EP2703507B1

EP2703507B1 - Ni base alloy and gas turbine blade and gas turbine utilizing the same

Info

Publication number: EP2703507B1
Application number: EP13180298.5A
Authority: EP
Inventors: Takashi Shibayama; Shinya Imano; Hironori Kamoshida
Original assignee: Mitsubishi Hitachi Power Systems Ltd
Current assignee: Mitsubishi Power Ltd
Priority date: 2012-08-30
Filing date: 2013-08-13
Publication date: 2016-01-20
Anticipated expiration: 2033-08-13
Also published as: EP2703507A1; US20140064982A1; JP2014047371A

Description

BACKGROUND OF THE INVENTION

The present invention relates to an Ni base alloy and a gas turbine blade utilizing the same and more particularly to the Ni base alloy having improved high-temperature strength while maintaining the weldablity and corrosion resistance and the gas turbine blade utilizing the same.
With high efficiency of a gas turbine, a request for material capable of resisting more severe operation conditions is increased. Particularly, it is an important subject that the high-temperature strength is improved and harmful phase precipitation is suppressed while ensuring the weldability and corrosion resistance and material for a turbine having these properties is required.
There is a nickel base alloy known generally as GTD-111 ( U.S. Pat. No. 6,416,596 ). Casting material of GTD-111 includes nominal composition containing Cr of 14 mass%, Co of 10 mass%, Mo of 1.5 mass%, W of 3.8 mass%, Ta of 3 mass%, A1 of 3 mass%, C of 0.10 mass%, Ti of 5 mass%, B of 0.02 mass%, Zr of 0.04 mass% and the remainder containing Ni. This alloy is excellent in the weldability and corrosion resistance, although the strength at higher temperature region is low and it is difficult to apply it to a high-efficient gas turbine.
JP-A-2004-197131 discloses a nickel base alloy known as RM02B, which contains Cr of 12.0 to 16.0, Co of 4.0 to 9.0, A1 of 3.4 to 4.6, Nb of 0.5 to 1.6, C of 0.05 to 0.16, B of 0.005 to 0.025 in mass %, Ti, Ta, Mo and W.
This alloy has the high-temperature creep strength, the corrosion resistance and the oxidation resistance in balance. However, the alloy is highly strengthened by addition of precipitation strengthening elements and solid solution strengthening elements, whereas the solid solution temperature of γ'-phase is high and there is a possibility that the weldability is not sufficient.
EP 2 169 087 A2 discloses a Ni-base alloy, which contains Cr of 13.1 to 16.0, Co of 1.0 to 6.8, Al of 3.0 to 3.4, Ti of 4.6 to 5.6, Ta of 2.0 to 4.4, W of 3.5 to 4.9, Mo of 0.1 to 0.9, Nb of 0.3 to 1.4, C of 0.05 to 0.20, and B of 0.01 to 0.03 in mass%, with a remainder being nickel and inevitable impurities.

SUMMARY OF THE INVENTION

It is an object of the present invention to provide an Ni base alloy using GTD-111 as a base and having improved high-temperature strength while maintaining the weldability and corrosion resistance and a gas turbine blade utilizing the same.
There is provided the Ni base alloy containing Al of 2.5 to 3.5%, Co of 1.5 to 5.5%, Cr of 11.8 to 13.8%, Mo of 0.4 to 1.4%, Ta of 3.0 to 5.0%, Ti of 5.1 to 6.1%, W of 3.3 to 4.3%, B of 0.01 to 0.02%, C of 0.08 to 0.12% and remainder containing Ni and inevitable impurities.
According to the present invention, there can be provided the Ni base alloy having improved high-temperature strength while maintaining the weldability and corrosion resistance and a gas turbine blade utilizing the same.
Other objects, features and advantages of the invention will become apparent from the following description of the embodiments of the invention taken in conjunction with the accompanying drawings.

