JP2008297579A - Nickel-based alloy excellent in structural stability and high tension strength, and method for producing nickel-based alloy material - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a Ni-based alloy having stable high temperature characteristic even at a using temperature of ≥700°C. <P>SOLUTION: This Ni-based alloy contains 0.005-0.1% C, 8-15% Cr, 1-7% Mo, 5-20% W, 0.5-1.0% Al 1.0-2.5% Ti and if desired, at least one or more kinds of ≤1.0% Nb, ≤0.015% B and ≤0.01% Mg, and the balance Ni with inevitable impurities. The structure is made stable even at high temperature of ≥700°C and the excellent high temperature characteristic is obtained. By regulating the balance of Mo and W to 0.1≤Mo/(Mo+W)≤0.5, this alloy is made more excellent in the high temperature strength, high temperature ductility and creep characteristic. By regulating the balance of Al and Ti to 0.2≤Al/Ti≤1.0, this alloy is made more excellent in the high temperature tensile characteristic and the creep characteristic. <P>COPYRIGHT: (C)2009,JPO&INPIT

Description

  The present invention is used, for example, as a material for a generator member exposed to high temperatures such as a turbine rotor, and has a Ni base having excellent structure stability, good high-temperature strength and ductility, and creep characteristics, particularly in a high temperature range. The present invention relates to a method for producing an alloy and a Ni-based alloy material.

  From the viewpoints of reducing the consumption of fossil fuels and preventing global warming, there are expectations for further improvement in the efficiency of USC (ultra-supercritical pressure) plants. Particularly in recent years, there has been a movement toward high-efficiency coal-fired power generation as a power plant in the 21st century, and development of turbine rotors and boiler components that support next-generation ultra-supercritical steam power generation with a main steam temperature exceeding 700 ° C. Is underway.

  The heat-resistant material used in the turbine rotor material exposed to high-temperature steam exceeding 700 ° C. can no longer be used in the conventional ferritic heat-resistant steel from the viewpoint of the service temperature, and a Ni-based alloy must be applied. I don't get it.

In order to obtain good high-temperature strength, a Ni-base heat-resistant alloy is added with a small amount of Ti, Al, or Nb, and a gamma prime phase (hereinafter referred to as γ) consisting of Ni ₃ (Al, Ti) in an austenite (hereinafter referred to as γ) matrix There are many precipitation strengthening type alloys that reinforce and precipitate a precipitation phase called a gamma double prime phase (denoted as γ ″) composed of Ni ₃ (Al, Ti, Nb) and / or Ni ₃ (Al, Ti, Nb). Inconel (trademark, the same applies hereinafter) 706 and 718 are examples of this.

Moreover, in addition to precipitation strengthening of the γ ′ phase, such as Waspaloy, alloys of the type that are strengthened in combination by solid solution strengthening and dispersion strengthening of M ₂₃ C ₆ carbide, and precipitation strengthening elements as represented by Inconel 617 There is also a so-called solid solution strengthening type alloy that does not contain Al and is strengthened by Co or Mo.

However, in any type of Ni-base alloy, the coefficient of linear expansion is higher than that of conventionally used ferritic heat-resistant steel. The problem of thermal fatigue at times has been pointed out.
For this reason, in Patent Document 1, Patent Document 2, or Patent Document 3, a precipitation-strengthened Ni-based alloy having a low coefficient of linear expansion equivalent to that of a ferritic heat resistant steel and exceeding the high temperature material properties of the ferritic heat resistant steel. Has been proposed.
Japanese Patent Laid-Open No. 9-157779 JP 2003-13161 A JP 2005-314728 A

On the other hand, the turbine rotor member must be a material with stable performance in high-temperature and long-time use so that a predetermined 100,000-hour creep rupture strength is required, not only strength and toughness in the operating temperature range, Stability is extremely important. The factor that has the greatest effect on the structural stability of the Ni-based alloy is the high-temperature stability of the strengthening precipitation phase itself, which is almost determined by the growth rate and transformation behavior of the strengthening precipitation phase.
The precipitation-strengthened Ni-based alloy of Patent Document 1 proposed above maintains a stable performance in a relatively short-time use environment at 600 to 700 ° C., which is conventionally assumed. However, when used for a long time of the order of tens of thousands to 100,000 hours in the temperature range of 700? Or higher as described above, precipitation proceeds during use, and overaging causes mechanical properties such as strength, ductility and toughness. There is a problem in that it can be used stably at 700 ° C. or more for a long period of time due to a problem caused by tissue stability, such as markedly lowering of mechanical properties.
In addition, in the precipitation strengthened Ni-base alloys of Patent Documents 2 and 3, the addition amounts of Mo and W are increased in order to reduce the linear expansion coefficient, and there is a problem from the viewpoint of structure stability.

