JP2862487B2 - Nickel-base heat-resistant alloy with excellent weldability - Google Patents

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a nickel-base heat-resistant alloy used as a material for forming turbine vanes and other high-temperature components of a gas turbine.

[0002]

2. Description of the Related Art Heat-resistant alloys conventionally used as materials for high-temperature components such as turbine stationary blades of gas turbines include an intermetallic compound Ni ₃ (Al, Ti), ie, precipitation strengthening by a γ ′ phase and Mo, W, etc. There is a Ni-based alloy which also has a solid solution strengthening by the above and a Co-based alloy which is precipitation strengthened by a carbide. In the case of Ni-base alloys, generally, the higher the precipitation amount of the γ 'phase is, the higher the high-temperature strength is, but the weldability tends to decrease. It is also evident from the fact that the weldability of JP-B-54-6968 is very poor, and the alloy having a reduced γ 'phase and improved weldability (JP-A-1-104738) has a remarkably low high-temperature strength. On the other hand, a Co-based alloy generally has good weldability, but has a low high-temperature strength and cannot expect remarkable improvement. From the above, since the high-temperature strength of the Co-based alloy has a limit, the weldability of the Ni-based alloy must be improved without impairing the high-temperature strength.

[0003]

SUMMARY OF THE INVENTION In order to simultaneously improve the weldability without impairing the high-temperature strength of a Ni-based alloy, A
l, Ti, and other elements such as W, C, and Zr are adjusted without decreasing the content of γ'-phase forming elements such as Ti, and the intended purpose is, for example, a welded structure and used at a high temperature. Alloys that can be used for stationary blades, welding structure equipment, etc. of gas turbines. The performance of such an alloy is 900 ° C. × 20 kgf /
mm ² creep rupture time is 110 hours or more and 5 × 6
TIG welding was performed using a 0 × 100 mm test piece under the welding conditions of a welding current of 100 A, a welding voltage of 12 V, and a welding speed of 1.67 mm / S, and an additional strain (total strain) of 0.25% or 0.77%. It is characterized in that the maximum crack length when performing a balest train test is 0.8 mm or less.

[0004]

Means for Solving the Problems As a result of earnest studies, the present inventors have added Cr and Co within a range in which no harmful phases such as σ phase and μ phase are formed as in the alloy composition described later, and γ ′ The addition of Al, Ti, Nb, Ta, etc., which is a phase forming element, and the addition of W, Mo, etc. as solid solution strengthening elements increase the high-temperature strength, while appropriately adding C, Zr, and B, which are easily segregated at grain boundaries. By adding to improve the weldability, an alloy having excellent high-temperature strength and weldability was obtained. Furthermore, it can also be used as a material for high-temperature components used under low-grade fuels such as heavy oil, that is,
The present inventors have found that a Ni-based alloy having excellent oxidation resistance and corrosion resistance can be adjusted, and have reached the present invention.

[0005] That is, the present invention relates to the method of
5 to 0.25% C, 18 to 25% Cr, 15 to 25
% Co, one or two of up to 3.5% Mo and up to 5-10% W in an amount where the value of W + 1 / 2Mo is 5-10%, 1.0-5.0% Ti, 1.0-4.0% Al, 0.5-4.5% Ta, 0.2-3.0% N
b, 0.005 to 0.10% Zr and 0.001 to 0.
Containing 0.1% of B, the balance consisting of Ni and unavoidable impurity elements, the amount of (Al + Ti) and (W + 1 / 2M)
o) The amount is the point A (Al + Ti: 3%, W + 1) in FIG.
/ 2Mo: 10%), point B (Al + Ti: 5%, W + 1)
/ 2Mo: 7.5%), point C (Al + Ti: 5%, W +
1/2 Mo: 5%), point D (Al + Ti: 7%, W + 1)
/ 2Mo: 5%), point E (Al + Ti: 7%, W + 1 /
(2Mo: 10%) having a composition within a range surrounded by a line connecting the points in sequence.
And 0.05-0.25% C, 10% by weight.
-20% Cr, 15-25% Co, W + 1 / 2Mo
One or two of Mo up to 3.5% and W up to 0.5-10% of the amount in which the value of
5.0% Ti, 1.0-4.0% Al, 0.5-
4.5% Ta, 0.2-3.0% Nb, 0.005
0.10% of Zr and 0.001 to 0.01% of B, with the balance being Ni and unavoidable impurity elements,
The amounts of (Al + Ti) and (W + 1 / 2Mo) are shown in FIG.
At point A (Al + Ti: 3%, W + ／ Mo:
10%), point B (Al + Ti: 5%, W + 1 / 2Mo:
7.5%), point C (Al + Ti: 5%, W + ／ M)
o: 5%), point F (Al + Ti: 4%, W + 1 / 2M)
o: 5%), point G (Al + Ti: 4%, W + 1 / 2M)
o: 0.5%), point H (Al + Ti: 7%, W + 1/2)
Mo: 0.5%) Point E (Al + Ti: 7%, W + 1/2)
(Mo: 10%) is a nickel-base heat-resistant alloy having a composition within a range surrounded by a line connecting points in sequence.