BRIEF DESCRIPTION OF THE DRAWINGS

Fig. 1 is a perspective view illustrating a gas turbine blade;
Fig. 2 is a side view illustrating an appearance of a gas turbine;
Fig. 3 is a diagram showing precipitation amounts (800°C) of γ'-phase;
Fig. 4 is a diagram showing precipitation amounts (1000°C) of γ'-phase;
Fig. 5 is a diagram showing solid solution temperatures of γ'-phase;
Fig. 6 is a diagram showing solid solution temperatures of harmful phase;
Fig. 7 is a diagram showing amounts of Cr in γ-phase; and
Fig. 8 is a diagram showing the relation of the solid solution temperature of the γ'-phase and the precipitation amount of the γ'-phase at 800°C.

DESCRIPTION OF THE EMBODIMENTS

The present invention is now described in detail.
An Ni base alloy of the present invention contains Al of 2.5 to 3.5%, Co of 1.5 to 5.5%, Cr of 11.8 to 13.8%, Mo of 0.4 to 1.4%, Ta of 3.0 to 5.0%, Ti of 5.1 to 6.1%, W of 3.3 to 4.3%, B of 0.01 to 0.02% and C of 0.08 to 0.12% in mass %. The Ni base alloy suppresses precipitation of harmful phase and improves high-temperature strength while maintaining the weldability and oxidation resistance.
Except the above components, an addition amount is substantially 0, although mixing of the following elements (Nb of 0 to 0.2% or less, Hf of 0 to 2.0% or less, Re of 0 to 0.5% or less, Zr of 0 to 0.05% or less, O of 0 to 0.005% or less, N of 0 to 0.005% or less, Si of 0 to 0.01% or less, Mn of 0 to 0.02% or less, P of 0 to 0.01% or less and S of 0 to 0.01% or less) is recognized as inevitable impurities mixed upon manufacturing of alloy within the above range.
The Ni base alloy of the present invention is an alloy of γ'-phase precipitation strengthening type, in which the γ'-phase having Ni₃Al which is intermetallic compound as representation is dispersedly precipitated in the γ-phase which is the matrix phase finely and innumerably. Ti and Ta are also solidly dissolved in the γ'-phase as alloy composition and Ni₃Ti and Ni₃GTa are formed.
The elements are now described. Composition ranges are all expressed by mass percentage (%).

Al: 2.5 to 3.5%

A1 is much contained in the γ'-phase contributing to the high strengthening and forms Ni₃Al which is an intermetallic compound. In order to attain this effect sufficiently, the content thereof larger than or equal to 2.5% is required. However, when 3.5% is exceeded, σ-phase which reduces the strength and the harmful phase such as α-Cr are precipitated to reduce the corrosion resistance. Accordingly, the component range of Al is set to be 2.5 to 3.5%. It if preferable that the content thereof falls within the range of 2.7 to 3.3% in consideration of balance of the precipitation amount of γ' and the solid solution temperature of the harmful phase and a component ratio to other elements.

Co: 1.5 to 5.5%

Co is effective in strengthening solid solution of the γ-phase and improving the corrosion resistance at high temperature. The effects require the content of Co larger than or equal to 1.5%, although in the alloy of the present invention since the solid solution temperature of the γ'-phase is increased with increased amount of Co, the weldability is reduced when Co is added excessively. Accordingly, the upper limit thereof is set to be 5.5% and the preferable component range is 3.0 to 4.0%.

Cr: 11.8 to 13.8%

Cr has the effect that Cr is solidly dissolved in the γ-phase to improve the corrosion resistance at high temperature. When the content thereof is larger than or equal to 11.8%, the sufficient effect is attained, although the solid solution amount of Cr in the γ-phase has an upper limit and when the upper limit is exceeded, there is a possibility that surplus Cr is stabilized as the harmful phase such as α-Cr. Since the harmful phase reduces the high-temperature strength, the content of Cr is set to be 11.8 to 13.8% in consideration of balance with phase fraction of the γ-phase. The preferable range of the content is 12.3 to 13.3%.

Mo: 0.4 to 1.4%

Mo is an element which is solidly dissolved in both of the γ-phase and the y'-phase and contributes to the high-temperature strength. This effect is attained when the content thereof is larger than or equal to 0.4%, although when this element is added excessively, the oxidation resistance and the corrosion resistance are reduced remarkably and accordingly the upper limit is set to be 1.4%. When the balance of these properties is considered, the preferable range is 0.6 to 1.2%.

