JP2014070230A - METHOD FOR PRODUCING Ni-BASED SUPERALLOY - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a method for producing a Ni-based superalloy having substantially improved creep rupture strength.SOLUTION: There is provided a method for producing a Ni-based superalloy which contains by mass%: 0.01-0.2% of C; 0.5% or less of Si; 0.5% or less of Mn; 10-24% of Cr; 5-17% of one or two of Mo and W in terms of Mo+0.5 W; 0.5-2.0% of Al; 1.0-3.0% of Ti; 10% or less of Fe; one or two of more than 0% and 0.02% or less of B and more than 0% and 0.2% or less of Zr; and the balance comprising Ni and impurities. The method for producing the Ni-based superalloy comprises: a step of providing a material for solution heat treatment having the above composition; a solution heat treatment step of performing solution heat treatment for 1.5-5 hours at 1,100-1,140°C (not performing multi-stage solution heat treatment) to obtain a solution heat treated material using the material for solution heat treatment; an aging treatment step of performing a first stage aging treatment at 820-880°C and a second stage aging treatment at 600-800°C using the solution heat treated material.

Description

  The present invention relates to a method for producing a Ni-base superalloy.

In recent years, in order to reduce CO ₂ emissions from the viewpoint of preventing global warming, higher efficiency of thermal power plants has been demanded. Currently, the steam temperature of thermal power plants has reached 600 to 630 ° C, but in order to further improve power generation efficiency, the steam temperature is expected to rise in the future. The use of a Ni-base superalloy having excellent breaking strength is considered promising. For example, as a Ni-based alloy that is optimal for use in a 650 ° C. class super-supercritical thermal power plant, Japanese Patent Application Laid-Open No. 9-157779 (Patent Document 1) describes Ni having excellent creep characteristics and a low thermal expansion coefficient. Inventions of basic superalloys have been proposed. In addition, in the re-published patent WO 10/038680 (Patent Document 2), an appropriate homogenization heat treatment is performed for the purpose of reducing the variation in creep characteristics due to the microsegregation of Patent Document 1 described above, and the Mo An invention of a Ni-base superalloy having a segregation ratio in the range of 1-1.17 has been proposed. Furthermore, in Japanese Patent Application Laid-Open No. 2011-84812 (Patent Document 3), for the purpose of increasing the strength of creep of a Ni-base superalloy and suppressing the decrease in ductility, the crystal grain size is adjusted by high-temperature solution treatment and low-temperature solution treatment. There has been proposed an invention for growing to a specific size and controlling the form of granular precipitates at grain boundaries.

JP-A-9-157779 Republished patent WO2010 / 038680 JP 2011-84812 A

The Ni-base superalloys described in Patent Document 1 and Patent Document 2 described above are manufactured by subjecting an ingot having a target composition to homogenization heat treatment, and then performing hot plastic working to finish it into a predetermined shape. By performing solution treatment and subsequent aging treatment on this, excellent creep characteristics that can be used for 700 ° C. class super supercritical thermal power plants can be obtained. The Ni-based superalloys described in Patent Document 1 and Patent Document 2 described above can be applied to members such as steam turbines and boiler tubes of 700 ° C. class super supercritical thermal power plants because of their excellent creep characteristics.
However, since the use environment of the steam turbine is extremely severe and higher reliability is required, it is required to maximize the creep characteristics.
Moreover, in patent document 3, the average grain of a crystal | crystallization is obtained by special solution treatment called the 1st solution treatment performed at 1100-1160 degreeC, and the 2nd solution treatment performed at 980-1080 degree after that. With a diameter of 72 to 289 μm and an average length of 0.5 to 2.5 μm, granular precipitates (carbides) are precipitated at the crystal grain boundaries, and it is possible to suppress a decrease in ductility while improving creep strength. Has been.
However, in consideration of productivity, the normal one-stage solution treatment is more advantageous than the two-stage solution treatment.
An object of the present invention is to provide a method for producing a Ni-base superalloy having significantly improved creep rupture strength while applying a normal one-step solution treatment.

The present inventor is superior in productivity and can exhibit higher creep rupture strength in the Ni-based alloys described in Patent Document 1 and Patent Document 2 as compared with the heat treatment conditions described in Patent Document 3. We studied various methods.
In the Ni-based alloys described in Patent Document 1 and Patent Document 2, the solution treatment is performed to control the solid solution of the alloy elements in the matrix and the crystal grain size, and the subsequent aging treatment mainly includes carbides and grains in the grain boundaries. A fine gamma prime phase is precipitated in it to obtain high strength.
As a control of the metal structure that improves the creep rupture strength, the present inventor has focused on the crystal grain size and the amount of carbide precipitated at the crystal grain boundary, and investigated the optimization of the C addition amount and the solution treatment conditions. As a result, carbides and gamma prime phases precipitated at the grain boundaries mainly depend on the amount of C added and the size of the crystal grains, and by optimizing the amount of C added and performing the solution treatment in a limited temperature range. The inventors have found that the creep rupture strength can be dramatically improved and have reached the present invention.