Next, the reason for limiting the action and addition amount (by weight) of each element in the alloy composition of the Ni-base heat-resistant alloy of the present invention will be described. C forms carbides, especially at grain boundaries,
Precipitates at dendrite boundaries and strengthens grain boundaries and resinous boundaries. If it is less than 0.05%, there is no strengthening effect,
%, Ductility and creep strength decrease. A particularly preferred range is from 0.09 to 0.23%. Cr is 18 to 25% in claim 1 and 10 to 20% in claim 2. However, Cr is an element that imparts oxidation resistance and corrosion resistance at high temperatures. If each exceeds the upper limit, there is a danger of σ phase formation during long-term high-temperature operation. In addition, the alloy of the first aspect particularly considers corrosion resistance and oxidation resistance, and the second aspect particularly considers high-temperature strength.

[0007] Co has the effect of increasing the limit of solid solution of Ti, Al, etc. that forms the γ 'phase in a substrate at a high temperature (solid solution limit). It is necessary to adopt 15.0% or more. on the other hand,
In order to avoid the danger of σ phase formation, the content is set to 25.0% or less. T
i is an element necessary for the precipitation of the γ 'phase for increasing the high-temperature strength. If the content is less than 1.0%, the required strength cannot be satisfied. Therefore, the content was set to 5.0% or less. Al, like Ti, forms a γ 'phase, increases the high-temperature strength, and contributes to imparting oxidation resistance and corrosion resistance at high temperatures. It is necessary that the amount is 1.0% or more. If too much is added, ductility and weldability are impaired.
It was as follows. In particular, Al + Ti is preferably in the range of 3.0 to 7.0%. W and Mo have the functions of solid solution strengthening and weak precipitation strengthening, and contribute to imparting high-temperature strength. In order to obtain the effect, W + 1 / 2Mo must be 0.5% or more, and if added too much, ductility is inhibited. Therefore, W is 10% or less, Mo is 3.5% or less, and W + 1 / 2Mo is 10% or less. It was as follows.

[0008] Ta and Nb contribute to the improvement of high-temperature strength by solid solution strengthening and γ 'phase precipitation strengthening, and Ta and Nb have an effect of 0.5% or more and Nb of 0.2% or more. On the other hand, if too much is added, the ductility decreases, so that Ta is 4.5% or less and N
b is 3.0% or less. Particularly, Ta is 1.0 to 4.2%.
Is preferable, and Nb is preferably in a range of 0.5 to 1.5%.
Zr has the effect of increasing the bonding force at the crystal grain boundaries to strengthen the grain boundaries, but if it is less than 0.005%, no improvement in creep strength is observed, and if it exceeds 0.10%, the weldability is reversed. Must be present in the range of 0.005 to 0.10%. A particularly preferred range is 0.01
０．１0.10%. B strengthens the grain boundary by increasing the bonding force of the grain boundary similarly to Zr, but if less than 0.001%, no improvement in creep strength is observed, and if it exceeds 0.01%, the weldability is reversed. , The B content is 0.001
-0.01%.