Ti: 5.1 to 6.1%

Ti is solidly dissolved in the γ'-phase in the form of Ni₃(A, Ti) and strengthens solid solution. Further, Ti is effective in improving the corrosion resistance at high temperature and accordingly the addition amount of Ti is set to be larger than or equal to 5.1%. However, since the excessive addition reduces the oxidation resistance, the upper limit of the content of Ti is set to 6.1%. The preferable content range is set to 5.3 to 5.9% in order to balance the high-temperature strength, the corrosion resistance and the oxidation resistance.

Ta: 3.0 to 5.0%

Ta is solidly dissolved in the γ'-phase which is the precipitation strengthening phase in the form of Ni₃(Al, Ti, Ta) and contributes to the high-temperature strength. In order to attain this effect sufficiently, the composition range thereof is set to 3.0 to 5.0%. However, when the total amount of Al, Ti and Ta which contribute to formation of the γ'-phase greatly is increased excessively, the solid solution temperature of the γ'-phase and the harmful phase is increased remarkably. These reduce the weldability and the high-temperature strength and accordingly the content of Ta is preferably set to 3.5 go 4.5% in consideration of balance with other alloy elements.

W: 3.3 to 4.3%

W is solidly dissolved in the γ-phase which is the matrix phase and the γ'-phase which is the precipitation phase and enhances the creep strength by solid solution strengthening. In order to attain such effects sufficiently, the content larger than or equal to 3.3% is required, although since the specific gravity of W is large, the mass of alloy is increased. Further, when W is added excessively, the corrosion resistance at high temperature is reduced and the strength and toughness by precipitation of α-W which is harmful phase are reduced. Accordingly, the upper limit is set to 4.3%. The preferable range is set to 3.6 to 4.0% in consideration of the high-temperature strength, the corrosion resistance and the phase stability.

B: 0.01 to 0.02%

B is segregated in the crystal grain boundary to improve the grain boundary strength and part of B forms boride such as (Cr, Ni, Ti, Mo)₃B₂ to be precipitated in the grain boundary of the alloy. In order to attain the effect of grain boundary strengthening, the addition amount of B larger than or equal to 0.01 % is required, although the boride generated has the melting point lower than that of the alloy, so that the temperature of the melting point of the alloy is reduced to narrow the temperature range of solution treatment. Accordingly, the upper limit is set to 0.02% and when balance of the strength and the temperature range of solution treatment is considered, the preferable range is 0.012 to 0.018%.

C: 0.08 to 0.12%

C is segregated in the crystal grain boundary to improve the grain boundary strength and part of C forms carbide such as TiC and TaC to be precipitated massively. In order to segregate C in the crystal grain boundary to increase the gain boundary strength, the addition amount of C larger than or equal to 0.08% is required. However, when C is added over 0.12%, excessive carbide is formed to reduce not only the high-temperature strength and ductility but also the corrosion resistance. Further, since the crystallization temperature of carbide upon solidification is increased, carbide is pinned between dendrites and porosities which are defects in casting are produced. The preferable component range is 0.9 to 0.11%.
The following component elements are inevitable impurities.

Nb: 0 to 0.2% or less

Nb is solidly dissolved in the γ'-phase in the form of Ni₃Nb similarly to Ti and strengthens the solid solution. However, in case of alloy having a lot of Ti amount as in the alloy of the present invention, even when Nb is added a little, the solid solution temperature of the harmful phase such as σ-phase is increased remarkably and η-phase which is embrittlement phase is also precipitated. Accordingly, Nb is not added and the content of Nb is set to be substantially 0%. When Nb is mixed, the mixed amount is suppressed to be smaller than or equal to 0.2%.

Hf 0 to 2.0 % or less

Hf does not almost contribute to improvement of strength but improves the corrosion resistance and the oxidation resistance at high temperature by improving adhesion properties of protective scales such as Cr₂O₃ and Al₂O₃ formed on the surface of alloy and when the addition amount of Hf is increased, the adhesion properties of protective scales are improved. However, in the present invention in which a lot of Ti amount is contained, when the addition amount of Hf exceeds 2.0%, a lot of eutectic of Ni₃(Hf, Ti) is formed to reduce the melting point of the Ni base alloy remarkably and make the solution treatment difficult. Accordingly, the upper limit thereof is required to be 2.0%. It is preferable that the upper limit is set to be smaller than or equal to 0.1 % and Hf is not substantially added.