  The present invention has been made in view of the above circumstances, and an object of the present invention is to provide a Ni-based alloy and a method for producing a Ni-based alloy material having stable high-temperature characteristics even at a use temperature of 700 ° C. or higher.

  Precipitation strengthening elements such as Al, Ti, and Nb added to Ni-base alloys and solid solution strengthening elements such as Mo and W form various intermetallic compounds depending on their combination and content, and also have a significant effect on mechanical properties. To do. The inventors of the present application proceeded with research on the formation of various intermetallic compounds, and by appropriately setting the content and balance of the above components, the stability of the structure was obtained even at a high temperature of 700 ° C. or higher, and an excellent high temperature The inventors have found that the present invention has characteristics and have completed the present invention.

  That is, among the Ni-based alloys having excellent structure stability and high temperature strength according to the present invention, the invention according to claim 1 is mass%, C: 0.005 to 0.1%, Cr: 8 to 15%, Mo: 1 to 7%, W: 5 to 20%, Al: 0.5 to 1.0%, Ti: 1.0 to 2.5%, the balance is made of Ni and inevitable impurities And

  The invention of the Ni-base alloy having excellent structure stability and high-temperature strength according to claim 2 is characterized in that, in the invention according to claim 1, it contains, in mass%, Nb: 1.0% or less. .

  The invention of the Ni-base alloy having excellent structure stability and high temperature strength according to claim 3 is the invention according to claim 1 or 2, wherein the invention further contains B: 0.015% or less by mass%. Features.

  The invention of the Ni-base alloy having excellent structure stability and high-temperature strength according to claim 4 is the invention according to any one of claims 1 to 3, and further contains Mg: 0.01% or less by mass. It is characterized by that.

The invention of the Ni-based alloy having excellent structure stability and high temperature strength according to claim 5 is the invention according to any one of claims 1 to 4, wherein the Mo content and the W content satisfy the following formula 1. It is characterized by satisfying.
0.1 ≦ Mo / (Mo + W) ≦ 0.5 Formula 1

The invention of the Ni-based alloy having excellent structure stability and high temperature strength according to claim 6 is the invention according to any one of claims 1 to 5, wherein the Al content and Ti content satisfy the following formula 2. It is characterized by that.
0.2 ≦ Al / Ti ≦ 1.0 Formula 2

  The invention of the method for producing a Ni-base alloy material excellent in structure stability and high-temperature strength according to claim 7 is a method of reusing a Ni-base alloy having the composition according to any one of claims 1 to 6 after hot forging. The solution treatment is performed at a temperature higher than the crystal temperature, and then the first aging treatment is performed in the range of 800 ° C. to 1000 ° C., and then the second aging treatment is performed in the range of 720 ° C. to 780 ° C. .

  The invention of the method for producing a Ni-based alloy material having excellent structure stability and high temperature strength according to claim 8 is the invention according to claim 7, wherein after the first aging treatment, the cooling rate is 20 ° C./h or more. It is characterized in that the aging treatment is continuously performed after cooling to the second aging treatment temperature.

  The reason for setting the alloy composition and production conditions of the present invention will be described below. In addition, all the following contents are shown by the mass%.

<Alloy composition>
C: 0.005-0.1%
C forms TiC with Ti, and forms M ₆ C and M ₂₃ C ₆ type carbides with Cr and Mo, suppressing coarsening of alloy crystal grains and improving high-temperature strength. Also contribute. Further, M ₆ C and M ₂₃ C ₆ are essential elements in the present invention in order to strengthen the grain boundary by precipitating an appropriate amount of carbides at the crystal grain boundary. If C is not contained in an amount of 0.005% or more, the above effect cannot be obtained. If it exceeds 0.1%, not only the amount of Ti required for precipitation strengthening is reduced but also Cr carbides precipitated at grain boundaries during aging treatment. Increases too much and grain boundaries become brittle and ductility decreases. Therefore, the addition amount of C is limited to a range of 0.005 to 0.1%. For the same reason, it is desirable that the lower limit is 0.01% and the upper limit is 0.08%.