The reason why the range is limited to the range surrounded by the line in FIG. 1 is as follows. Al and Ti increase the high-temperature strength by precipitating the γ 'phase, which is a strengthening factor of the Ni-based alloy, that is, Ni ₃ (Al, Ti). However, excessive addition lowers weldability and ductility, so that Al + Ti is 7% or less. And If the amount added is small, the effect of increasing the high-temperature strength is reduced,
As shown in FIG. Since the Cr content also affects the high-temperature strength, Al + T is considered in consideration of the Cr content.
The lower limit of i was 4% as shown in FIG. W and Mo have the effects of solid solution strengthening and precipitation strengthening by carbides, and have the effect of increasing high-temperature strength. W + 1 /
2Mo requires 0.5% or more, and if added too much,
The upper limit of W + 1 / 2Mo is set to 10% because it promotes precipitation of a harmful phase such as the σ phase and reduces ductility and strength.

[0010]

Now, the present invention will be described in further detail with reference to specific examples. Example 1 Tables 1 and 2 show chemical compositions (% by weight) of alloys invented for typical gas turbine stationary blades. Tables 3 and 4 show the chemical compositions of comparative alloys which are conventional alloys. For each composition, a 20 kg steel ingot was melted in a vacuum high-frequency melting furnace. The samples were lost wax precision cast using them as master ingots, 1160 ° C × 4hr + 1000 ° C × 6hr +
A heat treatment at 800 ° C. for 4 hours was performed. Thereafter, creep rupture test pieces of 6.25 φ × 25 mm in parallel portions and balest strain test pieces of 5 × 60 × 100 mm were produced by machining. Nos. 1 to 18 in Table 1 are the alloys of the present invention.
X, Y, Z, 19 to 36 are comparative alloys. Note that X,
Y is an example of the aforementioned Japanese Patent Publication No. 54-6968, and Z is an example of the Japanese Patent Application Laid-Open No. 1-14738.

[0011]

[Table 1]

[0012]

[Table 2]

[0013]

[Table 3]

[0014]

[Table 4]

FIG. 1 shows (Al + Ti) for each sample.
The relationship between the amount and the (W + 1 / 2Mo) amount is shown. In addition, in parentheses of each sample number, 900 ° C., 20 kgf / mm
² shows the creep rupture time. In FIG. 1, the alloy of the present invention is indicated by a white circle (○), and the comparative alloy is indicated by a black circle (●). Al + Ti and W in the line connecting points ABCDE
+ 1 / 2Mo high alloys of the present invention (1,4,11,12,
13, 14, 15, 16) show high strength,
In particular, No. 11 shows high strength. Also, Al + Ti, W +
ＭMo is in the line connecting the points FGHD and the alloy of the present invention (2,3,5,6,7,8,9,10,1) having a low Cr content
7, 18) show particularly high strength. FIG. 2 shows comparative samples Y, Z, and 20 of Table 2 which are representative samples, and alloys 9 and 11 of the present invention shown in Table 1 at 900 ° C., 20 kgf / mm ² and 98.
A comparison of creep rupture strength at 0 ° C. and 10 kgf / mm ² is shown. The horizontal axis is the Larson Miller parameter P = T
_k (20 + logt) × 10 ⁻³ , T _k : test temperature (°
K), t: Rupture time (Hr) is used. The test results at 900 ° C and 980 ° C show that the stress on the vertical axis is 20 kgf each.
/ Mm ² , 10 kgf / mm ² . The greater the parameter P on the horizontal axis, the greater the strength. The Larson Miller parameters of the alloys Nos. 9 and 11 of the present invention are larger than those of the comparative alloys No. Y, Z and No. 20 at the same test stress. This is the effect of increasing the amount of Al + Ti and the amount of W + 1 / 2Mo and decreasing the amount of Cr (No. 11). On the other hand, the comparative alloy No. Y, in which the amount of Al + Ti is slightly larger than that of No. 9 but has a larger amount of Cr, and the amount of Al + Ti is lower, but W + 1/2 Mo
Comparative alloy No. 20 with a large amount, Al + Ti amount, W + 1/2
The Larson-Miller parameters, such as the comparative alloy No. Z, which have a low Mo content, are on the lower side of the alloy of the present invention at the same test stress.