That is, the present invention is, by mass%, C: 0.01 to 0.2%, Si: 0.5% or less, Mn: 0.5% or less, Cr: 10-24%, one of Mo and W or Two types are Mo + 0.5W: 5 to 17%, Al: 0.5 to 2.0%, Ti: 1.0 to 3.0%, Fe: 10% or less, and B: 0.02% or less ( 0% is not included) and Zr: 0.2% or less (not including 0%) is included in one or two types, and the balance is Ni and an inevitable impurity Ni-base superalloy,
Preparing a solution treatment material having the composition;
A solution treatment step of obtaining a solution treatment material by performing a solution treatment for 1.5 to 5 hours at 1100 to 1140 ° C. (no multistage solution treatment is performed) using the solution treatment material; and ,
Using the solution treatment material, an aging treatment step of performing a first stage aging treatment at 820 to 880 ° C. and a second stage aging treatment at 600 to 800 ° C.,
A method for producing a Ni-base superalloy, comprising:

  According to the present invention, the creep rupture strength of a Ni-based alloy can be greatly improved. It becomes possible to improve the reliability of the steam turbine, boiler tube material, etc. of the thermal power plant using this.

It is an electron micrograph of the metal structure of the Ni base superalloy according to the manufacturing method of the present invention. It is a schematic diagram of the metal structure of the Ni-base superalloy according to the manufacturing method of the present invention.

First, in the present invention, a solution treatment material having the following components is prepared.
As the material for solution treatment, it is preferable to use a hot plastic working material which has been subjected to hot plastic working such as hot forging or hot pressing after melting. An example of the process for obtaining the solution treatment raw material is shown.
A conventional melting method may be applied for melting the Ni-base superalloy. In the present invention, in order to obtain high temperature and high strength, Al and Ti, which are constituent elements of the gamma prime phase, are indispensably added, so that precipitation of non-metallic inclusions such as harmful oxides and nitrides is prevented. It is preferable to perform some vacuum melting. It is preferable to perform remelting such as vacuum arc remelting or electroslag remelting after vacuum melting.
In addition, the ingot after vacuum melting, the electrode after vacuum melting, and the ingot after remelting are subjected to homogenization heat treatment at 1160 to 1220 ° C. for 1 to 100 hours to reduce component segregation, and for hot plastic working The material is preferred.

Next, in this invention, it is good to perform the process of obtaining the hot plastic work material by hot plastic working the raw material for hot plastic working which has the component mentioned above at 900-1200 degreeC.
Since the cast structure is a metal structure in which crystal grains grow in a specific direction during cooling after melting, anisotropy occurs in mechanical properties. Therefore, hot plastic working is performed to break the cast structure generated by melting and obtain a uniform recrystallized structure.
If the temperature of the hot plastic working is less than 900 ° C., the hot workability is poor, so that it is difficult to be deformed to a predetermined dimension and causes cracking. On the other hand, when the temperature of hot plastic working exceeds 1200 ° C., the strength of the crystal grain boundary where B or Zr segregates is greatly reduced, and hot workability is significantly impaired. Therefore, the hot plastic working is performed in a temperature range of 900 to 1200 ° C. A preferable upper limit of the temperature of the hot plastic working is 1180 ° C.
As described above, the solution processing material may be used by using the hot plastic working material described above.