Re: 0 to 0.5% or less

Re can be replaced by part of W if necessary and is an element which is solidly dissolved in the γ-phase to strengthen the solid solution and is effective in improvement of the corrosion resistance. However, Re is expensive and has the large specific gravity, so that the specific gravity of the alloy is increased. Accordingly, the upper limit is required to be 0.5% and it is preferable that it is smaller than or equal to 0.1%.

Zr: 0 to 0.05% or less

Zr is effective in segregating in the crystal grain boundary to enhance the grain boundary strength but Zr almost forms nickel and intermetallic compound Ni₃Zr which are main components of the alloy. This compound reduces the ductility of the alloy and has the remarkably low melting point. Accordingly, the solution processing of the alloy is made difficult and the harmful action is increased. Hence, the upper limit is set to 0.05% and it is preferable that the upper limit is set to be smaller than or equal to 0.01% and Zr is not substantially added.

O: 0 to 0.005% or less and N: 0 to 0.005% or less

O and N are impurities and both of them are often mixed from alloy material. O is mixed even from a melting pot and exists in the alloy as oxide Al₂O₃ and nitrides TiN and AlN massively. When these elements exist in an ingot, they are starting points of crack in creep deformation, so that creep rupture life is reduced and they are starting points of fatigue crack generation, so that fatigue life is reduced. It is preferable that the content of these elements is smaller, although when actual ingot is formed, these elements cannot be reduced to 0 and accordingly the upper limit of both elements is set to be 0.005% as the range in which properties are not deteriorated greatly.

Si: 0 to 0.01% or less

Si is brought from the alloy material. In the present invention, since this element is not an effective element particularly, it is preferable that this element is not contained and when it is contained, the content thereof is suppressed to be smaller than or equal to 0.01%.

Mn: 0 to 0.02% or less

Mn is also brought from the alloy material. Similarly to Si, in the present invention, since this element is not an effective element particularly, it is preferable that this element is not contained and when it is contained, the content thereof is suppressed to be smaller than or equal to 0.02%.

P: 0 to 0.01% or less

P is an impurity. It is preferable that the content thereof is as small as possible and it is necessary to suppress the content thereof to be smaller than or equal to 0.01%.

S: 0 to 0.01% or less

S is also an impurity. It is preferable that the content thereof is as small as possible similarly to P and it is necessary to suppress the content thereof to be smaller than or equal to 0.01%.
Fig. 1 illustrates a gas turbine blade cast by the Ni base alloy containing the above composition elements (Al, Co, Cr, Mo, Ti, Ta, W, B and C). The left of Fig. 1 is a perspective view of a first-stage blade and the right of Fig. 1 is a perspective view of a second-stage blade. The first-stage blade which is heated to highest temperature is formed of directional solidification alloy or single crystal alloy and the heating temperature is reduced as the blades go to second-stage and third-stage, so that the blade is formed of conventional casting alloy. The length of profile is longer as the profile approaches the lower temperature side. Fig. 2 illustrates a gas turbine in which first to fourth blades cast by the Ni base alloy are incorporated.
The following experimental data is based on the result of calculation simulation of phase equilibrium state using database of Ni base alloy.

EMBODIMENT

Table 1 shows chemical composition of the alloys (Al to A28) of the present invention and existing alloys (GTD-111: B1 to B5 and RM02B: C1 to C5) as comparison. The unit of numerical values is all mass percentage (%).