The weldability was evaluated by a balest train test (FIG. 6). That is, the welding current 100
A, TIG welding under welding conditions of a welding voltage of 12 V and a welding speed of 1.67 mm / s, applying a total strain of 0.25% or 0.77%, and an index of the embrittlement area at the time of welding generated at that time Was measured for the maximum crack length. The maximum crack length and creep rupture time (900 ° C x 20 kgf / m
FIG. 3 shows the relationship of m ² ). The vertical axis in the figure indicates that the smaller the value, the better the weldability. Therefore, the weldability is higher and the high-temperature strength is higher at the lower right of FIG.
The alloys No. 3, 7, 9, 10, 11, 12, and 15 of the present invention in which Zr is 0.1% or less and B is 0.01% or less are all small in valestrain maximum crack length. Especially No. 9,
Nos. 11 and 12 have a target value of 0.3 mm or less, and a creep rupture time of 185 hours or more, exhibiting excellent properties. On the other hand, in the comparative alloy, No. X, Y, 25, 27, 28,
Nos. 33 and 35 showed a creep rupture time of 110 hours or more, but in both cases, the length of the maximum valestrain crack was 0.3 hours.
The target was not satisfied with 8 mm or more. From the above results, Z
Even if the amount of r and the amount of B are reduced, Al + Ti and W + 1 / 2Mo
Is reduced by reducing the range of ABCDE or by lowering the amount of Cr.
By setting the relationship between l + Ti and W + 1 / 2Mo within the range of ABCFGHE, it is possible to improve weldability and increase creep strength.

Example 2 A stationary blade for a gas turbine shown in FIG. 4 was manufactured by the lost wax precision casting method using an alloy of No. 11 shown in Table 1 of Example 1 and solution-treated at 1160 ° C. × 4 hours. After that, a welding test was performed. This vane has a wing width of about 200 mm,
It has a height of about 200 mm and has a hollow structure with an air passage for cooling. The positions of 1, 2, 3 and 4 on the abdominal side of the wing and the positions 5 and 6 of the leading edge and the position 7 of the trailing edge shown in FIG. 4 are overlay welding, and the inner shroud 8 is shown in FIG. The shroud portion 8 (the alloy No. 11 of the present invention) and the cover plate 10 (the Hastelloy X alloy) are square-welded with Hastelloy W alloy 11 by TIG.
I went in. After welding, an appearance inspection at each position, a fluorescent penetration deep scratch test, and observation of a microstructure of a cross section at a position as shown in FIG. 5 were performed, but no crack was observed at any position. The same gas turbine vane as described above was manufactured using comparative alloy Y (Japanese Patent Publication No. 54-6968), and a welding test was conducted. As a result, a number of cracks were observed in the fluorescence penetrant inspection test. A crack having a length of about 1 mm was observed.

[0018]

As described above, according to the present invention, a Ni-based heat-resistant alloy having higher high-temperature strength and excellent weldability than conventional Ni-based heat-resistant alloys can be obtained. This Ni-base heat-resistant alloy is particularly suitable for a gas turbine stationary blade material that requires reliability as the temperature of the gas turbine increases.

[Brief description of the drawings]

FIG. 1 is a view showing a range of an alloy of the present invention and a test result of a creep rupture time.

FIG. 2 is a diagram showing a comparison of creep rupture strengths of match gold.

FIG. 3 is a diagram showing a relationship between a maximum crack length and a creep rupture time in a balest train test.

FIG. 4 is a perspective view of a gas turbine vane to which a welding test is performed by applying the alloy of the present invention.

FIG. 5 is an explanatory diagram of a welded portion in a welding test.

FIG. 6 is an explanatory view of a procedure of a balest train test performed for evaluating the weldability of the alloy of the present invention and a comparative alloy.