Next, the reasons for limiting the components defined in the present invention will be described. In addition, content of each element is the mass%.
C: 0.01% to 0.2%
C is the most important element for obtaining excellent creep rupture strength. In order to improve the creep rupture strength, it is effective to enlarge the crystal grains by solution treatment. However, if C is more than 0.2%, the effect of suppressing the crystal grain growth becomes excessive, and it becomes difficult to control the crystal grain size. In the present invention, the upper limit of C is set to 0.2%. Further, C has an effect of strengthening the crystal grain boundaries by carrying out an aging treatment to precipitate carbides at the crystal grain boundaries and suppressing creep deformation. In order to obtain a sufficient creep rupture strength, it is necessary to deposit about 0.2 to 0.4 mol% of carbides at the grain boundaries. The grain boundary carbide has an effect of strengthening the grain boundary and at the same time promoting the precipitation of the gamma prime phase to the grain boundary, and the effect of suppressing the grain boundary slip during creep is enhanced. However, if C is less than 0.01%, sufficient carbide cannot be precipitated at the crystal grain boundary, and the gamma prime phase is insufficiently precipitated at the crystal grain boundary. As a result, the grain boundary strengthening becomes insufficient and the creep rupture strength decreases. Therefore, the lower limit of C is defined as 0.01%. The preferable lower limit of C is 0.015%, and the preferable upper limit is 0.1%. A more preferred lower limit is 0.02%, and a more preferred upper limit is 0.05%.
Si: 0.5% or less Si is used as a deoxidizer during alloy melting. Moreover, Si has an effect of suppressing peeling of the oxide film. However, when it contains excessively, ductility and workability will fall, It limits to 0.5% or less.
Mn: 0.5% or less Mn is used as a deoxidizing agent or a desulfurizing agent during alloy melting. Deoxidation and desulfurization are performed using Mn because O and S as inevitable impurities segregate at the grain boundaries and cause hot brittleness. However, since ductility will fall when it adds excessively, it limits to 0.5% or less.

Cr: 10 to 24%
Cr combines with C to generate carbides at the grain boundaries, thereby strengthening the grain boundaries and improving the strength and ductility at high temperatures. In addition, it has the effect of improving the oxidation resistance and corrosion resistance of the alloy by dissolving in the matrix, and greatly reducing notch sensitivity. If the Cr content is less than 10%, the above-mentioned effects cannot be obtained with certainty, and excessive addition exceeding 24% causes a problem that the manufacturability, workability, and mechanical properties of the alloy deteriorate due to the formation of the brittle phase. For these reasons, Cr is limited to 10 to 24%. A preferable lower limit of Cr is 15%.
The amount specified by the formula “Mo + 0.5W”, 5 to 17% of one or two of Mo and W
Mo and W are effective in reducing the thermal expansion coefficient of the alloy as well as strengthening the matrix by dissolving in the matrix. Since the Ni-based alloy has a large coefficient of thermal expansion, there is a difficulty in that it is liable to cause thermal fatigue when used stably at high temperatures and lacks reliability. Since Mo is the most effective element for lowering the thermal expansion coefficient, Mo is essential, and Mo alone or Mo and W are added. In order to reliably obtain this effect, 5 + 17% is required at Mo + 0.5W. If the Mo + 0.5W amount is less than 5%, the above effect cannot be obtained reliably. If the Mo + 0.5W amount is more than 17%, the productivity and workability of the alloy tend to be difficult. Limit one or two to 5-17%.

Al: 0.5 to 2.0%
Al forms an intermetallic compound (Ni ₃ (Al, Ti)) called a gamma prime phase together with Ni and Ti, and is added to increase the high temperature strength of the alloy. The gamma prime phase precipitates in the grains and at the grain boundaries, and contributes to the strengthening of the grains and the grain boundaries. In particular, the precipitation at the grain boundary, like the grain boundary carbide, suppresses the slip of the grain boundary and contributes to the improvement of the creep rupture strength. If the content is less than 0.5%, the above effect cannot be obtained, and excessive addition deteriorates the manufacturability and workability of the alloy, so Al is limited to 0.5 to 2.0%.
Ti: 1.0-3.0%
Ti, like Ni and Al, has the effect of increasing the high temperature strength of the gamma prime phase (Ni ₃ (Ti, Al)) alloy. Since the atomic diameter of Ti is larger than that of Ni and gives an elastic strain to the matrix, it contributes to strengthening more than Ni ₃ Al. If less than 1.0%, the above effect cannot be obtained. If excessively added, the productivity and workability of the alloy deteriorate, so Ti is limited to 1.0 to 3.0%.

Fe: 10% or less Fe is not necessarily added, but it has an effect of improving the hot workability of the alloy, and can be added as necessary. When Fe is added excessively, the thermal expansion coefficient of the alloy increases, and there is a problem that cracks occur when used at high temperatures. Moreover, since oxidation resistance deteriorates, it limits to 10% or less. A preferable upper limit of Fe is 5%, and a more preferable upper limit of Fe is 3%.
One or two of B: 0.02% or less (not including 0%) and Zr: 0.2% or less (not including 0%) B and Zr are used for strengthening grain boundaries, It is necessary to add either or both of Zr and Zr. Since B and Zr are significantly smaller in size than Ni which is an atom constituting the matrix, B and Zr are segregated at the crystal grain boundary and have an effect of suppressing the grain boundary sliding at a high temperature. In particular, it has the effect of greatly reducing notch sensitivity. Therefore, the effect of improving the creep rupture strength and creep rupture ductility can be obtained, but if added excessively, the oxidation resistance deteriorates, so B and Zr are each 0.02% or less (excluding 0%), 0.2 % Or less (excluding 0%).