[TABLE 1]

	components (mass %)
	Ni	Al	Co	Cr	Mo	Nb	Ta	Ti	W	B	C
A1	66.285	3	3.5	12.8	0.9	0	4	5.6	3.8	0.015	0.1
A2	66.095	2.9	3.4	12.9	0.7	0	4.3	5.7	3.9	0.015	0.09
A3	66.185	3	3.5	12.8	0.9	0	4	5.7	3.8	0.015	0.1
A4	66.236	3	3.55	12.9	0.95	0	3.9	5.65	3.7	0.014	0.1
A5	66.485	3	3.5	12.8	0.9	0	3.8	5.6	3.8	0.015	0.1
A6	66.585	3	3.1	12.8	1.2	0	4	5.6	3.6	0.015	0.1
A7	66.285	3.1	3.45	12.8	0.9	0	4	5.55	3.8	0.015	0.1
A8	66.775	3.2	3.4	12.8	0.9	0	3.8	5.3	3.7	0.015	0.11
A9	66.085	3	3.5	13	0.9	0	4	5.6	3.8	0.015	0.1
A10	66.435	3.05	3.7	12.9	0.9	0	3.6	5.5	3.8	0.015	0.1
A11	66.785	3	3	12.8	0.9	0	4	5.6	3.8	0.015	0.1
A12	66.176	2.8	3.6	12.7	0.8	0	4.2	5.8	3.8	0.014	0.11
A13	65.785	3	4	12.8	0.9	0	4	5.6	3.8	0.015	0.1
A14	66.094	2.95	3.65	12.75	0.9	0	4.15	5.65	3.75	0.016	0.09
A15	66.185	3	3.5	12.8	1	0	4	5.6	3.8	0.015	0.1
A16	66.586	3.1	3.5	12.8	0.9	0	3.7	5.5	3.8	0.014	0.1
A17	66.385	2.9	3.5	12.8	0.9	0	4	5.6	3.8	0.015	0.1
A18	66.135	3	3.6	12.95	0.9	0	3.8	5.7	3.8	0.015	0.1
A19	66.085	3	3.5	12.8	0.9	0	4.2	5.6	3.8	0.015	0.1
A20	66.227	2.85	3.5	12.8	0.95	0	4.1	5.75	3.7	0.013	0.11
A21	66.385	3	3.5	12.8	0.8	0	4	5.6	3.8	0.015	0.1
A22	66.185	3	3.5	12.8	0.9	0	4	5.6	3.9	0.015	0.1
A23	66.435	13.05	3.4	12.85	0.85	0	3.95	5.5	3.85	0.015	0.1
A24	66.485	3	3.5	12.6	0.9	0	4	5.6	3.8	0.015	0.1
A25	66.185	2.9	3.5	12.8	1.1	0	4.05	5.75	3.6	0.015	0.1
A26	66.385	3	3.5	12.8	0.9	0	4	5.5	3.8	0.015	0.1
A27	66.335	3.05	3.45	12.85	0.9	0	3.95	5.55	3.8	0.015	0.1
A28	66.385	3	3.5	12.8	0.9	0	4	5.6	3.7	0.015	0.1
B1	59.588	3	10	14	1.5	0	3	5	3.8	0.012	0.1
B2	59.488	3.1	10	14.5	1.5	0	2.8	4.7	3.8	0.012	0.1
B3	60.238	2.9	9.5	13.5	1.55	0	3.3	5.2	3.7	0.012	0.1
B4	59.188	3	10.5	14	1.45	0	3.1	4.9	3.75	0.012	0.1
B5	59.688	3.15	10	14	1.5	0	2.75	4.75	3.85	0.012	0.1
C1	62.245	3.95	6.8	13.8	1.75	1.15	2.8	3.35	4	0.015	0.14
C2	62.095	4.05	6.8	13.8	1.7	1.15	2.9	3.3	4.05	0.015	0.14
C3	62.095	3.85	7	13.7	1.75	1.2	2.8	3.45	4	0.015	0.14
C4	61.695	3.95	7.2	13.8	1.8	1.15	2.9	3.35	4	0.015	0.14
C5	62.295	3.9	6.6	13.9	1.75	1.25	2.8	3.4	3.95	0.015	0.14

Table 2 and Figs. 3, 4, 5, 6, 7 and 8 show numerical values and graphs of calculation results of the phase equilibrium state in the composition of Table 1. Table 2 is a list of results. Evaluation of high-temperature strength uses the phase fraction of the γ'-phase as an index. Figs. 3 and 4 show the precipitation amounts of the γ'-phase at 800°C and 1000°C, respectively. Fig. 5 shows the solvus of the γ'-phase. Fig. 6 shows the solvus of the σ-phase and harmful phase such as α-Cr and α-W Fig. 7 shows amounts of Cr in the γ-phase at 800°C. Fig. 8 shows the relation of the solvus of the γ'-phase and the precipitation amounts of the γ'-phase at 800°C.