[Explanation of symbols]

 1-7 welding bead 8 inner shroud 9 outer shroud 10 cover plate 11 Hastelloy W alloy 12 balest strain test piece (before bending strain addition) 13 yoke 14 bead 15 welding torch 16 balest strain test piece (after bending strain addition) 17 bending block

──────────────────────────────────────────────────続 き Continuing on the front page (72) Inventor Takashi Kawabata Hirade Industrial Park 1, Utsunomiya City, Tochigi Prefecture Mitsubishi Steel Corporation Utsunomiya Works (72) Inventor Takemi Kokubu 1 Hirade Industrial Park, Utsunomiya City, Tochigi Prefecture Mitsubishi Inside Steel Utsunomiya Works (72) Inventor Toshio Mochizuki Tochigi Prefecture Utsunomiya City Hirade Industrial Park 1 Mitsubishi Steel Corporation Utsunomiya Works (72) Inventor Shuichi Sakashita Tochigi Prefecture Utsunomiya City Hirade Industrial Park 1 Mitsubishi Steel Co., Ltd. (72) Inventor Hisakataka Kawai 2-1-1 Shinama, Arai-machi, Takasago City, Hyogo Prefecture Inside Mitsubishi Heavy Industries, Ltd.Takasago Research Laboratory (72) Inventor Ikuo Okada 2-1-1, Araimachi Shinama, Takasago-shi, Hyogo Mitsubishi (72) Inventor Ichiro Tsuji 2-1-1 Shinama, Arai-machi, Takasago-shi, Hyogo Mitsubishi Heavy Industries, Ltd. Takasago Works (72) Inventor Koji Takahashi 2-1-1, Araimachi, Takasago, Hyogo Prefecture Mitsubishi Heavy Industries, Ltd. Takasago Works (72) Inventor Yasuharu Torikoshi 2-1-1, Araimachi Shinama, Takasago-shi, Hyogo Prefecture No. 1 Inside the Takasago Machinery Works, Mitsubishi Heavy Industries, Ltd. (56) References JP-A-56-81661 (JP, A) JP-A-1-104738 (JP, A) JP-A-51-18210 (JP, A) 7-286730 (JP, A) JP 54-6968 (JP, B2) (58) Fields investigated (Int. Cl. ⁶ , DB name) C22C 19/03

Claims (2)

(57) [Claims] (1) 0.05 to 0.25% by weight of C,
18-25% Cr, 15-25% Co, W + ／
One or two of Mo up to 3.5% and W up to 5-10%, in amounts where the value of Mo is 5-10%, 1.0-5.
0% Ti, 1.0-4.0% Al, 0.5-4.5
% Ta, 0.2-3.0% Nb, 0.005-0.
Containing 10% Zr and 0.001-0.01% B;
The balance consists of Ni and unavoidable impurity elements,
+ A) and (W + 1 / 2Mo) are points A (Al + Ti: 3%, W + 1 / 2Mo: 10%) in FIG.
Point B (Al + Ti: 5%, W + 1 / 2Mo: 7.5)
%), Point C (Al + Ti: 5%, W + 1 / 2Mo: 5)
%), Point D (Al + Ti: 7%, W + 1 / 2Mo: 5)
%), Point E (Al + Ti: 7%, W + 1 / 2Mo: 10)
%). A nickel-base heat-resistant alloy having excellent weldability, characterized by having a composition within a range surrounded by a line connecting each point of (%). 2. C. of 0.05 to 0.25% by weight,
10-20% Cr, 15-25% Co, W + ／
M up to 3.5% of the amount where the value of Mo is 0.5-10%
o and one or two of W up to 0.5-10%;
0-5.0% Ti, 1.0-4.0% Al, 0.5
-4.5% Ta, 0.2-3.0% Nb, 0.00
It contains Zr of 5 to 0.10% and B of 0.001 to 0.01%, and the balance consists of Ni and unavoidable impurity elements, and the (Al + Ti) amount and (W + 1 / 2Mo) amount are:
In FIG. 1, point A (Al + Ti: 3%, W + ／ M)
o: 10%), point B (Al + Ti: 5%, W + 1 / 2M)
o: 7.5%), point C (Al + Ti: 5%, W + 1/2)
Mo: 5%), point F (Al + Ti: 4%, W + ／ M)
o: 5%), point G (Al + Ti: 4%, W + 1 / 2M)
o: 0.5%), point H (Al + Ti: 7%, W + 1/2)
Mo: 0.5%) Point E (Al + Ti: 7%, W + 1/2)
(Mo: 10%). A nickel-based heat-resistant alloy having excellent weldability, characterized by having a composition within a range surrounded by lines sequentially connecting the respective points of (Mo: 10%).