  The remaining Ni is a main constituent element of the alloy according to the present invention, and is an element that forms a gamma phase of a matrix having an fcc structure while dissolving other elements. The gamma phase of the matrix has a large solid solubility limit of the alloy element, which is advantageous for the precipitation of the gamma prime phase, which is the key to precipitation strengthening. Moreover, since Ni is a main constituent element of the gamma prime phase that is a precipitation strengthening phase, it is necessary to have a sufficient amount with respect to the amounts of Al and Ti. In the present invention, the balance is Ni. Of course, impurities are included.

Solution Treatment Step In the present invention, a solution treatment material is obtained by performing a solution treatment at 1100 to 1140 ° C. using the hot plastic working material.
The solution treatment is performed for the purpose of adjusting the crystal grain size of the recrystallized structure obtained by hot plastic working and for once dissolving the constituent elements for precipitating the gamma prime phase and carbides in the matrix by aging treatment. In the present invention, in order to improve the creep rupture strength, the grains are coarsened to a certain extent, and at the same time, the carbide and gamma prime phase are continuously formed along the grain boundaries in the first stage aging treatment so as to form a rust-like shape. To precipitate. The grain boundary structure of the present invention can be obtained by sufficiently precipitating carbides at the grain boundaries by the first stage aging treatment and simultaneously precipitating the gamma prime phase at the grain boundaries starting from the carbides. That is, the grain boundary structure of the present invention is determined by the balance of crystal grain size, C addition amount, and aging treatment temperature. Therefore, for example, after the solution treatment as disclosed in Patent Document 3, creep rupture is performed without performing the second-stage solution treatment in which only carbides are precipitated at a temperature equal to or higher than the solid solution temperature of the gamma prime phase. A sufficient effect of strength improvement can be obtained.

The solution treatment temperature applied in the present invention necessary for obtaining the above grain boundary structure is set to 1100 to 1140 ° C., which is higher than the conventional temperature range and in a limited temperature range. Even when the temperature of the solution treatment is lower than 1100 ° C., the solid solution of the compound element in the matrix is not sufficient, and the effect of crystal grain growth sufficient to improve the creep rupture strength cannot be expected. On the other hand, when the temperature of the solution treatment is higher than 1140 ° C., although there is an effect of improving the creep rupture strength due to the coarsening of the crystal grains, the creep rupture ductility is lowered. Therefore, the solution treatment temperature is 1100 to 1140 ° C. The preferable lower limit of the solution treatment temperature is 1110 ° C., and the preferable upper limit is 1130 ° C.
In addition, the solution treatment time for obtaining the above effect is 1.5 to 5 hours. If the solution treatment time is less than 1.5 hours, the crystal grain size is not sufficiently grown, so that the proportion of carbides precipitated in the subsequent first stage aging treatment becomes insufficient, and as a result Thus, a continuous and wedge-shaped structure composed of carbides and gamma prime phases at grain boundaries cannot be obtained. In addition, since the purpose of the solution treatment is solid solution of the constituent elements of the precipitate and the coarsening of the crystal grains, the heat treatment in which the solution treatment time exceeds 5 hours is longer than necessary and is not economical.

Aging treatment step In the present invention, the solution treatment material is subjected to a first-stage aging treatment at 820 to 880 ° C and a second-stage aging treatment at 600 to 800 ° C.
In order to obtain sufficient creep rupture strength, it is necessary not only to adjust the crystal grain size by solution treatment, but also to strengthen the grain boundaries and crystal grains by aging treatment. As described above, in the first stage aging treatment, carbides and gamma prime phases are sufficiently precipitated continuously and in a wedge shape at the grain boundaries to strengthen the grain boundaries. When the temperature of the first stage aging treatment is less than 820 ° C., the amount of the precipitated compound is small and the strengthening of the crystal grain boundary becomes insufficient. On the other hand, when the temperature of the first stage aging treatment exceeds 880 ° C., the effect of precipitation strengthening in the grains is impaired due to the coarse growth of the gamma prime phase that precipitates in the grains. Therefore, the temperature of the first stage aging treatment is 820 to 880 ° C. The preferable lower limit of the temperature of the first stage aging treatment is 840 ° C., and the preferable upper limit is 860 ° C.
In addition, about 1 to 10 hours are sufficient for the time of the 1st stage aging treatment for obtaining the said effect.
In the second-stage aging treatment performed after the first-stage aging treatment described above, the gamma prime phase is finely precipitated in the crystal grains to strengthen the inside of the crystal grains. If the temperature of the second stage aging treatment is less than 600 ° C., the precipitation of the gamma prime phase is insufficient and the high temperature strength is insufficient. On the other hand, when the temperature of the second stage aging treatment exceeds 800 ° C., the gamma plum phase becomes coarse and the high-temperature strength tends to decrease. Therefore, in the present invention, the second stage aging treatment is set to 600 to 800 ° C. The preferable lower limit of the temperature of the second stage aging treatment is 700 ° C., and the preferable upper limit is 780 ° C.
In addition, the time of the second stage aging treatment for obtaining the above effect is preferably about 10 to 30 hours.