[TABLE 2]

	precipitation amounts of γ'-phase [%]		γ'-phase	harmful phase		Cr amount in γ-phase [%]
	800°C	1000°C	solvus [°c]	solvus [°C]	kind	800°C
A1	60.24	47.37	1167.91	757.15	α-Cr	26.91
A2	60.29	46.14	1165.67	774.8	α-Cr	27.41
A3	60.82	48.46	1169.64	769.17	α-Cr	27.28
A4	60.28	47.91	1168.47	762.37	α-Cr	27.2
A5	59.85	46.89	1168.05	746.95	α-Cr	26.64
A6	60.06	47.52	1165.82	768.44	α-Cr,α-W	26.85
A7	61.03	48.82	1170.7	770.18	α-Cr	27.35
A8	60.25	47.89	1169.75	737.74	α-Cr	26.53
A9	60.25	47.76	1166.82	771.91	α-Cr	27.41
A10	59.45	46.89	1168.4	739.74	α-Cr	26.6
A11	60.16	47.18	1166.22	762.2	α-Cr	26.82
A12	59.59	43.22	1164.06	734.95	α-Cr	26.12
A13	60.31	46.93	1169.62	715.04	σ,α-Cr	26.99
A14	60.26	46.73	1168.01	761.02	α-Cr	27.08
A15	60.24	47.68	1167.6	760.37	α-Cr	26.94
A16	60.16	47.83	1170.23	748.68	α-Cr	28.77
A17	59.14	44.56	1163.92	738.61	α-Cr	26.28
A18	60.38	47.87	1167.98	772.96	α-Cr	27.33
A19	60.64	47.83	1167.75	767.41	α-Cr	27.18
A20	59.61	44.8	1164.1	742.74	α-Cr	26.4
A21	60.24	46.89	1168.22	753.77	α-Cr	26.87
A22	60.3	47.57	1167.81	761.54	α-Cr	26.94
A23	60.13	47.49	1167.63	757.12	α-Cr	26.9
A24	60.23	46.67	1169.01	742.09	α-Cr	26.41
A25	60	46.83	1166.22	756.62	α-Cr,α-W	26.88
A26	59.66	46.27	1166.13	745.19	α-Cr	26.54
A27	60.39	47.93	1168.56	762.07	α-Cr	27.09
A28	60.19	47.16	1168.01	752.74	α-Cr	26.88
B1	54.96	42.09	1169.85	785.28	σ	26.94
B2	53.78	40.53	1164.5	791.76	σ	27.27
B3	55.72	42.76	1171.28	770.67	σ,α-W	26.33
B4	54.58	41.67	1169.44	775.73	σ	26.75
B5	54.49	41.22	1168.64	772.1	σ	26.57
C1	58.11	47.88	1171.49	814.89	σ-α-W	27.04
C2	58.64	48.46	1173.78	815.56	σ-α-W	27.26
C3	57.86	47.71	1171.43	809.82	σ-α-W	26.75
C4	58.25	48.28	1172.63	829.13	σ-α-W	27.23
C5	58.14	48.02	1170.77	822.23	σ-α-W	27.36

It is understood from the results shown in Table 2 that the precipitation amounts of the γ'-phase of the alloys (A to A) of the present invention are increased at both temperatures of 800°C and 1000°C as compared with GTD-111 (B to B) of existing alloys. Since the precipitation amounts of the γ'-phase in the Ni base alloy contribute to the precipitation strengthening, the strength by the precipitation strengthening is greatly improved in the alloys of the present invention. This reason is that the addition amounts of Ta and Ti are increased. The solvus in the γ'-phase is equivalent or a little low and the amount of Cr in the γ-phase is substantially equivalent. The lower the solvus in the γ'-phase is, the easier the welding is and the solution treatment temperature after welding can be also set to be low. Accordingly, recrystallization and crack are suppressed. The solvus is also increased as the γ'-phase is increased by great addition of Ti, although in the alloys of the present invention, the addition amount of Co is reduced, so that increase of the solvus is suppressed and the weldability is ensured. On the other hand, the amount of Cr in the γ-phase contributes to the corrosion resistance and the effect thereof is decided by the content of Cr. In the alloys of the present invention, since Cr is reduced as a whole but increased γ-phase of the γ'-phase is also reduced, the content of Cr in the γ-phase is substantially equal to GTD-111 and the corrosion resistance is maintained. Further, the main harmful phase precipitated at high temperature is the σ-phase in GTD-111 and α-Cr in the alloys of the present invention which are different in kind, although the solvus thereof is reduced slightly and precipitation of the harmful phase is suppressed.
When the alloys of the present invention are compared with RM02B (C to C) which is another existing alloy, the precipitation amount of the γ'-phase at 1000°C is substantially equal but is slightly increased at 800°C. The alloys of the present invention do not have much A1 amount but Ta and Ti are increased greatly to be solidly dissolved in the γ'-phase, so that the high-temperature strength is increased. The Cr amount in the γ-phase is substantially equal but the solvus in the γ'-phase of RM02B is higher clearly. These are based on the same principle as the comparison with GTD-111 and in the alloys of the present invention Co is reduced to make the solvus in the γ'-phase lower. There are two kinds of harmful phases of RM02B containing mainly the σ-phase and α-W and the solvus of α-W is higher. Since the kind and the temperature of the precipitated harmful phase depend on composition of alloys greatly, the solvus is reduced by about 50°C by adjustment of Cr amount and W amount in the alloys of the present invention.
Generally, in order to improve the high-temperature strength in the Ni base alloy, it is effective to increase the precipitation amount in the γ'-phase by adding Al, Ti, Ta, Nb or the like. However, the addition of these elements increases the solvus in the γ'-phase and the weldability is reduced. Accordingly, heretofore, the alloy is characterized as shown in Fig. 8 in that the strength is suppressed to obtain excellent weldability or the strength is high but the weldablity is a little low.
The alloy of the present invention improves the high-temperature strength while having the same weldability and corrosion resistance as GTD-111. As compared with RM02B, the precipitation amount in the γ'-phase at high temperature is equal to or larger than that of RM02B and the corrosion resistance is substantially equal and the weldability is improved. The solvus of harmful phase is reduced greatly and precipitation of harmful phase is suppressed.

Claims

An Ni base alloy containing Al of 2.5 to 3.5%, Co of 1.5 to 5.5%, Cr of 11.8 to 13.8%, Mo of 0.4 to 1.4%, Ta of 3.0 to 5.0%, Ti of 5.1 to 6.1%, W of 3.3 to 4.3%, B of 0.01 to 0.02%, C of 0.08 to 0.12% in mass % and remainder containing Ni and inevitable impurities, as well as Nb of 0 to 0.2%, Hf of 0 to 2.0%, Re of 0 to 0.5%, Zr of 0 to 0.05%, O of 0 to 0.005%, N of 0 to 0.005%, Si of 0 to 0.01%, Mn of 0 to 0.02%, P of 0 to 0.01% and S of 0 to 0.01% in mass % as optional components.
An Ni base alloy according to Claim 1, wherein Nb is not contained.
An Ni base alloy according to claim 1 or 2, wherein the Ni base alloy precipitates γ'-phase having Ni3Al which is intermetallic compound as representation in γ-phase which is matrix phase.
An Ni base alloy according to any of Claims 1 to 3, wherein the Ni base alloy of precipitation strengthening type contains Al of 2.7 to 3.3%, Co of 3.0 to 4.0%, Cr of 12.3 to 13.3%, Mo of 0.6 to 1.2%, Ta of 3.5 to 4.5%, Ti of 5.3 to 5.9%, W of 3.6 to 4.0%, B of 0.012 to 0.018% and C of 0.09 to 0.11% in mass %.
A casting product utilizing the Ni base alloy according to any of Claims 1 to 4.
A gas turbine blade (2, 11-14) utilizing the casting product according to Claim 5.
A gas turbine utilizing the gas turbine blade (2, 11-14) according to Claim 6.