Cr: 8-15%
Cr is an indispensable element for increasing the oxidation resistance, corrosion resistance, and strength of the alloy. Moreover, it couple | bonds with C and precipitates a carbide | carbonized_material and raises high temperature strength. In order to exhibit these effects, an addition amount of at least 8% is necessary. However, an excessive amount of addition inhibits the stability of the matrix, promotes the generation of harmful TCP phases such as σ phase and α-Cr, and adversely affects ductility and toughness. It is also known that the linear expansion coefficient increases as the Cr content increases. Therefore, the addition amount of Cr is limited to a range of 8 to 15%. For the same reason, it is desirable to set the lower limit to 9% and the upper limit to 14%.

Mo: 1-7%
Mo is mainly effective as a solid solution strengthening element for solid solution in the matrix to strengthen the matrix itself, and the stability of the γ 'phase by dissolving in the γ' phase and replacing the Al site of the γ 'phase. Is effective in increasing the strength at high temperatures and increasing the stability of the tissue. If Mo is less than 1%, the above effect is insufficient, and if it exceeds 7%, a TCP phase called a μ phase (Laves phase) is likely to be generated. Impairs tissue stability. Therefore, the addition amount of Mo is limited to a range of 1 to 7%. For the same reason, it is desirable to set the lower limit to 2% and the upper limit to 7%.

W: 5-20%
Like Mo, W is effective as a solid solution strengthening element for solid solution in the matrix and strengthening the matrix itself, and by dissolving in the γ ′ phase and replacing the Al site of the γ ′ phase, Since the stability is increased, it is effective to increase the strength at high temperature and the stability of the tissue. In addition, it has an effect of lowering the linear expansion coefficient. If the content is appropriate, the TCP phase does not precipitate, and the structural stability is not impaired. However, if it is added too much, not only α-W is precipitated and the structure stability is lowered, but also hot workability is remarkably deteriorated. Therefore, the addition amount of W is limited to a range of 5 to 20%. For the same reason, it is desirable to set the lower limit to 7% and the upper limit to 15%.

0.1 ≦ Mo / (Mo + W) ≦ 0.5
Mo and W can obtain the above-mentioned effects, but in order to ensure mechanical properties at high temperatures, particularly ductility, [Mo content] / [Mo content + W content] is set to 0.1 or more. Is desirable. On the other hand, when [Mo content] / [Mo content + W content] exceeds 0.5, the μ phase (Laves phase) precipitates, destabilizing the matrix structure at high temperature and increasing the temperature as described above. In order to deteriorate the tissue stability, it is desirable to make it 0.5 or less. For the same reason, it is more desirable to set the lower limit to 0.25 and the upper limit to 0.5.

Al: 0.5 to 1.0%
Al combines with Ni to precipitate a γ ′ phase and contributes to strengthening of the alloy. When Al is less than 0.5%, sufficient precipitation strengthening cannot be obtained, but too much addition can cause a concentrated region and no precipitation zone due to coarse aggregation at the γ 'phase grain boundary, resulting in deterioration of high temperature characteristics. In addition, the notch sensitivity is deteriorated, and the mechanical characteristics are greatly reduced. Therefore, the addition amount of Al is limited to the range of 0.5 to 1.0%. For the same reason, it is desirable to set the lower limit to 0.5% and the upper limit to 0.9%.

Ti: 1.0 to 2.5%
Ti mainly forms MC carbide and suppresses the coarsening of the crystal grains of the alloy, and like Al, it binds with Ni and precipitates a γ ′ phase, contributing to strengthening of the alloy. However, too much addition reduces the stability of the γ 'phase at high temperatures and precipitates the η phase, leading to a decrease in strength and ductility, and long-term tissue stability. Therefore, the addition amount of Ti is limited to the range of 1.0 to 2.5%. For the same reason, it is desirable to set the lower limit to 1.5% and the upper limit to 2.0%.

0.2 ≦ Al / Ti ≦ 1.0
Al and Ti combine with Ni as Ni ₃ (Ti, Al) to precipitate a γ ′ phase and contribute to strengthening of the alloy. However, it has been found that when the ratio of Al to Ti, that is, Al / Ti exceeds 1.0, the tensile ductility at high temperature is drastically lowered, and the creep rupture life and creep rupture ductility are markedly lowered. In particular, the decrease in creep rupture ductility must be avoided because it leads to weakening of the notch. On the other hand, when the ratio is less than 0.2, a large amount of η phase precipitates without forming a γ ′ phase, so in order to balance both the mechanical properties at high temperature and the structural stability, Al The ratio of Ti to Ti is preferably 0.2 or more and 1.0 or less. For the same reason, it is more desirable to set the lower limit to 0.4% and the upper limit to 0.8%.