  As described above, when the production method of the present invention to be described is applied, the creep rupture strength of the Ni-base superalloy can be improved. As a result, high reliability can be imparted as a member such as a steam turbine blade and a bolt of a 700 ° C. class ultra-supercritical thermal power plant.

A 10 kg ingot was produced by vacuum induction melting, and Ni-base superalloys having chemical components shown in Table 1 were obtained. The balance not shown in the table is Ni and impurities.
In order to clarify the relationship between the C content and the solution treatment, the contents of the elements specified in the present invention other than C were fixed.

A Ni-based superheat-resistant alloy ingot having the composition shown in Table 1 was subjected to homogenization heat treatment at 1200 ° C. to obtain a material for hot plastic working. Using the aforementioned hot plastic working material, hot plastic working was performed at 1150 ° C. to obtain a hot plastic work material.
Table 2 shows the conditions of the solution treatment step that affects the creep strength and the conditions of the aging treatment step.

The creep rupture properties of the Ni-base superalloys subjected to the solution treatment step and the aging treatment step described above were evaluated. The evaluation was made as a creep rupture test at 700 ° C. and 750 ° C. assuming use in a 700 ° C. class super supercritical thermal power plant.
The creep rupture test was performed under the conditions of a test temperature of 700 ° C., a load stress of 385 N / mm ² , a test temperature of 750 ° C., and a load stress of 242 N / mm ² . The creep rupture test results are shown in Table 3. In addition, the average crystal grain size shown in Table 3 is after the aging treatment step, and the amount of carbide is obtained by calculation using the CALPHAD method. In addition, alloy No. after the aging treatment step. The electron micrograph of the metal structure of 2 and its schematic diagram are shown in FIGS.

From Table 3, the results of the creep rupture test at 700 ° C. and 750 ° C. are shown for Alloy No. of the present invention. 1. Alloy no. 2 is Comparative Example No. Compared with 3, 4, and 5, the creep rupture life is improved. Regarding the creep rupture drawing, it can be seen that the alloy of the present invention satisfies the sufficient ductility although it is slightly inferior to the comparative example. It can be seen that the most important creep rupture strength can be improved within a range where satisfactory creep rupture ductility can be satisfied.
On the other hand, Comparative Example No. No. 5 shows the alloy No. Although the crystal grain size is about the same as that of Nos. 1 and 2, since the solution treatment temperature is low, undissolved carbides remain, and carbide precipitation at the grain boundaries is insufficient during aging treatment. Compared with 1 and 2, the creep rupture strength was low.
In addition, No. to which the manufacturing method of the present invention was applied. In the alloys 1 and 2, as shown in the metal structure photographs of FIG. 1 and FIG. 2 and the schematic diagram thereof, a large number of carbides and gamma prime phases are continuously and wedged precipitated along the grain boundaries, The creep rupture strength can be improved by suppressing the boundary slip.

1 Matrix 2 Gamma Prime Phase 3 Precipitate (Carbide and Gamma Prime Phase)

Claims (1)

By mass%, C: 0.01 to 0.2%, Si: 0.5% or less, Mn: 0.5% or less, Cr: 10 to 24%, and one or two of Mo and W are added to Mo + 0. 5W: 5 to 17%, Al: 0.5 to 2.0%, Ti: 1.0 to 3.0%, Fe: 10% or less, and B: 0.02% or less (0% not included) ) And Zr: 0.2% or less (not including 0%) or two of them, the balance being Ni and an inevitable impurity Ni-based superalloy,
Preparing a solution treatment material having the composition;
A solution treatment step of obtaining a solution treatment material by performing a solution treatment for 1.5 to 5 hours at 1100 to 1140 ° C. (no multistage solution treatment is performed) using the solution treatment material; and ,
Using the solution treatment material, an aging treatment step of performing a first stage aging treatment at 820 to 880 ° C. and a second stage aging treatment at 600 to 800 ° C.,
A method for producing a Ni-base superalloy, comprising: