characteristics at high temperatures. It has changed to a single crystal alloy.

In recent years, Ni-based heat-resistant alloys have been used as materials for heat-resistant alloys. Ni-based heat-resistant alloys have a γ phase (Ni
Γ ′ phase ((Ni ₃ (Al, Ti)) in the matrix)
And Ta, Mo, W in the γ phase and γ ′ phase.
It is an alloy strengthened by solid solution of Re and Re.

Further, various developments have been made on single crystal blade materials with the aim of further improving the creep strength.

The first generation single crystal alloy is Re (rhenium).
Alloy containing no CMSX-, which is described in, for example, U.S. Pat. No. 4,582,548.
2. Rene'N4 in U.S. Pat. No. 5,399,313, PWA1480 in U.S. Pat. No. 4,209,348 and British Patent 21281.
2A includes PWA1483 and the like.

In the second generation single crystal alloy, Re is 3%.
Added to a certain extent to improve creep rupture strength over the first-generation single crystal alloy, for example, US Pat. No. 4,643,782.
CMSX-4, US Pat. No. 4,71
PWA-1484 described in Japanese Patent No. 9,080 and Ren described in JP-A-5-59474.
e'N5, etc.

Further, as a third-generation single crystal alloy, an alloy containing Re in an amount of about 5 to 6% has been developed. For example, CMSX- which is disclosed in Japanese Patent Application Laid-Open No. 7-138683.
10.

On the other hand, the unidirectionally solidified alloy is also required to have an improved creep rupture strength, and various developments have been made in the same manner as the single crystal alloy.

The first generation unidirectionally solidified alloy is an alloy containing no Re (rhenium), and this alloy includes CM2 disclosed in US Pat. No. 4,461,659.
47LC etc.

Further, in the second generation unidirectionally solidified alloy, Re was added in an amount of about 3% to improve the creep rupture strength as compared with the first generation unidirectionally solidified alloy. For example, US Pat.
CM186LC and others listed in No. 9,873.

These single crystal alloys and unidirectionally solidified alloys are technologies that have made remarkable progress mainly in the fields of aircraft jet engines and small gas turbines, but their purpose is to improve combustion efficiency even in large industrial gas turbines. Due to the high temperature, the technology has been diverted.

[0012]

However, since a large-scale industrial gas turbine uses a low-priced fuel unlike an aircraft jet engine and a small-sized gas turbine, the fuel contains a large amount of corrosive sulfur components. Has been. For this reason, the industrial gas turbine is required to have excellent high-temperature corrosion resistance, but the alloy developed as a material for the aircraft gas turbine blade described above cannot obtain the required high-temperature corrosion resistance.

Further, the single crystal superalloy can exhibit excellent creep strength by completely solidifying the γ ′ phase into the γ phase by the complete solution treatment and then precipitating the γ ′ phase by the aging heat treatment. . However, in the case of large turbine blades for power generation, unlike the jet engine blade material for aircraft, the blade shape is large, a temperature gradient occurs in the furnace during heat treatment, and the temperature during the heating process becomes uneven. As a result, there are cases where the temperature rises partly and melts, or conversely the temperature falls partly and the γ'phase does not form a solid solution. Both of these cause a decrease in strength.
It is important that the temperature range between the melting temperature and the alloy matrix melting temperature is wide. However, in Ni-based single crystal alloys that have been developed for aircraft, the Al concentration is generally high because importance is attached to oxidation resistance, and thus the heat treatment temperature range is narrow,
It was difficult to perform heat treatment.

On the other hand, as for the unidirectionally solidified blades, the fuel used in the industrial gas turbine is highly corrosive as described above, and the creep rupture life and high temperature resistance required for the unidirectionally solidified alloy developed for aircraft are required. It was not possible to satisfy the corrosiveness at the same time.

The present invention has been made to solve these problems, and by considering the combination of alloying elements and the amount added, heat treatment is facilitated, and excellent creep strength and high temperature corrosion resistance are obtained. A first object of the present invention is to provide a Ni-based superalloy which is a Ni-based single crystal superalloy and a method for producing the same.

Further, by unidirectionally solidifying the above-mentioned single crystal alloy, a Ni-based superalloy which is a Ni-based unidirectionally solidified superalloy having excellent creep strength and high temperature corrosion resistance, and a method for producing the same are provided. This is the second purpose.

Further, by applying the obtained Ni-base single crystal superalloy and directionally solidified alloy to gas turbine parts, the gas turbine parts are provided with a longer life and improved reliability. The third purpose is to do so.

[0018]

The Ni-base superalloy according to claim 1 is, in% by weight, Co: 10% or more and 14% or less, C
r: 10% to 13%, Mo: 0.1% to 3%, W: 4% to 8%, Al: 4% to 6%,
Ti: 1% to 4%, Ta: 4% to 8%, H
It is characterized by containing f: 0.1% or more and 0.5% or less and Re: 2% or less, and the balance being Ni group and unavoidable impurities.

The Ni-based superalloy according to claim 2 has a weight percentage.
Then, Co: 11% or more and 13% or less, Cr: 11% or more 1
2.5% or less, Mo: 1% or more and 2% or less, W: 4.5%
Or more and 6.5% or less, Al: 4.5% or more and 5.5% or less,
Ti: 2% or more and 3% or less, Ta: 4.5% or more and 6.5%
Hereinafter, Hf: 0.1% or more and 0.3% or less and Re: 2
% Or less, with the balance being Ni-based and unavoidable impurities.

The reasons why the combination of alloying elements and the addition amount are defined as described above in the inventions of claims 1 and 2 will be described.

Co (cobalt) is an element necessary for lowering the solution temperature of the γ'phase and widening the solution temperature range. In the present invention, the addition amount of Co is 10% or more 14
% Or less, but to obtain the above effect sufficiently,
This is because the addition of 10% or more is necessary, and if the addition amount exceeds 14%, the creep strength in the high temperature range is lowered.
More preferably, the amount of Co added is 11% or more and 13% or less.

Cr (chromium) is an element having an action of improving the high temperature corrosion resistance of the alloy. In the present invention, the amount of Cr added is specified to be 10% or more and 13% or less, but it is necessary to add 10% or more in order to obtain the above effects sufficiently, and if the amount added exceeds 13%, sigma This is because the tendency of precipitation of harmful phases such as phases becomes stronger, leading to a decrease in creep strength. More preferably, the amount of Cr added is 11% or more and 12.5% or less.

Mo (molybdenum) is an element necessary for solid solution strengthening of the alloy and enhancement of creep strength by strengthening the matching interface due to negative lattice constant misfit. In the present invention, the amount of Mo added is specified to be 0.1% or more and 3% or less, but in order to obtain the above effects sufficiently, it is necessary to add 0.1% or more, and the amount added exceeds 3%. This is because the tendency of precipitation of harmful phases such as sigma phase becomes stronger, which leads to lowering of creep strength and deterioration of high temperature corrosion resistance. More preferably, the addition amount of Mo is 1% or more and 2% or less.

W (tungsten) is an element necessary for improving the creep strength by solid solution strengthening of the alloy. In the present invention, the amount of W added is defined as 4% or more and 8% or less, but at least 4% or more is required to be sufficiently obtained in order to sufficiently obtain the effect of improving the creep strength, and the addition amount exceeds 8%. This is because the tendency of precipitation of harmful phases such as sigma phase and the like increases, leading to a decrease in creep strength. Further, if W is added excessively like Mo, the high temperature corrosion resistance is deteriorated. More preferably, the amount of W added is 4.5% or more and 6.5% or less.

Al (aluminum) is an element necessary for increasing the precipitation amount of the γ'phase and improving the creep strength. In the present invention, the added amount of Al is 4% or more and 6
%, It is necessary to add at least 4% or more in order to sufficiently obtain the effect of improving the creep strength, and if the addition amount exceeds 6%, the creep strength in the high temperature range is lowered. This is because. More preferably, the amount of Al added is 4.5% or more and 5.5% or less.

Ti (titanium) is an element necessary for improving the creep strength by solid-solution strengthening the γ'phase and improving the high temperature corrosion resistance. In the present invention,
Although the amount of Ti added is specified to be 1% or more and 4% or less, it is necessary to add at least 1% or more in order to obtain the above effects sufficiently. If the amount of addition exceeds 4%, the melting point of the alloy decreases. The reason is that More preferably, the addition amount of Ti should be 2% or more and 3% or less.

Ta (tantalum) is an element having a function of solid solution strengthening γ'and improving creep strength. In the present invention, the addition amount of Ta is specified to be 4% or more and 8% or less, but in order to sufficiently obtain the effect of improving the creep strength,
This is because it is necessary to add 4% or more, and if the addition amount exceeds 8%, the creep strength in the high temperature range is lowered and the high temperature corrosion resistance is deteriorated. More desirable Ta
Is 4.5% or more and 6.5% or less.

Re (rhenium) is an element mainly improving the creep strength by forming a solid solution in the γ phase to strengthen it. In the present invention, the required creep strength can be obtained without adding Re, but by adding Re, the creep strength can be further improved. However, when the added amount of Re exceeds 2%, in the alloy of the present invention, the Re-Mo phase, Re-W phase and R
TCP (Topologica), which is a brittle phase such as e-Cr-W phase
lly Close-Packed phase) Accelerates the formation of phases and further narrows the solution temperature range. Therefore, the addition amount of Re is specified to be 2% or less.

Hf (hafnium) is an element which is not added at all in conventional Ni-based single crystal alloys because it lowers the melting point of the alloy and deteriorates the heat treatment characteristics. In order to strengthen the grain boundaries of recrystallization caused by the subsequent heat treatment and processing and improve the yield of turbine rotor blades and stator blades, the alloy of the present invention is added in the range of 0.1% or more and 0.5% or less. ing. Further, more preferably 0.1% or more and 0.3
The amount added is preferably not more than%.

The invention according to claim 3 is the invention according to claim 1 or 2.
In the described Ni-base superalloy, C: 0.5% by weight%
The feature is that Zr: 0.2% or less and B: 0.1% or less are added below.

C (carbon), Zr (zirconium) and B (boron) are grain boundary strengthening elements. Since the grain boundaries are sufficiently strengthened by adding these elements, it becomes possible to use the alloy material having the component composition described in claim 1 or 2 in the directionally solidified superalloy. However, when C is added in excess of 0.5%, Zr is in excess of 0.2%, and B is in excess of 0.1%, carbides and borides, which are the starting points of fatigue cracks, are formed and fatigue strength is reduced. In order to do so, the composition of C is 0.5% or less,
Zr was defined as 0.2% or less and B was defined as 0.1% or less.

The method for producing a Ni-base superalloy according to claim 4 is the method for producing Ni, Co, Cr, Mo, W, Al, Ti, Ta,
A material composed of Re and Hf is melted and solidified to obtain N
An i-based superalloy body is formed, and this Ni-based superalloy body is heated from 1200 ° C. to 128 ° C. in an environment of vacuum or an inert atmosphere.
After solution heat treatment in the temperature range up to 0 ° C, quenching, followed by 1-step aging heat treatment in the temperature range from 1000 ° C to 1200 ° C, then 2-step aging in the temperature range from 700 ° C to 900 ° C A feature is that heat treatment is performed.

In the alloy of the present invention, the alloy is strengthened mainly by precipitating the γ'phase in the Ni matrix. However, by prescribing the temperature and time for heat treatment as described above, excellent creep can be obtained. Ni-based superalloys with fracture properties can be produced.

The invention according to claim 5 is the N according to claim 4.
In the method for producing an i-based superalloy, Ni, Co, Cr, M
A material comprising o, W, Al, Ti, Ta, Re and Hf is used to obtain a Ni-based single crystal superalloy having the component composition according to claim 1 or 2.

In the present invention, the Ni-based superalloy, which is the Ni-based single crystal superalloy having the component composition according to claim 1 or 2, can be manufactured by using the manufacturing method described in claim 4.

The invention according to claim 6 is the N according to claim 4.
In the method for producing an i-based superalloy, Ni, Co, Cr, M
A material comprising o, W, Al, Ti, Ta, Re, Hf, C, Zr and B is used to obtain a Ni-based directionally solidified superalloy having the composition of claim 3.

In the present invention, since the grain boundaries are sufficiently strengthened by adding elements such as C, Zr and B, the component composition according to claim 3 is used by using the manufacturing method according to claim 4. It is possible to produce a Ni-based superalloy that is a Ni-based unidirectionally solidified superalloy.

The invention according to claim 7 is the method for producing a Ni-base superalloy according to any one of claims 4 to 6, wherein the solution heat treatment is performed within 10 hours and the aging heat treatment is performed within 3 hours.
It is characterized by being within 0 hours.

The invention according to claim 8 is the N according to claim 7.
In the method for producing an i-based superalloy, the solution heat treatment is characterized in that any temperature change from 2 to 4 steps is performed up to the final solution heat treatment temperature.

In the present invention, a Ni-base superalloy having excellent creep rupture properties can be obtained by performing a multi-step heat treatment while changing the temperature during the solution heat treatment and the aging heat treatment.

The invention according to claim 9 is the N according to claim 8.
In the method for producing an i-base superalloy, before the solution heat treatment, the temperature of the solution heat treatment from the final solution heat treatment temperature is 40 ° C. to 6 ° C.
It is characterized in that the preliminary heat treatment is carried out at a low temperature in the range of 0 ° C. for 1 hour or more and 2 hours or less.

In the present invention, before the solution heat treatment, the preliminary heat treatment is performed at a temperature of 40 ° C. to 60 ° C. lower than the temperature of the solution heat treatment to prevent local melting due to a rapid temperature rise. The Ni-based superalloy having excellent creep rupture strength can be obtained.

A gas turbine component according to a tenth aspect is made of a Ni-based superalloy of the Ni-based single crystal superalloy or the Ni-based directionally solidified superalloy according to any one of the first to third aspects. It was

The gas turbine component according to claim 11 is composed of the Ni-base superalloy produced by the manufacturing method according to any one of claims 4 to 9.

According to the tenth and eleventh aspects of the present invention, a Ni-base superalloy having excellent characteristics is applied to a gas turbine component such as a gas turbine moving blade and a stationary blade to extend the life of the gas turbine. As a result, the reliability of the gas turbine component can be improved.

[0046]

BEST MODE FOR CARRYING OUT THE INVENTION Embodiments of the present invention will be described below with reference to Tables 1 to 11 and FIGS.

First Embodiment (Tables 1 to 3; FIG. 1; Implementation)
Examples, Comparative Examples, Conventional Examples) In this embodiment, the fact that the alloy composition of the present invention has excellent creep rupture properties and high temperature corrosion resistance will be described with reference to Examples, Comparative Examples and Conventional Examples.

Example (Sample No. 1 to No. 11) In this example, the sample No. 1 shown in Table 1 was used. 1-No.
A Ni-based single crystal alloy having an ingredient composition range of 11 was used.

[0049]

[Table 1]

As shown in Table 1, the sample No. 1-No.
Alloy materials up to 11 have alloy compositions within the scope of the invention. Specifically, by weight%, Co: 10% or more 1
4% or less, Cr: 10% or more and 13% or less, Mo: 0.1
% To 3%, W: 4% to 8%, Al: 4% to 6%, Ti: 1% to 4%, Ta: 4% to 8%, Hf: 0.1% to 0 0.5% or less and R
e: 2% or less is contained, and the balance is composed of Ni group and inevitable impurities.

Comparative Examples (Sample Nos. 12 to 27) In this comparative example, the sample Nos. 12 to N
o. A Ni-based single crystal alloy having a composition range of 27 was used.

As shown in Table 1, the sample No. 12 to N
o. Alloy materials up to 27 have alloy compositions outside the scope of the present invention.

Conventional Example (Sample No. 28, No. 29) In this conventional example, the sample No. 28 as CMSX-
2 and sample No. IN738LC was used as 29. Although IN738LC is a normally cast alloy, it is added to evaluate the high temperature corrosion resistance.

Specifically, the alloy of CMSX-2 is, by weight%, Co: 4.6%, Cr: 8.0%, Mo: 0.6.
%, W: 8.0%, Al: 3.5%, Ti: 3.5%,
It contains Ta: 3.5% and Hf: 0.1%, and the balance is composed of Ni group and inevitable impurities.

Further, IN738LC is C in weight%.
o: 8.1%, Cr: 16.0%, Mo: 1.8%,
It contains W: 2.5%, Al: 3.5%, Ti: 3.6% and Ta: 1.9%, and the balance is composed of Ni group and inevitable impurities.

With respect to the alloys having the component compositions of the examples and the comparative examples, in order to prepare each test piece, the raw materials were scoured by vacuum melting in appropriate proportions so as to have the compositions shown in Table 1 and then redissolved. An ingot for use was made and cast into a melting stock having a diameter of about 100 × 1000 mm. Divide this melting stock into the required amount, and then use the drawing method to obtain a diameter of 9 × 100.
mm round bar-shaped single crystal alloy was cast.

Then, the sample Nos. In the examples, the comparative examples and the conventional example were used. 1 to Sample No. Each test piece having a composition up to 29 was etched with a corrosive solution prepared by diluting hydrochloric acid and hydrogen peroxide solution with glycerin to show that the entire test piece was single-crystallized, and that the growth direction was in the pull-out direction. It was visually confirmed that the angle was within 10 °.

After that, the sample Nos. 1 to Sample No. Each test piece having a composition up to 27 was subjected to the solution treatment and the aging treatment shown in FIG. 1 as follows.

FIG. 1 is a diagram showing a heat treatment sequence of Examples and Comparative Examples.

As shown in FIG. 1, sample Nos. 1 to No. Each alloy up to 27 was first preheated at 1260 ° C. for 1 hour and then
The solution heat treatment was performed at a temperature of 280 ° C. for 5 hours.

After the solution heat treatment, the test piece was sprayed with argon gas and rapidly cooled to room temperature to prepare a γ'phase precipitate.
The step aging heat treatment was performed in the temperature range of 1100 ° C. for 4 hours, and then the test piece was sprayed with argon gas to be rapidly cooled to room temperature. Subsequently, a two-step aging heat treatment for the purpose of stabilizing the γ'phase was carried out in a temperature range of 870 ° C for 20 hours, and then argon gas was similarly blown to the test piece to rapidly cool it to room temperature.

On the other hand, the CMSX-2 of the conventional example was subjected to a partial solution heat treatment and then subjected to a one-step aging heat treatment at a temperature of 1080 ° C. for 4 hours to air-cool the test piece to room temperature, and then at 870 ° C. for 20 hours. After the step aging treatment, the test piece was air-cooled to room temperature.

The IN738LC in the conventional example is 1
After the solution heat treatment at a temperature of 120 ° C. for 2 hours, the test piece was air-cooled to room temperature, and then subjected to a precipitation aging treatment at a temperature of 840 ° C. for 24 hours, and the test piece was air-cooled to a room temperature.

The test pieces thus obtained were subjected to a high temperature corrosion resistance test and a creep rupture test.

For the purpose of evaluating high temperature corrosion resistance, sample N
o. 1 to Sample No. Each test piece having a composition up to 29 was used as a high temperature corrosion test piece having a diameter of 6 × 4.5 mm. Then, the test piece for high-temperature corrosion test was treated with 75% Na ₂
After being immersed in a molten salt having a composition of SO ₄ + 25% NaCl and heated to a temperature of 900 ° C. for 20 hours, descaling was performed and the weight loss due to corrosion was measured. The results are shown in Table 2.
Shown in. The evaluation shows the mass reduction converted into the amount of corrosion erosion.

[0066]

[Table 2]

Table 2 is a table showing the results of high temperature corrosion tests of Examples, Comparative Examples and Conventional Examples.

As shown in Table 2, in the examples within the composition range of the present invention, the alloy CMSX-2 in the conventional example exhibits good high-temperature corrosion resistance, and is almost equivalent to the alloy IN738LC in the conventional example. Showed high temperature corrosion resistance.

Next, in order to evaluate the creep rupture characteristics, sample No. other than IN738LC was tested. 1 to Sample No.
After heat treatment of each test piece having a composition up to 28, the parallel portion 4
It processed into the creep test piece of mm x 30 mm and the total length of 60 mm. And in the atmosphere, temperature 1100 ° C, stress 137M
The creep rupture test was performed at Pa to determine the creep rupture life,
The elongation and the drawing were measured. The results are shown in Table 3.

[0070]

[Table 3]

Table 3 is a table showing the results of creep tests of Examples, Comparative Examples and Conventional Examples.

As shown in Table 3, in the alloys of Examples within the composition range of the present invention, the creep rupture life was 24.24 under the test conditions of the temperature of 1100 ° C. and the stress of 137 MPa in the atmosphere.
It became 5 to 28.2 hours, and showed the same creep rupture life as the alloy CMSX-2 of the conventional example. On the other hand, the sample N of the comparative example which is out of the range of the alloy composition of the present invention
o. In No. 12, since the amount of Co added was small, the phase stability was lowered and the creep rupture life was shortened. In addition, the sample No.
In No. 13, the solid solubility of γ'decreases due to excessive addition of Co, and the γ-γ 'eutectic precipitates, so the creep rupture life is reduced. Sample No. 16, No. 18 and No. 24, Mo, W and T which are solid solution strengthening elements
Since the amount of a added was small and it did not act effectively on the strength, the fracture life was shortened. Sample No. 20 and No. 20. In No. 22, the amount of Al and Ti was small, and the amount of γ'precipitation, which is the main factor for precipitation strengthening, decreased, so that the creep rupture life decreased.

Sample No. in the comparative example. 15, No. 1
7, No. 19, No. 25 and No. 26, C
Re-W, Re-Mo, Re-C which adversely affects the creep rupture life due to excessive addition of r, Mo, W and Ta
TCP (Topologica) such as r-W, α-W and α-Mo
l Closed Packed) phase and needle-like or plate-like precipitates were formed, and the creep rupture life was shortened. Sample No.
21 and No. In No. 23, γ-γ ′ eutectic was generated due to excessive addition of Al and Ti, and it became a crack generation point during creep, and the creep rupture life was shortened. Sample No. In No. 27, the melting point of the alloy was partially lowered by the excessive addition of Hf, and the creep rupture life was shortened.

Therefore, according to the present embodiment, the alloy composition within the range of the present invention makes it possible to obtain a Ni-based single crystal superalloy having excellent creep rupture characteristics and high temperature corrosion resistance.

Second Embodiment (Tables 4 to 6; FIG. 2; Example,
Comparative Example) In the present embodiment, it will be described using Examples and Comparative Examples that the Ni-based single crystal superalloy produced by performing the heat treatment of the present invention has excellent creep rupture properties.

Example (Heat Treatment Nos. 1 to 9) In this example, the heat treatment was carried out at a temperature and a time within the range of the present invention.

First, sample No. 1 shown in Table 1 in the first embodiment. Melting stock was prepared with the alloy composition of No. 1 as the target. Table 4 shows the results of the melting stock analysis. In addition, in Table 4, CMSX which is an alloy of the conventional example is shown.
-2, IN738LC and SC-16 are shown together.

[0078]

[Table 4]

Sample No. shown in Table 4 9 x 100m diameter by drawing method using 1 melting stock
A round bar-shaped single crystal test piece of m was prepared.

The single crystal test piece was etched with a corrosive solution of hydrochloric acid and hydrogen peroxide solution, and the entire test piece was single crystallized, and the growth direction was within 10 ° with respect to the pulling direction. It was visually confirmed.

Then, heat treatment was carried out by the heat treatment sequence shown in FIG. 2 and Table 5.

[0082]

[Table 5]

FIG. 2 is a diagram showing the heat treatment sequences of the example and the comparative example.

As shown in FIG. 2, sample N in the example
o. The alloy No. 1 was first preheated at I ° C. for 1) hour, and then solution heat treated at II ° C. for 2) hour. After the solution heat treatment, the test piece was rapidly cooled to room temperature and γ '
One-step aging heat treatment for the purpose of phase precipitation is performed in the temperature range of III ℃ for 3) hours, then the test piece is rapidly cooled to room temperature, and then two-step aging heat treatment for the purpose of γ'phase stabilization is performed at IV ℃. After carrying out in the temperature range of 4) hours, the test piece was rapidly cooled to room temperature.

As shown in Table 5, the heat treatment conditions of temperatures I to III ° C. and 1) to 4) are as follows. 1 to No. The conditions were up to 9. Specifically, the solution heat treatment temperature is set to 1200 ° C. to 1280 ° C., and 1
The stage aging temperature was 1000 ° C. to 1200 ° C., the total solution heat treatment time was 10 hours or less, and the total aging heat treatment time was 30 hours.

Comparative Example (Heat Treatment Nos. 10 to 13) In this comparative example, the sample No. shown in Table 1 is the same as the example. A melted stock prepared with the alloy composition of No. 1 as the target was used. Then, as shown in Table 5, the heat treatment conditions are heat treatment No. 10 to heat treatment No. 1
Instead of 3, the heat treatment was performed outside the scope of the present invention.

After the heat treatment of the examples and comparative examples,
A creep rupture test was conducted.

In order to evaluate the creep rupture characteristics, heat treatment No. Heat treatment No. 1 to No. 1 Each test piece that has been subjected to heat treatment up to 13 has a parallel part diameter of 4 mm x 30 mm and a total length of 60
It processed into the creep test piece of mm. Then, a creep rupture test was performed in the atmosphere at a temperature of 1100 ° C. and a stress of 137 MPa, and the creep rupture life, elongation and drawing were measured. The results are shown in Table 6.