Nb: 1.0% or less Nb is a precipitation strengthening element like Al and Ti, and precipitates the γ ″ phase and contributes to strengthening of the alloy, so it is contained as desired. The intermetallic compounds such as phases are liable to precipitate and the structure stability is remarkably impaired, so that the Nb content is optionally set to 1.0% or less. .2% or more is desirable, and for the same reason as described above, it is desirable to further limit the upper limit to 0.5%.

B: 0.015% or less B is segregated at the grain boundary and contributes to high temperature characteristics, so is contained as desired. However, too much addition tends to form a boride, and conversely causes grain boundary embrittlement. Therefore, if desired, the B content is 0.015% or less. In order to sufficiently obtain the above action, the content is preferably 0.0001% or more, and for the same reason as described above, the upper limit is preferably 0.01%, and the upper limit is 0.005%. Is more desirable.

Mg: 0.01% or less Mg is mainly combined with S to form a sulfide, and has the effect of improving hot workability. However, too much addition conversely causes grain boundary embrittlement and significantly reduces hot workability. Therefore, the Mg content is limited to a range of 0.01% or less.

First aging treatment: 800 ° C. to 1000 ° C.
By the first aging treatment, carbides are precipitated at the grain boundaries and the creep characteristics are improved. Here, when the aging treatment temperature is less than 800 ° C., the carbides are not effectively precipitated, and when heated above 1000 ° C., the agglomeration and coarsening of the carbides occur and the effect disappears. Therefore, the first aging treatment temperature is set to 800 ° C to 1000 ° C. The aging treatment time is appropriately set depending on the amount of C added and the shape and size of the member to be heat-treated.

Second aging treatment: 720-780 ° C
The γ ′ phase is precipitated by the second aging treatment to give a desired strength. Here, when the aging treatment temperature is less than 720 ° C., the γ ′ phase does not sufficiently precipitate, and aging precipitation further proceeds during use in an actual high temperature environment, and not only the structure but also various material properties change. It will end up. On the other hand, when heated above 780 ° C., the γ ′ phase becomes coarse, so that effective precipitation strengthening cannot be achieved. Therefore, the second aging treatment temperature is set to 720 ° C to 780 ° C. The aging treatment time is appropriately set depending on the addition amount of the precipitation strengthening element and the shape and size of the member to be heat-treated. For example, 10 to 100 hours is appropriate.

Step from the first aging treatment to the second aging treatment The first aging treatment can be followed by the second aging treatment continuously or once after passing through the cold material. However, it is desirable to perform the second aging treatment continuously at a cooling rate of 20 ° C./hour or more by cooling the fan from the first aging treatment. This is because carbide and γ ′ phase are inevitably deposited in the cooling from the first to the second aging treatment, or in the heating step, so that the carbide and γ ′ phase are not cooled at a cooling rate of less than 20 ° C./hour. This is because it becomes difficult to control the desired mechanical properties due to excessive precipitation.

As described above, the invention of the Ni-based alloy having excellent structure stability and high temperature strength according to the present invention is mass%, C: 0.005 to 0.1%, Cr: 8 to 15%, Mo: 1. -7%, W: 5-20%, Al: 0.5-1.0%, Ti: 1.0-2.5%, optionally containing Nb: 1.0% or less, Further, if desired, it contains B: 0.015% or less, and if desired, Mg: 0.01% or less, and the balance is composed of Ni and inevitable impurities. Excellent high temperature characteristics can be obtained.
Further, by defining the balance between Mo and W as 0.1 ≦ Mo / (Mo + W) ≦ 0.5, the high-temperature strength, high-temperature ductility, and creep properties are further improved.
Furthermore, by defining the balance of Al and Ti as 0.2 ≦ Al / Ti ≦ 1.0, the high-temperature tensile properties and creep properties are further improved.

  In addition, the manufacturing method of the Ni-based alloy material having excellent structure stability and high-temperature strength according to the present invention is a Ni-based alloy having the above composition, subjected to solution treatment at a recrystallization temperature or higher after hot forging, Since the first aging treatment is performed in the range of 800 ° C. to 1000 ° C., and then the second aging treatment is performed in the range of 720 ° C. to 780 ° C., the γ ′ phase is sufficiently precipitated so as not to become coarse, Thus, precipitation of the μ phase can be avoided, and thus a Ni-based alloy material having a stable structure state without causing overaging when used at high temperatures and having excellent high temperature characteristics can be obtained.

Hereinafter, an embodiment of the present invention will be described.
The Ni-based alloy of the present invention can be melted by a conventional method, and the manufacturing method is not particularly limited. However, it is desirable that the alloy of the present invention does not contain impurities such as Si, Mn, P, S, O, and N as much as possible. Therefore, preferably, the so-called double melt method using the VIM-ESR process, or VIM-ESR is preferably used. -A solubilization method such as a so-called triple melt method using a VAR process is desirable. Preferably, Si is 0.3% or less, Mn is 0.2% or less, P is 0.01% or less, S is 0.005% or less, O is 30 ppm or less, and N is 60 ppm or less.

  The melted Ni-base alloy is usually subjected to hot forging and the structure is adjusted by processing. In the present invention, hot forging conditions and the like are not particularly limited, and can be performed according to, for example, a conventional method.

  After the hot forging, a solution treatment is performed by heating above the recrystallization temperature. This solution treatment can be performed at, for example, 1050 to 1200 ° C. As the solution treatment time, an appropriate time is set according to the size and shape of the material. The solution treatment can be performed using a known heating furnace, and the heating method and the heating equipment are not particularly limited in the present invention. After the solution treatment, cooling is performed by water cooling, oil cooling, air cooling, or the like.

After the solution treatment, the first aging treatment is performed using a known heating furnace or the like. The aging treatment is performed at a temperature of 800 ° C to 1000 ° C. In the temperature increase to the aging treatment temperature, the temperature increase rate is not particularly limited in the present invention. After the first aging treatment, the second aging treatment is performed. However, the second aging treatment may be performed continuously or once after passing through the cooling material. In the second aging treatment after passing through the cold material, the same heating furnace or the like may be used, or another heating furnace or the like may be used.
It should be noted that the first aging treatment to the second aging treatment are preferably performed continuously by cooling with fan cooling or the like, and the cooling rate at that time is 20 ° C./hour or more. desirable.

  After the second aging treatment, the cooling rate is not particularly limited, and cooling can be performed by cooling or forced cooling. In the method of the present invention, the first aging process and the second aging process are defined as described above, but the subsequent aging process is not excluded, and the third and subsequent aging processes are performed as necessary. An aging treatment can also be applied.

  The Ni-based alloy material of the present invention can be used for a material such as a turbine rotor of a generator member. However, the application of the present invention is not limited to these, and it can be used for various applications requiring strength characteristics at high temperatures. Moreover, it is excellent also in long-term stability at high temperature, and can be used naturally even in a temperature range of a conventional generator member of, for example, about 600 to 650 ° C.

Examples of the present invention will be described below.
A Ni-based alloy having the composition shown in Table 1 (remainder Ni and other inevitable impurities) was melted into a 50 kg ingot by a vacuum induction melting furnace (VIM). After subjecting the test ingot to diffusion treatment, a test material was obtained as a plate having a thickness of 30 mm by hot forging. Each sample material was subjected to a solution treatment at a temperature equal to or higher than the recrystallization temperature, and then air-cooled to obtain a cold material. 840 ° C. × 10 hours of stabilization treatment, and then heat treatment under the conditions of 840 ° C. × 10 hours as the first aging treatment, and cooling by furnace cooling (cooling rate 50 ° C./h), The second aging treatment was continuously performed. In the second aging treatment, heat treatment was performed under conditions of 750 ° C. × 24 hours, and then cooled by furnace cooling (cooling rate 50 ° C./h) to prepare a test material.