[0089]

[Table 6]

As shown in Table 6, the heat treatment Nos. Heat treatment No. 1 to No. 1 Each of the test pieces subjected to the heat treatment up to 9 had a creep rupture life of 19.8 to 25.1 hours. 10 to heat treatment N
o. Good creep rupture life was exhibited for each of the test pieces that had been heat treated up to 13.

Therefore, according to the present embodiment, the heat treatment conditions are defined as in the present invention to produce an alloy,
A Ni-based single crystal superalloy having excellent creep rupture properties can be obtained.

Third Embodiment (Tables 4 to 8; FIG. 3; Example,
Conventional Example) In the present embodiment, a comparison between the conventional alloy and the Ni-based single crystal superalloy produced according to the present invention is evaluated particularly by the creep rupture property and the high temperature corrosion resistance.

Example (Table 4; Sample No. 1) In this example, the sample No. 1 shown in Table 4 was used. 1 melting stock was used. This melting stock was treated in the same manner as in the first embodiment, and a round bar-shaped single crystal test piece with a diameter of 9 × 100 mm was produced by a drawing method. Subsequently, first, a preliminary heat treatment was performed at 1260 ° C. for 1 hour, and then a solution heat treatment was performed at a temperature of 1280 ° C. for 5 hours. After the solution heat treatment, the test piece was sprayed with argon gas and rapidly cooled to room temperature to aim for γ'phase precipitation.
The step aging heat treatment was performed in the temperature range of 1100 ° C. for 4 hours, and then the test piece was sprayed with argon gas to be rapidly cooled to room temperature. Subsequently, a two-step aging heat treatment for the purpose of stabilizing the γ'phase was carried out in a temperature range of 870 ° C for 20 hours, and then argon gas was similarly blown to the test piece to rapidly cool it to room temperature.

Thereafter, a cylindrical high temperature corrosion test piece having a diameter of 6 mm × 4.5 mm and a creep test piece having a parallel portion of 4 mm × 20 mm and a total length of 60 mm were processed.

Conventional Example (Table 4; Sample Nos. 28 to 3)
0) In this conventional example, the sample No. 28 CMs
SX-2, sample No. 29 IN738LC and sample no. 30 SC-16 was used.

In order to compare and evaluate the high-temperature corrosion resistance, a cylindrical test piece having a diameter of 6 mm × 4.5 mm was prepared for the conventional alloys IN738LC and CMSX-2 in the same manner as in the examples.

A creep rupture test and a high temperature corrosion resistance test were conducted for each of the alloys of Examples and Conventional Examples.

Sample No. in the examples. In order to evaluate the creep rupture property of 1, the test temperature is 900 ° C in the atmosphere.
A creep rupture test was performed with a test stress of 392 MPa to measure the rupture life. Also, the test temperature is 1100 ° C.
Then, the creep rupture test was conducted by changing the test stress to 137 MPa, and the creep rupture life, elongation and drawing were measured. The measurement results are shown in Table 7.

[0099]

[Table 7]

As shown in Table 7, the test condition is the test temperature 9
At 00 ° C. and a test stress of 392 MPa, the creep rupture life was 118.0 hours, the elongation was 21.8% and the drawing was 37%. Also, the test condition is test temperature 11
At 00 ° C. and a test stress of 137 MPa, the creep rupture life was 24.9 hours, the elongation was 23.1% and the drawing was 49%.

In order to compare the result of the example with the conventional example, the data of CMSX-2 and SC-16 were used as the conventional example. For CMSX-2, "DS AND S
XSUPERALLOYS FOR INDUSTRIAL GAS TURBINE ”: GLEric
kson and K. Harris: Materials for Advanced Power Eng
The value described in innering 1994 was read and used. For SC-16, see "SINGLE-CRYSTAL BLADES
”: D.GOLDSCHMIDT: Materials for Advanced Power En
The value described in ginnering 1994 was read and used.

FIG. 3 is a diagram comparing the creep characteristics of the example and the conventional example. The horizontal axis represents the Larson-Millar parameter, which is a parameter of temperature and time, and the vertical axis represents the stress.

As shown in FIG. 3, the alloys of the examples have a significantly higher creep rupture life than the conventional single crystal alloys CMSX-2 and SC-16 containing 14 to 16% by weight of Cr. Was confirmed to improve.

Next, a high temperature corrosion resistance test was conducted on each of the alloys of Examples and Conventional Examples.

In order to evaluate the high temperature corrosion resistance, the temperature 90
Each test piece was immersed in a molten salt containing 75% Na ₂ SO ₄ and 25% NaCl for 20 hours at 0 ° C., the generated corrosion product was removed with a metal wire brush, and then the weight change was measured. Then, it was converted into the amount of corrosion erosion. The results are shown in Table 8.

[0106]

[Table 8]

As shown in Table 8, the alloys of the examples have the same creep rupture life as the single crystal alloy CM.
Alloy 738L of the conventional example, which has good hot corrosion resistance against SX-2 and also shows good hot corrosion resistance.
It became clear that it has high-temperature corrosion resistance equivalent to C.

Therefore, according to the present embodiment, it is possible to obtain a Ni-based single crystal alloy in which the creep rupture properties and the high temperature corrosion resistance are improved in a better balance than the conventional alloys.

Fourth Embodiment (Tables 9 to 10; FIG. 4; Implementation)
Examples, Comparative Examples) In the present embodiment, the Ni-based unidirectionally solidified alloy produced by unidirectionally solidifying the alloy material having the alloy composition of the present invention has excellent creep rupture properties and high temperature corrosion resistance. Will be described with reference to Examples and Comparative Examples.

Example (Table 9; Sample No. 1a) In this example, the sample No. 1 shown in Table 9 was used. The melting stock of 1a was used.

[0111]

[Table 9]

Sample No. 1a is the sample No. Hf: 0.1% in a melting stock having an alloy composition of 1,
C: 0.1%, Zr: 0.05% and B: 0.05%
Is added. Using this melting stock, a unidirectionally solidified test piece with a round bar shape having a diameter of 9 × 100 mm was prepared by a drawing method. Subsequently, solution heat treatment was performed at 1220 ° C. for 4 hours. After the solution heat treatment, the test piece was blown with argon gas and rapidly cooled to room temperature, and the one-step aging heat treatment for the purpose of γ'phase precipitation was performed at a temperature range of 1100 ° C for 4 hours, and then the test piece was blown with argon gas. And rapidly cooled to room temperature. Subsequently, after performing a two-step aging heat treatment for the purpose of stabilizing the γ'phase in a temperature range of 870 ° C for 20 hours, the test piece was similarly sprayed with argon gas and rapidly cooled to room temperature.

Then, a high temperature corrosion test piece having a cylindrical shape with a diameter of 6 mm × 4.5 mm and a creep test piece with a parallel portion of 4 mm × 20 mm and a total length of 60 mm were processed.

Conventional Example (Table 9: Sample Nos. 31 to 3)
3) In this conventional example, the sample No. 31 CM
247LC, sample no. 32 IN738LC and sample no. 33 IN6203 was used.

In order to compare and evaluate the high temperature corrosion resistance, cylindrical test pieces having a diameter of 6 mm × 4.5 mm were prepared for the conventional alloys IN738LC and CM247LC in the same manner as in the examples.

A creep rupture test and a high temperature corrosion test were carried out for each of the alloys of Examples and Conventional Examples.

Sample No. in the examples. In order to evaluate the creep rupture property of 1a, a test temperature of 900
A creep rupture test was performed at 0 ° C. and a test stress of 392 MPa to measure the rupture life. In addition, the test temperature is 110
The creep rupture test was performed at 0 ° C. and the test stress was changed to 137 MPa to measure the cleave rupture life, elongation and drawing. The measurement results are shown in Table 10.

[0118]

[Table 10]

As shown in Table 10, when the test conditions were a test temperature of 900 ° C. and a test stress of 392 MPa, the creep rupture life was 25.0 hours, the elongation was 24.0% and the drawing was 41%. Also, the test condition is test temperature 11
At 00 ° C. and a test stress of 137 MPa, the creep rupture life was 6.7 hours, the elongation was 28.1% and the drawing was 51%.

In order to compare the result of the example with the conventional example, as a conventional example, a directionally solidified alloy, CM247LC.
And the data of IN6203 was used. In addition, CM24
For 7LC, DS AND SX SUPERALLOYS FOR INDUSTRI
AL GAS TURBINE: GLErickson and K. Harris: Materia
The value described in ls for Advanced Power Engineering l994 was read and used. For IN6203,
"DSIN6203: FROM BIRTH TO APPLICATIONS": Barnard
PM, Allen DH: ThirdInternational Charles Pa
The value described in rsons Turbine Conferences was read and used.

FIG. 4 is a diagram comparing the creep characteristics of the example and the conventional example. The horizontal axis is the Larson mirror parameter, which is a parameter of temperature and time, and the vertical axis is the stress.

As shown in FIG. 4, the alloy of the embodiment is 22
A conventional directionally solidified alloy containing I wt% Cr
It was confirmed that the creep rupture life was significantly improved compared to N6203.

Next, a high temperature corrosiveness test was conducted on each of the alloys of the working row and the conventional example.

In order to evaluate the high temperature corrosion resistance, the temperature 90
Each test piece was immersed in a molten salt containing 75% Na ₂ SO ₄ and 25% NaCl for 20 hours at 0 ° C., the generated corrosion product was removed with a metal wire brush, and then the weight change was measured. Then, it was converted into the amount of corrosion erosion. The results are shown in Table 11
Shown in.

[0125]

[Table 11]

As shown in Table 11, the alloys of the examples have good hot corrosion resistance and good hot corrosion resistance as compared with the unidirectionally solidified alloy CM247LC having the same creep rupture life. Alloy IN of the conventional example showing the property
It became clear that it has high-temperature corrosion resistance equivalent to 738LC.

Therefore, according to this embodiment, it is possible to obtain a Ni-based unidirectionally solidified alloy in which the creep rupture property and the high temperature corrosion resistance are improved in a better balance than the conventional alloy.

[0128]

As described above, the N according to the present invention is
The i-base superalloy and the Ni-base superalloy produced by the method have excellent creep rupture characteristics and high-temperature corrosion resistance in a well-balanced manner, and the Ni-base superalloy can be used for gas turbine rotor blades and stator blades. By applying it to a gas turbine component, it can greatly contribute to the improvement of its efficiency and reliability.

[Brief description of drawings]

FIG. 1 is a diagram showing a heat treatment sequence of an example and a comparative example in the first embodiment of the present invention.

FIG. 2 is a diagram showing a heat treatment sequence of an example and a comparative example in the second embodiment of the present invention.

FIG. 3 is a diagram comparing the creep characteristics of an example and a conventional example in the third embodiment of the present invention.

FIG. 4 is a diagram comparing the creep characteristics of an example and a conventional example in the fourth embodiment of the present invention.

─────────────────────────────────────────────────── ─── Continuation of the front page (51) Int.Cl. ⁷ Identification code FI theme code (reference) C22F 1/00 651 C22F 1/00 651B 682 682 691 691B 691C 691Z 692 692Z (72) Inventor Hiroaki Yoshioka Kanagawa 2-4 Suehiro-cho, Tsurumi-ku, Yokohama-shi Toshiba Keihin Plant (72) Inventor Takao Suzuki 1-1-1 Shibaura, Minato-ku, Tokyo Toshiba Head Office Co., Ltd. (72) Toshiharu Kobayashi Tsukuba, Ibaraki Prefecture 1-2-1, Sengen Ichi, Metal Materials Research Laboratory, Agency for Science and Technology (72) Inventor Tadaharu Yokokawa 1-2-1, Sengen, Sengen, Tsukuba, Ibaraki Prefecture (72) Inventor Yutaka Koizumi, Ibaraki 1-2-1 Sengen, Tsukuba-shi, Japan Metal Science and Technology Research, Agency for Science and Technology In-house (72) Inventor Hirofumi Harada 1-2-1 Sengen, Tsukuba-shi, Ibaraki F-Term (Reference), Institute for Materials Research, Science and Technology Agency 3G002 EA06

Claims (11)

[Claims] 1. By weight%, Co: 10% or more and 14% or less, Cr: 10% or more and 13% or less, Mo: 0.1% or more and 3% or less, W: 4% or more and 8% or less, Al: 4 % Or more 6%
Below, Ti: 1% to 4%, Ta: 4% to 8%, Hf: 0.1% to 0.5% and Re: 2%
A Ni-based superalloy containing the following, the balance consisting of Ni-based and unavoidable impurities. 2. By weight%, Co: 11% or more and 13% or less, Cr: 11% or more and 12.5% or less, Mo: 1% or more and 2% or less, W: 4.5% or more and 6.5% or less. , Al: 4.
5% or more and 5.5% or less, Ti: 2% or more and 3% or less, T
a: 4.5% or more and 6.5% or less, Hf: 0.1% or more and 0.3% or less, and Re: 2% or less, with the balance being N
N characterized by comprising i group and unavoidable impurities
i-based superalloy. 3. The Ni-based superalloy according to claim 1 or 2, wherein C: 0.5% or less and Zr: 0.2% by weight.
The following and B: Ni-based superalloys characterized by being added by 0.1% or less. 4. Ni, Co, Cr, Mo, W, Al, T
A material consisting of i, Ta, Re and Hf is melted and solidified to form a Ni-base superalloy body, and the Ni-base superalloy body is heated at 1200 ° C. in an environment of vacuum or an inert atmosphere.
Solution heat treatment in the temperature range from 1 to 1280 ° C, followed by quenching, followed by 1 stage aging heat treatment in the temperature range from 1000 ° C to 1200 ° C, 2 stages in the temperature range from 700 ° C to 900 ° C N characterized by aging heat treatment
Method for producing i-based superalloy. 5. The method for producing a Ni-based superalloy according to claim 4, wherein Ni, Co, Cr, Mo, W, Al, Ti,
A material comprising Ta, Re and Hf is used.
Alternatively, a method for producing a Ni-base superalloy is characterized in that a Ni-base single crystal superalloy having the composition described in 2 is obtained. 6. The method for producing a Ni-based superalloy according to claim 4, wherein Ni, Co, Cr, Mo, W, Al, Ti,
A method for producing a Ni-base superalloy, characterized in that a Ni-base unidirectionally solidified superalloy having a composition according to claim 3 is obtained using a material consisting of Ta, Re, Hf, C, Zr and B. 7. The method for producing a Ni-base superalloy according to claim 4, wherein the solution heat treatment is 10 times.
The method for producing a Ni-base superalloy is characterized in that the heat treatment is performed within an hour and the aging heat treatment is performed within 30 hours. 8. The method for producing a Ni-base superalloy according to claim 7, wherein the solution heat treatment is performed by changing the temperature to any of 2 to 4 steps up to the final solution heat treatment temperature. Superalloy manufacturing method. 9. The method for producing a Ni-base superalloy according to claim 8, wherein before the solution heat treatment, 1 at a low temperature in the range of 40 ° C. to 60 ° C. with respect to the final solution heat treatment temperature of the solution heat treatment temperature. A method for producing a Ni-base superalloy, which comprises performing a preheating treatment for not less than 2 hours and not more than 2 hours. 10. A gas turbine component whose constituent material is composed of the Ni-based single crystal superalloy or the Ni-based unidirectionally solidified superalloy as described in any one of claims 1 to 3. 11. A gas turbine component made of a Ni-base superalloy manufactured by the manufacturing method according to any one of claims 4 to 9.