  About the obtained test material, the room temperature tensile test and the high temperature (700 degreeC) tensile test were done, and also the creep test which changed load conditions was done. A list of test results is shown in FIGS. As is apparent from FIG. 1, the test materials (Nos. 1 to 14) of the present invention have both good strength and high ductility at both room temperature and 700 ° C. in short-time tensile properties. As shown in the graph, the specimen of the present invention shows a good rupture time and rupture ductility even in the creep rupture characteristics at 700 ° C. On the other hand, the comparative materials (Nos. 15 to 24) had the test results that the creep rupture properties were insufficient or vice versa even though the short-time tensile properties were good. That is, the short-time strength at room temperature and 700 ° C. shows a sufficient strength, but the ductility is low, particularly a remarkable decrease in creep rupture ductility is observed, or the short-time tensile ductility and creep rupture ductility are good values. However, the result was that the creep rupture ductility was completely insufficient due to insufficient strength. That is, the material of the present invention was excellent in the balance between strength and ductility at room temperature and 700 ° C., and was able to obtain extremely high creep rupture time and creep rupture ductility at 700 ° C.

  3 and 4 show the short-time tensile properties and creep rupture properties at 700 ° C. of the specimens with varying Al / Ti ratios. The relative evaluation result compared with 4 as 1 is shown. The test material No. 4 is Al / Ti = 0.8 / 1.8≈0.3. 3 is Al / Ti = 0.8 / 1.6≈0.5. 1 is Al / Ti = 0.8 / 1.2≈0.7. 20 is Al / Ti = 1.4 / 0.8≈1.8. 19 is Al / Ti = 1.8 / Free (Ti-free material), and the contents of elements other than Al and Ti are almost constant. In the short-time tensile properties at 700 ° C, there is no correlation between strength and Al / Ti ratio, or the effect of no Ti addition, but the test material (No. 20) with an Al / Ti ratio exceeding 1 and no Ti addition In the test material (No. 19), the ductility tended to decrease slightly. On the other hand, in the creep rupture characteristics at 700 ° C., the creep rupture time tended to clearly decrease as the Al / Ti ratio increased and in the test material without addition of Ti. In particular, in the creep rupture ductility, a remarkable decrease occurs in the specimen (No. 20) having an Al / Ti ratio exceeding 1 and the specimen without addition of Ti (No. 19), and the Al / Ti ratio is 0. The difference with the test material of less than .8 was obvious.

  The above results are shown in No. 1 shown in FIG. 3 and no. This can be explained by the EPMA analysis results in the vicinity of 20 grain boundaries. In the case of the stabilization treatment + aging treated material (840 ° C. × 10 h + 750 ° C. × 24 h), no concentration of components is observed at the grain boundary regardless of the Al / Ti ratio, but the stabilization treatment + aging treatment material is further added to 700. It was revealed that when aging treatment was performed at a temperature of 1000 ° C. for 1000 hours, Al was concentrated at the grain boundaries. This is because the γ 'phase is preferentially precipitated at the grain boundaries during the long-term aging treatment, and is agglomerated and coarsened as the aging time elapses. That is, in a material with a high Al / Ti ratio, the γ 'phase is likely to precipitate at the grain boundaries, and the speed of precipitation-aggregation coarsening is significantly increased. In the creep test that becomes longer, the fracture ductility is significantly reduced. Therefore, in order to obtain good characteristics at high temperatures, it is necessary to keep the upper limit of the Al content low, and it has become clear that it is desirable to appropriately determine the Al / Ti balance.

  Furthermore, in the test materials shown in Table 1, the test material Nos. With different Mo / W ratios were used. 22, no. 1, no. 2, and no. FIG. 6 shows the result of comparison of the creep rupture characteristics of No. 23 at 700 ° C. The test material No. No. 22 is a Mo-free material, test material No. 1 is Mo / (Mo + W) = 0.3; 2 is Mo / (Mo + W) = 0.5. 23 is a W additive-free material, and the contents of elements other than Mo and W are almost constant. As is apparent from FIG. 6, the test material No. 15 and W additive-free material No. No. 16 significantly reduced the creep rupture ductility, and it was found that good creep rupture characteristics and rupture ductility can be obtained by the combined addition of Mo and W.

  Furthermore, the test material No. having a varied Mo / W ratio was used. 22, no. 1, no. 2, and no. When the material was subjected to a long-term aging treatment of 700 ° C. × 1000 h on the material (840 ° C. × 10 h + 750 ° C. × 24 h) as it was, the 14% Mo material (W Free) as shown in FIG. Specimen No. 23, a large number of needle-like precipitates presumed to be the μ phase were confirmed at the grain boundaries and within the grains. On the other hand, test material No. with Mo amount of 7% or less. 22, no. 1, no. In No. 2, the microstructure after the long-term aging treatment at 700 ° C. × 1000 h did not change at all from the structure of the material as it was in the stabilization treatment + aging treatment, and it was found that the high-temperature structure stability was very excellent.

  In addition, the test material No. having a varied Mo / W ratio was used. 22, no. 1, no. 2, and no. For No. 23, a tensile test was performed at room temperature and 700 ° C. under the four aging conditions shown in FIG. The result is shown in FIG. Under any aging condition, the room temperature strength decreases as the value of Mo / (Mo + W) increases, and in particular, the T.V. at Mo / (Mo + W) = 1. S. The decrease of was remarkable. The room temperature ductility maintained a good value up to Mo / (Mo + W) = 0.5, but when Mo / (Mo + W) = 1, it decreased abruptly as with the strength. On the other hand, the strength at 700 ° C. is 0.2% Y.V. S. And T. S. And showed different behavior. That is, 0.2% Y.P. S. Was decreased as the value of Mo / (Mo + W) was increased as in the case of room temperature strength. S. Tended to be maximum at Mo / (Mo + W) = 0.3 and 0.5. Further, the ductility at 700 ° C. takes a minimum value at Mo / (Mo + W) = 0, and shows a tendency to improve as the value of Mo / (Mo + W) increases. The tendency was prominent. Therefore, regarding the Mo amount and the W amount, from the viewpoint of these mechanical properties and structure stability, the amount of Mo is limited to 1 to 7%, and 0.1 ≦ Mo / (Mo + W) ≦ 0.5. Obviously it is desirable to do so.

It is a figure which shows the test result of the tensile property in room temperature and high temperature of the test material in one Example of this invention. Similarly, it is a figure which shows the result of the creep test in 700 degreeC of a test material. Similarly, it is a figure which shows the test result of the short-time tensile characteristic in 700 degreeC of the test material which changed Al / Ti ratio. Similarly, it is a figure which shows the result of the creep test in 700 degreeC of the test material which changed Al / Ti ratio. Similarly, it is the drawing substitute photograph which shows the EPMA result of the test material which changed Al / Ti ratio. Similarly, it is a figure which shows the creep rupture test result in 700 degreeC of the test material which fluctuated Mo / W ratio. Similarly, it is the drawing substitute photograph which shows the EPMA result of the test material which changed Mo / W ratio. Similarly, it is a figure which shows the heat pattern of aging conditions AD employ | adopted as an Example. Similarly, it is a figure which shows the test result of the short-time tensile property in room temperature and 700 degreeC of the test material which changed Mo / W ratio and aging conditions.

Claims (8)

In mass%, C: 0.005 to 0.1%, Cr: 8 to 15%, Mo: 1 to 7%, W: 5
Excellent structure stability and high temperature strength, characterized by containing ~ 20%, Al: 0.5-1.0%, Ti: 1.0-2.5%, the balance being made of Ni and inevitable impurities Ni-based alloy.   The Ni-based alloy having excellent structure stability and high-temperature strength according to claim 1, further comprising Nb: 1.0% or less by mass%.   The Ni-based alloy having excellent structure stability and high temperature strength according to claim 1 or 2, further comprising B: 0.015% or less in mass%.   The Ni-based alloy having excellent structure stability and high-temperature strength according to any one of claims 1 to 3, further comprising Mg: 0.01% or less in terms of mass%. The Ni content alloy excellent in structure stability and high temperature strength according to any one of claims 1 to 4, wherein the Mo content and the W content satisfy the following formula 1.
0.1 ≦ Mo / (Mo + W) ≦ 0.5 Formula 1 The Ni-based alloy having excellent structure stability and high-temperature strength according to any one of claims 1 to 5, wherein the Al content and the Ti content satisfy the following formula 2.
0.2 ≦ Al / Ti ≦ 1.0 Formula 2   The Ni-based alloy having the composition according to any one of claims 1 to 6 is subjected to a solution treatment at a recrystallization temperature or higher after hot forging, and then a first aging treatment in a range of 800 ° C to 1000 ° C. And then performing a second aging treatment in the range of 720 ° C. to 780 ° C., and a method for producing a Ni-based alloy material having excellent structure stability and high temperature strength.   The tissue stability and high temperature according to claim 7, wherein after the first aging treatment, the aging treatment is continuously performed by cooling to a second aging treatment temperature at a cooling rate of 20 ° C./h or more. A method for producing a Ni-based alloy material having excellent strength.